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Creep-resistant steels

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Creep-resistantsteels

Edited byFujio Abe, Torsten-Ulf Kern and R. Viswanathan

Woodhead Publishing and Maney Publishingon behalf of

The Institute of Materials, Minerals & Mining

CRC PressBoca Raton Boston New York Washington, DC

W O O D H E A D P U B L I S H I N G L I M I T E DCambridge England

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Woodhead Publishing Limited and Maney Publishing Limited on behalf ofThe Institute of Materials, Minerals & Mining

Woodhead Publishing Limited, Abington Hall, AbingtonCambridge CB21 6AH, Englandwww.woodheadpublishing.com

Published in North America by CRC Press LLC, 6000 Broken Sound Parkway, NW,Suite 300, Boca Raton, FL 33487, USA

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Contents

Contributor contact details xiii

Preface xix

Part I General

1 Introduction 3

F. ABE, National Institute for Materials Science (NIMS), Japan

1.1 Definition of creep 31.2 Creep and creep rate curves 31.3 Creep rupture data 71.4 Deformation mechanism map 91.5 Fracture mechanism map 111.6 References 14

2 The development of creep-resistant steels 15

K.-H. MAYER, ALSTOM Energie GmbH, Germany and F. MASUYAMA,Kyushu Institute of Technology, Japan

2.1 Introduction 152.2 Requirements for heat-resistant steels 182.3 Historical development of ferritic steels 192.4 Historical development of austenitic steels 422.5 Historical development of steel melting and of the purity

of heat-resistant steels 642.6 Summary 672.7 References 70

3 Specifications for creep-resistant steels: Europe 78

G. MERCKLING, RTM BREDA Milano, Italy

3.1 Introduction 78

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3.2 Specifications and standards 813.3 The European Creep Collaborative Committee (ECCC) 853.4 European Pressure Equipment Research Council

(EPERC) 923.5 The latest generation of CEN standards for creep-resistant

steels 953.6 Future trends 1503.7 References 151

4 Specifications for creep-resistant steels: Japan 155

F. MASUYAMA, Kyushu Institute of Technology, Japan

4.1 Introduction 1554.2 Types of heat-resistant steels in Japan 1554.3 Specifications for high temperature tubing and piping

steels 1584.4 Specifications for steam turbine steels 1694.5 Heat-resistant super alloys 1694.6 Summary 1694.7 References 173

5 Production of creep-resistant steels for turbines 174

Y. TANAKA, Japan Steel Works, Japan

5.1 Introduction 1745.2 Overview of production technology of rotor shaft

forgings for high temperature steam turbines 1755.3 Production and properties of turbine rotor forgings for

high temperature applications 1925.4 Future trends 2075.5 References 212

Part II Behaviour of creep-resistant steels

6 Physical and elastic behaviour of creep-resistant steels 217

Y. YIN and R.G. FAULKNER, Loughborough University, UK

6.1 Introduction 2176.2 Elastic behaviour 2196.3 Thermal properties of creep-resistant steels 2256.4 Electrical resistivity and conductivity of creep-resistant steels 2346.5 Implications for industries using creep-resistant steels 2386.6 Future trends 2396.7 References 239

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7 Diffusion behaviour of creep-resistant steels 241

H. OIKAWA and Y. IIJIMA, Tohoku University, Japan

7.1 Introduction 2417.2 Diffusion and creep 2417.3 Diffusion characteristics 2437.4 Roles of atom/vacancy movement in creep 2487.5 Influence of some factors on creep through their effects

on diffusion 2507.6 Diffusion data in iron and in some iron-base alloys 2557.7 Concluding remarks 2607.8 References 263

8 Fundamental aspects of creep deformation anddeformation mechanism map 265

K. MARUYAMA, Tohoku University, Japan

8.1 Introduction 2658.2 Stress–strain response of materials 2658.3 Temperature and strain rate dependence of yield stress 2678.4 Deformation upon loading of creep test 2698.5 Creep behavior below and above athermal yield stress 2708.6 Change in creep behavior at athermal yield stress σa 2718.7 Deformation mechanism maps 2758.8 Concluding remarks 2788.9 References 278

9 Strengthening mechanisms in steel for creep andcreep rupture 279

F. ABE, National Institute for Material Science (NIMS), Japan

9.1 Introduction 2799.2 Basic ways of strengthening steels at elevated temperature 2799.3 Strengthening mechanisms in modern creep-resistant steels 2879.4 Loss of strengthening mechanisms in 9–12Cr steels

during long time periods 2959.5 Future trends 3019.6 References 301

10 Precipitation during heat treatment and service:characterization, simulation and strength contribution 305

E. KOZESCHNIK and I. HOLZER, Graz University of Technology, Austria

10.1 Introduction 305

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10.2 Microstructure analysis of the COST alloy CB8 30610.3 Modelling precipitation in complex systems 31210.4 Computer simulation of the precipitate evolution in CB8 31510.5 Microstructure–property relationships 32010.6 The back-stress concept 32210.7 Loss of precipitation strengthening during service of CB8 32410.8 Summary and outlook 32510.9 References 326

11 Grain boundaries in creep-resistant steels 329

R.G. FAULKNER, Loughborough University, UK

11.1 Introduction 32911.2 Ferritic steels 33011.3 Austenitic steels 34111.4 Grain boundary properties and constitutive creep design

equations 34511.5 Future trends 34611.6 References 347

12 Fracture mechanism map and fundamental aspectsof creep fracture 350

K. MARUYAMA, Tohoku University, Japan

12.1 Introduction 35012.2 Fracture mechanisms and ductility of materials 35112.3 Stress and temperature dependence of rupture life 35212.4 Fracture mechanism maps 35512.5 Influence of fracture mechanism change on creep rupture

strength 35612.6 Influence of microstructural degradation on creep rupture

strength 35812.7 Change in creep rupture properties at athermal yield stress 35912.8 Multi-region analysis of creep rupture data 36112.9 Summary 36212.10 References 364

13 Mechanisms of creep deformation in steel 365

W. BLUM, University of Erlangen-Nuernberg, Germany

13.1 Introduction 36513.2 Initial microstructure 36613.3 Creep at constant stress 36813.4 Transient response to stress changes 370

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13.5 Cyclic creep 37413.6 Microstructural interpretation of creep rate 37513.7 Dislocation models of creep 38513.8 In situ transition electron microscope observations of

dislocation activity 38913.9 Discussion and outlook 39313.10 Acknowledgments 39513.11 References 39513.12 Appendix: Microstructural model Mikora 401

14 Constitutive equations for creep curves andpredicting service life 403

S.R. HOLDSWORTH, EMPA – Materials Science & Technology,Switzerland

14.1 Introduction 40314.2 Constitutive equations 40514.3 Constitutive equation selection 40514.4 Predicting service life 41214.5 Future trends 41614.6 Concluding remarks 41614.7 Nomenclature 41614.8 References 417

15 Creep strain analysis for steel 421

B. WILSHIRE and H. BURT, University of Wales Swansea, UK

15.1 Introduction 42115.2 Creep-induced strain 42215.3 Patterns of creep strain accumulation 42715.4 Practical implications of creep strain analysis 43315.5 Future data analysis options 44115.6 References 442

16 Creep fatigue behaviour and crack growth of steels 446

C. BERGER, A. SCHOLZ, F. MUELLER and M. SCHWIENHEER, DarmstadtUniversity of Technology, Germany

16.1 Introduction 44616.2 Creep–fatigue experiments 44716.3 Stress–strain behaviour 44916.4 Creep–fatigue interaction, life estimation 44916.5 Multiaxial behaviour 45616.6 Creep and creep–fatigue crack behaviour 459

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16.7 Concluding remarks 46816.8 Acknowledgements 46916.9 References 469

17 Creep strength of welded joints of ferritic steels 472

H. CERJAK and P. MAYR, Graz University of Technology, Austria

17.1 Introduction 47217.2 Influence of weld thermal cycles on the microstructure of

ferritic heat-resistant steels 47417.3 Weld metal development for creep-resistant steels 48217.4 Creep behaviour of welded joints 48317.5 Selected damage mechanism in creep-exposed welded

joints 48417.6 Implications for industries using welded creep-resistant

steels 49517.7 Future trends 49617.8 References 498

18 Fracture mechanics: understanding inmicrodimensions 504

M. TABUCHI, National Institute for Materials Science (NIMS),Japan

18.1 Introduction 50418.2 Non-linear fracture mechanics 50418.3 Effect of mechanical constraint 50718.4 Effect of microscopic fracture mechanisms 50918.5 Type IV creep crack growth in welded joints 51318.6 References 517

19 Mechanisms of oxidation and the influence ofsteam oxidation on service life of steam powerplant components 519

P. J. ENNIS and W. J. QUADAKKERS, Forschungszentrum Juelich GmbH,Germany

19.1 Introduction 51919.2 Mechanisms of enhanced steam oxidation 52019.3 Steam oxidation rates 52519.4 Oxidation and service life 53019.5 Development of steam oxidation-resistant steels 53219.6 Outlook 533

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19.7 Sources of further information 53419.8 References 534

Part III Applications

20 Alloy design philosophy of creep-resistant steels 539

M. IGARASHI, Sumitomo Metal Industries, Japan

20.1 Introduction 53920.2 Creep-resistant steels for particular components in power

plants and the properties required 53920.3 Alloy design philosophies of creep-resistant steels 54120.4 References 570

21 Using creep-resistant steels in turbines 573

T.-U. KERN, Siemens AG Power Generation Group, Germany

21.1 Introduction 57321.2 Implications for industries using creep-resistant steels 57421.3 Improving the performance and service life of steel

components 58321.4 Next steps into the future 59121.5 Summary 59321.6 References 593

22 Using creep-resistant steels in nuclear reactors 597

S.K. ALBERT, Indira Gandhi Centre for Atomic Research, India andS. SUNDARESAN, Maharaja Sayajirao University, Baroda, India

22.1 Introduction 59722.2 Radiation damage 59822.3 Embrittlement caused by ageing 61122.4 Use of heat-resistant steels in major reactor types 61322.5 Fabrication and joining considerations 62922.6 Summary 63122.7 References 632

23 Creep damage – industry needs and future researchand development 637

R. VISWANATHAN and R. TILLEY, Electric Power ResearchInstitute, USA

23.1 Introduction 63723.2 Calculational methods for estimating damage 638

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23.3 Non-destructive evaluation methods 64323.4 Accelerated destructive tests 65323.5 High temperature crack growth 65823.6 Future trends 66223.7 References 663

Index 667

Contentsxii

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Contributor contact details

Editors

Fujio Abe*Structural Metals Center,National Institute for Materials

Science1-2-1 SengenTsukuba 305-0047Japan

Email: [email protected]

T.-U. KernSiemens AG, Power Generation

GroupDept. MaterialsRheinstr. 100D-45478 MuelheimGermany


R. (Vis)ViswanathanTechnical ExecutiveElectric Power Research Institute3420 Hillview AvePalo AltoCA 94304USA


Chapter 1

Fujio AbeStructural Metals Center,National Institute for Materials

Science1-2-1 SergenTsukuba 305–0047Japan


Chapter 2

K.-H. Mayer*Am Kirchbühl 1D-90592 SchwarzenbruckGermany

Fujimitsu MasuyamaDept. Applied Science for

Integrated System EngineeringKyushu Institute of Technology1-1 Sensui-choTobataKitakyushu 804-8550Japan

Email: [email protected];[email protected]

(* = main contact)

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Chapter 3

Gunther MercklingVia Po 8420032CORMANO MIMilanoItaly

Email:[email protected]

Chapter 4

Fujimitsu MasuyamaDept. Applied Science for

Integrated System EngineeringKyushu Institute of Technology1-1 Sensui-choTobataKitakyushu 804-8550Japan


Chapter 5

Yasuhiko TanakaThe Japan Steel Works4 Chatsu,Muroran,Hokkaido 051-8505Japan


Chapter 6

Y.F. Yin and R.G. FaulknerIPTMELoughborough UniversityAshby RoadLoughboroughLE11 3TUUK


Chapter 7

Hiroshi Oikawa*2-2 Kagitori-3Sendai 982-0804Japan


Yoshiaki Iijima37-2 Kamo-1Sendai 981-3122Japan


Chapter 8

Kouichi MaruyamaGraduate School of Environmental

StudiesTohoku University6-6-02 AobayamaAoba-kuSendai 980-8579Japan


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Chapter 9

Fujio AbeStructural Metals Center,National Institute for Materials

Science1-2-1 SergenTsukuba 305–0047Japan


Chapter 10

Ernst Kozeschnik* and Ivan HolzerInstitute for Materials Science,

Welding and FormingGraz University of TechnologyKopernikusgasse 24A-8010Austria

Email: [email protected];[email protected]

Chapter 11

R.G. FaulknerIPTMELoughborough UniversityLoughboroughLE11 3TUUK


Chapter 12

Kouichi MaruyamaGraduate School of Environmental

StudiesTohoku University6-6-02 AobayamaAoba-kuSendai 980-8579Japan

Email:[email protected],jp

Chapter 13

Wolfgang BlumDepartment of Materials Science

EngineeringInstitute I: General Materials

Properties WWIUniversity of Erlangen-NuernbergMartensstr. 591058 ErlangenGermany


Chapter 14

Stuart HoldsworthEMPA – Materials Science &

TechnologyÜberlandstrasse 129CH-8600 DübendorfSwitzerland


Contributor contact details xv

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Chapter 15

B.Wilshire* and H. BurtMaterials Research CentreSchool of EngineeringUniversity of Wales SwanseaSingleton ParkSwanseaSA2 8PPUK


Chapter 16

Christina Berger*,Alfred Scholz, F. Mueller,Michael SchwienheerInstitute for Materials TechnologyDarmstadt University of TechnologyGrafenstr 264283 DarmstadtGermany

Email: [email protected]; [email protected];[email protected]

Chapter 17

Horst Cerjak* and Peter MayrInstitute for Materials Science,

Welding and FormingGraz University of TechnologyKopernikusgasse 24A-8010Austria

Email: [email protected];[email protected];

Chapter 18

Masaaki TabuchiMaterials Reliability CentreNational Institute for Materials

Science (NIMS)1-2-1 SengenTsukuba 305-0047Japan


Chapter 19

P. J. Ennis and W. J. QuadakkersForschungszentrum Juelich GmbHIEF-2D 52425 JuelichGermany


Chapter 20

M. IgarashiCorporate Research and

Development LaboratoriesSumitomo Metal Industries Ltd.1-8 Fuso-choAmagasakiHyogo 660-0891Japan


Contributor contact detailsxvi

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Chapter 21

T.-U. KernSiemens AG Power Generation

GroupDept. Materials Rheinstr. 100D-45478 MuelheimGermany


Chapter 22

S. Sundaresan*L&T Visiting Welding ChairDept. of Metallurgical EngineeringFaculty of Technology and

EngineeringMaharaja Sayajirao UniversityKala BhavanBaroda-390001India

S.K. AlbertIndira Gandhi Centre for Atomic

ResearchKalpakkam 603102India

Email: [email protected]@igcar.gov.in

Chapter 23

R. (Vis) Viswanathan* and RichardTilley

Technical ExecutiveElectric Power Research Institute3420 Hillview AvePalo AltoCA 94304USA


Contributor contact details xvii

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Creep-resistant steel that can be used for a long time at elevated temperatureis the key to the construction of thermal and nuclear power generation plants,chemical plants and petroleum plants. During the last decade, great progresshas been made in developing creep-resistant steels of high strength andcorrosion resistance at ever increasing temperatures and in evaluating thesteels in terms of the weld characteristics, creep strength and corrosionresistance necessary for constructing plants. Although in the past the drivingforce for these developments has been primarily to achieve higher efficiencies,the focus has shifted more recently to the reduction of emissions of CO2,dioxins and other environmentally hazardous gases.

In the field of thermal power generation, the maximum allowabletemperature was about 565°C for conventional low alloy ferritic steels.However, progress in recent years has led to the development of high-strength9–12% Chromium ferritic steels capable of operating in ultra super critical(USC) power plants at metal temperatures approaching up to 650°C. Thecreep strength of austenitic creep-resistant steels has been enhanced to enableoperation up to temperatures of 675–700°C through the development of highCr, high nickel steels. In the field of nuclear power, creep-resistant steels,which are excellent both in high-temperature creep strength and in irradiationresistance, have been developed for cladding tubes for 650°C fast breederreactors. The temperature and pressure used were 454°C and 17 MPa,respectively in the early 1990s for hydrogen refining equipment in chemicalplants, when reaction chambers were made of 2.25Cr–1Mo steel, but thesubsequent development of high-strength 3Cr–1Mo–V steel and 2.25Cr–1Mo–V steel raised the limiting temperature and pressure to 482°C and 24MPa, respectively, by 1995. These figures are now about to reach 510°C and24 MPa. For power generation from wastes, the development of austeniticcreep-resistant steels that have high corrosion resistance enabled the boilersteam temperature to be raised from about 300°C in conventional plants upto about 500°C in more modern plants. In the automotive field, exhaustmanifolds used to be made of cast iron to withstand exhaust heat. However,as the exhaust gas temperature rose with improved engine performance,

Preface

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Prefacexx

higher strength was required and so 18Cr–2Mo–Nb and other steels weredeveloped, raising the exhaust gas temperature to 900°C or higher.

Recent research on enhancing the creep strength of 9–12Cr steels for650°C operation has revealed that the formation of even a partially weakmicrostructure near a grain boundary promotes local creep deformation andcauses premature fracture. This suggests the importance of taking into accountmicrostructural evolution phenomena during creep such as precipitation andcoarsening of carbonitrides and intermetallic compounds, dynamic recoveryand dynamic recrystallization, in the matrix as well as in the vicinity of grainboundaries.

Recently, some high-strength 9–12Cr steels have been found to sufferpremature loss of creep strength at 550°C or higher often after prolonged useup to relevant times. Therefore, efforts have been made to clarify themechanisms of creep strength loss, using modern transmission electronmicroscopy studies. Extrapolation of short duration laboratory data usingtime–temperature parameter (TTP) methods, such as the Larson–Millerparameter, have been used widely in the past to predict long-term life. However,it has now become clear that conventional TTP methods tend to overpredictthe long term strength because of microstructural degradation phenomena.To address this issue, new analysis techniques have been proposed taking themechanisms of creep deformation and creep rupture into account.

Welded structures made of ferritic creep-resistant steels used under hightemperature and low stress (about 600°C and 100 MPa or less) are subject topremature brittle creep fracture by the so-called type IV fracture in the fine-grained heat-affected zone (HAZ). Therefore, 9–12Cr steels are beinginvestigated to clarify the mechanisms and the means of preventing thisform of fracture. Operation of thick section components under thermallycyclic conditions further exacerbates the cracking problem by creep–fatigueinteraction. Thus, as plant temperatures are raised to improve energy efficiency,it is becoming increasingly important to establish the foundation of creep-resistant steels that can be used safely for a long time without showingdeterioration of creep strength and creep ductility.

The aim of this book is to consolidate and review the current state ofknowledge of creep resistant steels, summarizing the information which isnow scattered throughout voluminous scientific journals and a large numberof proceedings of international conferences. Each chapter of the book hasbeen written for engineers and researchers in particular by a world renownedexpert in the field. Therefore, the book contains not only background onmaterials but also recent progress from an engineering and technology pointof view. It also can be used as a reference source by graduate level students.It is hoped that the book will serve as an authoritative source of informationrelating to creep of steels.

This book consists of three parts: a general Part I on specifications and

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manufacture, Part II on the behaviour of creep-resistant steels and Part III onspecific applications. The introductory Part I includes the introductorydescription of creep and rupture (Chapter 1) and the historical developmentof creep-resistant steels (Chapter 2). Part I also includes the specifications ofcreep-resistant steels in Europe (Chapter 3) and in Japan (Chapter 4) and theproduction of creep-resistant steels for turbines (Chapter 5). Part II on thebehaviour of creep-resistant steels covers physical and elastic behaviour(Chapter 6), diffusion behaviour (Chapter 7), fundamental aspects of creepdeformation (Chapter 8), strengthening mechanisms (Chapter 9), precipitation(Chapter 10), grain boundaries (Chapter 11), fracture mechanisms and creepfracture (Chapter 12), mechanisms of creep deformation (Chapter 13),constitutive equations for creep curves and the prediction of service life(Chapter 14), creep strain analysis (Chapter 15), creep crack growth andcreep-fatigue behaviour (Chapter 16), creep strength of welded joints (Chapter17), fracture mechanics (Chapter 18), and oxidation and corrosion (Chapter19). Part III on specific applications includes the alloy design philosophybehind creep-resistant steels (Chapter 20), creep-resistant steels in turbines(Chapter 21), creep-resistant steels in nuclear reactors (Chapter 22), andindustry needs and future research trends in understanding creep damage(Chapter 23).

We are grateful to all the contributors for their willing participation andfor the cooperation they have extended to us in producing this book. We arealso grateful to Mr Robert Sitton, Mr Ian Borthwick, Mrs Lynsey Gathercoleand Ms Laura Bunney of Woodhead Publishing for their help in the publicationof this book.

F. AbeT.-U. Kern

R. Viswanathan

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Part I

General

1

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Creep-resistant steels2

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3

1Introduction

F . A B E, National Institute for Materials Science(NIMS), Japan

1.1 Definition of creep

Plastic deformation is irreversible and it consists of time-dependent andtime-independent components. In general, creep refers to the time-dependentcomponent of plastic deformation. This means that creep is a slow andcontinuous plastic deformation of materials over extended periods underload. Although creep can take place at all temperatures above absolute zeroKelvin, traditionally creep has been associated with time-dependent plasticdeformation at elevated temperatures, often higher than roughly 0.4Tm, whereTm is the absolute melting temperature, because diffusion can assist creep atelevated temperatures. For detailed description of mechanical equation ofstate, creep behavior of metals and alloys, dislocation motion during creep,mechanisms of creep, creep damage and fracture, the reader is referred tostandard text books on creep.1–6

1.2 Creep and creep rate curves

Creep tests can be conducted either at constant load or at constant stress. Forexperimental convenience, most frequently the creep tests of engineeringsteels are conducted at constant tensile load and at constant temperature. Thetest results can be plotted as creep curves, which represent graphically thetime dependence of strain measured over a reference or gauge length. Figure1.1 shows schematically three types of creep curves under constant tensileload and constant temperature conditions and also their creep rates ε = dε/dt, where ε is the strain and t the time, as a function of time. Textbooks oncreep of metals and alloys generally describe that three stages of creep,consisting of primary or transient, secondary or steady-state and tertiary oracceleration creep that appear after instantaneous strain ε0 upon loading asshown in Fig. 1.1(a), when the test temperature is high enough or at a highhomologous temperature. The homologous temperature is defined as theratio T/Tm, where T is the test temperature in absolute Kelvin and Tm the

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absolute melting temperature. The instantaneous strain ε0 contains elasticstrain and possibly plastic strain depending on the stress level.

In the primary creep stage between ε0 and ε1, the creep rate, ε , decreaseswith time, as shown in Fig. 1.1(d). The decreasing creep rate in the primarycreep stage has been attributed to strain hardening or to a decrease in free ormobile dislocations. In the secondary creep stage between ε1 and ε2, thecreep rate remains constant. This creep rate is designated as a steady-statecreep rate, ε s , which is given by ε s = (ε2 – ε1)/(t2 – t1) and is commonlyattributed to a state of balance between the rate of generation of dislocationscontributing to hardening and the rate of recovery contributing to softening.At high homologous temperatures, creep mainly involves diffusion and hencethe recovery rate is high enough to balance the strain hardening and resultsin the appearance of secondary or steady-state creep. In the tertiary creepstage, the creep rate increases with time until rupture at rupture time tr andrupture strain, εr. It should be remembered that under the constant tensileload, the stress continuously increases as creep proceeds or as cross-sectiondecreases and a pronounced effect of increase in stress on the creep rateappears in the tertiary creep stage.

Necking of the specimens before rupture causes a significant increase instress. The increase in creep rate with time in the tertiary creep stage canfollow from increasing stress or from microstructure evolution including

Time

log

(st

rain

rat

e o

r cr

eep

rat

e)

εmin.

εs.

t2t1 tr

trtm

(d)

(e)

(f)

Steady-statecreep rate

Minimumcreep rate

Time

Str

ain

ε0

ε0

εm

εr

ε0

ε1

ε2

εr

t1 t2 tr

trtm

(a) Three stages creep curve

(b) Two stages creep curve

(c) Logarithmic creep curve

1.1 (a), (b) and (c) Creep curves of engineering steels under constanttensile load and constant temperature and (d), (e) and (f) their creeprate curves as a function of time.

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Introduction 5

damage evolution taking place during creep. Microstructure evolution usuallyconsists of dynamic recovery, dynamic recrystallization, coarsening ofprecipitates and other phenomena, which cause softening and result in adecrease in resistance to creep. Damage evolution includes the developmentof creep voids and cracks, often along grain boundaries. The extent andshape of the three creep stages described above can vary markedly dependingon test conditions of stress and temperature, as shown schematically in Fig.1.2, where the final point in each curve represents creep rupture. With increasingstress and temperature, the time to rupture and the extent of secondary creepusually decrease but the total elongation increases.

Under certain conditions, the secondary or steady-state creep stage maybe absent, so that immediately after the primary creep stage the tertiary creepstage begins at tm, as shown in Fig. 1.1(b) and 1.1(e). In this case, theminimum creep rate, ε min, can be defined instead of the steady-stage creeprate, ε s. Similar to the steady-stage creep rate, ε s, the minimum creep rate,ε min, can be explained by the process where hardening in the primary stageis balanced by softening in the tertiary stage. In many cases, there is substantiallyno steady-state stage in engineering creep-resistant steels and alloys. Manyresearchers have shown that there is an ever-evolving microstructure duringcreep for engineering creep-resistant steels and alloys. This suggests thatthere is no dynamic microstructural equilibrium in engineering creep-resistantsteels and other alloys during creep, which characterizes steady-state creepof simple metals and alloys. Therefore, the term ‘minimum creep rate’ hasbeen favored by engineers and researchers who are concerned with engineeringcreep-resistant steels and alloys.

The stress dependence of minimum or steady-state creep rate is usuallyexpressed by a power law as:

Time

Str

ain

Arrows: increasing stress andtemperature

1.2 Schematic creep curves varying with stress and temperature.

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˙ ˙ε ε σmin s or = A n [1.1]

A = A′ exp (– Qc/RT) [1.2]

where n is the stress exponent, Qc the activation energy for creep, R the gasconstant and T the absolute temperature. The parameter A′ includesmicrostructure parameters such as grain size and so on. Equation [1.1] isoften referred to Norton’s law. It is well known that the minimum or steady-state creep rate is inversely proportional to the time to rupture tr as:

˙ ˙ε ε σmin s r c or = /( ) = exp (– / )C t A Q RTm n′ [1.3]

where C is a constant depending on total elongation during creep and m is aconstant often nearly equal to 1. Equation [1.3] is often referred to as theMonkman–Grant relationship, which has been experimentally confirmed notonly for simple metals and alloys but also for a number of engineering creep-resistant steels and alloys. Equation [1.3] suggests that the minimum orsteady-state creep rate and the time to rupture vary in a similar manner tostress and temperature.

At low homologous temperatures, with T/Tm often less than roughly 0.3,where diffusion is not important, only the primary stage appears. Usuallyonly limited strains well below 1% occur that do not lead to final rupture, asshown in Fig. 1.1(c) and 1.1(f). This deformation process is designated aslogarithmic creep.

Considerable efforts have been made to describe the creep curves, namely,the time dependence of creep strain. There are several model equationsavailable for characterizing the primary, secondary and tertiary creep stagecharacteristics, ranging in complexity from simple phenomenological tophysically based constitutive. Recent progress on the suitability of some ofthese to specific materials classes and analytical applications is reviewed byHoldsworth et al. [7].

Although Fig. 1.1 shows the idealized creep and creep rate curves,engineering creep-resistant steels sometimes exhibit complicated behavior,especially under low stress and long time conditions, reflecting complexmicrostructural evolution during creep. Complicated behavior is clearlydemonstrated by creep rate curves rather than creep curves. Figure 1.3 showsan example of complicated creep rate curves of 1Cr–0.5Mo steel at 550°C.8

At high stresses above 108 MPa, the creep rate curves are relatively simpleand consist of the primary and tertiary stages but there is no substantialsteady-state stage, similar to Fig. 1.1(e). The shape of creep rate curve becomesgradually complicated with decreasing stress. At low stresses below 88 MPa,two minima appear in the creep rate curves. This suggests that newstrengthening effects such as the precipitation of new phases seem to operateafter an extended period, causing a decrease in creep rate again after angrowing the previous acceleration creep. The subsequent loss of the existing

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Introduction 7

strengthening effects by microstructural evolution such as the coarsening ofnew phases causes an increase in creep rate again after reaching a secondminimum. Eventually the creep rate versus time curves exhibit oscillatedshapes under low stress and long time conditions, reflecting complexmicrostructural evolution during creep. Similar oscillated shapes havesometimes been observed in other low alloyed steels.

In fundamental investigations of creep, creep tests are often conducted atconstant stress. The applied stress does not change during the creep testprovided that the reduction in cross-sectional area is uniform along the wholegauge length. The stress can be kept constant during creep using properloading mechanisms. When we need to avoid any influence of oxidation,creep tests are usually conducted in vacuum or in an inert atmosphere. Otherwisethe influence of oxidation in reducing the cross-sectional area has to beconsidered, especially at higher temperatures and longer times for low alloyedsteels.

1.3 Creep rupture data

Elevated-temperature components used under creep conditions are designedusing allowable stress under creep conditions, which is usually determinedon the basis of 100 000 h creep rupture strength at the operating temperature,and sometimes also for 200 000–300 000 h creep rupture strength. The100 000 h creep rupture strength at a temperature T is defined as the stressat which creep rupture, the last point in Fig. 1.1(a) and (b), occurs at 100 000h. Generally creep rupture data are represented in graphic form showing the

Time/h10610510410310210110010–1

10–8

10–7

10–6

10–5

10–4

10–3

10–2

Cre

ep/h

–1

265MPa 216MPa108MPa

88MPa

74MPa

61MPa

53MPa

137MPa

1Cr–0.5Mo steel(JIS STBA22)

550°C

1.3 Creep rate versus time curves of 1Cr–0.5Mo steel at 550°C (823K).

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relationship between the stress σ and the time to rupture tr. Figure 1.4 showsan example of creep rupture data for 1Cr–0.5Mo steel from the NationalInstitute for Materials Science (NIMS) Creep Data Sheet.9 This figure contains309 data points for 11 heats. The material specification defines the chemicalcomposition, heat treatment conditions and so on. In terms of the chemicalcomposition of 1Cr–0.5Mo steel (JIS STBA 22), the Cr concentration isspecified as the range 0.80–1.25%, the concentration of Mo in the range0.45–0.65%, and so on. Practically, the melting of steels causes a differencein the concentration of alloying elements, so that No. 1 ingot contains 1.0%Crand 0.50%Mo, but No. 2 ingot contains 0.90%Cr and 0.60%Mo and so on,in which the two ingots satisfy the materials specification of 1Cr–0.5Mosteel (JIS STBA 22). Usually, such a small variation in chemical compositioncauses a difference in creep strength.

The 100 000 h creep rupture strength is evaluated to be, for example, 61MPa at 550°C. The creep rate curves shown in Fig. 1.3 were obtained for oneheat of the 1Cr–0.5Mo steel shown in Fig. 1.4. The creep rupture data in Fig.1.4 exhibit rather complicated curves showing inverse sigmoidal bending atintermediate stress levels of about 130 MPa. It should be noted that twominima appear in the creep rate curves at intermediate stress levels andbelow, while only one minimum appears at higher stress levels, Fig. 1.3.

500°C550°C600°C650°C

Time to rupture (h)10610510410310210

Str

ess

(MP

a)

500

400

300

200

100

80

60

50

40

30

20

n = 309 650°C 600°C

550°C

500°C

1.4 Creep rupture data for 1Cr–0.5Mo steel at 500–650°C.

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Introduction 9

Recently, long-term creep rupture test data and creep strain data beyond 100000 h have become available for a number of creep-resistant steels in severalmaterials test institutions in the world, for example, in NIMS, Japan. Forlong-term creep and creep rupture data, the reader is referred to the NIMSCreep Data Sheets, for example.10

NIMS Creep Data Sheets contain a full set of data, such as creep rupturedata, often exceeding 100 000 h, minimum creep rates, short-time tensiledata, evaluation of short-time tensile strength and long-term creep rupturestrength by curvilinear regression analysis and optical micrographs, togetherwith the details of materials production procedures and chemical compositions.Microstructure Data Sheets, have also been published as the MetallographicAtlas of Long-Term Crept Materials,11 another series of NRIM Creep DataSheets. The Metallographic Atlas not only contains series micrographs thatshow microstructural evolution during creep for up to 100 000 h, but thatalso show related data such as time–temperature–precipitation (TTP) diagrams,histograms describing the distributions of precipitates and creep-voids, andcreep damage parameters, using specimens in the Creep Data Sheets.Furthermore, the Atlas of Creep Deformation Property12 was published asCreep Strain Data Sheets for Grade 91 steel (9Cr–1Mo–V–Nb), providingcreep curves, creep rate curves and related data.

As can be seen from Equation [1.3], stress and temperature are importantvariables that influence creep rate and time to rupture. In addition, creep andcreep rupture properties are markedly affected by not only microstructurevariables but also by external variables. The external variables include pre-straining (cold-working), additional heat treatments, oxidation and corrosion,stress mode such as uniaxial or multiaxial loading, and superimposition ofcyclic loading (creep–fatigue mode). High-temperature structure componentsin plants are usually used under the complicated conditions described aboveover long duration up to 300 000 h or longer.

1.4 Deformation mechanism map

Ashby13 proposed the concept of a deformation mechanism map, based onthe assumption that all six deformation mechanisms concerned are mutuallyindependent and operate in a parallel way. The six deformation mechanismsinclude (1) defect-less flow, (2) glide motion of dislocations, (3) dislocationcreep, (4) volume diffusion flow, (5) grain boundary diffusion flow and (6)twinning. The twinning can supply only a limited amount of deformationand usually does not appear in the deformation mechanism map. It should benoted that Ashby considered steady-state flow only but no fracture. As illustratedschematically in Fig. 1.5, the deformation mechanism map is constructedwith axes of normalized stress σ/G, where G is the shear modulus and T /Tm

is the homologous temperature. The map is divided into fields. Within a

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field, one mechanism is dominant, that is, it supplies a greater strain ratethan any other mechanisms. The upper limit of the boundary is set by atheoretical or ideal strength of roughly G/20 to G/30. At stresses lower thanthe ideal strength, the deformation takes place by dislocation glide, as inshort-time tensile tests. At stresses lower than yield stress, dislocation creepcan take place with the aid of diffusion: probably dislocation core diffusionat low homologous temperatures and volume diffusion at high homologoustemperatures.

Sometimes the dislocation creep field is further divided into two fields:low- and high-temperature dislocation creep fields. At further low stresses,volume diffusion creep (Nabarro–Herring creep) and grain boundary diffusioncreep (Coble creep) dominate. The boundaries between adjacent fields in thecreep region indicate the conditions under which two mechanisms contributeequally to the overall creep rate. Using an appropriate constitutive equationfor creep rates as functions of stress and temperature, we can calculate thecreep rates and can draw the boundaries. This also allows us to plot the contoursof constant creep rate onto the map, as shown schematically in Fig. 1.5.

The locations of the boundaries between adjacent creep fields differ fordifferent materials and also depend on microstructure valuables such asgrain size. Experimentally, the deformation mechanism map can be constructedby the measurements of stress and temperature dependence of strain rates orcreep rates caused by the individual mechanisms. It should be also noted thatthe time dependence is not included in the deformation mechanism map. As

Ideal strength

Homologous temperature, T/Tm

Defect-free flow

Dislocation glide

1.00.80.60.40.20

No

rmal

ized

str

ess,

σ/G

10–7

10–5

10–3

10–1

Dislocationcreep

10–2

10–4

10–6

10–8Creep rate

10–10/S

Nabarrocreep

Coble creep

1.5 Schematic deformation mechanism map with contours ofconstant creep rate.

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Introduction 11

already shown in Fig. 1.2, the creep rate of engineering creep-resistant steelsvaries in complex manner with time because of complicated microstructureevolution during creep exposure at elevated temperatures. Therefore, in thecase of engineering creep-resistant steels, the deformation mechanism mapcan be applied to predict a dominant deformation mechanism at the beginningof creep under specific stress and temperature conditions.

1.5 Fracture mechanism map

Ashby14 also proposed the concept of fracture mechanism map for facecentred cubic (fcc) metals and alloys with axes of normalized stress σ/G andhomologous temperature T/Tm, which provides us with information aboutthe dominant mechanism resulting in fracture in a shorter time than anyother mechanisms. The fracture mechanism map is more important than thedeformation mechanism map in practice, because the former relates to damageand fracture processes, which provide us with useful guidelines for assessmentof damage evaluation and the remaining life estimation of components inplants. Because the minimum or steady-state creep rate and the time torupture vary in a similar manner stress and temperature, as suggested by Eqn[1.3], approaches similar to those employed in the construction of a deformationmechanism map can be adopted for the construction of a fracture mechanismmap.

Figure 1.6 shows schematically the fracture mechanism map for fcc metals,where a cleavage fracture field does not appear. The ideal strength appears

Homologous temperature, T/Tm

1.00.80.60.40.2010–7

10–5

10–3

10–1

No

rmal

ized

str

ess,

σ/G

Ideal strengthDuctile fracture

Dynamic fracture

Transgranularcreep fracture

Intergranularcreep fracture

105 h rupturestrength

Rupture

104 h 103 h

1.6 Schematic fracture mechanism map with contours of constanttime to rupture for fcc metals.

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as the upper limiting fracture strength which will overcome interatomicforces in defect-free materials. At stresses lower than the ideal strength,fracture takes place in a ductile, transgranular way, designated ductile fracture,and often designated ductile transgranular fracture. In the creep regime, twofields of transgranular creep fracture and intergranular creep fracture appearat high and low stresses, respectively. At high temperature and relativelyhigh strain rate, dynamic recrystallization can allow materials to deformextensively so that deformation becomes localized in a neck and failureeventually occurs by the specimen necking until the cross-sectional area hasgone to zero, usually called the field of rupture. Because grain boundariesbecome highly mobile under conditions of dynamic recrystallization, thedevelopment of creep voids and cavities is suppressed.

Figure 1.7 shows schematically the three fracture mechanisms in creepregime: intergranular creep fracture, transgranular creep fracture and rupture.14

Contours of constant time to rupture can be also plotted onto the map, asshown schematically in Fig. 1.6. Although the axes of most of the fracturemechanism maps are stress and temperature, axes of stress and time are alsoused.

Figure 1.8(a) and 1.8(b) show examples of fracture mechanism maps for1Cr–1Mo–0.25V steel for a turbine rotor, plotted for stress–time to ruptureand for stress–temperature coordinate systems, respectively.15 In these figures,the stress–time to rupture plots and stress–temperature plots for constanttimes to rupture of 100–100 000 h are superimposed. It should be noted thatthe axes of stress and temperature but not those of normalized stress σ/G andhomologous temperature T/Tm are used in these figures because the objectiveof constructing fracture mechanism maps is primarily for use in assessmentof reliability, such as in the design and remaining life prediction of a steamturbine rotor. The intergranular creep fracture field is located in a long time-

Rupture due todynamic recoveryor recrystallization

(c)

Intergranular creepfracture

(voids) (wedge cracks)

(a)

Growth of voids bypower-law creep

(transgranular) (intergranular)

(b)

1.7 Schematic drawing of three fracture mechanisms in a high-temperature creep regime.

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Introduction 13

to-rupture region at 500–575°C. The rupture field appears at temperatureshigher than 600°C. The region of practical importance for 1Cr–1Mo–0.25Vsteel turbine rotor in power plants is the low stress and long time-to-ruptureregion at temperatures of 550°C or lower, which belongs to the intergranularcreep fracture field. This suggests that precise measurements of the developmentof creep voids at grain boundaries during creep contributes to the improvementin the reliability of the remaining life estimation.

1.8 Fracture mechanism maps for 1Cr–1Mo–0.25V steel, as functionsof time to rupture and of temperature.

Temperature (°C)650600550500450

600

400

200

100

40

Str

ess

(MP

a)

Tensile strength

Rupture(recrystallization)

10 5 h

10 4 h

10 3 h

10 2 h

Transgranularcreep rupture

Intergranular creepfracture (cavitation)

Ductilityminimum

Rupture

strength

Time to rupture (h)105104103102

600

400

200

100

40

Str

ess

(MP

a)

Rupture(recrystallization)

Transgranularcreep fracture

Ductilityminimum

Intergranularcreep fracture

(cavitation)

450°C500°C

625°C

675°C

650°C

600°C

575°C

550°C525°C

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1.6 References

1. Finnie I. and Heller W. R., Creep of Engineering Materials, McGraw-Hill, NewYork, 1959.

2. Garofalo F., Fundamentals of Creep and Creep-Rupture in Metals, The MacmillanCompany, New York, 1965.

3. Penny R. K. and Marriott D. L., Design for Creep, McGraw-Hill, London, 1971.4. Evans R. W. and Wilshire B., Creep of Metals and Alloys, The Institute of Metals,

London, 1985.5. Cadek J., Creep in Metallic Materials, Elsevier, Amsterdam, 1988.6. Viswanathan R., Damage Mechanisms and Life Assessment of High-Temperature

Components, ASM International, Ohio, 1989.7. Holdsworth S. R., Baker A., Gariboldi E., Holmstrom S., Klenk A., Merckling G.,

Sandstrom R., Schwienheer M. and Spigarelli S., ‘Factors influencing creep modelequation selection’, Proceedings of ECCC Creep Conference, 12–14 September2005, The Institute of Materials, London, UK, 2005, 380–393.

8. Kushima H., Kimura K., Abe F., Yagi K., Irie H., Maruyama K., ‘Effect ofmicrostructural change on creep deformation behaviour and long-term creep strengthof 1Cr–0.5Mo Steel’, Tetsu-to-Hagane, 2000, 86, 131–137.

9. NIMS (formerly NRIM) Creep Data Sheets No.1. Tokyo, Tsukuba, National Institutefor Materials Science, 1996.

10. Series of NIMS (formerly NRIM) Creep Data Sheets No. 1–48. Tokyo, Tsukuba,National Institute for Materials Science, 2007.

11. Series of NIMS Metallographic Atlas of Long-Term Crept Materials No. M1-M6.Tokyo, Tsukuba, National Institute for Materials Science, 2007.

12. NIMS Atlas of Creep Deformation Property, No. D-1. Tokyo, Tsukuba, NationalInstitute for Materials Science, 2007.

13. Ashby M. F., ‘A first report on deformation-mechanism maps’, Acta Metallurgica,1972, 20, 887–897.

14. Ashby M. F., Gandhi C. and Taplin D. M. R., ‘Fracture-mechanism maps and theirConstruction for FCC Metals and Alloys’, Acta Metallurgica, 1979, 27, 699–729.

15. Shinya N., Kyono J. and Kushima H., ‘Creep fracture mechanism map and creepdamage of Cr–Mo-V turbine rotor steel’, ISIJ International, 2006, 46, 1516–1522.

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15

2The development of creep-resistant steels

K.-H . M AY E R , ALSTOM Energie GmbH, Germanyand F . M A S U YA M A , Kyushu Institute of

Technology, Japan

2.1 Introduction

The development of creep-resistant steels is a result of continuous technologicalprogress throughout the 20th century. The urgent need to improve the creepstrength of steels was based on endeavours by the power station industry toimprove the thermal efficiency of steam power plant by raising the steamtemperature and steam pressure in order to reduce the cost of fuel and reduceuse of fuel resources. Since roughly 1900, as shown for instance by Fig. 2.1,the heat rate of thermal power plant in Germany has been reduced followinga step-by-step increase in the steam parameters from 275°C/12 bar to 620°C/300 bar.1,2

A major contribution to the increase in power plant efficiency consistedof the development of heat-resistant steels with a higher creep strength at anacceptable creep ductility level (see for example Kallen).3 The significanceof these material properties was not recognised until early damage wassuffered by steam turbine bolts in the 1930s, which pointed to the fact thatthe strength of steels used in power stations operating at higher temperaturesdepends significantly on the creep behaviour of the material over the fullperiod of operation.4 Based on this experience it was concluded that thestrength values should no longer be determined in short-term tests,5,6 forexample the ‘durability strength’ according to the DVM (Deutscher Verbandfur Material prüfung) creep rate limit test. The procedure to be adoptedshould be to determine the fracture strength, the creep elongation and creepductility of the heat-resistant steel in a creep test extending over a period ofroughly 100 000 h (see for example Siebel).7 For the DVM creep rate limittest established in Germany in 1930, the ‘durability strength’ was defined tobe the stress at the test temperature at which a creep rate of 10 × 10–4 %/hwas reached between the 25th and 35th hour.5 Typical results of both tests,which were commenced at the end of the 1930s, are shown by Fig. 2.2.The tests were performed at 500°C on a steel containing 0.30%C–1.61%Cr–1.28%Mo–0.10%V. The DVM creep rate limit test using smooth specimens

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and the creep rupture test using smooth and notched specimens (notch factor4.3) are compared.

Creep rupture tests were even continued up to roughly 300 000 h at theend of the 1970s.8 The ‘durability strength’ in the short-term test was determinedat a strength level of 306 MPa. At this stress level, the rupture of the creeprupture tests was reached after about 3000 h, whereas the 100 000 h rupturestrength of the smooth specimens lies at 190 MPa. The notched specimensstressed at the same level of 190 MPa failed after roughly 30 000 h, distinctlyearlier than the smooth specimens owing to a significant notch-weakening

2.1 Heat rate of steam power plants in Germany as a function ofsteam parameters since the year 1900.

40

30

20

10

0

Sp

ecif

ic h

eat

rate

(kJ

/kW

h) 12 bar/275°C

15 bar/350°C

35 bar/450°C

100 bar/500°C100 bar/540°C

280 bar580/600°C

300 bar580/600°C

With reheat (540/540°C)Supercritical (250 bar)

1900 1910 1920 1930 1940 1950 1960 1970 1980 1990 2000 2010Year

2.2 Creep rupture strength as a function of time to rupture and‘durability strength’ of a 1.6%CrMoV steel at 500°C. Test steel:0.30%C–1.6%Cr–1.3%Mo–0.1%V; heat treatment: 950°C/air + 680°C/air;tensile strength: 893 MPa.

500°C

Smooth specimens

313192 h

Notched specimens

294000 h

‘Durability strength’ of DVMcreep rate limit test carried

out at about 1936

10–1 100 101 102 103 104 105 106

Time to rupture (h)

Cre

ep r

up

ture

str

eng

th (

MP

a)

1000

600

400

200

100

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The development of creep-resistant steels 17

behaviour. It was also recognised that the tendency to notch-weakeningbehaviour of the steels was a wrong turn in the development of heat resistantsteels on the basis of the DVM creep rate limit test, because the aim ofraising the ‘durability strength’ of the heat-resistant steels as high as possibleinvolves the risk of increasing the susceptibility of the steels to embrittlement.

To provide a further example of the influence of creep processes on thestrength of heat resistant steels, Fig. 2.3 demonstrates the dependence of thecreep rupture strength and the strength for 1% creep strain on the testtemperature and test period for a carbon steel and a 1%Cr–0.5%Mo steel incomparison with the 0.2% yield limit determined in the short-term tensiletest (see for example Wellinger).9 In comparison with the 0.2% yield limitdetermined in the short-term tensile test, the 100 000 h creep rupture strengthis lower for the carbon steel at higher than about 410°C and is also lowerfor the 1%Cr–0.5%Mo steel higher than about 480°C. The crossovertemperatures between the results of the short-term tensile test and the creepstrength values are distinctly lower if the 0.2% or the 1% permanent creepstrain determined in the 100000 h test are decisive for the design of powerstation components.

Forming influences marking the development of heat resistant steels overthe past 100 years are:

• long-term operational experience• experience gained from long-term creep rupture tests• improvements in melting technology• systematic investigations into the influence of heat treatment on creep

behaviour

2.3 0.2-limit, 100000 h creep rupture strength and 100 000 h 1%-creep strength of a carbon steel and 1%Cr–0.5%Mo steel as afunction of test temperature.

300

200

100Str

eng

th (

MP

a)

0.2-Limit1% Cr 0.5%

Mo steel

100 000 hcreep rupture

100 000 h1%-creep

strainC steel

200 300 400 500 600 700Test temperature (°C)

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• examination of the microstructure of specimens in the virgin conditionfollowing long-term thermal and creep loading

• systematic investigations into the influence of alloying elements• computer-aided alloy design methods (e.g. Thermocalc, DICTRA)• modelling of creep processes• development of metallographic methods and equipment for the

identification of precipitates (e.g. transmission electron microscopy (TEM),energy dispersive X-ray spectrometry (EDS), energy filtered transmissionelectron microscopy (EFTEM), atom probe field ion microscopy (APFIM),field emission Auger electron spectroscopy (FE-AES) and secondaryion mass spectroscopy (SIMS).

• national and international joint research activities and research projectsrelated to the development of advanced creep resistant steels and long-term tests under creep stress conditions7–16

• testing of newly developed heat-resistant steels on the basis of largepilot components and welds fabricated under normal workshopconditions

• investigations into the oxidation behaviour of advanced heat-resistantsteels in the laboratory and in test fields of steam power stations

• international exchange of experience at conferences and in workshopse.g. EPRI (Electric Power Research Institute USA), EPDC (ElectricPower Development Center/Japan), COST (Community of Science andTechnology of the European Communities), ECCC (European CreepCollaboration Committee), NIMS (National Institute for Materials Science/Japan).

2.2 Requirements for heat-resistant steels

Heat-resistant steels for use in thermal power stations must be capable ofsatisfying the specific requirements established for dependable and economicoperation. All phases of development and testing must therefore be specificallyaligned to the following requirements:

• high thermal efficiency• operational capability in the medium and peak load ranges• life expectancy of at least 200 000 h• high availability• long intervals between overhauls• short overhaul periods• short manufacturing times• competitive production costs for the steam plant and electric power.

These requirements mean that the application of newly developed steelsmust not involve any additional risks, implying:

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• that long time creep testing up to 100000 h is needed to predict reliablythe creep strength for 200 000 h (which means that the tests must bestarted with a large number of specimens, because at the outset of thetests a prediction cannot be made about the stress level which will bereached after a test period of 100000 h. A long test period should also bescheduled if a research project is only due to last for 3–5 years);

• satisfactory oxidation resistance;• high ductility of the steels under conditions of creep stressing;• high fracture toughness of the steels in a new condition and following

prolonged operational stressing;• satisfactory production of the new steels in terms of melting, casting,

forging, hot forming and welding.

2.3 Historical development of ferritic steels

2.3.1 Carbon steels

Up to the 1920s it was general practice to use non-alloyed steels for componentsin the steam admission zone exposed to maximum temperatures of 350°Cand pressures of about 15 bar. The components were designed according tothe material requirements established in a hot tensile test. In these short-termtests it was not possible to recognise that the elements N, Al and Mn exerciseda major influence on the creep strength of carbon steels. Figure 2.3 hasalready shown the 0.2-limit and the creep rupture strength obtainable withpresent-day standard non-alloyed steels as a function of the test temperaturein comparison with the 1%Cr–1%Mo steel.9

2.3.2 Low alloy steels

At the beginning of the 1920s, operation at steam temperatures of 450°C andpressures of 35 bar called for the development of low-alloyed heat-resistantsteels. Developments were limited to individual steel works which at thattime were not yet coordinated in joint research programmes. The steels wereidentified by the trade name of the steel works. The basic test in the developmentof low-alloyed steels was a hot tensile test which later on was followed bya short-term test, for example the DVM creep rate limit test in Germany.5 Inthe USA in 1933, a guideline was prepared between the ASME and ASTMfor tests covering periods of 500–2000 h to determine the creep strain limitsfor a permanent creep strain of 0.01%, 0.1%, 1% and the ultimate rupturelimit.6 The results were extrapolated at a double logarithmic scale on astraight line pattern up to 104 h and further up to 105 h.

Based on the multiplicity of investigations of test steels carried out withdifferent Mo, Cr, Ni, V, CrMo, CrV, MnSi, MoMnSi, CrSiMo, CrNiMo,

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CrMnV, CrMoV contents, worldwide developments in the manufacture ofsteam boilers and small forgings for steam turbines produced steels withchemical compositions of 0.15%C–0.3–0.5%Mo, 0.13%C–1%Cr–0.5%Mo17

and 0.10%C–2.25%Cr–1%Mo20 which are still in use today. In addition, inthe 1950s, a MoV steel with a composition of 0.14%C–0.5%Mo–0.3%Vwith an even higher creep strength was developed in Europe for gas turbinesand later also qualified in long-term creep tests for steam plants. In the fieldof turbine manufacturing since the 1950s a steel with a composition ofapproximately 0.25%C–1.25Cr–1%Mo–0.30%V is in use worldwide for turbinerotors, casings, bolts and small forgings.

Systematic investigations into the creep strength of the steels developedin short-term tests between the 1920s and 1940s were followed in the 1950sby long-term creep tests.10–16 In Germany, for instance, a joint research projectwas established for this purpose in 1949 between steel and power plantmanufacturers and plant operators.10 Long-term creep tests of individualmelts have actually been performed by research bodies in Germany since themid-1930s (see for example Diehl and Granacher).8 The activities of theindividual national creep groups operating within Europe were coordinatedin 1990 and culminated in the establishment of the European CreepCollaborative Committee (ECCC) in December 1991.11

Molybdenum was recognised as an important element for increasing hightemperature strength. Mo steels developed in the USA and the UK are alloyedwith a Mo-content of about 0.5%. The Mo-content of the steel developed inGermany is roughly 0.3% at a C-content of about 0.15%.

Figure 2.4 illustrates the influence of molybdenum on the 100 000 h creeprupture strength at 450°C as opposed to an unalloyed steel with roughly0.15%C.18 By the addition of approximately 0.5%Mo, the 100 000 h creepstrength of the unalloyed steel of roughly 70 MPa is increased to about260 MPa. The alloying effect of Mo is the result of solution hardening andMo2C precipitation.9,18

A drawback of Mo alloying to over about 0.35% is amarked decline in ductility under creep stress conditions as well as graphiteprecipitation. Consequently, steels with an Mo-content of 0.5% should notbe used in temperature environments over 400°C. However, the strength-increasing effect of higher Mo contents, without an unacceptable decrease inductility, can be utilised by the addition of Cr as in the case of the steels witha composition of 0.13%C–1%Cr–0.5%Mo and 0.10%C–2.25%Cr–1%Mo.

Figure 2.5 shows the influence of Mo and Cr on the 100 000 h creeprupture strength of the three steels 0.3%Mo, 1%Cr–0.5%Mo and 2.25%Cr–1%Mo at 500°C and 550°C.18 The highest creep rupture strength is alreadyachieved by the 0.13%C–1%Cr–0.5%Mo steel at 500°C. At 550°C, subjectedto an increase in the Mo and Cr-contents as in the case of the 0.10%C–2.25%–Cr1%Mo steel, a further increase in the creep rupture strength isobtained. Microstructure investigations on the initial condition of the 0.13%C–

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1%Cr–0.5%Mo steel revealed M3C, M7C3 and M23C6 precipitations whereasfor the 0.10%C–2.25%Cr–1%Mo steel Mo2C and M23C6 precipitations werefound (see for example Florin).19 An excellent literature survey about themicrostructure formation of the CrMo steels after heat treatment and longterm creep is given by Orr et al.20

The 0.14%C–0.6%Mo–0.3%V steel, which in view of its higher creeprupture strength (Fig. 2.8) is given preference for live steam pipes and pipeswith superheated steam, features higher strength than the 0.10%C–2.25%Cr–1%Mo steel owing to finely distributed and thermally very stable V4C3

precipitation and Mo2C.19,21 A drawback for this steel is its tendency to typeIV cracking in the intercritical area of the heat affected zone of welds (seefor example Schüller et al).22

Amongst the numerous steel versions developed in the 1930s and 1940sfor the manufacture of rotors, casings, valves and bolts for steam turbines, a1%CrMoV steel has found worldwide acceptance which, depending oncomponent size and the location of the site of development, is alloyed witha composition of roughly 0.20–0.30% C, 1–1.5% Cr, 0.70–1.25% Mo, 0.25–0.35% V and 0.50–0.75% Ni.15,23,24

Figure 2.6 shows schematically the relationship of the 100 000 h creeprupture strength and the fracture toughness FATT50 as a function of themicrostructure for the 1%CrMoV steel.26 The highest creep rupture strengthof this steel type is achieved with an upper bainite structure.25 Thedisadvantage of the upper bainite structure is the lower toughness25 so that

2.4 100000 h creep rupture strength of a C-steel as a function of Mocontent at 450°C.

100

000

h c

reep

ru

ptu

re s

tren

gth

(M

Pa) 300

200

100

C-steel

450°C

0.1 0.2 0.3 0.4 0.5 0.6Mo content (mass%)

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the individual alloying elements of the steel as well as heat treatment mustbe aligned to the specific operational properties of the components.26 Insome cases the procedure for turbine rotors is to adapt the heat treatmentcontour to the operational properties and/or to apply a method of sprayhardening with different quenching rates for the specific regions of thecomponents.26 Investigations into the microstructure in the initial state revealedV4C3, Mo2C and M23C6 (see for example Smith).27 With regard to theductility and toughness of the 1%CrMoV steel, experience of operationalstressed components has emphasised the significance of the austenitisingand tempering temperature.28–32

Figure 2.7 illustrates the behaviour of smooth and notched specimens at

100

000

h c

reep

ru

ptu

re s

tren

gth

(M

Pa)

200

160

120

80

40

0.2 0.4 0.6 0.8 1.0Mo content (mass%)

Steel1%Cr–0.5%Mo

Steel2.25%Cr–1%Mo

500°C

500°CSteel0.3Mo

100

000h

cre

ep r

up

ture

str

eng

th (

MP

a)

200

160

120

80

40

0.5 1.0 1.5 2.0Mo content (mass%)

Steel1%Cr–0.5%Mo

Steel2.25%Cr–1%Mo

500°C

500°CSteel

0.3%Mo

2.5 100000 h creep rupture strength of low alloyed steels as afunction of Mo and Cr content at 500°C and 550°C.

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500°C in a creep rupture test for two different heat treatment temperatures oftwo 1%CrMoV melts over test periods up to about 120000 h.28 Austenitisingat 1050°C in connection with a tempering temperature of 700°C was foundto cause substantial notch weakening and very low ductility of the smoothspecimens (case 17c). An acceptable deformation behaviour is obtained withan austenitising temperature of 980°C and tempering treatment at 670°C. Afurther disadvantage of heat treatment at an excessive austenitising temperatureis the initiation of long-term embrittlement which results in a remarkablereduction of toughness at low temperatures. This loss of toughness has been

Tem

per

atu

re (

°C)

1200

800

400

TTT-diagram 1%CrMoV steel

F Perlite

B

M

102 104 106

Time (s)Martensite (M)

LowerBainite

(B)

Upper Ferrite (F)

FATT

FAT

T

Lon

g t

erm

cre

ep s

tren

gth

Creep strength

2.6 Creep rupture strength and toughness FATT of 1%CrMoV steel asa function of cooling rate after austenisation respectively formartensite, bainite and ferrite microstructures (schematic). TTT is thetime temperature transformation.

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24

2.7 Creep rupture strength of smooth and notched specimens of 1%CrMoV steels as a function of heat treatment andtime to fracture at 500°C.

17a: 0.19%C–1.32%Cr–1.05%Mo–0.54%V980°C/oil + 2h 670°C/air

17c: 0.17%C–1.10%Cr–1.16%Mo–0.35%V1050°C/oil + 3h 700°C/air

Smooth specimen

Notched specimen

500°C

Smooth specimen

Notched specimen

500°C

Cre

ep r

up

ture

str

eng

th (

MP

a)

500

400

300

200

100102 103 104 105 102 103 104 105

102 103 104 105

Time to fracture (h)102 103 104 105

Time to fracture (h)

Red

uct

ion

of

area

(%

) 100

80

60

40

20

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the cause of brittle failure of turbine and valve bolts in the past.32,33 Thelong-term embrittlement also increases the risk of brittle failure of HP (highpressure) and IP (intermediate pressure) turbine rotors.34,35 In one investigatedcase, a FATT50 of 340°C was found after long-term service of a turbine rotorwith a component temperature of about 380°C.36 However, in this connection,attention must be drawn to the fact that in accordance with technologicalprogress in the early 1950s, the trace elements, owing to the melting process,were still at a relatively high level (e.g. phosphorus up to 0.028%). Thesehigh contents of trace elements also contributed to the embrittlement.36 Goodlong-term experience was gained in Germany with components of 0.20%C–1%Cr–1%Mo–0.3%V steels in the mid-1950s when the austenitisingtemperature was limited to a maximum of 950°C and tempering treatmentwas fixed at 680–740°C. The maximum permissible tensile strength wasspecified at 835 MPa.

Figure 2.8 provides an overview of the 100 000 h creep rupture strengthas a function of the test temperature for the non-alloyed and low-alloyedheat-resistant steels currently established for the temperature range belowabout 565°C. Two new low-alloyed heat-resistant steels have been developedover the past 15–20 years predominantly for the manufacture of water wallsfor advanced steam power stations. The steels are named HCM2S (0.06%C–2.25%Cr–2%Mo–1.6%W–0.25%V–0.05%Nb–0.02%N–0.003%B)37 and7CrMoVTiB (0.07%C–2.4%Cr–1.0%Mo–0.25%V–0.07%Ti–0.01%N–0.004%B).38 Both steels lend themselves well to welding and do not requirepost-weld heat treatment. Their creep rupture strengths in comparison with

105

h c

reep

ru

ptu

re s

tren

gth

(M

Pa) 200

100

0

(b) 0.3%Mo

(a) C-steel

(f) 1%CrMoV*

*Upper bainite

(e) 0.6%MoV

(d) 2.25%CrMo

(c) 1%CrMo

450 500 550 600Temperature (°C)

Steel C Cr Mo V (mass%)(a) 0.18 – – –(b) 0.15 – 0.3 –(c) 0.13 1.0 0.5 –(d) 0.10 2.25 1.0 –(e) 0.14 0.5 0.6 0.3(f) 0.28 1.0 1.0 0.3

2.8 100000 h creep rupture strength of a C-steel and low alloyedheat-resistant steels as a function of test temperature.

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the conventional used 0.15%C–0.5%Mo and 0.13%C–1%Cr–0.5%Mo steelsare given in Fig. 2.9.

2.3.3 9–12%Cr steels

The development of heat-resistant 9–12%Cr steels was strongly motivatedby two major events. During the 1950s it was the development of thermalpower stations for public power supply, operating at steam temperaturesranging from 538°C to 566°C and during the 1980s the target was set todevelop low-pollution power stations operating at steam admission temperaturesof 600–650°C and supercritical pressures up to 350 bar. Figure 2.10 presentsa summary of current national and international research projects in progresssince the 1980s in Japan, USA and Europe.

An overview of the historical development of heat-resistant ferritic–martensitic 9–12%Cr steels from the 1950s to the 1990s is given in the upperpart of Fig. 2.11. The lower part shows recent values of the 100 000 h creeprupture strength at 600°C, extrapolated from long-term test data. Table 2.1illustrates the chemical composition of the steels. As a rule, the steels are anonward development of steels already applied over extended periods of timeby using the trial-and-error method.

Development of 9–12% Cr steels for steam temperatures up to 620°C

The steel X22CrMoV 12 1 was developed in the 1950s for thin-walled andthick-walled power station components. Its creep strength is based on solution

105

h c

reep

ru

ptu

re s

tren

gth

(M

Pa) 200

100

0

(a) 0.3%Mo

(c) HCM 2 S

(d) 7 CrMoVTiB 10 10

(b) 1%CrMo

450 500 550 600Temperature (°C)

Steel C Cr Mo W V Nb Ti N B (mass%)(a) 0.16 – 0.3 – – – – – –(b) 0.13 1.0 0.5 – – – – – –(c) 0.06 2.25 0.2 1.6 0.25 0.05 – 0.02 0.003(d) 0.07 2.4 1.0 – 0.25 – 0.07 0.01 0.004

2.9 100000 h creep rupture strength of heat-resistant low alloyedsteels used for water walls of steam plant boilers as a function oftest temperature.

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hardening and on the precipitation of M23C6 carbides. The steel has beenapplied successfully in power stations over several decades.

The steels H46, FV448 and 56T5 (nos. 2 and 3 in Fig. 2.11) exhibitadditional alloying of 0.30–0.45% Nb and roughly 0.05 N. The targetedincrease in strength is obtained by secondary MX precipitations of the typeVN and Nb (C,N). However, a distinct improvement in creep strength at600°C, which is of primary interest for components for the aviation industry,is only obtainable in the short-term range. In view of the high Nb-content,these steel grades are only suitable for the manufacture of small-sizecomponents because the relatively high Nb-content results in pronouncedsegregations in ingots used in manufacturing thick-walled components.

TAF steel (no. 4) developed in Japan by Fujita39 for small components isan onward development of European Nb-containing steels (no. 2: H46 andFV 448). In addition to an improved balance of the alloying elements–basedon a very extensive investigation of the influence of all alloying elements onthe creep strength – it also features a high boron contents up to 0.040%,which permits the steel only to be used for small components. According toFujita’s investigations, boron stabilises the M23C6 carbides by forming M23

(C,B)6. At the end of 1999, Fujita40 gave a report on the actual results ofcreep tests on specimens of this steel which were carried out at 550°C up toabout 70 000 h, at 600°C up to about 20 000 h and at 650°C up to about125 000 h (Fig. 2.12). The results show that this steel has an extremely high

2.10 International research projects for the development of heat-resistant steels for advanced steam power plants since 1978.

International projects of advanced power plants

Japan USA Europe

R & D : DPDC R & D: EPRI Cost 50/501

Manufacturers, utilities, EPDC Manufacturers Manufacturers. steelworks,utilities and R & D institutes

1983–1997Study 1978–19801981–1991

316 Bar 566/566/566°C314 Bar 593/593/593°C343 Bar 649/593/593°C

310 Bar 566/566/566°C310 Bar 593/593/593°C345 Bar 649/549/549°C

300 bar 600/600/600°C300 bar 600/620°C

Steels and components for Boiler + Turbine50 MW pilot power plant EPRI-RP 1403-15300–900 MW COST 522

1989–19901991–19931994–2000300 bar 630/630°C Steels and components for boiler + turbine

(USA, Japan, Alstom+Man)Steels and components for boiler + turbine

1998–2003300 bar 620/650°C

COST 536

2004–2008

Steels and components for boiler + turbine300 bar max. 650°C

NRIM-STX 21Project:USC 650°C/350 bar boiler

1997–2012Thick wall boiler components Thick wall pipes: P 92 + P 122

(USA, Japan, UK + Denmark)

R&D: 1986–1993

EPRI-RP 1403-50-WO9000-38

1990–1999

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28

Historical development

X 22 CrMo(W)V 12 1/rotors, casings, bolts, blades, pipes

56 T 5/bolts, blades,

H 46; FV 448/bolts, blades, gas turbine discs

Japan : 4 TAF/blades, small forgings

USA : 5 11% CrMoVNbN/rotors (GE)

USA : 6 X 10 CrMoVNbN 9 1 (P 91)/pipes, pressure vessels, casings

Development forfast breeder

(1) Casing EPRI 1403–15(2) Rotors and casings COST 501–2

Japan : 7 + 8 HCM 12/Tubes; TMK1, TMK2/rotors

Cost 501 : 9 X 18 CrMoVNbB 9 1/rotors

Cost 501 : 10 X 12 CrMoWVNbB 10 11/E911

Japan : 11 + 12 NF 616/HCM 12 A/pipes

1950 1960 1970 1980 1990 Year

TMK 1 + TMK 2

100000 h creep strength at 600°C

Improvedpowerplants (600 °C)

MPa160

120

80

40

MPa

120

80

40

1950 1960 1970 1980 1990 Year

2,3

1

5

4

6 8

7E911

9

10

1112

2.11 Overview of the historical development of heat-resistant 9–12% Cr steels within the time range 1950–1995 and the100000 h creep rupture strength of these steels at 600°C.

USA, Germ 1UK: 2

France: 3

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The development of creep-resistant steels

29

Table 2.1 Chemical composition and creep rupture strength at 600°C of the steels in Figure 2.11 (mass%)

Country Rupturestrength at

Steel Chemical composition (weight%) 600°C (MPa)

Basic steels C Cr Mo Ni W V Nb N B 104 h 105 h

Germany 1. X22CrMoV 12 1 0.22 12.0 1.0 0.50 – 0.30 – – – 103 59UK 2. H46 0.16 11.5 0.65 0.70 – 0.30 0.30 0.05 – 118 62

FV448 0.13 10.5 0.75 0.70 – 0.15 0.45 0.05 – 139 64France 3. 56T5 0.19 11.0 0.80 0.40 – 0.20 0.45 0.05 – 144 64Japan 4. TAF 0.18 10.5 1.5 0.05 – 0.20 0.15 0.01 0.035 216 (150)USA 5. 11%CrMoVNbN 0.18 10.5 1.0 0.70 – 0.20 0.08 0.06 – 165 (85)

Advanced steelsUSA 6. P 91 0.10 9.0 1.0 <0.40 – 0.22 0.08 0.05 124 94Japan 7. HCM 12 0.10 12.0 1.0 1.0 0.25 0.05 0.03 75Japan 8. TMK 1 0.14 10.3 1.5 0.60 – 0.17 0.05 0.04 170 90

TMK 2 0.14 10.5 0.5 0.50 1.8 0.17 0.05 0.04 185 90Europe 9. X18CrMoVNbB 91 0.18 9.5 1.5 0.05 – 0.25 0.05 0.01 0.01 170 122Europe 10. X12CrMoWVNbN 0.12 10.3 1.0 0.80 0.80 0.18 0.05 0.06 – 165 90

E911 0.11 9.0 0.95 0.20 1.0 0.20 0.08 0.06 – 139 98Japan 11. P92 0.07 9.0 0.50 0.06 1.8 0.20 0.05 0.06 0.003 153 113Japan 12. P122 0.10 11.0 0.40 <0.40 2.0 0.22 0.06 0.06 0.003 156 101

1.CuJapan 13. HCM 2S 0.06 2.25 0.20 – 1.6 0.25 0.05 0.02 0.003 80Germany 14. 7CrMoTiB 0.07 2.40 1.0 – – 0.25 – 0.01 0.004 60

0.07Ti

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creep strength. Furthermore they demonstrate the creep strength potentialsof ferritic–martensitic 9–11% Cr steels subject to optimum alloying.

The rotor steel 11%CrMoVNbN (no. 5), patented in 1964 by the GeneralElectric Company, USA, is also an onward development of the Nb alloyedsteels no. 2.41 In particular, the Nb content was greatly reduced (0.08%) inorder to prevent harmful segregation at the centre of a rotor. Furthermore,the alloying elements were balanced in order to avoid the formation of deltaferrite. The published creep strength of about 85–90 MPa for 600°C and100 000 h was extrapolated on the basis of tests at 620°C up to times of16 195 h duration.42

The steel referred to in the literature as mod. 9Cr1Mo or P91 (no. 6) is asteel of the newer generation. It was developed under a huge Americannational project in the late 1970s for manufacturing pipes and vessels for afast breeder reactor. It is tough, readily weldable and, as shown by creeptests at 593°C up to about 80 000 h, has a high creep strength at 600°C and100 000 h of about 94 MPa.43 In comparison with earlier steels it ischaracterised, for example, by a lower C content of only about 0.10% and areduced Cr content of about 9%. This steel has meanwhile found wideapplication in all new Japanese and European power stations for themanufacture of pipes and small forgings. It is also used for the manufactureof valve chests and turbine casings.44,45

Steel HCM 12 (no. 7) is a newly developed Japanese 12% Cr steel with

Cre

ep r

up

ture

str

eng

th (

MP A

)

500

400

300

200

100

80

60

40

30100 1000 10000 100000

Time to fracture (h)

550°C

600°C

650°C

700°C

2.12 Creep rupture strength of the TAF steel (0.18%C–10.5%Cr–1.5%Mo–0.2%V–0.15%Nb–0.035%B) as a function of temperature andtime to fracture.

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0.10%C–1%Mo–1%W–0.25%V–0.05%Nb–0.03%N and a duplex structureof delta-ferrite and tempered martensite with improved weldability and creepstrength.37,53 The stability of the creep strength of this steel has been obtainedprimarily by precipitation strengthening with very fine VN precipitates andhigh-temperature tempering at over 800°C.37 Experience with this steel hasalready been accumulated over a period of more than 20 years. The steel hasbeen extensively used for superheater tubes in chemical recovery boilersexposed to severe high-temperature corrosion attack. The good resistance tocorrosion is due to the high Cr-content of 12%.

The Japanese rotor steels TMK 1 and TMK 2 (no. 8), developed in the1980s, were based on the known properties of steels no. 1–6.46 Comparedwith the GE rotor steel (no. 5), the C content in particular was reduced andthe sum total (C + N) was selected at around 0.17%. Based on the researchwork of Fujita,47 the Mo content was raised to 1.5% in TMK 1, whereasTMK 2 is additionally alloyed with about 1.8%W. The Mo content issimultaneously balanced to 0.50% in keeping with Fujita’s result that thehighest solution hardening was obtained with (Mo + 0.5W) = 1.5%. Bothrotor steels have been used for new power stations in Japan.

Rotor steels nos. 9 and 10 are primarily the result of research work performedin the 1980s under the European cooperation programme COST 501.48 Steelno. 9 is a 9%CrMoVNb steel which is additionally alloyed with about 0.01%boron. It is an onward development of the TAF steel for large componentswith reduced contents of Cr, Nb and B and with an increase in V. Creep testswith specimens from a 900 mm diameter pilot rotor, which have so farreached about 100 000 h, indicate a probable creep rupture strength of about120 MPa at 100 000 h and 600°C. Creep tests up to 94 000 h with specimensfrom a 600 mm diameter pilot rotor with a similar composition confirm thisvalue. A larger pilot rotor with a diameter of about 1200 mm is underinvestigation.

The creep strength of the steel X12CrMoWVNbN 10 1 1 (steel no. 10)containing 0.06%N instead of 0.01%B and only 1% Mo instead of 1.5% Moand, in addition, 0.8% W, is about 20% lower, see Fig. 2.13. So far, thelongest test period is about 100 000 h. This rotor steel has primarily beenused for the new advanced European steam power plants. This steel type hasalso successfully been applied for many valve chests and turbine casings inthe new plants. The pipe steel, E 911, developed under COST 501, featuresa somewhat comparable chemical composition. The Cr content was reducedto roughly 9%.49 The Ni content is also distinctly lower since because of thelower Cr content there is no risk of the occurrence of delta-ferrite. Based onresults up to roughly 100 000 h, the creep rupture strength of this steel isestimated to be 98 MPa at 600°C and 100 000 h.

Steel no. 11 is also based on the research work of Fujita. It is a 9% Cr pipesteel specifically alloyed with 0.10%C–1.8%W–0.5%Mo–0.2%V–0.06%Nb–

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0.05%N and about 0.003%B and tempered at about 750°C. It was developedin the second half of the 1980s under the designation NF 616 (P92).50 Basedon creep tests up to about 100000 h, the creep rupture strength is estimatedto be 113 MPa at 600°C and 100 000 h.51

A similar pipe steel HCM 12A (P122), steel no. 12, has also been developedin Japan alloyed with a higher Cr content of 11% in order to improve oxidationresistance. About 1% copper has been added to reduce the tendency to delta-ferrite formation caused by the higher Cr-content.52 The newest evaluationof the available results of long term creep tests led to an estimated creeprupture strength of 101MPa at 600°C and 100 000 h.53 Figure 2.14 comparesthe 100 000 h creep rupture strength versus the temperature for the new pipesteels. In addition to the pipe steel P91, all three pipe steels (E 911, NF 616and HCM 12A) are in successful use for the new advanced steam powerplants.

Development of 10–11%Cr steels for steam turbines up to 650°C

Based on the experience with the steels developed for temperatures up to620°C, the alloying principles needed to obtain a further increase in thecreep strength of 100 MPa for 100 000 h at the application temperature andto improve the resistance to oxidation are very similar in the Japanese andEuropean research programmes:

• increase in Cr-content of 10–11% to improve the resistance to steamoxidation;

ca. 70°C

ca. 20°C

200

100

0

105

h

cree

p r

up

ture

str

eng

th (

MP

a)Steel C Cr Mo W V Nb N B (weight%)(a) 0.28 1.0 0.9 – 0.30 – – –(b) 0.21 12.0 1.0 – 0.30 – – –(c) 0.12 10.0 1.5 – 0.20 0.05 0.05 –(d) 0.12 10.0 1.0 1.0 0.20 0.05 0.05 –(e) 0.18 9.0 1.5 – 0.25 0.05 0.02 0.010

(c/d) X12CrMo(W)VNbN101(1)

(a) 1%CrMoV

(b) 12%CrMoV(e) X18CrMoVNbB91

500 550 600 650Temperature (°C)

2.13 100000 h creep rupture strength of steam turbine rotor steels asa function of test temperature.

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• reduction in Si, Mn- and Ni-content to the lowest possible level to improvethe creep strength;

• addition of 3–6% Co to reduce the diffusion rate and to prevent theformation of delta-ferrite;

• addition of 0.002–0.018% boron to stabilise the M23C6 carbides by formingM23(C,B)6;

• addition of 3%W (Japan) or 1.5%Mo (Europe) to improve the solutionhardening and to stabilise the M23C6 and M23(C,B)6 carbides;

• study of the influence of C content on diffusion rate and on M23C6precipitates.

Table 2.2 provides an overview of the chemical composition of Japaneseand European test melts. None of the referenced analysis versions meets theneeds for use at 650°C. A direct comparison of the creep strength of thetested versions is only possible in part because, as a rule, the results are onlyshown in the form of Larson and Miller diagrams. The investigations withthe MTR and TOS/JSW test melts resulted in the development of rotor steelsMTR 10 A and TOS 110 which were also tested in large pilot rotors.55,57 Forthese two steels a maximum application temperature of 630°C is given.

By way of example, Fig. 2.15 shows the determined creep strength of theweakest version (FB6) and the strongest version (FB8) of the COST programmeat 650°C in comparison with the master steel FB2.58,59 At test periods below10 000 h, the test melts FB6 and FB8 undercut the creep strength curve of the9% Cr master steel FB2 which is suitable for use up to about 625°C. The

40°C25°C

200

100

0

105

h

cree

p r

up

ture

str

eng

th (

MP

a)Steel C Cr Mo W V Nb N B (mass%)(a) 0.10 2.50 1.0 – – – – –(b) 0.20 12.0 1.0 – 0.30 – – –(c) 0.10 9.0 1.0 – 0.20 0.05 0.05 –(d) 0.10 9.0 1.0 1.0 0.20 0.05 0.07 –(e) 0.10 9.0 0.45 1.8 0.20 0.06 0.05 0.002

(a) 2.25%CrMo

500 550 600 650Temperature (°C)

(b) 12%CrMoV

(c) P91(d) E911

(e) P92

2.14 100000 h creep rupture strength of steam plant piping steels asa function of test temperature.

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34Table 2.2 Chemical composition of rotor test melts investigated in Europe and Japan (mass%)

Type C Si Mn Cr Mo W Co Ni V Nb N B

COST ProgrammeMaster steel:FB2 (625°C) 0.13 0.05 0.80 9.3 1.50 – 1.0 0.15 0.20 0.05 0.020 0.010Test Melts:FB5 0.13 0.09 0.09 10.1 1.46 – 2.85 0.15 0.20 0.06 0.017 0.010FB6 0.14 0.10 0.10 11.2 1.45 – 2.93 0.15 0.22 0.08 0.016 0.009FB7 0.16 0.12 0.08 10.2 1.54 – 2.97 0.17 0.20 0.07 0.019 0.009FB8 0.17 0.09 0.09 11.1 1.46 – 2.94 0.20 0.21 0.07 0.023 0.010FB9 0.18 0.08 0.08 11.0 1.48 – 6.03 0.13 0.20 0.06 0.017 0.009FB10 0.12 0.03 0.12 10.3 0.90 0.42 3.02 0.17 0.21 0.07 0.030 0.010FB11 0.19 0.02 0.11 11.1 0.90 0.42 3.02 0.17 0.21 0.07 0.020 0.009

JapanHR1200/20t ingot 0.09 0.02 0.50 11.0 0.23 2.60 2.50 0.51 0.22 0.07 0.020 0.018HR1200/80t ingot 0.10 0.06 0.46 10.2 0.14 2.50 2.40 0.25 0.21 0.07 0.020 0.013HR1200/C 0.12 0.03 0.50 10.3 0.16 2.51 1.18 0.04 0.19 0.07 0.018 0.011HR1200/D 0.13 0.03 0.18 10.5 0.14 2.40 1.17 0.11 0.23 0.07 0.023 0.011HR1200/E 0.14 0.03 0.15 10.1 0.18 2.41 1.45 0.18 0.20 0.07 0.018 0.011MTR various 0.10– 0.04– 0.06– 10.1– 0.33– 0– 0– 0.02– 0.17– 0.05– 0.021– 0–

test melts 0.12 0.10 0.50 11.6 2.50 1.96 6.0 0.30 0.21 0.06 0.076 0.008MTR10A pilot rotor 0.11 0.05 0.06 10.2 0.71 1.76 3.7 0.06 0.20 0.05 0.029 0.004TOS/JSW 0.11 0.04 0.08 10.3 0.70 1.80 3.02 0.20 0.20 0.06 0.018 –

test melts 0.11 0.03 0.08 10.2 0.70 1.76 2.98 0.21 0.20 0.06 0.020 0.0090.11 0.04 0.08 10.2 1.51 – 2.95 0.21 0.20 0.06 0.018 –0.11 0.05 0.08 10.1 1.54 – 2.98 0.19 0.20 0.06 0.018 0.0080.11 0.03 0.08 10.1 0.69 1.81 2.95 0.20 0.20 – 0.022 –0.11 0.03 0.08 10.1 0.72 1.78 2.97 0.20 0.20 – 0.019 0.0090.11 0.03 0.08 10.2 0.70 1.81 2.97 0.20 – 0.06 0.019 –0.11 0.03 0.08 10.2 0.70 1.78 2.97 0.20 – 0.06 0.020 0.008

TOS110 pilot 0.11 0.08 0.10 10.0 0.65 1.80 3.0 0.20 0.02 0.05 0.020 0.010rotor nominal

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650°C creep strengths of the HR1200/20t ingot, the HR1200/80t ingot andthe HR1200/C are shown in Fig. 2.16.60

After 10 000–20 000 h the HR1200/20t and the HR1200/80t versions(variants A and B) reveal a distinct reduction in creep rupture strength.Based on published test progress at 15 000 h, the long-term behaviour of theHR 1200/C version cannot be evaluated. However, the target of 100 MPa at100 000 h will not be reached even after a further stable pattern of the creeprupture strength curve. According to the results of investigations into themicrostructure of versions FB6 and FB8 of the COST programme, the markeddecline in creep strength is mainly caused by precipitation of the Z-phaseCr2(Nb,V)2N2 which is formed at the expense of MX precipitations.54,61

Afurther striking fact is a remarkable coarsening of M23C6 precipitations andthe Laves phase. A significant coarsening of the M23C6 precipitations and theLaves phase was also determined in investigations of versions HR1200/20tand HR1200/80t.60 This publication, however, does not provide any informationon MX and Z-phase precipitations.

It is significant when considering the interpretation of the investigationresults into the microstructure of these rotor test melts that the long-termcreep-stressed specimens of the thermally very stable steels TAF, B2 andFB2 revealed no Z-phase precipitations, but fine MX precipitations and onlyminor coarsening of the M23C6 precipitations.61,62 The influence of 8.56–11.59%Cr on the creep strength at 650°C for a steel which in addition isalloyed with 0.10%C–3%W–3%Co–0.70%Mo–0.15%V–0.06%Nb–0.02%N–0.01%B has been reported.63 Figure 2.17 presents an overview of the results.The summary includes details of the creep rupture time as a function of the

Cre

ep r

up

ture

str

eng

th (

MP

a)

200

100

80

60

40

100 1000 10000 100000Time to fracture (h)

Test temperature 650°C

FB8

FB6

650°C FB2

2.15 Creep rupture strength of the COST test steels FB2, FB6 and FB8as a function of time to rupture at 650°C.

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Cr content and test stresses in the range 300–98 MPa. In the long-term range,that is, at lower test stresses, the alloying procedure with approximately9%Cr exhibits the highest creep strength whereas in the short-term range thealloy with about 11.5%Cr reveals the highest creep strength.

Investigations into the microstructure of the 157 MPa stressed creepspecimens of the 9.5%Cr and 11.5%Cr versions showed the following resultsafter creep rupture periods of 5289 h and 1928 h respectively. For the firstset result the 11.5%Cr steel, the dislocation density decreased, subgrainswere formed and only a minor fine Laves phase could be observed within thelath structure. For the second set result the 9.5%Cr steel on the other hand,the lath width remained narrow and the dislocation density was higher thanfor the 11.5% Cr steel and in addition, a significant amount of fine Lavesphase had still been precipitated within the lath structure. The conclusiondrawn was that the coarsening of precipitates, such as the Laves phase, wasincreased and that recovery of the microstructure could not be suppressed inthe higher Cr steels.

Development of 11% Cr steels for pipes and headers for 650°C steamplants

Based on experience with the steels developed for temperatures up to 620°C,the alloying principles needed to obtain a further increase in the creep strength

Ru

ptu

re s

tren

gth

(M

Pa)

300

200

100

80

60

40

30



Ingot sizes:20 t (A)80 t (B)

small (C)small (E)

HR1200 variantsA B C E

newdesign

C 0.09 0.09 0.10 0.14Si 0.02 0.06 0.03 0.03

Mn 0.50 0.47 0.50 0.15Cr 11.0 10.2 10.3 10.1Mo 0.23 0.14 0.16 0.18Ni 0.51 0.25 0.04 0.10W 2.6 2.59 2.51 2.41V 0.22 0.21 0.19 0.20

Nb 0.07 0.07 0.07 0.07N 0.02 0.018 0.018 0.018B 0.018 0.012 0.011 0.011

Co 2.53 2.46 1.18 1.45Al 0.007 <0.005 0.005 0.001

2.16 Creep rupture strength of the rotor test steels HR1200 (VariantsA-B-C) as a function of time to rupture at 650°C.

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to reach 100 MPa for 100 000 h at the application temperature and to improvethe resistance to oxidation are similar in the research programmes undertakenin Japan and Europe:64–66

• increase in Cr-content to about 11% to improve the resistance to steamoxidation,

• addition of 1–3%Co to reduce the diffusion rate and to prevent theformation of delta-ferrite,

• addition of 0.003–0.012% boron to stabilise the M23C6 carbides by formingM23(C,B)6,

• addition of 0.3–3%W and/or 0.15–1.5%Mo (Europe) to improve solutionhardening and to stabilise the M23C6 and M23(C,B)6 carbides,

Heat treatment5 h 1070°C+20 570°C

+20 h 680°C

Metallographicinvestigation

8 9 10 11 12Cr content (%)

98MPa

118MPa

137MPa

157MPa

177MPa

210MPa

240MPa

270MPa

300MPa

Tim

e to

ru

ptu

re (

h)

104

103

102

101

2.17 Creep rupture strength of 8.5–12%Cr–3.5%W–3%CoVNbB teststeels as a function of Cr content, applied stress and time to fractureat 650°C.

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• addition of 0.07%Ta and 0.04%Nd to improve the creep strength by fineand stable nitrides.

Table 2.3 provides an overview of the chemical composition of Japanese andEuropean test melts. None of the referenced analysis versions meets theneeds for use at 650°C. Figure 2.18 shows a comparison of the publishedresults for steels VM12, NF12 and SAVE 12 with regard to the target creepstrength at 650°C.67–69 The behaviour is comparable with the creep strengthestablished for the rotor steels in Table 2.2, that is, with 10–11%Cr and Co-contents of roughly 3%. Results of the investigations into the microstructurehave hitherto only been published for version NF12. After creep stressing for15 000 h at 650°C it was similarly discovered that Z-phase precipitation tookplace at the expense of MX precipitation.70

Japanese NIMS research project STX 21 for the development of thick-section boiler components

The Japanese research project STX 21, started in 1995,71 embraces a mostcomprehensive 15-year investigation into the influence of the alloying elementson the creep strength and oxidation resistance of the ferritic–martensiticsteels in thick-section boiler components for advanced power stations. Thesystematic investigation into all relevant elements in regard to material propertyrequirements for main steam pipes and headers covers:

• 105 h creep rupture strength at 650°C• oxidation resistance in steam• weldability and creep rupture strength of welded joints• thermal fatigue• impact properties and hot workability.

Figure 2.19 demonstrates the design philosophy of a 9%Cr steel72 withthe alloying partners W, Mo, Ni, Cu, Co, Si, V, Nb, Ta, C, N and B and theirspecific property profile. In addition, further analysis versions were tested.Table 2.4 provides an overview of investigated element variations within theproject. Reports on test progress were issued on a regular basis at internationalconferences. The best version was established to be a steel with 0.08%C–9%Cr–3%W–3%Co–0.20%V–0.05%Nb–0.008%N–0.014%B.73 It was foundthat the addition of boron at more than 0.01% to the 9%Cr steel remarkablyimproves the long-term creep strength. Boron stabilises the lath martensiticmicrostructure of 9%Cr–3%W–3%Co steels during creep deformation at650°C through the stabilisation of M23C6 in the vicinity of prior austenitegrain boundaries by an enrichement of boron in the M23C6 carbides. Furtherimprovement of creep strength in this high boron-containing steel isaccomplished by the addition of 0.008% nitrogen enhancing precipitation offine MX. The steel also exhibits good creep–rupture ductility.

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39

Table 2.3 Chemical composition of pipes and header test melts investigated in Europe and Japan (mass%)

Type C Si Mn Cr Mo W Co Ni V Nb N B

COST programme 63

Master steelP91: 0.10 0.35 0.45 8.75 1.00 – – 0.20 0.21 0.080 0.050 –F1 0.16 0.15 0.09 11.0 1.04 0.32 2.12 0.14 0.26 0.065 0.059 0.009F2 0.16 0.30 0.22 11.1 1.04 0.32 2.16 0.43 0.27 0.067 0.060 0.011Master steelP 92: 0.10 0.40 0.50 9.0 0.45 1.8 – 0.15 0.20 0.060 0.045 0.003F 0.11 0.40 0.50 11.5 0.28 1.4 1.30 0.23 0.24 0.065 0.056 0.003F3 0.15 0.30 0.22 11.2 0.50 2.1 2.15 0.44 0.26 0.068 0.063 0.009Master steelB2: 0.17 0.07 0.06 9.34 1.55 – – 0.12 0.27 0.064 0.015 0.010A 0.19 0.10 0.52 11.6 1.48 – – 0.20 0.26 0.058 0.015 0.004B 0.18 0.10 0.51 11.5 1.48 – – 0.25 0.26 0.058 0.047 0.004Master steelFB2: 0.13 0.05 0.82 9.32 1.47 – 0.96 0.16 0.20 0.05 0.019 0.009C 0.16 0.09 0.53 11.3 1.46 – 0.90 0.26 0.25 0.047 0.063 0.006D 0.16 0.10 0.50 11.4 1.48 – 1.06 0.26 0.25 0.058 0.063 0.012E 0.13 0.10 0.49 11.3 1.44 – 3.06 0.28 0.25 0.057 0.051 0.008G 0.16 0.10 0.50 11.4 1.46 – 1.49 0.27 0.27 0.044 0.050 0.008Pilot pipes 0.12 0.50 0.35 11.5 0.30 1.5 1.6 0.30 0.25 0.050 0.065 0.005VM1264

JapanResults of investigationsof various test meltsNF1265 0.10 0.35 0.50 10.2 0.15 2.5 2.0 <0.10 0.22 0.07 0.02 0.005SAVE 1265 0.10 0.30 0.20 11.0 – 3.0 3.0 <0.10 0.20 0.07 0.04 0.07Ta

0.04Nd

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2.18 Creep rupture strength of piping test steels NF12, SAVE 12 andVM12 as a function of test time at 650°C in comparison to thebehaviour of the 9% Cr piping steel P92.


NF12SAVE12

VM12

P92 (mean ECCC)


Ru

ptu

re s

tren

gth

(M

Pa)

300

200

100

80

60

40

30

Cr < 8 Cr = 9 Cr > 11

W + 2Mo < 2 W + 2Mo = 3 W + 2Mo > 3

Ni = 0 Ni = 0.5(Cu, Co) Ni > 0.5(Cu, Co)

Si < 0.3 Si = 0.5 Si > 0.6

V, Nb, Ta

C, N, B

Oxidation resistancedecreaseAddition of Si

δ ferrite formationAddition of austenitestabilising elements

Creep strength decrease Toughness decrease

δ ferrite formation Creep strength decreaseAc1 temperature decreaseAdditon of Ir, Rh, Pd

Carbide agglomerationAcceleration of Fe2Wprecipitaton

Fine dispersion of MCby TMCPIncrease in MC volumefraction

Carbide agglomerationWeldability decrease

C = 0.08–0.1, N = 0.03 –0.06, B = 0.005–0.008

Creep strength decrease

Oxidation resistancedecrease

Addition of Y

V = 0.2, Nb = 0.05, Ta = 0.1–0.2

2.19 NIMS alloy design philosophy of high Cr ferritic steels for 650°Cultrasupercritical (USC) boilers.

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41Table 2.4 Chemical composition of test steels for pipes and boilers of the NIMS STX 21 project (mass%)

Type C Si Mn Cr Mo W V Nb N Co B

9CrW 0.10 0.30 0.50 9.0 – – 0.02124

9CrWMo 0.08 0.30 0.50 9.0 3 0.003 0.22 0.05 0.052.8 0.102.5 0.331.9 0.62<0.01 1.54

9CrWN 0.10 0.30 0.50 9.3 – 3.3 0.20 0.05 0.054 3.0 0.0050.0260.001

9CrWB 0.08 0.30 0.50 9.0 – 3.0 0.20 0.05 0.02 3.0 0+0.0149CrCoTiAl 0.10 0.01 0.01 9.0 – – 0.20 0.05 0.05 3.0 0.014

(Ti+Al)9–12Cr –0.01 <0.01 Var. – 2.0 0.20 0.05 0.05 – Var. Var.Y Var.WSiTiY 0.14 0.30 8.5 Ti 0 Al(Oxidation– 0.50 10 0 0.05 0.001tests)

0.80 12 0.05 0.10 0.0081.0 0.10

9CrWCoPd 0.08 0.30 0.50 9.0 – 3.0 0.20 0.05 0.05 0+3 0.005 0Pd1Pd3Pd

9CrWVTa 0.10 0.30 0.50 9.0 – 3.0 0.20 0.05 0.05 Ta0.02

ODS 0.065 9 – 2 Ti Y2O310CrTiY2O3 –0.12 –13 0.12 0.08

+.22 –0.35ODS <0.05 12.3 – 0–3 Ti Y2O312CrWY2O3 0–0.39 0.24

–0.25

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Figure 2.20 presents an overview of the influence of B and N on creepstrength of 9Cr3W3CoVNbNB steel at 650°C in comparison with the TAFand P92 steels. The creep strength of the best 9%Cr version with 0.014%Band 0.008%N correlates in the long-term range with the creep strength of theTAF steel. Considering the results hitherto available, the authors of thispublication conclude that, based on the present chemical composition, a100 000 h creep strength of 100 MPa at 650°C is achievable. Protective Cr-rich scale is formed on the surface of 9%Cr by a combination of Si additionand pre-oxidation treatment in argon gas at 700°C for 20 h. This significantlyimproves oxidation resistance of 9%Cr steels at 650°C.74

2.4 Historical development of austenitic steels

2.4.1 History of austenitic steels

Nickel alloyed austenitic steel originated from 25%Ni–Fe alloy and 25%Ni–5 to 8%Cr–Fe alloy melted by Krupp in Germany in 1893 and 1894,respectively.75 Krupp also produced 35%Ni–13 to 14%Cr–Fe and 25%Ni–8to 15%Cr–Fe alloys for use in thermocouple sheaths and molten glass mouldmaterial in 1910.75 Krupp continued the development of a series of Ni–Cr–Fe steels and identified a martensitic 10%Cr–2%Ni steel and an austenitic

0.0139B/0.0034N 1080/800°C0.0092B/0.0016N 1050/790°C0.0048B/0.0011N 1050/790°C0.0135B/0.0079N 1150/770°C0.0100B/0.0200N 1080/800°C

9Cr3W3CoVNbB(NIMS)


TAF (Fujita)

P92(ECCC)

100 1000 10000 100 000Time to fracture (h)

Ru

ptu

re s

tren

gth

(M

Pa)

300

200

100

80

60

40

30

2.20 Creep rupture strength of 9Cr3W3CoVNbB NIMS test steel as afunction of B and N content, heat treatment and time to fracture incomparison to the behaviour of the 9% Cr piping steel P92 and theTAF steel.

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20%Cr–5%Ni steel in 1912 as stainless steels.76 The latter was designated asV2A in 1922 (V for Versuchstahl, respectively test steel, 2 as the developmentnumber, and A for austenitic) and served as an original austenitic steel foruse after solution annealing for corrosion resistance in vessel piping, machinery,and so on in acid environments.77

Conventional 18%Cr–8%Ni austenitic steel originated from V2A in themid-1920s, optimising the content of Cr and Ni to maximise the economicalbenefit by keeping the structure austenitic.78 Since these austenitic steelssuffered severe intergranular corrosion in the weld heat affected zone, extensiveinvestigation into the corrosion mechanism and improvement of the steelswas carried out at Krupp around 1930. Consequently, stabilised austeniticsteels were developed employing alloying elements Ti, V, Nb and Ta, as wellas reducing the carbon content to a maximum of 0.07% and achieving a finegrained steel with finer distribution of chromium carbide.79 These steels areoriginal versions of the conventional Type 321 and Type 347 steels. Type316 originated from 18%Cr–8%Ni steel containing 3% molybdenum for usein ammonium chloride and sulphuric acid environments as a corrosion-resistantalloy. At that time (early 1930s), addition of the combination of Mo and Cuwas found to deliver improved corrosion resistance against sulphuric acidand 18%Cr–8%Ni–2%Mo–2%Cu was thus developed by Krupp. From theviewpoint of better cold workability of austenitic steels, austenite formingelements such as Mn and Cu were increasingly used to develop several newsteels, namely 19-9LW(19%Cr–9%Ni–1.25%Mo–1.25%W–NbTi), 19-9DX(19%Cr–9%Ni–1.5%Mo–1.2%W-Ti) and 17-14CuMo81 modified fromKrupp V6A by Armco Steel in USA.80 These steels with stabilised austeniticstructures exhibited not only corrosion resistance to sulphuric acid but alsoresistance to heat and creep.

Figure 2.2182 shows the elevation of creep rupture strength of heat-resistantsteels for boilers, viewed in terms of change in 100 000 h creep rupturestrength at 600°C, for materials developed during the 20th century. AfterWorld War II, 18%Cr–8%Ni steels previously developed in Germany beforethe war were used for heat-resistant purposes as well as in chemical plantsworldwide, thereby increasing the steam pressure and temperature of fossil-fired power plants. Also, ultra-supercritical pressure power plants constructedin the late 1950s were realised by further applying these austenitic steels tothick walled components. For example, TP316H was used for boiler headersand steam piping and 17-14CuMo and TP321H for superheater and reheatertubes at Eddystone unit No. 1. Esshete 1250 (15%Cr-10%Ni-6%Mn-1%MoVNbB)83 austenitic steel was also used for thick section componentsin the UK.

In the 1950s high nickel Alloy800H was developed and put in service inthe USA as a high strength and anti-corrosive steel substitutable for Ni-based alloys. We must remember that these austenitic steels were basically

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developed for corrosion-resistant materials and unexpectedly showed goodcreep-resistant properties. From the perspective of the improved creep resistanceof the 18%Cr–8%Ni system, austenitic steel alloy development studies startedin the 1960s. Tempaloy A-184 was developed through the optimisation ofcarbon stabilising elements of Ti and Nb in 18%Cr–8%Ni steel in the early1970s, followed by thermo-mechanically strengthened TP347H with finergrain structure developed in the early 1980s, designating TP347HFG85 foruse in superheaters and reheaters as a steam oxidation-resistant and creep-resistant material. Subsequently, several other 18%Cr–8%Ni austenitic steelswith highly improved creep strength, such as Super304H,86 XA70487 andTempaloy AA-188 were developed.

Within the group of high Cr and high Ni austenitic steels, several new 20–25%Cr austenitic steels with nitrogen and relatively lower Ni content weredeveloped from the 1980s to the 1990s. As seen in Fig. 2.21 (for reference),the improved creep strength of ferritic steels is notable, along with theimprovement of austenitic steels. Regarding ferritic steels, low alloy steels

105

h C

reep

ru

ptu

re s

tren

gth

at

600 °

C (

MP

a)

200

150

100

50

01900 1920 1940 1960 1980 2000 2020

Year

SAVE25HR6WNF709

RH3C AA-1 XA-704Spr304H

A-3 Next generaton12Cr-WCoVNb

3rdgenerationA-1 TP347HFG17-17CuMo

E1250800H

Stable austenitic

TP347HTP316HTP321H

TP304H

Meta-stable austenitic

Ferritic

(2Cr)T22

410 T9(9-12Cr)Cr-Mo

Ni-Cr

NF616HCM12A

E9112nd generation

HCM12T91

HCM2ST24 (2Cr)

1st generationF-9EM12 HCM9M

HT9

HT91

2.21 Historical improvement in creep rupture strength of boilersteels.

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or 9–12% Cr steels with about 40 MPa of 100 000 h creep rupture strength(half or less than that of 18%Cr–8%Ni steels) had been used for many yearsand the problem of cost increases existed because, especially in the case ofsuperheater and reheater applications, there was an alloy ‘gap’ between thelow alloy steels and the 18%Cr–8%Ni steels, arising as a result of temperatureelevation. Accordingly, research work for high strength 9–12%Cr steels wasinitiated in order to fill this gap and 60 MPa class materials (first generation)were developed over the period from 1950 to 1970. Further efforts have beenmade and creep strength reached the 100 MPa class (second generation) inthe 1980s, whereas the 130 MPa class (third generation) was achieved in the1990s. Materials for 150 MPa at 600°C or 100 MPa at 650°C, as the nextgeneration, are expected to emerge. In general, the outside diameter and wallthickness of pipes and tubes can be greatly reduced through elevation ofcreep rupture strength. Thermal stresses can accordingly be lowered andconstruction of fossil-fired power plants capable of load sliding operationunder further elevated pressure and temperature steam conditions will bepossible.

2.4.2 Alloy design of creep-resistant austenitic steels

Heat-resistant steels for practical application must be designed by takingtheir service conditions and environments into consideration and by examiningtheir various properties.89 However when alloy design is performed based onmodification of existing steels, both oxidation and corrosion resistance aswell as their general material properties are expected to be nearly equivalentto those of the original materials. Hence, chemical compositions and heattreatment conditions are examined with particular consideration of creepstrength improvement. Figure 2.2289 shows the concept of alloy design forheat-resistant austenitic steels to improve creep strength through themodification of existing steels. In the case of austenitic steels, chemicalcomposition can be largely classified into the four categories shown in thefigure, and solution strengthening and precipitation strengthening are designedspecifically for each of these categories.

18%Cr–8%Ni steels based on Type 304 steels include Type 316 steelssolution-strengthened through the addition of Mo, as well as Type 321 steelsand Type 347 steels precipitation-strengthened through the addition of Ti orNb. However, these materials were originally developed for chemical equipmentas mentioned above, placing emphasis on corrosion resistance, but were notdesigned from the standpoint of creep strengthening. Accordingly, furtherenhancement of precipitation strengthening by means of ‘under-stabilising’90

C and/or composition design for improved creep strength is used. 15%Cr–15%Ni or 21%Cr–30%Ni steels with a full austenitic phase structure arecapable of high creep strength in the ‘as-is’ condition, although they are

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46Austenitic steel

18%Cr-8%Ni Type 304 Type 316 (Mo) Cr carbide

Solutionstrengthening

Precipitationstrengthening

Stabilising

Under stabilising

Alloy design for creep

Type 321 (Ti), TiC

Type 347 (Nb), NbC

Lower (Ti + Nb)/C

Micro alloying ofNb, Ti, B

High strength18%Cr-8%Ni steel

Cu, Mo, W addition

Creepresisting

steelN addition

Creep/corrosion-resistant

steel

Low-costcreep/

corrosion-resistant

steel

Type 17-14CuMo

Type 310

Type alloy800H

15%Cr-15%Ni(15%Cr-10%Ni)

25%Cr-20%Ni

21%Cr-32%Ni

2.22 General concept of alloy design for austenitic heat-resistant steels.

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costly because of their high Ni content. Steels containing 20% and more Crare likely to have excellent oxidation and corrosion resistance, but a costlyNi content of at least 30% is required to maintain a full austenitic structure.Nevertheless, low-cost, high-strength, highly corrosion-resistant austeniticsteel can be designed by adding about 0.2% N to reduce the Ni content andby combining the strengthening mechanism as described above.

The features of alloy design of austenitic heat-resistant steels are discussedbelow. As previously noted, ‘under-stabilising’ is a technique for improvingthe creep strength of 18%Cr–8%Ni steels. This method enhances creep strengththrough the improvement of precipitation morphology by fixing C in alloysand decreasing carbide forming elements such as Ti and Nb, which hinder Crcarbide formation, to the point where their contents are insufficient for Cfixation. Figure 2.2390 shows this, and the peak point of the creep rupturestrength against the ratio of (Ti + 0.5Nb)/C is at position far away from thepeak point of conventional Type 321 or Type 347 steels, indicating thatreducing additions of Ti and Nb relative to the C content can be useful.Figure 2.2486 shows the effect of the Cu and Nb additions on the creeprupture strength of 18%Cr–9%NiNbN steel and 18%Cr–9%NiCuN steel,respectively. Although Cu addition does not show a major change up toabout 2%, a substantial enhancement in creep strength is observed owing tothe very finely dispersed Cu-rich particle precipitates by means of Cu addition

321H347H

0.01 0.1 0.5 1.0 1.5 2.0(Ti + Nb)/C Atomic ratio

0.01 0.1 0.5 1.0 1.5 2.0 3.0 4.0 5.0 6.0 7.0 8.0(Ti + 0.5Nb)/C Mass ratio

105

h C

reep

ru

ptu

re s

tren

gth

at

650 °

C (

MP

a)

130

120

110

100

90

80

70

60

50

40

30

2.23 Effect of (Ti + Nb)/C ratio on creep rupture strength of18Cr10NiNbTi steel.

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of about 3% or more. However, because the strength tends to be saturated,while a decline in creep rupture ductility can occur when the Cu additionexceeds 3%, the addition of Cu at 3% should be suitable. On the other hand,the effect of Nb addition on the creep rupture strength of 18%Cr–9%NiCusteel is significant when Nb content exceeds 0.2% and is saturated at aconcentration of above 0.4%, which is not sufficient to stabilise the 0.08%Cusually contained in 18%Cr–8%Ni H grade steels. This was mentioned aboveas an effect of under-stabilising on the creep strength enhancement. Super304Hand other Cu-alloyed and under-stabilised steels were thus developed.

2.4.3 Development of austenitic heat-resistant steels

Austenitic steels for heat exchanger and boiler tube applications

Chemical compositions of austenitic heat resistant steels for boiler applicationsare given in Table 2.5, with development progress in austenitic boiler steelspresented in Fig. 2.25.91 Considering steels for boiler applications, variousimprovements have been made to enhance creep strength in the grade of18%Cr–8%Ni system steel because this grade is used in the highest temperatureregion of superheater and reheater sections of boilers with conventionalsteam parameters. The steam oxidation resistance of some 18%Cr–8%Nisteels is also improved at their inner surface through finer grains. Furthermore,new austenitic steels with a Cr content of 20% or more have been developedfor the purpose of improving creep strength and corrosion resistance.

700°C/137MPa750°C/108MPa

167MPa137MPa

137MPa108MPa

Tim

e to

ru

ptu

re (

h)

103

102

10

18Cr-9NiNbN 18Cr-9NiNuN

0 2 4 6Cu (%)

Tim

e to

ru

ptu

re (

h)

103

102

100 0.2 0.4 0.6 0.8 1.0

Nb (%)

700°C 750°C

2.24 Effect of Cu and Nb on creep rupture strength of 18Cr–8Nisteels.

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49

Table 2.5 Nominal chemical composition of austenitic steels for boilers

Steels Chemical composition (mass%)

C Si Mn Ni Cr Mo W V Nb Ti B Others

18%Cr– TP304H 18Cr–8Ni 0.08 0.6 1.6 8.0 18.0 – – – – – – –8%Ni TP316H 16Cr–12NiMo 0.08 0.6 1.6 12.0 16.0 2.5 – – – – – –

TP321H 18Cr–10NiTi 0.08 0.6 1.6 10.0 18.0 – – – – 0.50 – –TP347H 18Cr–10NiNb 0.08 0.6 1.6 10.0 18.0 – – – 0.80 – – –TP347HFG 18Cr–10NiNb (FG) 0.08 0.6 1.6 10.0 18.0 – – – 0.80 – – –Tempaloy A-1 18Cr–10NiNbTi 0.12 0.6 1.6 10.0 18.0 – – – 0.10 0.08 – –Super304H 18Cr–9NiCuNbN 0.10 0.2 0.8 9.0 18.0 – – – 0.4 – – 3.0Cu,

0.1NXA704 18Cr–9NiWVNb 0.03 0.3 1.5 9.0 18.0 – 2.5 0.3 0.3 – – 0.2NTempaloy AA-1 18Cr–10NiCuTiNb 0.10 0.3 1.5 10.0 18.0 – – – 0.3 0.2 0.02 3.0Cu

15%Cr– 17–14CuMo 17Cr–14NiCuMoNbTi 0.12 0.5 1.7 14.0 16.0 2.0 – – 0.40 0.30 0.008 3.0Cu15%Ni– 15–15N 15Cr–15NiMoWNbN 0.12 0.7 1.5 15.0 15.0 1.5 1.5 – 1.0 – – 0.1N

AN31 15Cr–13NiMoNbN 0.10 0.5 1.5 14.0 16.0 1.5 – 0.5 1.0 – – 0.1NEsshete1250 15Cr–10Ni6MnVNbTi 0.12 0.5 0.6 10.0 15.0 1.0 – 0.2 1.00 0.06 – –12R72 15Cr–10NiMoNbVB 0.10 0.4 2.0 15.0 15.0 1.0 – – – 0.3 0.006 –

20%– TP310 25Cr–20Ni 0.08 0.6 1.6 20.0 25.0 – – – – – – –25%Cr HR3C 25Cr–20NiNbN 0.08 0.4 1.2 20.0 25.0 – – – 0.45 – – 0.2N

Alloy 800H 21Cr–32NiTiAl 0.08 0.5 1.2 32.0 21.0 – – – – 0.50 – 0.4AlTempaloy A-3 22Cr–15NiNbN 0.05 0.4 1.5 15.0 22.0 – – – 0.70 – 0.002 0.15NNF709 20Cr–25NiMoNbTi 0.15 0.5 1.0 25.0 20.0 1.5 – – 0.20 0.10 – –SAVE25 22.5Cr–18.5NiWCuNbN 0.10 0.1 1.0 18.0 23.0 – 1.5 – 0.45 – – 3.0Cu,

0.2NSanicro25 22Cr–25NiWCuNbN 0.08 0.2 0.5 25.0 22.0 – 3.0 – 0.3 – – 3.0Cu,

0.2NHigh Cr– CR30A 30Cr–50NiMoTiZr 0.06 0.3 0.2 50.0 30.0 2.0 – – – 0.20 – 0.03ZrHigh Ni HR6W 23Cr–43NiWNbTi 0.08 0.4 1.2 43.0 23.0 – 6.0 – 0.18 0.08 0.003 –

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50

18Cr-8Ni

(AISI 302)

–C 18Cr-8Ni,C<0.08

(AISI 304)+Ti

18Cr-10NiTi

(AISI 321)

+Nb18Cr-10NiNb

(AISI 347)+Mo

16Cr-12NiMo

(AISI 316)

+Cr–Ni

22Cr-12Ni

(AISI 309)

25Cr-20Ni

(AISI 310)

(37) (65)

25Cr-20NiNbN 23Cr-18.5NiWCuNbN

SAVE25(SUS310J3TB)

HR3C ASME TP310CbN SUS310J1TB

( ) Designates 105 h Creep rupture strength (MPa) at 700°C.

22Cr-25NiWCuCoNbN

(90)

(91)

Sanicro 25

(40–60)

H Grade0.04 – 0.10C

AISI 304HAISI 321HAISI 347HAISI 316H

(75)

17Cr-14NiCuMoNbTi

W Addition (74)

18Cr-9NiWVNb

XA704(ASME TP347W)(SUS347 J1TB)

HeatTreatment (58)

18Cr10NiNb

(ASME TP347HFG)

ChemistryOptimization (58)

18Cr10NiNbTi

Tempaloy A-1(SUS321J1HTB)

18Cr9NiCuNbN

Cu Addition (70)

Super 304H

ASME TP304CuCbNSUS304J1HTB

(53)

21Cr-32NiTiAl

(Alloy 800H)

Cu Addition (76)18Cr10NiCuTiNb

Tempaloy AA-1ASME TP 321HCuCb

SUS321J2HTB

High Cr-High Ni

20Cr-25NiMoNbTi

22Cr-15NiNbN

23Cr-43NiWNbTi

30Cr-50NiMoTiZr

(85)

(66)

(89)

(73)

NF 709(SUS310J2TB)

Tempaloy A-3(SUS309J4HTB)

HR6W

Cr30A

2.25 Development progress of austentic boiler steels.

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18Cr–8Ni steels such as TP304H, TP321H, TP316H and TP347H gradesare still used for fossil-fired power plants operating under conventional steamconditions. TP347H, which has the highest allowable stress among thesefour types, was improved in order to provide a fine-grained structure (grainsize number 8 and finer) for steam oxidation resistance and creep strengthening,designated as TP347HFG by ASME. The effect of solution temperature onthe creep rupture strength and grain size in thermo-mechanical processes isschematically illustrated in Fig. 2.26.85 This steel is highly effective inimproving the reliability of superheater tubes, being applicable to ultra-super critical pressure power plants up to 600°C class and already fullyemployed in the superheater tubes of a substantial number of these plants.

Since the 15%Cr–15%Ni system steel is stable austenitic and high strengthsin creep are likely to be obtained, their allowable stresses are very high. Themechanism of the strengthening of these steels is strong solution hardeningby Mo and W, and precipitation hardening effects by inter-metallic compoundsas well as by MC and M23C6 carbides make a contribution. Table 2.692

shows the nominal chemical compositions of typical elements forming inter-metallic compounds and carbides, Nv – Nc (Nv – Nc is the excess electronvacancy number by PH ACOM which is defined in ASTM A567, v = vacancy,

DevelopedConventional

10

8

6

4

2

120

100

80

60

40

20

Gra

in s

ize,

AS

TM

No

.10

5 h

cre

ep r

up

ture

stre

ng

th (

MP

a)

ASMEA S./0.67

ASMEA.S./0.67

700°C × 105 h

650°C × 105 h

1100 1200 1300 1400Solution temperature (°C)

A.S.: Allowable stress

Developed TP347HFG

Anneal AnnealNbCPrecipitated

NbCResolved

Colddrown

Coldcrown

Solutiontreatments

Fine graln Coarce graln

Solutiontreatments

Tem

per

atu

re

2.26 Effect of solution temperature on creep rupture strength andgrain size in a thermo-mechanical process.

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c = critical) values, and schematic microstructures. The creep rupture propertiesof five 15%Cr–15%Ni system steels are presented in Fig. 2.27.92 Amongthese alloys 17-14CuMo exhibits the highest creep rupture strength for thelongest time and at the highest temperature. In Fig. 2.28, the allowablestresses for 18%Cr–8%Ni steels and for 17-14CuMo and Esshete1250representing 15%Cr–15%Ni steels are compared. The allowable stress ofSuper304H is much higher than that of 17-14CuMo, which is thought tohave the highest stress at temperatures up to 670°C among conventionalmaterials. TP347HFG and Tempaloy A-1 also have allowable stresses higherthan those of existing conventional steels. Because these steels have beendeveloped on the basis of TP304H or TP321H, their cost effectiveness isexcellent. They are also advantageous from the standpoint of resistance tosteam oxidation owing to their fine grains.

Although newly developed 20%–25%Cr austenitic steels and high Cr–high Ni austenitic steels such as CR30A93 and HR6W94 have excellentresistance to high temperature corrosion and steam oxidation compared toother austenitic steels, their drawback lies in the fact that they are toocostly for their allowable stresses. However, as shown in Table 2.5, thematerials developed in the 1980s, particularly N-alloyed 20%–25%Cr steel,have excellent high temperature strength as well as being relativelyinexpensive. They are practically applied as high strength steels taking hightemperature corrosion resistances into account. Allowable stresses for HR3C,95

M23C6

TiC

Fe2MoNbC

Fe2Mo

NbC(Nb, Ti)C

Schematic microstructuresafter ageing at 700°C for 104 hSteel

17–14CuMo

15–15N

AN31

Esshete 1250

12R72

W

–

1.49

–

–

–

Nb

0.67

1.07

1.07

0.96

–

Ti

0.28

–

–

–

0.30

NV –NC (700°C)

0.117

0.096

0.162

0.197

0.093

Mo

2.12

1.58

1.53

1.04

1.16

Table 2.6 Comparison of major alloying elements and schematic microstructures in15Cr–15Ni austenitic steels

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NF70996 and Tempaloy A-397 are far higher than that of Alloy 800H, asshown in Fig. 2.29 and they can be used in higher steam conditions and incorrosive environments.

As an example, Fig. 2.3095 shows the effect of the content of solute N onHR3C, with the changing Nb and Ni content in the alloy. Nitrogen as analloying element acts as a creep strengthening agent in 25%Cr austeniticsteel and the combination of 0.4%Nb and 20%Ni addition is an effectivemeasure of the enhancement of creep strength. Alloy 800H has a stableaustenite structure, stabilised by using a large addition of Ni, but hightemperature strength is insufficient in relation to cost.

Although there is currently no choice but to use austenitic steels forsuperheater and reheater tubes for ultra-supercritical pressure boilers, certainmaterials have already been developed that are sufficient to meet the steamconditions of 650°C class boiler superheater and reheater tubes, as notedpreviously. From the mid-1990s to date, materials taking cost effectivenessinto consideration have also been developed. SAVE2598 and Sanicro2599 are

2.27 Comparison of creep rupture strength of 15Cr–15Ni austeniticsteels.

17–14CuMo15–15NAN31Esshete 125012R72

700°C

102 103 104 105

Time to rupture (h)

700°C

102 103 104 105

Time to rupture (h)

Str

ess

(MP

a)

300

200

100

70

Str

ess

(MP

a)

300

200

100

70

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examples that use 0.25% addition of N to stabilise the austenitic structurebased on HR3C, in addition to a small amount of Nb addition aimed atprecipitation strengthening by means of ‘under-stabilising’. Furthermore,comprehensive strengthening techniques covering a wide range of temperatureshave been employed by introducing the concept of Cu addition in Super304Hand W addition in HR6W. Here, cost effectiveness has been secured bystabilising the austenitic structure through addition of N and Cu, whiledecreasing the Ni addition to 18% and reducing the Cr content to a levelslightly below that of HR3C. As shown in Fig. 2.25, austenitic steels thathave the highest creep rupture strength at 700°C are HR6W, SAVE25 andSanicro25, exhibiting around 90 MPa at 100 000 h. Figures 2.3194 and 2.32100

show creep rupture strength in the temperature range 650–800°C. Stablecreep rupture properties are confirmed in these figures.

From the standpoint of realising 700°C class ultra-super critical pressurepower plants, achievement of 100 MPa of 100 000 h creep rupture strength(70 MPa in allowable stress) is considered to be required, so that the limitsfor application of 18%Cr–8%Ni steels and 20–25%Cr austenitic steels with

2.28 Allowable stresses for 18Cr–8Ni and 15Cr–15Ni austenitic steels.

Allo

wab

le s

tres

s (M

Pa)

150

100

50

0500 550 600 650 700

Temperature (°C)

Super 304H

XA704

TP347HFG

TP347H

17–14CuMo

Tempaloy AA-1

TP316H

TP304HTP321H

Tempaloy A-1

E1250

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Allo

wab

le s

tres

s (M

Pa)

150

100

50

0500 550 600 650 700

Temperature (°C)

NF709

HR3C

SAVE25

Tempaloy A-3

Alloy 800H

HR6WCR30A

TP310H

2.29 Allowable stresses for 20–25Cr and high Cr–high Ni austenitic steels.

700°C 750°C Nb Ni

0.4/0.5

20/23

17

20

23

103

h C

reep

ru

ptu

re s

tren

gth

(M

Pa)

200

180

160

140

120

100

80

60

40

20

00.05 0.10 0.15 0.20 0.25 0.30

Soluble N (%)

2.30 Effect of content of solute N and Nb and Ni on creep rupturestrength of HR3C.

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the highest creep rupture strength in 350 bar plants would be approximately660°C and 680°C, respectively, as shown in Fig. 2.33.91 In the accompanyingfigure, allowable stresses at higher temperatures for numerous austeniticsteels developed to date are presented together with those for Ni-based alloys.However the allowable stresses in Ni-based alloys are not yet well establishedowing to insufficient databases and inaccuracy of predicted strength at100 000 h. Reviewing the austenitic steels developed beyond the strength ofconventional steels to date, materials can be classified in the strength levelsas shown in Table 2.7 at a temperature of 700°C. Austenitic steels are graded

650°C700°C750°C800°C

HR6W:0.07C-23Cr-43Ni-6W-0.1Ti-0.2Nb

101 102 103 104 105

Time to rupture (h)

Str

ess

(MP

a)

400

300

200

100

70

50

40

2.31 Creep rupture properties of HR6W.

650°C700°C750°C800°C

100 101 102 103 104 105

Time to rupture (h)

Str

ess

(MP

a)

500

400

300

200

70

50

2.32 Creep rupture properties of SAVE25 steel.

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in five levels except for the 80 MPa and 85 MPa classes in which strengthsare separated in increments of 5 MPa. Table 2.891 lists austenitic steels andNi-based alloys whose maximum application temperature is categorised by620–660°C, 620–680°C and 680–770°C. The table also shows the ASMECode Case number assigned and applicable product forms, tubes and/orpipes.

Austenitic steels for thick-section pipe and turbine componentapplications

Thick section components in boilers consist of headers and steam piping andunder conventional steam conditions, CrMo low alloy steel (for example,2.25%Cr–1%Mo steel, 0.5%Cr–0.5%MoV steel or martensitic 12%Cr stainlesssteel X20CrMoV121)101 has been used extensively up to 566°C for manyyears. From 1989 onwards, P91102 has been selected for headers and pipingoperating up to 600°C and from the mid-1990s P92103 has also been employedfor recently constructed 600°C class ultra-super critical power plants. However,under steam conditions above the temperature at which ferritic steels can beapplied, generally above 620°C or 630°C, austenitic steels must presently beused from the standpoint of creep strength and oxidation resistance.

Among austenitic steels, ASTM TP316 has actual application experiencefor the header and main steam piping in the Eddystone Station unit No. 1 inthe USA under steam conditions of 350 bar and 650°C. Being an austeniticsteel, this material has a low thermal conductivity meaning that thermalstresses will be higher than for ferritic steels. Further, as the high temperatureproperties of TP316 are greatly influenced by its chemical composition, itshigh temperature strength will be highly unstable if the chemical composition

Table 2.7 Strength classification of austenitic steels

105 h creep rupture Allowable Steels Categorystrength at 700°C (MPa) stress (MPa)

60 40 TP347HFG 18CrTempaloy A-1 18Cr

65 43 HR3C 20/25CrNF709 20/25CrTampaloy A-3 20/25Cr

70 50 Super304H 18CrXA704 18CrTempaloy AA-1 18Cr

75 53 CR30A High Cr80, 85 55 – –90 60 HR6W High Cr

SAVE25 20/25CrSanicro25 20/25Cr

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58

Table 2.8 Candidate alloys for USC power boilers

Trade designation Nominal composition ASME Code/ Preferred Temperature ofCode Case application application

(metal)/°C

Austenitic 347HFG 18Cr–10Ni–Nb 2159 T 620–660steels Super304H 18Cr–8Ni–Cu–Nb–N 2328 T(18%Cr) XA704 18Cr–9Ni–W–V–Nb 2475 T

Temp. AA-1 18Cr–10Ni–Nb–Ti 2512 T

Austenitic HR3C 25Cr–20Ni–Nb–N 2115 T 620–680steels NF709 20Cr–25Ni–Nb–Ti–N 2581 T

Temp. A-3 22Cr–15Ni–N6–N – T(20–25%Cr) 800HT 21Cr–32Ni–Al–Ti 1325 T

SAVE25 23Cr–18Ni–W–Cu–Nb–N – TSanicro25 23Cr–18.5Ni–W–Cu–Nb–N – THR120 25Cr–42Ni–N – THR6W 23Cr–43Ni–6W–Nb–Ti – P.T

Ni-base Haynes230 22Cr–5Co–3Fe–14W–2Mo–La 2063 P.T 680–770allows Inco617 22Cr–13Co–9Mo–Al–Ti 1956 P.T

Inco625 22Cr–5Fe–9Mo–Nb–Al–Ti 1409 P.TInco740 25Cr–20Co–Mo–Nb–Al–Ti – P.T45TM 27Cr–23Fe–2.75Si 2188 P.T

T: superheater/reheater tubes, P: pipes and headers

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is not carefully optimised. This must be given special consideration becauseTP316 is very suitable for use in the heavy sectioned components operatingat steam temperatures above the maximum limit for ferritic steels. In thecase of TP316, the formation of the sigma phase can be suppressed in theoryby having the Ni-valance (= 11.6 + Nieq – 1.35Creq, where Nieq = 0.5Mn +30(C+N) + Ni, Creq = 1.5Si + Cr + Mo + 0.5 Nb) above 0.104 In practice,however, the sigma phase may precipitate after a long term of service causinga major drop in the creep rupture strength. Figure 2.34105 shows creep rupturetest results for the main steam pipes produced at the end of the 1950s andused in the Eddystone No. 1 unit over a period of 130 250 h. It can be seenthat a low Ni balance greatly reduces the creep rupture strength. The materialsindicated as 17U and 17D in the graph were produced at the start of the1980s using modern steel-making techniques to keep Ni balance high andwere used for only 6000 h under the same conditions, with other samplesshowing high strength. Among the steels listed in Table 2.5, HR6W is beingconsidered for thick section components of 700°C class ultra supercriticalpressure power boilers, taking advantage of its very high creep rupture strengthwhich is on a par with Ni-based alloys.

2.33 Comparison of maximum allowable stresses at hightemperatures.

Allo

wab

le s

tres

s (M

Pa)

120

110

100

90

80

70

60

50

40

30

20

10

0600 650 700 750 800 850 900 950 1000

Temperature (°C)

Inconel 740N

imonic 263

Inconel 617

Hynes230

Sanicro25

SA

VE25Tem

paloy A-3

Alloy 800H

TP310S Super304HXA704

HR3CHR6W

NF709CR30A

Hastelloy XHastelloy XR

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A typical thick section component in a steam turbine is the rotor. Forquite some time, 1%CrMoV steel and conventional 12%Cr steel have beenused for steam temperatures of 566°C and below. The maximum usetemperatures for these steels are considered to be 570°C and 590°C,respectively. 12%Cr steel, for which 1.5% W is added in substitution for theequivalent amount of Mo, can be used up to approximately 600°C, buthigher temperatures require 12%Cr steel with the addition of W and Co.While the limit has not been clearly defined, safety considerations suggest atemperature of around 630°C. The arrival of ferritic steel that can withstand650°C is awaited, but advanced materials development is needed since rotorsrequire not only high-temperature strength but also good tensile strength andtoughness.

Austenitic steels such as Discaloy (13%Cr–25%Ni–3%MoAlTi) and A286(15%Cr–26%Ni–1.5%MoAlTi) have demonstrated performance in EddystoneNo. 1 unit at 650°C and 350 bar, and in a 50 MW demonstration plant inWakamatsu, Japan also at 650°C. But temperatures over 700°C would requireNi-based alloys such as Inco617 (Ni–22%Cr–13%Co–9%MoAlTi) or Waspaloy(Ni–19%Cr–14%Co–4.5%MoAlTi), or Fe–Ni-based alloys with improvedcreep strength. From the perspective of practical application, fabricabilitywould need to be improved to enable the manufacture of large components.

Figure 2.35 presents the relationship between 100 000 h creep rupturestrength and temperature for various materials based on the results of researchto date. Along with consideration of Inco617, efforts are also being madeworldwide to improve Inco706 (Ni–16%Cr–36%Fe–3%NbAlTi) and to developalloys modifying the low thermal expansion Inver-type Ni alloy. Of these,LTES700 (Ni–12%Cr–18%MoAlTi)106 is an austenitic alloy featuring a low

2.34 Creep rupture properties at 650°C for as-exposed TP316 steel.

8D (–0.134)8U (–1.167)7D (–0.017)7U (–0.961)

18D (1.195)18U (2.306)19D (0.425)19U (0.045)

Sample (Ni-bal ) 25U (2.134) 25D (2.126)

20D (–0.370)20U (–0.367)17D (1.536)17U (1.339)

102 103 104 105

Time to rupture (h)

NIMSdata (ave.)

650°C

Str

ess

(MP

a)400

300

200

100

70

50

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thermal expansion coefficient similar to ferritic steels and having strengthequivalent to that of Refractaloy26 (Fe–18%Cr–38%Ni–20%Co–3%MoAlTi).Furthermore FENIX-700 (Ni–16%Cr–36%Fe–2%NbAlTi),107 modifiedInco706 and TOS-1X (Ni–23%Cr–10%Mo–15%CoAlTiB)108 have beendeveloped aiming at 700°C class steam turbine rotor application.

For temperatures above around 630°C, austenitic steels must be usedbecause ferritic steels are presently not strong enough. The Discaloy used inEddystone No. 1 unit weighed 1.6 tons and was made by melting at atmosphericpressure. However, rotors in 1000 MW class ultra supercritical pressurepower plants will require a steel ingot weighing nearly 30 tons, and althoughit is possible to cast ingots of over 800 mm in diameter by the electro slagremelting (ESR) melting process, there will be a greater tendency for frecklingto occur and this may impair the mechanical properties.109 Tests on a 1000mm diameter A286 steel ingot with a rationalised chemical composition castusing large-size ESR equipment have shown that even if freckling occurs, ithas no effect on high cycle fatigue or fracture toughness at room temperature,or on creep rupture strength, high or low cycle fatigue at high temperatures,although it may reduce tensile elongation at temperatures below 200°C.110 Tiand C contents influence freckling to the greatest extent and it has beenreported that there was no freckling at all in a 1350 mm diameter ESR A286steel ingot in which the Ti and C contents had been reduced to 1.13% and0.02%, respectively. Although reducing the Ti content lowered the creeprupture strength of an A286 ingot, as shown in Fig. 2.36,110 for the level ofreduction mentioned above, it satisfies the tensile property and creep rupture

2.35 Comparison of 105 h creep rupture strength of turbine rotoralloys.

105

h C

reep

ru

ptu

re s

tren

gth

(M

Pa)

700

500

400

300

200

100

70450 500 550 600 650 700 750

Temperature (°C)

Waspaloy

Refractaloy 26

Inco706

LTES

Inco617

Mod. A286

12CrWCo12CrMoW12CrMoV

1CrMoV

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strength requirements for rotor material to be used at 650°C, with no reductionin creep notch strength. The Ti and C contents in A286 can also be reducedto 1.52–1.81% and 0.023–0.039%, respectively with no loss in toughness at650°C confirming in theory the excellent suitability of the chemically improvedA286 as a 650°C class rotor material.

Austenitic steels for chemical plant applications

The typical heat-resistant austenitic steels in chemical plants are used forreformer tubes in hydrogen treating equipment and for cracking tubes inethylene polymerisation equipment in thermal cracking of naphtha attemperatures above about 800°C. In the case of high pressure equipmentsuch as reformer tubes, higher creep rupture strength is required and anti-carbonisation is also required for cracking tubes which are used at highertemperatures. Table 2.9111 shows the chemical composition of these austenitictube materials for chemical plants made by centrifugal casting. The basealloy is HK40, 25%Cr–20%Ni with a high carbon content of 0.4%. Thecreep rupture strength has been improved through investigation into theeffect of chemical elements. Nb and Ti are used to form eutectic carbidewhich strongly contributes to improved creep rupture strength, and W andMo are used as solution strengthening elements in the austenitic matrix.

In order to prevent the precipitation of the sigma phase and to stabilise theaustenitic matrix as well as to improve anti-carbonisation, the Ni content isincreased from 20% to 30–50%, and Si is added (also for the anti-carbonisation).

100 kg ingot2 ton ingot16–19 ton ingot

Ti (mass%)1.8 1.91.61.41.21.00.9

105 h

Cre

ep r

up

ture

str

eng

th a

t 65

0°C

160

150

140

130

120

110

2.36 Effect of Ti content on 105 h creep rupture strength of A286 alloyat 650°C.

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63

Table 2.9 Nominal chemical composition of austenitic steels for chemical plants


C Cr Ni Nb Co W Mo Ti Others

Cr–Ni– HK-40 0.40 25.0 20.0 – – – – – –based Mod. HK40 0.43 23.5 22.5 – – – – – –

High Si-HK40 0.45 25.0 20.0 – – – – – Si:2.0HP 0.50 25.0 35.0 – – – – – –High Si-HP 0.50 25.0 35.0 – – – – – Si:2.0HO 0.50 27.0 29.0 – – – – – –

Ti addition HiKa-1A 0.40 25.0 20.0 – – – – 0.20 R.E.:0.3Nb addition IN-519 0.30 24.0 24.0 1.5 – – – – –

Manaurite 36 0.50 25.0 35.0 1.5 – – – – –Manaurite 36XS 0.50 25.0 35.0 1.0 – 1.5 – – Si:2.0Hika-1B 0.40 25.0 20.0 0.2 – – – 0.20 R.E. :0.3BST 0.45 25.0 20.0 0.6 – – – 0.15 –Mod. BST 0.50 24.0 23.0 0.7 – – – 0.15 –HP-BST 0.50 25.0 35.0 0.7 – – – 0.15 –Manaurite 900 0.15 22.0 30.0 1.2 – – – – –

Mo addition HOM 0.40 25.0 35.0 – – – 1.2 – –MoRe-1 0.45 25.0 35.0 – – 2.0 – – –NA-22H 0.50 28.0 48.0 – – 5.0 – – –MoRe-2 0.40 33.0 50.0 – – 17.0 – – –

Co addition HOM-3 0.45 25.0 45.0 – 3.0 3.0 3.0 – –Super 22H 0.40 28.0 48.0 – 5.0 5.0 – – –Tenex 0.50 22.0 30.0 – 13.0 4.5 – – –Supertherm 0.50 25.0 35.0 – 15.0 5.0 – – –

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High temperature oxidation and corrosion resistance are improved by theaddition of rare earth metal material. Figure 2.37112 demonstrates the creeprupture strength of these centrifugally cast austenitic steel tubes. The creeprupture strength of HK40 is lowest, while the other steels are substantiallyimproved by alloy development. HP-BST whose chemistry is a nominal0.5%C–25%Cr–35%NiNbTi exhibits the highest creep rupture strength. Inthe case of centrifugally cast tubes, small diameter sizes are difficult toproduce and the extrusion process is also applied to produce tubes. However,the eutectic structure is decomposed by the extrusion process, resulting in aweakening of creep rupture strength. In this case, creep strength improvementby heat treatment is used in reforming the eutectic type carbide structures inthe austenitic matrix.

2.5 Historical development of steel melting and of

the purity of heat-resistant steels

The increasing demand of the steam plant manufacturer for turbine rotorforgings with improved cleanliness, soundness, higher ductility and toughnesshas mainly led to improvement in steelmaking processes for heat-resistantsteels within the last century. Figure 2.38 gives a rough overview of thehistorical development of steelmaking processes throughout the 20th centuryand the achievable reduction in the phosphorus and sulphur contents as themain indicators of the effectiveness of these steelmaking processes.15,113–117

At the beginning of the last century only steels of the Thomas and open

1000°C

Mod. BST

HP-BST

Super themNA22HMoRe-1

Manaurite 36XHPHK40IN519

102 103 104 105

Time to rupture (h)

Str

ess

(MP

a)

50

40

30

20

10

2.37 Creep rupture strength of centrifugally casted tubes for chemicalplants.

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652.38 Historical development of steel melting processess and their achievable P and S contents.

(*) Oxidizing refining in EF reladle

(1) S < 0.075%/P < 0.10%

1900 1920 1940 1960 1980 2000Year of application

Ladle refining furnace LRFVacuum induction melting VIMVacuum arc remelting VARVacuum oxygen decarburization VODArgon oxygen decarburization AOD

Electric slag remelting ESRVacuum carbon desoxydation VCD

Vacuum degassing VAD (H2 <2ppm)Oxygen blown steel BOF

Electric-arc steel EFThomas steel (1877)Basic open hearth steel (1964) BOH

Ste

el m

elti

ng

pro

cess

es

(1)

Max

. P

an

d S

(%

)

0.02

0.01

P,S

P,S

P,S

P P

P*P* S S S

Thomas steel BOH EF BOF EF + ESR EF + AOD EF + LRFSteel melting processes

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hearth processes, which possess high P and S contents were available. Theelectric-arc furnace steel, first introduced in 1908, provided very cost effectiveand much lower levels of P and S and better overall cleanliness but hydrogenflaking was still a problem of larger forgings. This problem was solved bythe introduction of vacuum degassing introduced in the 1950s. The productionof oxygen blown steels started in 1952. This basic oxygen steelmaking converterprocess, which works with a high amount of pig iron (nearly 70%), offers avery economic way to produce clean steels with lower P and S contents andwith lower tramp elements, such as As, Sb and Sn, as long as cleaner pig ironis used. A further improvement in vacuum treatment was established in 1961with the adoption of the vacuum carbon deoxidisation (VCD) process. VCDallows for the removal of oxygen through reaction with carbon to form COas a gas which is then removed from the pouring stream by a vacuum system.This provided cleaner steels than the earlier practice in which oxygen wastied-up with solid deoxidants (Si, Al) which resulted in undesirable silicateand Al-oxide inclusions.

Demands for cleaner and more uniform chemical homogenity were providedby the secondary remelting methods: electroslag-remelting (ESR) in themiddle of the 1960s and vacuum-arc remelting (VAR) at the end of the1960s. The advantage of the ESR process is also a low S content of about0.002%. The VAR process provides no reduction of the trace elements butvery low gas contents. Very effective secondary refining steelmaking processesare also the decarburisation technologies Argon-oxygen decarburisation (AOD)and vacuum-oxgen decarburisation (VOD) were introduced at the end of the1960s. The advantages of the AOD process are lower S contents (about0.002%) and lower N2 and H2 levels. Vacuum-induction melting (VIM) alsoprovides very low gas contents. For this process very clean iron and alloyores, selected scraps and very clean pig iron are used. The risk of shrinkagesin the VIM ingot requires a remelting (ESR or VAR) after the VIM process.VIM+ESR is used, for example, for the Ni-alloys Inconel 617 and Inconel625, but also for special steels, like 17-4PH. For the very high stressedInconel 706 gas turbine discs, a triple melting of VIM+ESR+VAR is common.

In the middle of the 1970s, very large forgings for nuclear and fossil firedplants with higher output demanded larger ingots. This challenge led to thedevelopment of ladle refining furnaces and subsequently commited the electric-arc furnace only to the role of melting. With the developed ladle refiningfurnaces, various combinations of heating, degassing, desulphurising,dephosphorising and methods for inclusion shape control are possible inorder to provide large forging with high quality.114 This ladle metallurgyallows delivery P and S contents in the range of about 0.002%. Low P andlow Mn contents are also often made by oxidising refining processes in theelectric furnace (see for example Azuma et al.)118 In addition to phosphorus,antimony, arsenic, tin and copper are elements which lead to embrittlement

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in service and to lower creep strength in the heat-resistant steels. Each ofthese elements is non-oxidisable and therefore controlled through scrapselection for the electric-arc furnace melting process in order to avoid steeldegradation. A further weakness in steelmaking practice in the past was thedeoxidisation of the steels with aluminium. In consequence of the discoverednegative influence of Al on creep strength and creep ductility in long terminvestigations, Al is nowadays limited to levels lower than 0.010%.

2.6 Summary

The development of creep-resistant steels is a result of continuous technologicalprogress throughout the 20th century. The urgent need to improve the creepstrength of steels was based on endeavours by the power station industry toimprove the thermal efficiency of steam power plant by raising the steamtemperature and steam pressure in order to reduce the cost of fuel and goeasy on fuel resources. The major contribution to the increase in power plantefficiency consisted in the development of heat-resistant steels with a highercreep strength at an acceptable creep ductility level. The significance ofthese material properties was not recognised until early damage was sufferedby steam turbine bolts in the 1930s, which pointed to the fact that the strengthof steels used in power stations operating at higher temperatures dependssignificantly on the creep behaviour of the material over the full period ofoperation. Based on this experience, it was concluded that the strength valuesshould no longer be determined in short-term tests but that a procedureshould be adopted to determine the rupture strength, the creep strain andcreep ductility of the heat-resistant steel in a creep test extending over aperiod of roughly 100000 h.

Based on the multiplicity of investigations of test steels carried out withdifferent Mo, Cr, Ni, V, CrMo, CrV, MnSi, MoMnSi, CrSiMo, CrNiMo,CrMnV and CrMoV contents, worldwide developments in the manufactureof steam boilers and small forgings for steam turbines brought forth lowalloyed steels with chemical compositions of 0.15%C–0.3% to 0.5%Mo,0.13%C–1%Cr–0.5%Mo and 0.10%C–2.25%Cr–1%Mo which are still inuse today. In addition, in the 1950s, a MoV steel with 0.14%C–0.5%Mo–0.3%V with an even higher creep strength was developed in Europe for gasturbines and afterwards also qualified in long term creep tests for steamplants. In the field of turbine manufacture, a steel with approximately 0.25%C–1.25Cr–1%Mo–0.30%V is in worldwide use for turbine rotors, casings, boltsand small forgings.

The development of heat-resistant 9–12%Cr steels was strongly motivatedby two major events: during the 1950s by the development of thermal powerstations for public power supply operating at steam temperatures rangingfrom 538–566°C and during the 1980s by the target to develop low-pollution

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power stations operating at steam admission temperatures of 600–650°Cwith supercritical pressures up to 350 bar. The steel X22CrMoV 12 1 wasdeveloped in the 1950s for thin-walled and thick-walled power stationcomponents. Its creep strength is based on solution hardening and on theprecipitation of M23C6 carbides. The steel has been applied successfully inpower stations over several decades up to temperatures of about 566°C.

The steel referred to in the literature as mod. 9Cr1Mo or P91, is a steel ofthe newer generation. It was developed under a huge USA project in the late1970s for the manufacture of pipes and vessels for a fast breeder. This steelhas meanwhile found wide application in all new Japanese and Europeanpower stations with steam temperatures up to 600°C for pipes and smallforgings. It is also used for the manufacture of castings, like valve chests andturbine casings. The increase of creep strength in comparison to the 12%CrMoVsteel is caused by 0.05%Nb and 0.05%N which form thermal stable VN andNb(C, N) precipitates. A lower Cr content of about 9% also contributes to thehigher creep strength. Similar creep strengths are exhibited by new steelsdeveloped for rotors, castings and pipes which are in addition alloyed with1%W. A further increase of about 10% in creep strength reveals the steeltype P92 which is in addition alloyed with about 0.003%B and has an increasedW content of about 1.8%. Addition of boron gives thermally stabile M23(C,B)6precipitates whereas the higher W content leads to a higher amount of theLaves phase. Higher B contents in the range of 0.010–0.014% give steels forrotors, casings and pipes which can be used up to temperatures of about630°C. Further ferritic 9–10%Cr steels are under development for steamtemperatures up to 650°C.

A great challenge for the development of further improved 9–10%Crheat-resistant steels is, first, the avoidance of the Z-phase precipitationCr2(Nb,V)2N2 which precipitates at the expense of the MX particles, andsecond, the avoidance of BN and metallic borides which precipitate at theexpense of M23(C,B)6 and VN. An important contribution to the developmentof advanced steam power plants was made in the last century with a continuousimprovement in the melting technology which improved significantly thepurity of the steels and the manufacturability of large heat-resistant steelcomponents.

Austenitic steel itself has originally been developed for chemical plantequipments used in various corrosion and oxidation environments.Economically to keep the structure austenitic, Cr and Ni contents wereoptimised at 18% and 8%, respectively (Type 304 steel) in the 1920s; latercompletely stabilised austenitic steels such as Types 321 and 347 steels weredeveloped. Type 316 steel was also developed containing 3% Mo for use inammonium chloride and sulphuric acid environments in 1930s. At the sametime, the addition of combination of 2%Mo and 2%Cu to 18%Cr–8%Ni steelwas found to provide excellent corrosion resistance against sulphuric acid.

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Fortunately the austenitic steels exhibit not only good corrosion resistancebut also very high creep rupture strength, such as above 100 MPa of 100000 hstrength at 600°C for 18%Cr–8Ni% steels.

The steam temperature of fossil power plant in the USA and Europe roseto around 538–566°C from the end of the 1940s to the beginning of the1950s. The same situation reached Japan at the end of the 1950s. This meantthat low alloyed Cr–Mo steels could no longer be used in boilers in superheaterand reheater tubes whose metal temperature exceeded 580°C, from the viewpoint of oxidation and corrosion as well as creep strength behaviour. Austeniticsteels or high Cr ferritic steels needed to be used. Four types of 18%Cr–8%Ni system austenitic steels could be used for such high temperature steamplants. Also, ultra-supercritical pressure power plants constructed in the latterhalf of the 1950s were made possible by further application of these austeniticsteels to thick walled components. For example, TP316H was used for boilerheaders and steam piping, 17%Cr–14%NiCuMoNbTi (categorised in the15%Cr–15%Ni system austenitic steel), TP321H for superheater and reheaterand Discaloy (13%Cr–25%Ni–3%MoAlTi) for a very high pressure turbinerotor at the Eddystone unit No. 1 of Philadelphia Electric operated at atemperature of 649°C in the 1960s.

After 1980, numerous studies on the development of austenitic steels foruse in power plants operating at temperatures above 650°C were startedaiming at the enhancement of creep strength rather than using conventionalaustenitic steels originally developed for chemical applications. The alloydesign concepts for austenitic steels were (1) enhanced precipitationstrengthening by means of ‘under-stabilising’ C composition, (2) loweringthe material cost by addition of N to reduce Ni content to maintain a fullaustenitic structure of high Cr above 20% containing steel, (3) utilisation ofCu and W for precipitation hardening and solution strengthening, and (4)precipitation hardening by inter-metallic compounds. Presently, the greatest100 000 h creep rupture strengths at 700°C available are 70 MPa for 18%Cr–8%Ni system steels and 90 MPa for 25%Cr austenitic steels. As 100 MPa isthought to be required for high-temperature and high-pressure power plants,the limits for application of 18%Cr–8%Ni system steels and 25%Cr austeniticsteels in 350 bar pressure plants would be approximately 660°C and 680°C,respectively. For thick wall components experience of Type 316 steel isalready available at the 650°C steam temperature. However, for temperaturesabove 650°C, development of new austenitic steels or improvement of Ni-based alloy is needed. For steam turbine rotor applications, austenitic superalloys such as Dicaloy (13%Cr–25%Ni–3%MoAlTi) and A286 (15%Cr–26%Ni–1.5%MoAlTi) have demonstrated a restricted performance at the650°C level, but temperatures of 700°C and greater would require Ni-basedalloys such as Inco617 (Ni–22%Cr–13%Co–9%MoAlTi) or Waspaloy(Ni19%Cr–14%Co–4.5%MoAlTi), or Fe–Ni based alloys with improved creep

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strength. The limit for austenitic steels is approximately 680°C for 350 barpressure plants. On the other hand, Ni-based alloys can be applied attemperatures in excess of 680°C or 700°C, but the material cost is quite highcompared with austenitic steels. In future, continued development is neededfor austenitic steels that can be used at temperatures above 700°C aiming atthe 100 000 h creep rupture strength of 100 MPa and greater at 700°C.

2.7 References

1 Schult H, ‘Grundlage der zukünftigen Entwicklung der Stromwirtschaft’, Kohle inder Elekrizitätswirtschaft Essen, 1952, 6–11.

2 Scarlin B, Alstom Switzerland, personal communication 2006.3 Kallen H, ‘Der Werkstoff Stahl in der technischen Entwicklung der letzten 100

Jahre’, Stahl und Eisen, 1960, 80, 1864–1877.4 Ruttmann W, ‘Untersuchungen an Schraubenbolzen’, Mitt. VGB, 1937, issue 65,

395–396.5 Schmitz H, ‘Vereinheitlichung des Dauerstandversuches mit Stahl’, Stahl und Eisen,

1935, 1523–1534.6 ASTM Tentative Standards, ASTM, Philadelphia, 1934, 1138–1156.7 Siebel E, ‘Neuzeitliche Werkstoffprüfung im Hochdruckkessel’, Mitt. VGB, 1938,

Issue 67, 74–79.8 Diehl H and Granacher J, ‘Ergebnisse aus Zeitstandversuchen bei 500°C mit einer

Beanspruchungsdauer bis über 300 000 h’, Arch. Eisenhüttenwesen, 1979, 50, (7),299–303.

9 Wellinger K, ‘Beanspruchung und Werkstoffe’, VGB-Werkstofftagung, Essen,Germany, 1969, Tagungsband, 9–17.

10 Ewald J, Bendick W, Granacher J, Maile K, Melzer B, Mayer K H, Rhode W andTolksdorf E, ‘50 Years of joint german activities in the area of creep resistantmaterials’, Proceedings of International Colloquium on the Occasion of the 50thAnniversary of the German Creep Committee, Düsseldorf, Germany, 25 November,1999, 1–31.

11 Thornton D V, ‘UK power industry collaboration high temperature materials testingprogramme’, Proceedings of International Colloquium on the Occasion of the 50thAnniversary of the German Creep Committee, Düsseldorf, Germany, 25 November1999, 32–37.

12 Masuyama F, ‘Industry view of creep study activities in Japan’, Proceedings ofInternational Colloquium on the Occasion of the 50th Anniversary of the GermanCreep Committee, Düsseldorf, Germany, 25 November 1999, German Iron andSteel Institute, Düsseldorf 1999, 38–41.

13 Abe F, Irie H and Yagi K, ‘Recent activities in NRIM creep data sheet project’,Proceedings of International Colloquium on the Occasion of the 50th Anniversaryof the German Creep Committee, Düsseldorf, Germany, 25 November 1999, 42–46.

14 Eberle F, ‘Einige Ergebnisse amerikanischer Langzeit-Standversuche anRöhrenstählen, ihre Berücksichtigung bei der Festlegung zulässiger Spannungen,International Meeting about Long Term Creep Tests’, 31.05./01.06.1957 Düsseldorf,Germany, Archiv für das Eisenhüttenwesen, 1957, 28, 702–706.

15 Curran R M, ‘The development of improved forgings for modern steam turbines’,

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Proceedings of ASTM Symposium on Steel, Stainless Steel and Related Alloys,Williamsburg, VA, USA, 28–30 November 1984, ASTM Special Technical Publication903, Code No. 04-903000-02, 9–32.

16 Mitteilung der Schweizerischen Arbeitsgemeinschaft (BBC, Gebr. Sulzer, G. FischerAG, Escher Wyss und von Roll); K.N. Melton et al, Zeitstanduntersuchungen an21/4 Cr-1Mo-Stahlguss GS-18Cr Mo910, Material und Technik 1982/4, December1982, 190–197.

17 Baumann K and Gramberg U, ‘260 000 Stunden einer Löffler-Kesselanlage mit 130bar und 510°C’, VGB Kraftwerkstechnik, 1977, 57, (10), 699–706.

18 Fabritius H, Entwicklungsstand von warmfesten und korrosionsbeständigen Stählenfür die Erdöl- und Erdgasindustrie, Mannesmannforschungsbericht, 1973, 610.

19 Florin C, ‘Ferritische warmfeste Stähle’, in Werkstoffkunde der gebräuchlichenStähle, Teil 2, Verlag Stahleisen, Düsseldorf 1977, 96–105.

20 Orr J, Beckitt F R, Met A and Fawkes G D, ‘The physical metallurgy of chromium–molybdenum steels for fast reactor boilers’, Proceedings of BNES (British NuclearEnergy Society) Conference on Ferritic Steels for Fast Reactor Steam Generator,London, 1977.

21 Foldyna F, Purmensky J, Prnka T and Kadulova M, ‘Einfluß des Molybdängehaltesauf die Zeitstandfestigkeit von Chrom-Molybdän-Vanadinstählen mit niedrigemKohlenstoffgehalt’, Arch. Eisenhüttenwesen, 1971, 42, 927–932.

22 Schüller H J, Hagn L, Woitscheck A, Kober A and Christian H, ‘Schweißnahtrissein Formstücken an Heissdampfleitungen-Werkstoffuntersuchungen’, VGB-KonferenzWerkstoffe und Schweißtechnik im Kraftwerk, VGB Essen, Germany, 1973, 163–194.

23 Stahl-Eisen-Werkstoffblatt SEW 555, January 2001, Steels for Large Forgings forComponents in turbine and generator installations, Verlag Stahleisen GmbH,Düsseldorf.

24 EN 10269:1999, Steels and Nickel Alloys for Fasteners with Specific Elevated and/or Low Temperature Properties, Beuth Verlag GmbH, Berlin, Germany, 1999.

25 Norton F and Strang A, ‘Improvement of creep and rupture properties of large1%CrMoV steam turbine forgings’, Journal Iron and Steel Institute, 1969, February,193–203.

26 Mayer K H and Rieß W, ‘The influences of the microstructure on the operationalcharacteristics of steam turbine components subjected to high stresses’, VGBKraftwerkstechnik, 1976, No 3, March, 138–142.

27 Smith R, PhD Thesis, Sheffield University, 1959, cited in detail in Ref. 31.28 Thum A and Richard K, ‘Ergebnisse von Langzeit-Dauerstandversuchen bei 500°C’,

Schweizer Archiv, 1953, 19, (8), 235–245.29 Schinn R, ‘Die deutsche Entwicklung von Schmiedstücken für Wellen von

Turbogeneratoren’, VGB Technisch-wissenschaftlicher Bericht, WärmekraftwerkeVGB-TW503, VGB Essen Germany.

30 Conrad J D and Mochel N L, ‘Operating experiences with high temperature steamturbine rotors and design improvements in rotor blade fastening’, TransactionsASME, 1958, 80, No 6 1210–1224, presented in Allentown,USA, 21–23 October 1957.

31 Cress T, Zum Zeitstandverhalten warmfester Chrom-Molybdän-Vandin Stähle undderen Neigung zu verformungslosen Zeitstandbrüchen, Dissertation Techn. HochschuleDarmstadt, Germany, 1967.

32 Ruttmann W, ‘Untersuchungen an Schraubenbolzen’, Mitt. VGB, 1937, issue 65,395–396.

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33 VGB Guidelines for Bolts in the Range of High Temperatures, VGB-R 505M, 3rdedition, VGB Essen, Germany, 1977.

34 Kramer L D, Randolph D D and Weisz D A, ‘Analysis of the Tennessee ValleyAuthority Gallatin Union No 2 Turbine Rotor Burst’, Winter Meeting of ASME,New York, December 5–10, 1976, ASME New York NY, USA 1976.

35 Toughness of CrMoV-Steels for Steam Turbine Rotors, EPRI Special Report RD-2357, April, 1982.

36 Viswanathan R and Gehl S, ‘Temper embrittlement of rotor steels’, Robert I. JaffeeMemorial Symposium on Clean materials technology, ASM/TMS Materials Week,2–5 November 1992, Chicago, Illinois, USA, ASM International, Materials Park,Ohio 44073-0002, USA 1992.

37 Masuyama F, ‘Steam plant material development in Japan‘, 6th Liège COSTConference, Materials for Advanced Power Engineering 1998, Liège, Belgium,September 1998, Forschungszentrum Jülich GmbH, Germany 1998.

38 Husemann R U, ‘Werkstoffe und Werkstoffentwicklung für die KomponentenMembranwände und Überhitzerrohre für zukünftige Dampferzeuger’, Proceedingsof TAM-Fachtagung Kohlekraftwerke im Jahre 2000/2015, 30–31 März Dresden,expert verlag Renningen-Malmsheim, Germany 1995.

39 Fujita T and Takahashi N, ‘The effects of boron on the long period creep rupturestrength of the modified 12% chromium heat-resisting steel’, Transactions ISIJ,1976, 16, 606–613.

40 Fujita T, ‘Twenty-first century electricity generation plants and materials’, Proceedingsof International Workshop on Development of Advanced Heat Resisting Steels,Yokohama, Japan, 8 November, NIMS Tsukuba, Japan 1999.

41 Boyle C J and Newhouse D L, United States Patent 3, 139, 337, June, 1964.42 Curran R M, ‘Progress in the development of large turbine rotor forgings’, Proceedings

of the Fifth International Forgemasters Meeting, 6–9 May Terni, Italy, 1970.43 Brinkmann C R, Gieseke B, Alexander J and Maziasz P J, ‘Modified 9Cr-1Mo Steel

for Advanced Steam Generator Applications’, ASME/IEEE Power GenerationConference, 21–25 October, Boston, MA, USA, ASME/IEEE, ASME New York,NY, USA, 1991.

44 Mayer K H and Bakker W T, ‘New ferritic steels increase the thermal efficiency ofsteam turbines’, International Conference Joint Power Generation, Houston, Texas,USA, October 1996, ASME New York.

45 Thornton D V and Mayer K H, ‘New materials for advanced steam turbines’,Proceedings of the Fourth International Charles Parsons Conference, Advances inTurbine Materials, Design and Manufacturing, 4–6 November, Newcastle uponTyne, UK, Institute of Materials London, 1997.

46 Hizume A, Takeda Y, Yokota H, Takano Y, Suzuki A, Kinoshita S, Koono M andTsuchiyama T, ‘The Probability of a new 12%Cr Rotor Steel applicable for SteamTemperature above 593°C’, Journal of Engineering Materials and Technology/Transaction of ASME Chicago, USA, (109) 319–325, 1987.

47 Fujita T, Sato T and Takahashi N, ‘Effect of Mo and W on long term creep strengthof 12%Cr heat-resisting steel containing V, Nb and B’, Transaction Iron and SteelInstitute of Japan 1978 18, 115–124.

48 Berger C, Mayer K H, Scarlin B and Thornton D, ‘Improved ferritic rotor and caststeels for advanced steam power plants – a collaborate European effort in COST501’, 4th International EPRI Conference on Improved Coal-Fired Plants, 1–4 March,Washington DC, USA, Electric Power Reseach Institute Palo Alto CA, 1993.

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49 Orr J, Buchanan L W and Everson H‚ ‘The commercial development and evaluationof E 911, a strong 9%CrMoWVNbN steel for boiler tubes and headers’, InternationalConference Advanced Heat Resistant Steels for Power Generation, 27–29 April,San Sebastian, Spain, Electric Power Research Institute Palo Alto CA, 1998.

50 Masumoto H, Sakakibara M, Takahashi T, Sakurai H and Fujita T, ‘Development ofa 9%Cr–Mo–W steel for boiler tubes’, First International EPRI Conference onImproved Coal-Fired Plants, 19–21 November, Palo Alto/USA, Electric PowerResearch Institue Palo Alto CA, 1986.

51 Hald J, ‘ECCC E911-P92 assessment, ECCC Data Sheet’, published in September2005 at the ECCC Creep Conference, 12–14 September, London, UK, ERATechnology Ltd, Leatherhead, Surrey KT22 7SA, 2005.

52 Iseda A, Sawaragi Y, Kato S and Masuyama F, ‘Development of a new 0.1C–11Cr–2W–0.4Mo–1Cu steel for large diameter and thick wall pipe for boilers’, 5thInternational Conference on Creep of Materials, 18–21 May 1992, Lake BuenaVista, Florida, USA, ASM International, Materials Park, Ohio, 1992.

53 Kimura K, Assessment of Long-Term Creep Strength and Review of AllowableStress for High Cr Ferritic Creep Resistant Steels, ASME PVP2005-71039, ASMENew York NY, USA 2005

54 Strang A and Vodarek V, ‘Modelling and experimental verification of minor phasecomposition changes in creep resistant 12%CrMoVNb steels’, 5th InternationalCharles Parsons Turbine Conference, Advanced Materials for 21st Century Turbinesand Power Plants, 3–7 July 2000, Cambridge, UK, IOM Communication Ltd,London UK 2000.

55 Kagawa Y, Tamura F, Ishiyama O, Matsumoto O, Honjo T, Tsuchiyama T, ManabeY, Kadoya Y, Magoshi R and Kawi H, ‘Development and manufacturing of the nextgeneration of advanced 12%Cr steel rotor for 630°C steam temperature’, 14thInternational Forgemaster Meeting, Wiesbaden, Germany, September 3–8, GermanIron and Steel Institute, Düsseldorf, Germany, 2000.

56 Azuma T, Miki K and Tanaka Y, ‘Effect of boron on microstructural change duringcreep deformation in 12%Cr heat resistant steel’, 3th EPRI Conference on Advancesin Materials Technology for Fossil Power Plants, Swansea, UK, April 2001, ElectricPower Research Institue Palo Alto, CA.

57 Fukuda M, Tsuda Y,Yamashita K and Shinozaki Y, Takahashi, ‘Materials and designfor advanced high temperature steam turbines’, 4th EPRI International Conferenceon Advanced in Material Technology for Fossil Power Plants, Hilton OceanfrontResort, Hilton Head Island, SC, USA, October 25–28, 2004, Electric Power ResearchInstitute Palo Alto CA, 207–221.

58 Kern T U, Staubli M, Mayer K H, Donth B, Zeiler G and DiGianfrancesco A, ‘TheEuropean effort in development of new rotor materials – COST 536’, 8th LiègeConference on Materials for Advanced Power Engineering, 18–20 September, Liège,Belgium, Forschungszentrum Jülich, Germany, 2006.

59 Kern T U, Mayer K H, Berger C, Zies G and Schwienheer M, ‘Stand derEntwicklungsarbeiten in COST 522 für Hochtemperatur-Dampfturbinen 27’,Vortragsveranstaltung FVHT, 26 November, German Iron and Steel Institute,Düsseldorf, Germany, 2004.

60 Arai M, Doi H, Azuma T and Fuita T, ‘Development of high WCoB-containing12Cr Rotor Steels for use at 650°C in USC Power Plants’, 15th InternationalForgemasters Meeting IFM 2003, Kobe, Japan, October 26–29, 2003, Steel Castingsand Forgings Association of Japan, 261–268.

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61 Kauffmann F, Mayer K H, Straub S, Zies G, Scheu C, Willer D, Ruoff H and MaileK, ‘Characterization of the precipitates in modern boron containing 9–12%Cr steelsdeveloped in the frame of the COST program’, 8th Liège Conference on Materialsfor Advanced Power Engineering, 18–20 September, Liège, Belgium,Forschungszentrum Jülich, Germany, 2006.

62 Kager F, Böck N, Spiradek-Hahn K, Höfinger S, Brabetz M and Zeiler G, ‘Superiorlong-term creep behaviour and microstructure evolution of 9%Cr-steels with boron’,8th Liège Conference on Materials for Advanced Power Engineering, 18–20September, Liège, Belgium, Forschungszentrum Jülich, Germany, 2006.

63 Miki K, Azuma T and Ishiguro T, ‘Improvement of long term creep strength in highCr heat resistant steel’, 15th International Forgemasters Meeting IFM 2003, Kobe,Japan, October 26–29, 2003, Steel Castings and Forgings Association of Japan,269–275.

64 Scarlin B and Stamatelopoulos G N, ‘New boiler materials for advanced steamconditions’, Proceedings of the 7th Liège COST Conference on Materials for AdvancedPower Engineering 2002, Part III, Liège Belgium, Forschungszentrum Jülich,Germany, 2002, 1091–1108.

65 Bendick W, Vaillant J C, Vandenberghe B and Lefebre Bo, ‘VM12 – a new 12%Crsteel for boiler tubes, headers and steam pipes in USC power plants’, ‘Developmentof a new 12%Cr-steel for tubes and pipes in power plants up to 650°C’, EPRIInternational Conference on Materials and Corrosion Experiences for Fossil PowerPlants, 18–21 November, Wild Dunes Resort, Isle of Palms, SC, USA, HiltonOceanfront Resort, Hilton Head Island, SC, USA, October 25–28, 2004, ElectricPower Research Institute Palo Alto CA, 2003.

66 Masuyama F, ‘Steam plant material development in Japan’, 6th Liège COSTconference, Materials for advanced power engineering 1998, Conference proceedingsIII, Liège, Belgium, September 1998, Forschungszentrum Jülich GmbH, Germany,1807.

67 Bendick W, ‘Neue Werkstoffentwicklung für moderne Hochleistungskraftwerke’,VGB Tagung Werkstoffe und Qualitätssicherung 2004, 10–11 März, 2004.

68 Blum R, Vanstone R W and Messlier-Gouze C, ‘Materials development for boilersand steam turbines operating at 700°C’, 4th EPRI International Conference onAdvanced in Material Technology for Fossil Power Plants, Hilton Oceanfront Resort,Hilton Head Island, SC, USA, October 25–28, 2004, Electric Power ResearchInstitute Palo Alto CA, 118–138.

69 Masuyama F in Landolt–Boernstein, Numerical Data and Functional Relationshipsin Science and Technology, Group VIII, Vol. 2 Materials Subvol. B, Creep Propertiesof Heat Resistant Steels and Superalloys, ed by Yagi K, Merckling G, Kern T U, IrieH, Warlimont H, Springer Verlag, Berlin Germany 2004.

70 Danielsen H and Hald J, Z-phase in 9-12%Cr Steels, VÄRMEFORSK Service AB,Stockholm, Sweden, April 2004.

71 Sato A, ‘Research project on innovative steels in Japan’, International workshop onInnovative Structural Materials for Infrastructure in 21th Century, Tsukuba, NIMSTsukuba, Japan, 12–13 January 2000.

72 Abe F, Igarashi M, Fujitsuna N and Muneki S, ‘Alloy design of advanced ferriticsteels for 650°C USC boilers’, Conference on Advanced Heat Resistant Steels forPower Generation, 27–29 April San Sebastian, Spain, Electric Power ResearchInstitute Palo Alto CA, 1998.

73 Semba H and Abe F, ‘Development of creep resistant 9Cr-3W-3Co steel containing

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high boron for USC boilers’, First International Conference Super-High StrengthSteels, 2–4 November, Centro Sviluppo Metallurgia, Rome, Italy, 2005.

74 Abe F, ‘High performance creep resistant steels for the 21st century power plants’,First International Conference Super-High Strength Steels, 2–4 November, CentroSviluppo Metallurgia, Rome, Italy, 2005.

75 Krainer H, ‘50 Jahre nichtrostender Stahl’, Stahl und Eisen, 1962, 82, 1527.76 Strauses B, ‘Non-rusting Cr-Ni steels’, Proceedings American Society Test Materials,

1924, 24, 208.77 Schottky H, ‘German version of birth of stainless steels’, Iron Age, 1929, December

5, 1512.78 Mitchel W M, ‘18-8 and related stainless steels’, Metals and Alloys, 1940, January,

14.79 Houdremont E and Schafmeister P, ‘Verhütung der Korrosion bei Stählen mit 18%Cr

und 8%Ni’, Archiv. Eisenhuettenwesen, 1933, 7, 187.80 Suzuki T, Historical Invent of Stainless Steels (in Japanese), AGNE Technical

Center, Tokyo, 2000, 110.81 Eberle F, Ely E G and Dillon J S, ‘Experimental superheater for steam at 2000Psi

and 1250F – Progress report after 12 000 h of Operation’, Trans. ASME, 1956, 76,665.

82 Masuyama F, Materials for Advanced Power Engineering 1998, Part III, J. Lecomte-Beckers et al. (eds), Juelich GmbH, Forschungszentrum, Juelich, 1998, 1807.

83 Murray J D, ‘Welding Trails on Esshete 1250’, Welding and Metal Fabrication,1962, 9, 350.

84 Minami Y, Kimura K and Tanimura M, ‘Creep rupture properties of 18%Cr-8%Ni-Ti-Nb and Type 347H Austenitic Stainless Steels’, ASM International ConferenceNew Development in Stainless Steel Technology, Detroit, Michigan, September 17–21, ASM International, Materials Park, Ohio, 1984.

85 Yoshikawa K, Fujikawa H, Teranishi H, Yuzawa H and Kubota M, ‘Fabrication andprogress of corrosion-resistant TP347H stainless steel’, ASM Conference Coatingsand Bimetallic for Aggressive Environments, ed. by R. D. Sission Jr.1985 99, ASMInternational, Materials Park, Ohio 44073 1984.

86 Sawaragi Y, Ogawa K, Kan S, Natori A and Hirano S, ‘Development of the economical18-8 Stainless Steel (SUPER 304) having elevated Temperature Strength for FossilPower Boilers’, Sumitomo Search, 1992, 48, 50.

87 Ishituka T and Mimura H, ‘Development of New 18Cr-9Ni austenitic stainless steelboiler tube’, 7th Liège COST Conference Materials for Advanced Power Engineering2002, September 2002, Liège Belgium, J. Lecomte-Beckers et al. (eds), JuelichGmbH, Forschungszentrum, Juelich, 2002, 1321.

88 Minami Y, Tohyama A and Hayakawa H, ‘Properties and experiences with a newaustenitic stainless steel (Tempaloy AA-1) for Boiler Tube Application’, 7th LiègeCOST Conference Materials for Advanced Power Engineering 2002, September2002, Liège Belgium, J. Lecomte-Beckers et al. (eds), Juelich GmbH,Forschungszentrum, Juelich, 2002, 1445.

89 Masuyama F, ‘History of power plants and progress in heat resistant steels’,Transactions of Iron Steel Institute Japan, ISIJ International, 2001, 41, 612.

90 Shinoda T and Tanaka R, ‘Role of carbide precipitations on creep rupture strengthof austenitic stainless steels (in Japanese)’, Bulletin Japan Institute Metals, 1972,11, 180.

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91 Masuyama F, ‘Advanced power plant development and material experiences inJapan’, 8th Liège conference on Materials for Advanced Power Engineering, J.Lecomte-Beckers et al. (eds), Juelich GmbH, Forschungszentrum, Juelich, 2006,Liège, Belgium, 175.

92 Sawaragi Y, ‘Elevated temperature strength and microstructure of several highstrength austenitic steels’, JSPS, 123 Committee Report on Heat Resistant MetallicMaterial, 1985, 26, 369.

93 Tamura M, Yamaguchi N, Tanimura M and Murase S, ‘Processing alloys for theheat exchangers of advanced coal fired boilers’, 1985 Exposition and Symposium,Industrial Heat Exchanger Technology, Pittburgh, PA, 1985.

94 Sawaragi Y and Yoshikawa K, ‘High temperature strength and microstructure ofhigh strength and corrosion resistant austenitic steel for boiler (in Japanese)’, Tetsu-to-Hagane, 1986, 72, S672.

95 Sawaragi Y, Teranishi H, Makiura H, Miura M and Kubota M, ‘The development ofHR3C steel with high elevated temperature strength and high corrosion resistancefor boiler tubes (in Japanese)’, Sumitomo Metals, 1985, 37, 166.

96 Kikuchi M, Sakakibara M, Otoguro M, Mimura H, Araki S and Fujita T, ‘Anaustenitic heat resisting steel tube developed for advanced fossil-fired steam plants’,International Conference High Temperature Alloys, Petten, Netherlands, 1985.

97 Toyama A, Minami Y and Yamada T, ‘Effect of alloying elements on high temperatureproperties in an austenitic stainless steel containing high Cr (in Japanese)’, CAMP-ISIJ, 1988, 1, 928.

98 Semba H, Igarashi M and Sawaragi Y, ‘Development of 23Cr–18Ni–3Cu–1.5W–Nb–N Austenitic Steel for USC Boilers’, Proceedings International ConferencePower Engineering–97, Volume 2, JSME, Tokyo, 1997, 125.

99 Rautio R and Bruce S, ‘Sandvik Sanicro 25, a new material for ultrasupercriticalcoal fired boiler’, Advances in Material Technology for Fossil Power Plants, R.Viswanathan et al. (eds), 2005, ASM International, Metals Park, OH, 274.

100 Masuyama F, in Landolt-Boernstein, Numerical Data and Functional Relationshipsin Science and Technology, Group VIII, Volume 2, Materials Subvolume B, CreepProperties of Heat Resistant Steels and Superalloys, Yagi K, Merckling G, KernT U, Irie H, Warlimont H, Springer Verlag, Berlin Germany 2004.

101 Kalwa G, Haarmann K and Janssen J K, ‘Experiences with ferritic and martensiticsteels tubes and piping in nuclear and non-nuclear applications’, Topical ConferenceFerritic Alloys for Use in Nuclear Energy Technology, Metallurgical Society AIME,Snowbird, ed. by J W Davis and D J Michel, 1983.

102 Sikka V K, Ward C T and Thomas K C, ‘Modified 9Cr-1Mo Steel – An ImprovedAlloy for Steam Generator Application’, ASM International Conference Production,Fabrication, Properties and Applications of Ferritic Steels for High-TemperatureApplications, Warrendale, PA, ASM International, Materials Park, Ohio, 1981.

103 Sakakibara M, Masumoto H, Ogawa T, Takahashi and Fujita T, ‘High Strength 9Cr-0.5Mo-1.8W(NF616) Steel for Boiler Tube (in Japanese)’, Thermal and NuclearPower, 1987, 38, 841.

104 DeLong W T, Ostrom G A, Szumachowski E R, ‘Mearsurement and Calculation ofFerrite in Stainless Steel Weld Metal’, Weld Journal, 1956, 35, 521S.

105 Masuyama F, Nishimura N, Haneda H, Ellis F V and DeLong J F, ‘Investigation ofthe Deterioration plus Restoration Behaviour of fourteen Heats of TP316 StainlessSteel removed from Eddystone Unit No 1 Main Steam Lines after 130 520 h

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Service’, Properties of Stainless Steels in Elevated Temperature Service, MPC-Vol.26/PVP-Vol 32, 1987, 173.

106 Yamamoto R, Kadoya Y, Ueta S, Nobe T, Magoshi R, Nishimoto S and Nakano T,‘Development of wrought Ni-based Superalloy with low Thermal Expansion for700°C Steam Turbines’, Advances in Material Technology for Fossil Power Plants,R. Viswanathan et al. (eds), 2005, ASM International, Metals Park, OH 623.

107 Imano S, Doi H and Iijima E, ‘Modification of Ni-Fe base Superalloy for SteamTurbine Applications’, Advances in Material Technology for Fossil Power Plants,R. Viswanathan et al. (eds), 2005, ASM International, Metals Park, OH, 575.

108 Imai K, private communication, 2006.109 Takeda Y and Masuyama F, ‘Heat Resistant Materials for Ultra Super Critical

Power Plants (in Japanese)’, JSPS 123 Committee Report on Heat Resistant MetallicMaterial, 1987, 29, 399.

110 Furuya K, Hizume A, Takeda Y, Fujikawa T, Fujita A, Kinoshita S, Kohono M,Honjo T and Suzuki A, ‘Austenitic Steel Rotor Forging for EPDC’s WAKAMATSU50MW High Temperature Turbine Project’, Proc. 2nd Int. Conf. Improved Coal-Fired Power Plants, Palo Alto CA, USA 1989, Electric Power Research Institute,1989, 59–1.

111 Ohta S, ‘Heat Resistant Steels for Chemical Industries (in Japanese)’ JSPS 123Committee Report on Heat Resistant Metallic Material, 1977, 18, 383.

112 Koori M, Yosida T and Ohta S, ‘Analysis and Prevention of Failures in SteamReforming Furnace (in Japanese)’, Kobe Steel Technical Bulletin, 1983, 33, 65.

113 Weißbach W, Werkstoffkunde und Werkstoffprüfung, Vieweg & Sohn, Braunschweig/Wiesbaden, Germany 1979.

114 Bodnar R L and Cappepellini R F, ‘Effects of residual elements in heavy forgings,past, present, future’, Research Department Publication Bethlehem Steel Corporation,Bethlehem, PA, Philadelphia, USA, 1986.

115 Schmollgruber F, Verfahrenswege zur Herstellung grosser Schmiedestücke und derenqualitative und wirtschaftliche Auswirkungen, Dissertation TH Aachen, February1974.

116 Nutting J and Viswanathan R, ‘Clean steel: superclean steel’, Conference on CleanSteel: Superclean Steel, Conference Proceedings, 6–7 March 1995, London, UK,1995.

117 Viswanathan R, ‘Clean steel technology’, Proceedings of the R.I. Jaffee MemorialSymposium, 2–5 November 1992, Chicago, Illinois, USA, ASM International,Materials Park Ohio, 1997.

118 Azuma T, Tanaka Y, Ikeda Y and Yoshida H, ‘Production and properties of superclean3.5%NiCrMoV rotor forgings for low pressure steam turbine’, Proceedings of theR.I. Jaffee Memorial Symposium, 2–5 November 1992, Chicago, Illinois, USA,1992, ASM International, Materials Park Ohio, 213–220.

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78

3Specifications for creep-resistant steels:

Europe

G . M E R C K L I N G, RTM BREDA Milano, Italy

3.1 Introduction

Creep problems with metallic materials emerged simultaneously in the USAand in Europe in the 1920s. In both cases, efforts to increase efficiency hadled to higher service temperatures in pressure and other load-bearing systems.Unfortunately, on both sides of the Atlantic ignoring creep effects resulted inseveral casualties and in some cases severe injuries were sustained.

As a consequence, researchers in the USA and the UK started to investigatethe effect and the first publications describing creep appeared in themid-1920s.1–3 The search for a simple answer to creep, encouraged bythe decreasing strain rate of the primary creep stage, suggested to earlyresearchers that there was a creep threshold stress, that is a stress levelbeyond which no creep, or at least no creep fracture, would occur. A particulartesting type4 was created in Germany in the early 1930s. Criteria that helpedidentify ‘creep-resistant’ materials at a given temperature were defined byassuming that materials fulfilling these properties (creep strain rate lowerthan 10–3% h–1 between 25 and 35 h of testing and creep strain smallerthan 0.2% in 45 h) would not fail and were applied extensively in parts ofEurope. Again, it was casualties that proved this assumption wrong. It thusbecame apparent that expensive and difficult long-term creep investigationswere needed to guarantee operational safety and plant reliability at hightemperatures.

All the major European countries developed activities to study, understandand predict creep strength and deformation behaviour. Building on efforts inthe 1930s, research was re-activated after World War II and has continuedever since. Although aimed at the same target – for the best use of materialsat the best safety levels with the available technology – different approachesto the investigation of creep were developed. Essentially three differentdirections were taken, although other interpretations and sub-fields grewwith time:

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Specifications for creep-resistant steels: Europe 79

1. German industry developed a common effort approach from the outset,which, coordinated by an umbrella organisation, pooled experiences anddata from most industrial companies, research and university institutions,including very long-term creep results. As a result, a huge number ofsteels were characterised, optimised and developed for application increep regime, and creep strength data containing DIN standards wereissued. From this activity, perhaps the only raw long-term creep datacollation was published in 19695 (see also Table 3.1), and this formedthe basis of the publicly available data pool throughout Europe for along period. In parallel, technical rules for material properties, weldingand component design were issued in the TRD and AD series.

2. In the United Kingdom, creep investigation efforts were more or lessconfined to big industrial companies operating independently, with lessnational coordination. Nevertheless, here also a large series of materialswere developed on a company-to-company basis for use in creep regimes,accompanied by the related BS and PD standards and rules for creepstrength data assessment, welding and component design.

3. A third approach was followed in the 1950s by some of the lessindustrialised countries, like Italy, who had fewer resources to developtheir own materials and/or rules and so relied on external sources. Italiandesign and material standards in creep regimes were based on a uniquemixture of European (German and British) and US research, assessedto suit the local circumstances. In these countries, US materialnomenclature, design criteria and codes were used as widely as Europeanspecifications.

Other countries, like Sweden and France, also developed a significant amountof technical documentation for material applications in creep regimes.

Further developments followed the introduction of nuclear power to Europe,mainly concerning austenitic materials, nuclear pollution and the effects ofirradiation. As this a very specialised area, the findings will not be discussedhere, although it is acknowledged that nuclear research produced valuabledata, some of which was made available for the development of new EuropeanStandards and joint assessment activities in the European Creep CollaborativeCommittee (ECCC) (see Section 3.2.5).

With the introduction of the Common Market in Europe in the 1960s andthe increasing importance of exports within and outside Europe, several pan-European institutions began attempting to harmonise the various differentstandards. The European Carbon and Steel Collaboration programmes CECA/ECSC6 had already begun this process in the 1950s and continued into the1960s. Since 1971, most major work has been coordinated through theprogrammes of the organisation for European Cooperation in the field ofScientific and Technical Research (COST)7 and these have contributed

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significantly to innovation and the development and characterisation of newmaterials for high temperature applications, such as rotor steels, advancedpiping materials, welding techniques, and so on. More focused front-lineresearch has been concentrated in the EC funded Brite-Euram-Projects sincethe 1980s and technical, pre-competitive development of future technologicalsolutions has been through various Thermie-Projects,8 which began in 1993.All these activities have helped European industry and research specialists toexchange their experience and knowledge and have encouraged the formationof a strong core of organisations collaborating while competing in the hightemperature field.

On the standards development side, right from the beginning the EuropeanCommunity tried to find common denominators for relevant technical issuesthat could be published as Euronorms (abbreviated as EU). However, thefirst series of Euronorms were not very well accepted, did not gain generalapplication and made little difference to the use of independent nationalstandards.

The situation changed when the now European Union started to issue‘directives’, which the member countries had agreed would become law andso would formalise voluntary industrial standardisation. For the hightemperature sector, the relevant event was the publication in 1997 andmandatory introduction in 2002 of the Pressure Equipment Directive (PED)97/23/EC,9 which required guarantees for the in-service performance ofmaterials and, as a consequence, required harmonisation of the relevant ENstandards. Although not dealing specifically with pressure vessels under acreep regime, the PED required a new series of standards harmonised at alllevels for high-temperature materials, for which ‘European’ creep strengthvalues were to be defined. The PED came from the European Committee forStandardisation (CEN) and its related technical committee, the EuropeanCommittee for Iron and Steel Standards (ECISS) and, under the frameworkof the European Community funded research programmes, encouraged alarge group of industries, research and inspection organisations to form:

• the European Pressure Equipment Research Council (EPERC), whichstrongly contributed to implementation and clarification of the PEDdirective itself; and

• the European Collaborative Creep Committee (ECCC), which developedcreep strength data for some of the materials under the new harmonisedEN standards, using new data assessment approaches.

The following sections and tables will try to give a concentrated overview ofthe most relevant standard publications related to creep-resistant steels suitablefor structural applications, beginning with the original national grades up tothe current PED-harmonised standards.

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3.2 Specifications and standards

3.2.1 Institutions, publishing rules and standards

Each European country publishes norms for voluntary industrial standardisation,which are issued by various organisations, as follows:

• The national standardisation body (DIN in Germany, BS in the UnitedKingdom, AFNOR in France, UNI in Italy, etc.) is responsible for preparingand issuing voluntary industrial standardisation documents.

• Some of the national standardisation bodies, their sub-organisations orumbrella organisations also issue ‘second-level standardisation documents’,such as PD documents in the United Kingdom, SEP documents in Germany,and so on.

• Ad hoc industrial committees, in some cases affiliated to the nationalstandardisation body, issue design codes for particular areas of technicalinterest. For example, for pressure application in creep regimes, designrules are issued by TRD and AD in Germany, CODAP in France, CTIwith ANCC/ISPESL in Italy, and so on. Often these codes do not justspecify a national standard for materials, but add additional prescriptions,controls or requirements.

• In some countries, inspection authorities have issued additional material-related specifications (e.g. the ‘Raccolta M/S’ of the Italian AuthorityISPESL) or have required qualifications for suppliers of materials to beused in creep regime (e.g. VdTÜV-Richtlinien in Germany).

• A few major users or builders of equipment operating under creep regimeshave developed design and material qualification guidelines (e.g. NationalElectricity Boards, refinery and petrochemical groups, turbine and boilermanufacturers, etc).

Since the PED Directive was issued in 1997, standards across Europehave tended towards unification. Standards issued by CEN, the EuropeanStandardisation Body, are jointly prepared by specialists from all CEN membercountries and some which are not yet members, and are quickly substitutedfor the corresponding national standards. Under the CEN framework10 forsteels ECISS11 is responsible for issuing harmonised standards. Experts fromindustry and research, delegated by national standards bodies, are recruitedinto Technical Committees and advisory organisations provide specificinformation as required, or, as is often the case, via their specialists beingmembers of both the relevant CEN/ECISS Technical Committee issuing thestandard and the advisory organisation (e.g. ECCC, EPERC, etc).

The European Commission itself is ultimately in charge of defining EU-wide regulations in the form of directives, with the objective of standardisingapproaches and methods, both technical and administrative, within the membercountries to simplify and encourage transport, trade, exchange and

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implementation of goods. EU directives normally become law in all EUmember countries within 12 to 18 months, superseding previous locallegislation. Technical directives often call for common standardisation orrefer directly to EN standards, significantly increasing the importance ofCEN documents and their influence on industry and administration. Formetallic materials used in creep regimes, the issue of the PED in 1997followed by mandatory application starting in 2002 was a critical step.

3.2.2 The pressure equipment directive

The PED9 deals with all relevant safety issues in the design, manufactureand use of components designed to resist stresses induced mainly by pressure.The directive is aimed at simplifying intra-European trade and exchange ofpressure handling equipment and harmonising safety prescriptions in orderto allow intra-European exchangeability and serviceability of pressure devices.There are many types of boilers, pipelines, furnaces, heat exchangers,incinerators, chemical reactors, etc within Europe, and all have to conformto the directive. Explicitly excluded are all equipment regulated by otherdirectives (e.g. pressure equipment for transportation on rail or road, ‘simplepressure equipment’, etc), and all types of rotating machines and motors(e.g. combustion engines, turbines and pumps, etc.).

The PED is a general law (not a design code) and does not directlyconsider specific design and safety factors for components operating undercreep regimes, although creep is mentioned and regarded as a non-negligibleissue in designing. For dimensioning, including equipment operating in creepregime, PED allows the use of any internationally recognised standard andaccepts fracture mechanics approaches and finite element simulation, but formaterials, the rules are more strict:

• All materials must have a guaranteed minimum impact energy at thelowest possible service temperature that the equipment for which theyare used may encounter under pressure, e.g. including hydraulic tests.

• All materials must demonstrate ‘sufficient’ ductility in service conditionsto guarantee ‘leak before break’.

• Materials need to be selected from harmonised standards, i.e. materialstandards that in their intention and scope consider and include thesafety criteria as set forth by the PED. This requirement strongly enhancedand accelerated the production of EN standards incorporating Europe-wide agreed strength values. If materials cannot be selected from ENstandards (typically because they come under non-European design codes,e.g. where ASME Boiler and Pressure Vessel Code or API standards areused), they must either pass a general acceptance procedure conductedby an accredited, notified body, or they must undergo a ‘particular material

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appraisal’ to check their suitability for the application and their conformityto the PED safety criteria, on a case-by-case basis.

• A final, and in the creep regime very relevant, PED requirement is thatall material producers involved in the production chain of a givencomponent must declare and take responsibility that the material supplied(possibly with provisos relating to particular handling, assembly or designoperations) is suitable for the intended application and its operatingconditions. Alternatively, in a somewhat milder interpretation, producersmust certify that the material is indeed fully compliant with the purchasespecifications and the design-relevant material properties so that it canbe guaranteed for safe service. For materials and welds in the creepregime this requires the material, shape, fitting and so on producer toguarantee:

– long term strength of the base material– long term strength of the construction welds– long term strength of the assembly welds– long-term strength of any repair (if allowed by the design code) of

any of the items listed above.

The easiest way to comply is to select materials that conform with theharmonised standards, for which long-term strength values are indicatedin the material norm. Materials subjected to the ‘particular materialappraisal’ (typically ASTM or API grades, sometimes Japanese materials)have to demonstrate the points listed above, which can be achieved bytesting the actual material or by presenting previous data for the samegrade and producer. Owing to the expected long service duration ofmodern plants (generally 200 000–250 000 h, i.e. 25 to 35 years), thesedemonstrations often become extremely complex for welds, repairs andmultiple repairs.

Another route that is applied by some notified bodies and/or end userinspectorates is to qualify material producers as PED-compliant usingstringent procedures when they first supply materials under the aegis ofthis notified body. This first-time qualification should include creepstrength verification and demonstration through a reduced creep campaign,which, depending on the notified body/inspectorate, the type of materialand the scope of supply, may include 2–3 isotherms made or composedfrom of three to five points each and durations of generally not less than10 000 h, and often not less than 30 000 h.

3.2.3 Non PED creep applications

Equipment not subjected to PED regulations or the particular EU directivesfor dangerous goods and fluids, transportation or general safety in the work

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place, is built under the sole responsibility of its manufacturer. For hightemperature applications, this is mainly the case for manufacturers of gasand steam turbines, compressors, pumps, geothermics, some types of well-drilling equipment, combustion engines, aircraft engine components and, insome cases, electrical or railway equipment.

Creep data and strength determination in Europe in the past were stronglydriven by turbine manufacturers. Based on their own large data collections,turbine manufacturers imposed creep testing and, less frequently, creep strengthproperties on components they bought from forges, steel works and foundries.Typical specifications included quite severe prescriptions and sometimesrange limitations in chemical composition, stringent heat treatment detailsand control and mechanical testing, but generally only specifying ‘creepquality testing’ by means of one or a small group of short-term (24–100 h)stress rupture or time-to-rupture (see ASTM E292) tests. The turbinemanufacturers themselves then often carried out large-scale, long-durationtests, so that the assessment of the creep properties of a single material batchwas fully under their control.

3.2.4 European national standards reporting creep data

Before the era of CEN and the PED directive, creep rupture strengths weregenerally:

• assessed within the user companies (turbine and boiler manufacturers,or big end-users such as utilities, refineries and the chemical industry);

• supplied to support sales between companies (typically from steel makeror tube/pipe manufacturer to user);

• assessed by national associations (often under the control of a nationalelectricity board acting as the main national customer).

Generally such assessments were not made available to the wider publicand were confined to the members of the working group collating and preparingthe strength values. In some cases, these were one-off activities, that is thecollated or produced data were assessed by the working group and thendisappeared into the databases of the companies. In other cases, fortunately,centralised databases were created which allowed data to be updated andstrength values upgraded as the information available on creep expanded(for instance AGW within VDEh in Germany).

National standardisation bodies made several attempts to include creepvalues into standards and rules, but their success was quite limited becauselong-term data were expensive to obtain and considered commercially sensitive,so they were often not made available to larger groups for common assessment.Nevertheless some of the older standards include creep strength information,sometimes as an informative appendix, sometimes as a true part of the standard.

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Table 3.1 gives some of the older standards, now mainly discontinued,including creep strength information. Table 3.2 gives examples of creep datafor tubes and pipes in the main steel grades included in the older standards.This table also shows the difficulties encountered in achieving commonstandardisation, because the correlation between steel grades of differentcountries is not always obvious and their application and heat treatmentconditions are sometimes quite different. It is relatively obvious that in spiteof the national preferences for using particular steel grades as recommendedby national producers, the main national standards also include grades for‘foreign’ materials, although sometimes with quite significant differences inguaranteed/expected creep strength values.

3.3 The European Creep Collaborative Committee

(ECCC)

3.3.1 Motivation and history

With advancing integration within Europe, in 1991 a large number of industrialand research companies from the United Kingdom, Germany, Italy, France,Denmark, Switzerland, Austria, The Netherlands and Belgium founded theEuropean Creep Collaborative Committee (ECCC).12 The common purposesof this organisation were laid down in a Memorandum of Understanding(MoU) that was jointly agreed, signed by all members and is still the foundationfor cooperation. The MoU established that ECCC is a voluntary associationof organisations representing their nations, led by industry, with the aim ofimproving, reinforcing and enhancing the position of European industry inthe high temperature application market. The companies participating inECCC cover the whole scope of industries and research organisations dealingwith high temperature problems: steel makers, steel product manufacturers,boiler, turbine, plant and equipment builders, utilities and end users, inspectionbodies, research institutes, technical universities, testing houses.

Currently, the ECCC represent 14 nations, which include, besides thosealready mentioned above, Sweden, Finland, Portugal, Czech Republic andSlovakia. More than 50 organisations from the 14 countries contribute toECCC activities. The MoU defines detailed targets to be addressed by ECCCto support European standardisation and research in creep. ECCC and itsmembers:

• collate, exchange and jointly assess creep data;• support European standardisation;• coordinate European creep data generation;• develop common creep research and testing programmes;• mutually exchange information on HT material development; and• define common procedures for data generation and assessment.

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86Table 3.1 Former national European Standards including creep strength values (examples)

Designation Nation Title Date of issue Status

BS 1501 UK Plates, Specification for carbon, alloy and austenitic stainless steel tubes 1988/90 Discontinuedwith specified elevated temperature properties

BS 1502 UK Section and bars: Specification for carbon, alloy and austenitic stainless 1982 Discontinuedsteel tubes with specified elevated temperature properties

BS 1503 UK Forgings: Specification for carbon, alloy and austenitic stainless steel 1989 Discontinuedtubes with specified elevated temperature properties

BS 1506 UK Bars for bolting: Specification for alloy and austenitic stainless steel 1990 Discontinuedtubes with specified elevated temperature properties

BS 3059-2 UK Steel boiler and superheater tubes. Specification for carbon, alloy and 1990 Discontinuedaustenitic stainless steel tubes with specified elevated temperatureproperties

BS 3602 UK Specification for carbon steel pipes and tubes with specified room 1987 Discontinuedtemperature properties for pressure purposes

BS 3604 UK Steel pipes and tubes for pressure purposes: ferritic alloy steel with 1991 Current (onlyspecified elevated temperature properties. Specification for longitudinally part 2 –arc welded tubes welded tubes)

BS 3605 UK Austenitic Stainless steel pipes and tubes for pressure purposes 1991 DiscontinuedDIN 17175 D Nahtlose Rohre aus warmfesten Stählen; Technische Lieferbedingungen, 1979 DiscontinuedDIN 17176 D Nahtlose kreisförmige Rohre aus druckwasserstoffbeständigen 1990 Discontinued

Stählen; Technische LieferbedingungenDIN 17177 D Elektrisch pressgeschweißte Rohre aus warmfesten Stählen; Technische 1979 Discontinued

LieferbedingungenDIN 17240 D Heat resisting and highly heat resisting materials for bolts and nuts 1976 DiscontinuedDIN 17245 D Ferritic Steel Castings Creep Resistant at elevated Temperatures: 1977 Discontinued

technical Conditions for DeliveryDIN 17458 D Seamless circular austenitic stainless steel tubes subject to special 1985 Discontinued

requirements: Technical delivery conditions

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Specifications for creep-resistant steels: E

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DIN 17459 D Nahtlose kreisförmige Rohre aus hochwarmfesten austenitischen 1992 DiscontinuedStählen: Technische Lieferbedingungen

DIN 17460 D Hochwarmfeste austenitische Stähle. Blech, kalt- und warmgewalztes 1992 DiscontinuedBand, Stabstahl und Schmiedestücke

DIN 17465 D Heat resistant steel castings: Technical conditions for delivery 1977 DiscontinuedDIN 17470 D Heating conductor alloys : Technical delivery conditions for round and 1984 Current

flat wireNF A36-209 F Produits Sidérurgiques : Tôles en aciers inoxydables austénitiques et 1990 Discontinued

austéno-ferritiques pour chaudières et appareils à pressionNF A49-213 F DiscontinuedNF A49-215 F DiscontinuedNF A49-219 F DiscontinuedBS PD 6525 UK Elevated temperature properties for steels for pressure purposes 1990 CurrentUNI 5462 I Tubi di acciaio senza saldatura: Tubi per caldaie, per apparecchi, per 1964 Discontinued

tubazioni di impianti termici ad alte temperature ed alte pressioni –qualità, prescrizioni e prove

UNI 6904 I Tubi senza saldatura di acciaio legato speciale inossidabile resistente alla 1971 Discontinuedcorrosione e al calore.

UNI 7660 I Prodotti finiti di acciaio fucinati, per recipienti a pressione. Qualità, 1977 Discontinuedprescrizioni e prove.

VDEh D Ergebnisse deutscher Langzeitversuche 1961 na

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88Table 3 2 Example of creep data availability in some national European Standards for carbon, low and high alloyed ferritic steels for boiler tubes and pipes (all mentioned national

standards discontinued in the meantime)

EN 10216 Heat BS 3059-2 United Kingdom NF A49-213, France DIN 17175 Germany UNI 5462 Italy

Material Treat- BS 3604

grade ment Creep rupture strength 215,219 Creep rupture strength DIN 17176 Creep rupture strength Ceep rupture strength

Designation T range Duration Designation T range Duration Designation T range Duration Designation T range Duration

(°C) (h) (°C) (h) (°C) (h) (°C) (h)

P235GH N 33 350–500 100k St35.8 380–480 10k, 100k C14 390-520 100k200k

P265GH N 45 350–500 100k St45.8 380–480 10k, 100k C18 390-520 100k8 CrMo 4 5 NT 200k8 CrMo 5 5 NT TU 10CD5-05 500–575 10k, 100k13 CrMo 4 5 N 620–460 450–630 10k, 30k,

50k, 100k,150k, 200k,250k*

13 CrMo 4 5 NT 620–440 450–630 10k, 30k,50k, 100k,150k, 200k,250k* TU 13CD4-04 450–560 10k 100k 13 CrMo 4 4 450–570 10k, 100k 14 CrMo 3 450–600 100k

13 CrMo 5 5 NT 621 450–630 10k, 30k, 200k50k, 100k,150k, 200k,250k*

X 11 CrMo 5 A 625 450–650 10k, 30k,50k, 100k,150k, 200k,250k* TU Z12CD5-05 500–600 10k, 100k 12 CrMo 19 5 480–650 10k, 100k,

200kX 11 CrMo 5 NT QT TU Z12CD5-05 500–600 10k, 100k 12 CrMo 19 5 350–650 10k, 100k,

200k

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16 Mo 3 N 243 450–550 10k, 30k,50k, 100k,150k, 200k,250k* TU 15 D3 450–550 10k, 100k 15 Mo 3 450–500 10k, 100k 16Mo5 450–550 100k

11 CrMo 9 10 NT 622 460–610 10k, 30k, 200k50k, 100k,150k, 200k,250k* TU 10CD9-10 500–625 10k, 100k 10 CrMo 9 10 450–600 10k, 100k, 12 CrMo 470–620 100k

200k 9 1011 CrMo 9 10 QT 12 CrMo 9 10 400–520 10k, 100k,

200k12 CrMoV 6 2 NT 660 450–600 10k, 30k,

50k, 100k,150k, 200k,250k* 14 MoV 6 3 480–580 10k, 100k, 12CrMoV 6 2

200kX 11 CrMo 9 1 A 629 450–640 10k, 30k,

50k, 100k,150k, 200k,250k* TU Z10CD09 550–625 10k, 100k X 12CrMo9 1 460–600 10k, 100k,

200kX 11 CrMo 9 1 NT 629–590 440–670 10k, 30k,

50k, 100k,150k, 200k,250k* TU Z10CD09 450–675 10k, 100k X 12CrMo9 1 400–650 10k, 100k,

200k,X 10 NT ‘T91’ 490–690 10k, 30k, TU X 470–650 10k, 100k X 10CrMoVNb9 1 50k, 100k, Z10CD 10CrMoVNb CrMoNiVNb

150k, 200k, VNb09-01 9 1 9 1250k*

X 20 CrMoNiV NT 762 480–680 10k, 30k, X20 CrMoV 470–650 10k, 100k X 2011 1 50k, 100k, 12 1 200k CrMoNiV

150k, 200k, 11 1 1250k*

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3.3.2 Organisation

Being a voluntary group, ECCC has no basic resources or funding to rely on.Nevertheless, there are several sources which, owing to the multi-nationalconfiguration of the ECCC, provide at least some limited support:

• Over the last ten years, ECCC has run European Community fundedprojects.13–15

• Under certain circumstances, ECCC or ECCC sponsored sub-groupshave applied for and obtained European research projects, during whichdetailed problems identified through ECCC activities could be discussed,tested and satisfactorily solved. In some cases these SMT-Projects16,17

formed the basis for new European Standards (e.g. EN 10291 (creeptesting) or 10319 (stress relaxation testing)).

• Each member nation contributes a balanced mixture of testing activityeach year, so that testing resources are available each year for commonlyagreed research programmes.

This strategy compensates for the recent reductions in national and ECresearch budgets and the loss of companies and related specialists dealingwith high temperature resulting from globalisation. It has also allowed theECCC to continue to contribute to debate on safety-related technical issues,which are becoming more and more stringent as requirements for reliabilityand foreseeability of component behaviour increase. Included in this are thenew flexible service conditions, which introduce new hazards by pushingplants into not previously experienced operating conditions.

The ECCC is led by a management committee, in which all membernations have a seat and a vote. Technical work is done in:

• Working Group 1: Creep Data Assessment and Generation Procedures,in which specialists from all over Europe can take part

• Working Group 2: Results Dissemination and Exchange Forum18

• Working Group 3: Creep Data Assessment, divided into sub-groups A(ferritic low and high alloyed steels), B (austenitic steels), C (nickelbase alloys) and

• Working Group 4: Creep of Components and Features, which can bothonly be attended by specialists belonging to organisations that contributeto their country’s fee to the organisation.

3.3.3 ECCC’s contribution to standardisation

The ECCC WG3 subgroups were charged with producing the strength valuesto be proposed to the CEN Technical Committees working on the new EuropeanStandards in 1991/92. To comply with this task, the WG3 groups:

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• identified the steel grades which were relevant to industry and should beconsidered for standardisation;

• collated raw creep data from all available, reliable sources in Europe(and in some cases, the world);

• assessed the credibility of the raw data collated;• identified gaps within the available data;• assessed creep rupture and creep strain strength values for the

standardisation body; and• promoted and led ECCC harmonised testing programmes relying – at

least partially – on common testing resources. Since 1997, more thanfive million hours of stress rupture and creep rupture testing has beenadded to the collated data bases of most relevant steel grades (includingX10CrMoVNNb 9 1 (grade 91), X 10 CrWMoVNb 9 2 (grade 92), X 19CrWMoVNb9 1 1 (grade 911), X 6 NiCrMoTiBN 25 20 (grade 709),NiCr22Co12Mo (alloy 617), NiCr20TiAl (alloy 80A), 9%Cr welds,dissimilar welds (2,25%Cr to 9%Cr steels), aged materials (nickel basealloys), etc).

A strategically essential point in this was the need to establish:

• a modern system for collating and exchanging creep data in order toguarantee completeness of information and the fair involvement of allcontributors;

• an objective system to judge and to guarantee quality of experimentaldata for old and future (including those jointly produced) creep tests;and

• a reliable, commonly agreed procedure on how to derive creep strengthvalues from the large data population collated all over Europe.

A dedicated Working Group (WG1) was established to deal with thiscentral problem and agreed in 1993/9419 on a common approach which wassubsequently applied to all ECCC strength value determinations and wasenhanced and enlarged several times up to 2005.20

Significant innovations included:

• a detailed written procedure by which the whole process, from setting upthe detailed specifications of the material to be assessed for creep strengthto the presentation of the results, is defined;

• defined criteria for evaluating statistical significance and material propertiesof a data set (Table 3; source: ECCC Recommendations, Reference 20;Vol 5 part Ia) and EN 12952-2 Annex B);

• strict requirements for strength assessment:

– strength values for standards to be assessed by two independent assessors

— whose results may not diverge by more than 10%,

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— who have to use at least one of the methods whose procedure isgiven in (Reference 20: Vol 5 part Ia – app. D);

— both assessments to be positively verified by objective postassessment tests (Reference 20: Vol 5, References 21, 22), checkingthe strength prediction physical credibility, data description qualityand repeatability/stability of extrapolation (see Chapter 14,Constitutive equations for creep curves and predicting service life).

Owing to the time needed to prepare all these agreed guidelines, the veryearly data assessments recommended by ECCC did not completely followthe WG1 procedures, but since 1995 all creep strength values published bythe ECCC have been validated by this approach. A first collation of creepstrength values was presented in 1996,23 and an upgraded and enlargedversion in 2005.24 Further upgrades and assessments on other grades areexpected over the coming years.

ECCC-recommended values have been made available to the relevant CENTechnical Committees and their sub-committees by common members, discussedand sometimes iteratively agreed. ECCC strength values formed the basis fora large part of the creep strength values included in the new EN standards.

3.3.4 ECCC’s future

In spite of the difficult funding situation in industrial research in Europe,ECCC will continue to support CEN in providing reliable creep strengthvalues. In order to avoid duplication, to enhance the interaction with PEDtopics and to optimise available resources, ECCC will tighten its collaborationwith the European Pressure Equipment Research Council (EPERC).

3.4 European Pressure Equipment Research Council

(EPERC)

3.4.1 Motivation and history

The European Pressure Equipment Research Council (EPERC)25 was foundedin the early 1990s and established officially in 1995 by a group of Europeanindustry and research organisations. It was set up to collaborate in harmonisingand unifying European laws, regulations, codes and standards on all pressure-bearing equipment, including those operating under high temperature andcreep regimes. In 2005, EPERC became the EPERC-Technological Platform,in the light of the new EC Framework VII research funding strategy.

According to the EPERC Member Agreement,25 the main objectives ofEPERC are the establishment of:

• a European Network to support the pressure equipment industry andsmall and medium enterprises (SMEs), in particular;

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Table 3.3 Recommended data size requirements according to Reference 20, Vol. 5 part I

Interim-minimum Target-minimum requirementsrequirements Original (TM1) TM2 TM3

For datasets with ≥300 For datasets with ≥500observations, originating observations, originating fromfrom ≥10 casts, at ≥5 ≥20 casts, at ≥5 temperaturestemperatures covering the covering the range TMAINrange TMAIN ± ≥50°C ± ≥50°C

For ≥3 casts, there should be For ≥6 casts, there should For ≥5 casts, there should be For ≥5 casts, there shouldtu(T, σ0) observations from: be tu(T, σ0) observations tu(T, σ0) observations from: be tu(T, σ0) observations from:

from:• ≥3 tests at each of ≥3 • ≥5 tests at each of ≥3 • ≥5 tests at each of ≥2 • ≥5 tests at ≥1 temperature(s)

temperatures, at temperatures in the temperatures in the in the design applicationintervals of 50–100°C design application range design application range range (at intervals of

at intervals of 25–50°C at an interval(s) of 25–50°C)25–50°C

– ≥3 tests per temperature – ≥4 tests per – ≥4 tests per – ≥4 tests per temperature(different σ0) with temperature temperature (different σ0) withtu,max≥ 10 kh (different σ0) with (different σ0) with tu ≤ 35 kh

tu,max≥ 40 kh tu,max ≤ 35 kh – ≥1 test per temperature – ≥1 test per temperature – ≥1 test per temperature with tu,max≥ 35 kh

with tu,max≥ 40 kh with tu,max≥ 35 khPredicted strength values determinedfrom an Interim minimum datasetshall be regarded as tentative untilthe data requirements defined inone of the target-minimumcolumns are obtained

tu,max is the longest test duration available, σ0 is the initial test stress, I is the temperature.

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• short- and long-term research priorities for the European pressureequipment industry;

• coordination of cooperative research in the domain of pressure equipmentand identification of funding sources for this research;

• dissemination of research results to the European industry andstandardisation bodies; and

• joint European attitudes to pressure equipment safety and reliability.

EPERC was particularly successful in interconnecting European cooperativeresearch with current design needs, construction, inspection and in-serviceactivities and especially in trying to harmonise contradictory European standards,rules and laws. For this purpose, EPERC was and probably will remain the majordiscussion forum and investigation coordinator providing technical and scientificdepth to support the ‘New approach to Technical Harmonisation and Standards’,which will be the privileged means of complying with the ‘Essential SafetyRequirements of the new Pressure Equipment Directive (PED 97/23/EC)’.

Another important issue for EPERC is maintaining permanent contactswith standardisation bodies (CEN) and the relevant international organisations,PVRC (USA) and JPVRC (Japan), and so on. This cooperation is supportedby a common Memorandum of Understanding (2001) and involves a stronginteraction with all the committees and notified bodies’ umbrella organisationsdealing with the Pressure Equipment Directive (PED 97/23/EC)9 and itsamendments, guidelines and sometimes conflicting interpretations.

Since 2002, EPERC has entered several new areas, including support forthe new European Design Codes (EN 13445, EN 13480, families etc), alsoin the area of creep. In this context, EPERC and ECCC have collaboratedand will develop an even closer relationship in the future.

3.4.2 Organisation

EPERC has relied on in-kind contributions and financial support via EC-funded projects. It is led by a steering committee in which nationalrepresentatives from the countries with member organisations have a seatand a vote. The original EPERC structure can be found in Reference 25. Theactual organisation of EPERC-TP is structured into five task groups, dealingwith (i) Materials & fabrication, (ii) Design, (iii) Damage & rupture analysis,(iv) Operation, inspection, testing & maintenance, (v) Special and emergingissues, and Support to Standardisation.

3.4.3 EPERC’s contribution to creep-relatedstandardisation

EPERC has collaborated on and influenced a huge number of standards andpre-standardisation projects in all areas related to pressure equipment. In the

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high temperature area, essential contributions were made to the harmonisationof the non-destructive testing (NDT) codes and practices used for creep-serviced component inspection, creep crack growth, minimum invasive testing(small punch) and the new philosophies founded on risk base analysis (RIMAP)and ‘fitness for service’ concepts (FITNET). Further EPERC contributionscentred on the design codes and the related CEN TCs, especially for the newEN 13445 standards which include design issues in creep regime in Part 3.This particular project required information on parent material and weldcreep strengths, hence the interaction with ECCC.

3.5 The latest generation of CEN standards for

creep-resistant steels

3.5.1 General principle

The CEN10 standards are prepared by the relevant CEN Technical Committeeand its sub-committees assisted, in some cases, by additional working groups.

The new EN standards attempt to harmonise the often contrasting nationalnorms and rules, focusing on those materials which are either present inmost of the national standards, are not yet included in national standards butare of immediate technical interest, or are evidently commonly used throughoutEurope. In some cases, similar grades codified in Partner National Standardscan be grouped into one common type.

In the case of standards related to creep, the framework covers revisedmethods for testing, specifications for steels, extensive use of creep strengthdata (although this is generally indicated as being informative), welds and –still mainly under development – design rules.

3.5.2 Material specifications

Table 3.4 includes an overview of the latest available CEN standards (up toNovember 2006, although several were expected in 2007) that relate tocreep-resistant steels and reporting strength values. The materials includedin these standards are listed in Tables 3.5 to 3.15, along with some informationon the available creep strength values.

A particular feature is the use of informative annexes reporting meancreep rupture strengths, which are determined from quality experimentaldata collations representing European and sometimes worldwide research.

• Values marked * were obtained by extended extrapolation in time, i.e.by more than a factor of 3, beyond the scope of the available experimentaldata.

• Values in brackets ( ) were calculated by an extrapolation beyond 80%of the minimum stress of the experimental data set.

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96Table 3.4 European specifications for creep-resistant steels

Designation Title Date of Statusissue (Nov/2006)

EN10028-2 Flat products made of steels for pressure purposes – non-alloy and alloy steels 2004 Currentwith specified elevated temperature properties

EN10028-7 Flat products made of steels for pressure purposes – stainless steels 2002 CurrentEN10095 Heat resisting steels and nickel alloys 2001 CurrentEN10213-2 Technical delivery conditions for steel castings for pressure purposes – steel grades 1998 Current

for use at room temperature and elevated temperaturesEN10213-4 Technical delivery conditions for steel castings for pressure purposes – austenitic 1998 Current

and austeno-ferritic steel gradesEN10216-2 Seamless steel tubes for pressure purposes – Technical delivery conditions – Part 2: 2002 Current

Non-alloy and alloy steel tubes with specified elevated temperature propertiesEN10216-5 Seamless steel tubes for pressure purposes – Technical delivery conditions - Part 5: 2002 Current

Stainless steel tubesEN10217-2* Welded steel tubes for pressure purposes – Technical delivery conditions - Part 2: 2002 Current

Electric welded non-alloy and alloy steel tubes with specified elevated temperatureproperties

EN 10217-5* Welded steel tubes for pressure purposes – Technical delivery conditions - Part 5: 2002 CurrentSubmerged arc welded non-alloy and alloy steel tubes with specified elevatedtemperature properties

EN10217-7* Welded steel tubes for pressure purposes – Technical delivery conditions – stainless 2005 Currentsteel tubes

EN10222-2 Steel forgings for pressure purposes – ferritic and martensitic steels with specified 2001 Currentelevated temperature properties

EN10222-5 Steel forgings for pressure purposes – martensitic, austenitic and austeno-ferritic 2001 Currentstainless steels

EN 10253* Butt welding pipe fittings 1999 CurrentEN 10269 Steels and nickel alloys for fasteners with specified elevated and/or low 2001 Current

temperature properties

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EN 10272* Stainless steel bars for pressure purposes 2003 CurrentEN 10273 Hot rolled weldable steel bars for pressure purposes with specified elevated 2002 Current

temperature propertiesEN 10295 Heat resistant steel castings 2003 CurrentEN 10302 Creep resistant steels, nickel and cobalt alloys 2002 CurrentEN 12952-2* Water tube boilers and auxiliary installations – Part 2: Materials for pressure 2001 Current

parts of boilers and accessoriesEN 14532-2 Welding consumables: Test methods and quality requirements Part 2: 2004 Current

Supplementary methods and conformity assessment of consumables for steelnickel and nickel-alloys

* These standards do not include creep strength values themselves, but the materials included can be referred back to other standards in thetable, where strength values are listed.

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98Table 3.5 Chemical composition of carbon, low and high alloyed ferritic steels in European Standards

Material grade Gross chemical composition

Heat Concentration in mass% of the following elementstreatment

C Cr Mo Ni V Other

GP240GH N/ QT 0.18–0.23 – – – –GP280GH N–QT 0.18–0.25 – – – –P195 N <0.13 <0.30 <0.08 <0.3 <0.02P235GH N <0.16 <0.30 <0.08 <0.3 <0.02P245GH A, NT, QT 0.08 –0.2 – – – –P250GH N 0.18–0.23 <0.30 – <0.3 <0.02P265GH N <0.2 <0.3 <0.08 <0.3 <0.02P280GH N, NT, QT 0.08 –0.2 – – – –P295GH N 0.08–0.2 <0.3 <0.08 <0.3 <0.02P305GH N, NT, QT 0.15–0.2P355GH N 0.1–0.22 <0.3 <0.08 <0.3 <0.02C7–C24 N <0.24 – – – –C35E N, QT 0.32–0.39 <0.4 <0.1 <0.4 –15 MnCrMoNiV 5 3 NT QT <0.17 0.5–1 0.2–0.35 0.3–0.7 0.05–0.1 Mn 1–1.515 MnMoV 4 5 NT QT <0.18 – 0.4–0.6 – 0.04–0.08 Mn 0.9–1.418 MnMo 4 5 NT QT <0.2 <0.3 0.45–0.6 <0.3 – Mn 0.9–1.520 Mn 5 N 0.17–0.23 <0.4 <0.1 <0.4 – Mn 1–1.520 MnMoNi 4 5 QT 0.15–0.23 <0.2 0.45–0.6 0.4–0.8 <0.02 Mn 1–1.58 MoB 5 4 N 0.06 –0.1 <0.2 0.4–0.5 – – B 0.002–0.00612 MoCrV 6 2 2 NT QT14 MoV 6 3 NT QT 0.1–0.18 0.3–0.6 0.5–0.7 – 0.22–0.2816 Mo 3 N, NT, QT 0.12–0.2 <0.3 0.25–0.35 <0.3 –17 Mo 5 N, NT, QT20 MnNb 6 N <0.22 – – – – Mn 1–1.5; Nb

0.015–0.1(Cr+Cu+Mo+Ni)>0.7

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G 20 Mo5 QT 0.15–0.23 – 0.4–0.6 – –10 CrMo 5 5 NT <0.15 1–1.5 0.45–0.65 <0.3 –10 CrMo 9 10 NT QT 0.08–0.14 2.0–2.5 0.9–1.1 – –11 CrMo 9 10 IA 0.08–0.15 2–2.5 0.9–1.1 – –11 CrMo 9 10 NT QT 0.08–0.15 2–2.5 0.9–1.1 – –11 CrMo 9 10 QT (!) 0.08–0.15 2–2.5 0.9–1.1 – –12 CrMo 9 10 NT QT 0.10–0.15 2.0–2.5 0.9–1.1 <0.3 –12 CrMoV 12 10 NT QT 0.10–0.15 2.75–3.25 0.9–1.1 <0.25 0.2–0.312 CrMoV 6 2 2 NT 0.09–0.17 0.3–0.55 0.51–0.7 <0.22 0.21–0.34 N 0.0015–0.016G 12 MoCrV 5 2 QT 0.1–0.15 0.3–0.5 0.4–0.6 – 0.22–0.313 CrMo 4 5 NT QT 0.08–0.18 0.7–1.15 0.4–0.6 – –13 CrMo 4 5 N 0.08–0.18 0.7–1.15 0.4–0.6 – –13 CrMoSi 5 5 NT QT <0.17 1–1.5 0.45–0.65 <0.3 – Si 0.5–0.813 CrMoV 9 10 NT QT 0.11–0.15 2–2.5 0.9–1.1 <0.25 0.25–0.35G 17 CrMo 5 5 QT 0.15–0.2 1–1.5 0.45–0.65 – –G 17 CrMo 9 10 QT 0.13–0.2 2–2.5 0.9–1.2 – –G 17 CrMoV 5 10 QT 0.15–0.2 1.2–1.5 0.9–1.1 – 0.2–0.320 CrMoV 13 5 NT QT 0.17–0.23 3–3.3 0.5–0.6 – 0.45–0.5520 CrMoV 13 5 5 QT 0.17–0.23 3.0–3.5 0.5–0.6 <0.3 0.45–0.5520 CrMoVTiB 4 10 QT 0.17–0.23 0.9–1.2 0.9–1.1 <0.2 0.6–0.8 B 0.001–0.010;

Ti 0.007–0.1521 CrMoV 5 7 QT 0.17–0.25 1.2–1.5 0.55–0.8 <0.6 0.2–0.3525 CrMo 4 QT 0.22–0.29 0.9–1.2 0.15–0.3 – –40 CrMo 5 6 NT 0.39–0.45 1.2–1.5 0.5–0.7 – –40 CrMoV 4 6 QT 0.26–0.44 0.9–1.2 0.5–0.65 – 0.25–0.3540 CrMoV 4 6 NT42 CrMo 5 6 QT 0.39–0.45 1.2–1.5 0.5–0.7 – –15 NiCuMoNb 5 6 4 NT QT <0.17 <0.3 0.25–0.5 0.1–1.3 – Cu 0.5–0.8; Nb

0.015–0.045X 3 CrAlTi 18 2 A <0.04 17.0–18.0 – – – Si <1; Al 1.7–2.1;

Ti 0.2+4*(C+N)– 0.8

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X 8 CrCoNiMo 10 6 QT 0.05–0.12 9.8 –11.2 0.5 –1 0.2–1.2 0.1–0.4 Nb 0.2–0.5;. W<0.7;B 0.005–0.015; Co 5–7

X 10 CrAlSi 13 A <0.12 12.0–14.0 – – – Si 0.7–1.4; Al 0.7–1.2X 10 CrAlSi 18 A <0.12 17.0–19.0 – – – Si 0.7–1.4; Al 0.7–1.2X 10 CrAlSi 25 A <0.12 23.0–26.0 – – – Si 0.7–1.4; Al 1.2–1.7X 10 CrAlSi 7 A <0.12 6.0–8.0 – – – Si 0.5–1. Al 0.5–1X 10 CrMoVNb 9 1 NT QT 0.08–0.12 8–9.5 0.85–1.05 <0.3 0.18–0.25 Nb 0.06–0.1X 11 CrMo 5 A 0.08–0.15 4.0–6.0 0.45–0.65 – –X 11 CrMo 5 NT1 0.08–0.15 4.0–6.0 0.45–0.65 – –X 11 CrMo 5 NT2 0.08–0.15 4.0–6.0 0.45–0.65 – –X 11 CrMo 5 IA 0.08–0.15 4.0–6.0 0.45–0.65 – –X 11 CrMo 9 1 A1 0.08–0.15 8.0–10.0 0.9–1.1 – – –X 11 CrMo 9 1 IA 0.08–0.15 8.0–10.0 0.9–1.1 – – –X 11 CrMo 9 1 NT 0.08–0.15 8.0–10.0 0.9–1.1 – – –X 11 CrMoWVNb 9 1 1 NT QT 0.09–0.13 8.5–9.5 0.9–1.1 0.1–0.4 0.18–0.25 W 0.9–1.1; Nb

0.06–0.1; N0.05–0.09;B0.0005–0.005

X 11 CrWMoVNb 9 1 NT 0.07–0.13 8.5–9.5 0.3–0.6 <0.4 0.15–0.25 W 1.5–2; Nb0.04–0.09; N0.03–0.07; B0.001–0.006

GX 12 CrMoVNbN 9 1 NT QT 0.1–0.14 8.0–9.5 0.85–1.05 <0.4 0.18–0.25 Nb 0.06–0.1; N0.03–0.07

X 12 CrMo 5 NT QT 0.10–0.15 <4–6 0.45–0.65 <0.3 –

Table 3.5 (Cont’d.)


Heat Concentration in mass% of the following elementstreatment

C Cr Mo Ni V Other

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X 12 CrNiMoV 12 3 QT 0.08–0.15 11–12.5 1.5 –2 2 –3.0 0.25–0.4GX 15 CrMo 5 QT 0.12–0.19 4–6 0.45–0.65 – –X 15 CrMo 5 1 NT QT <0.18 4.0–6.0 0.45–0.65 – –X 16 CrMo 5 1 NT QT <0.18 4–6 0.45–0.65 – –X 16 CrMo 5 1 A <0.18 4–6 0.45–0.65 – –X 18 CrN 28 A 0.15–0.20 26.0–29.0 – – – Si <1. N 0.15–0.25X 19 CrMoNbVN 11 1 QT 0.17–0.23 10–11.5 0.5–0.8 0.2–0.6 0.1–0.3 Nb 0.25–0.55;

N 0.05–0.1X 20 CrMoV 11 1 QT 0.17–0.23 10–12.5 0.8–1.2 0.3–0.8 0.2–0.35X 20 CrMoWV 12 1 QT 700 0.17–0.24 11–12.5 0.8–1.2 0.3–0.8 0.2–0.35 W 0.4–0.6X 20 CrMoWV 12 1 QT800 0.17–0.24 11–12.5 0.8–1.2 0.3–0.8 0.2–0.35 W 0.4–0.6X 22 CrMoV 12 1 QT 0.18–0.24 11–12.5 0.8–1.2 0.3–0.8 0.25–0.35GX 23 CrMoV 12 1 QT 0.2–0.26 11.3–12.2 1–1.2 <1.0 0.25–0.35GX 30 CrSi7 A 0.2–0.35 6.0–8.0 <0.15 <0.5 – Si 1–2.5GX 40 CrSi 13 A 0.3–0.5 12.0–14.0 <0.5 <1.0 – Si 1–2.5GX 40 CrSi 17 A 0.3–0.5 16.0–19.0 <0.5 <1.0 – Si 1–2.5GX 40 CrSi 24 as cast 0.3–0.5 23.0–26.0 <0.5 <1.0 – Si 1–2.5GX 40 CrSi 28 as cast 0.3–0.5 27.0–30.0 <0.5 <1.0 – Si 1–2.5GX 130 CrSi 29 as cast 1.2–1.4 27.0–30.0 <0.5 <1.0 – Si 1–2.5GX 160 CrSi 18 as cast 1.4–1.8 17.0–19.0 <0.5 <1.0 – Si 1–2.5

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102Table 3.6 Available creep data for carbon, low and high alloyed ferritic steels: plates; heat-resistant steels and Ni alloys; castings

Product form Plates Heat-resisting steels and Ni alloys CastingsEN 10028-2 EN 10095 EN 10213-2

1% strain creep Creep rupture 1% strain creep Creep rupture Creep ruptureMaterial grade strength strength strength strength strength

T range Duration T range Duration T range Duration T range Duration T range Duration(°C) (h) (°C) (h) (°C) (h) (°C) (h) (°C) (h)

GP240GH 400–500 10k, 100k,200k

GP280GH 400–500 10k, 100k,P195P235GH 380–480 10k–100k 380–480 10k, 100k,

200kP245GHP250GHP265GH 380–480 10k–100k 380–480 10k, 100k,

200kP280GHP295GH 380–500 10k–100k 380–500 10k, 100k,

200kP305GHP355GH 380–500 10k–100k 380–500 10k, 100k,

200kC7–C24C35E15 MnCrMoNiV5 315 MnMoV 4 5

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18 MnMo 4 5 425–525 10k–100k 425–525 10 k, 100k20 Mn 520 MnMoNi 4 5 Na na 450–490 10 k, 100k8 MoB 5 412 MoCrV 6 2 214 MoV 6 316 Mo 3 450–530 10k–100k 450–530 10k, 100k

200k17 Mo 520 MnNb 6G 20 Mo5 400–550 10k, 100k,

200k10 CrMo 5 510 CrMo 9 10 450–600 10k–100k 450–600 10k, 100k

200k11 CrMo 9 1011 CrMo 9 1011 CrMo 9 1012 CrMo 9 10 na na 400–520 10k, 100k12 CrMoV na na 400–550 10k, 100k12 1012 CrMoV 6 2 2G 12 MoCrV 5 2 450–600 10k, 100k13 CrMo 4 5 450–570 10k–100k 450–570 10k, 100k,

200k13 CrMo 4 513 CrMoSi 5 5 450–570 100k 450–570 100k13 CrMoV 9 10 na na 400–550 10k, 100kG 17 CrMo 5 5 400–550 10k, 100k,

200kG 17 CrMo 9 10 400–600 10k, 100k,

200k

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104Table 3.6 (Cont’d)




G 17 CrMoV 400–600 10k 100k5 10 200k20 CrMoV 13 520 CrMoV 13 5 520 CrMoVTiB4 1021 CrMoV 5 725 CrMo 440 CrMo 5 640 CrMoV 4 640 CrMoV 4 642 CrMo 5 615 NiCuMoNb 400–500 10k, 100k 400–500 10k, 100k5 6 4X 3 CrAlTi 18 2 500–900 1k, 10k 500–900 1k, 10k,

100kX 8 CrCoNiMo10 6X 10 CrAlSi 13 500–900 1k, 10k 500–900 1k, 10k,

100kX 10 CrAlSi 18 500–900 1k, 10k 500–900 1k, 10k,

100k

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X 10 CrAlSi 25 500–900 1k, 10k 500–900 1k, 10k,100k

X 10 CrAlSi 7 500–900 1k, 10k 500–900 1k, 10k,100k

X 10 CrMoVNb na na 500–670 10k, 100k,9 1 200kX 11 CrMo 5X 11 CrMo 5X 11 CrMo 5X 11 CrMo 5X 11 CrMo 9 1X 11 CrMo 9 1X 11 CrMo 9 1X 11 CrMoWVNb 9 1 1X 11 CrWMoVNb9 1GX 12 CrMoVNNb 9 1X 12 CrMo 5 460–625 10k 460–625 10kX 12 CrNiMoV12 3GX 15 CrMo 5 470–600 10k 100kX 15 CrMo 5 1X 16 CrMo 5 1X 16 CrMo 5 1X 18 CrN 28 500–900 1k, 10k 500–900 1k, 10k,

100kX 19 CrMoNbVN 11 1X 20 CrMoV11 1

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X 20 CrMoWV12 1X 20 CrMoWV12 1X 22 CrMoV12 1GX 23 CrMoV 400–600 10k 100k12 1 200kGX 30 CrSi7GX 40 CrSi 13GX 40 CrSi 17GX 40 CrSi 24GX 40 CrSi 28GX 130 CrSi 29GX 160 CrSi 18

Table 3.6 (Cont’d)




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Table 3.7 Available creep data for carbon, low and high alloyed ferritic steels; seamless tubes/pipes; forgings; bars for fasteners and bolts; bars

Product form Seamless tubes/pipes Forgings Bars for fasteners and bolting BarsEN 10216-2 EN 10222-2 EN 10269 EN 10273

Creep rupture 1% strain Creep Creep rupture 1% strain creep Creep rupture 1% strain Creep Creep rupture

Material gradestrength strength strength strength strength strength strength

T range Duration T range Duration T range Duration T range Duration T range Duration T range Duration T range Duration(°C) (h) (°C) (h) (°C) (h) (°C) (h) (°C) (h) (°C) (h) (°C) (h)

GP240GHGP280GHP195P235GH 400–500 10k 100k 380–480 10k– 380–480 10k, 100k,

200k 250k 100k 200kP245GH 380–480 10k–100k 380–480 10k, 100k,P250GH 200k 380–480 10k– 380–480 10k, 100k,

100k 200kP265GH 400–500 10k 100k 380–480 10k – 380–480 10k, 100k,

200k 250k 100k 200kP280GH 380–480 10k–100k 380–480 10k, 100k,P295GH 200k 380–500 10k– 380–500 10k, 100kP305GH 380–480 10k–100k 380–480 10k, 100k, 100k 200kP355GH 200k 380–500 10k– 380–500 10k, 100k,C7–C24 100k 200kC35E 350–500 10k 100k 350–500 10k 100k15 MnCr na na 400–500 10k, 100k,

200kMoNiV 5 315 MnMoV 4 5 450–500 10k–100k 430–500 10k–100k18 MnMo 4 520 Mn 5 na na 380–500 10k 100k20 MnMoNi 4 5 200k8 MoB 5 4

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12 MoCrV 6 2 214 MoV 6 3 450–600 10k, 100k, 480–600 10k 100k 450–600 10 k, 100k,

200k, 250k 200k16 Mo 3 450–550 10k, 100k, 450–530 10k–100k 450–530 10k, 100k, 450–530 10k – 450–530 10k, 100k,

200k, 250k 200k 100k 200k17 Mo 520 MnNb 6 400–500 10k, 100k

200k* 250k*G 20 Mo510 CrMo 5 5 450–600 10k, 100k

200k, 250k10 CrMo 9 10 450–600 10k, 100k 450–600 10k, 450–600 10k, 100k,

200k, 250k 100k 200k11 CrMo 9 1011 CrMo 9 10 400–520 10k, 100k 450–600 10k–100k 450–600 10k, 100k,

200k11 CrMo 9 10 400–520 100k12 CrMo 9 1012 CrMoV12 1012 CrMoV6 2 2G 12 MoCrV5 2






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13 CrMo 4 5 450–600 10k, 100k, 450–570 10k–100k 450–570 10k, 100k, 450–570 10k 450–570 10k, 100k,200k 250k 200k 100k 200k

13 CrMo 4 513 CrMoSi 5 513 CrMoV 9 10G 17 CrMo 5 5G 17 CrMo 9 10G 17 CrMoV5 1020 CrMoV 13 520 CrMoV 420–550 10k, 100k13 5 520 CrMoVTiB na na 450– 600 10k, 100k,4 10 200k21 CrMoV 5 7 420–550 10k, 100k 420–550 10k, 100k,

200k25 CrMo 4 420–550 10k, 100k 420–550 10k, 100k40 CrMo 5 640 CrMoV 4 6 na na 450–550 10k, 100k

200k40 CrMoV 4 642 CrMo 5 6 na na 450–550 10k, 100k,

200k15 NiCuMoNb 400–500 10k, 100k5 6 4X 3 CrAlTi 18 2X 8 CrCoNiMo10 6X 10 CrAlSi 13X 10 CrAlSi 18X 10 CrAlSi 25X 10 CrAlSi 7X 10 CrMoV 500–650 10k, 100k na na 500–670 10k, 100k,Nb 9 1 200k* 200kX 11 CrMo 5

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110Table 3.7 (Cont’d.)





X 11 CrMo 5 450–600 10k, 100k200k, 250k

X 11 CrMo 5 450–600 10k, 100k200k, 250k

X 11 CrMo 5 450–630 10k, 100k200k, 250k

X 11 CrMo 9 1X 11 CrMo 9 1 460–600 10k, 100kX 11 CrMo 9 1 450–650 10k, 100k,

200kX 11 CrMoWVNb 9 1 1X 11 CrWMoVNb 9 1GX 12 CrMoVNbN 9 1X 12 CrMo 5X 12 CrNiMoV12 3GX 15 CrMo 5X 15 CrMo 5 1 490–570 10k 100k 450–600 10k, 100k,

200kX 16 CrMo 5 1 490–570 10k, 100k 450–600 10k, 100k,

200k

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X 16 CrMo 5 1 450–600 10k, 100k 450–600 10k, 100k,200k

X 18 CrN 28X 19 CrMoNbVN 11 1 450–600 10k, 100k 450–600 10k, 100k,X 20 CrMoV 480–650 10k 100k 470–650 10k 100k 480–630 10k, 100k, 200k11 1 200k 200kX 20 CrMoWV12 1X 20 CrMoWV12 1X 22 CrMoV12 1 450–600 10k, 100k 450–600 10k, 100kGX 23 CrMoV12 1GX 30 CrSi7GX 40 CrSi 13GX 40 CrSi 17GX 40 CrSi 24GX 40 CrSi 28GX 130 CrSi 29GX 160 CrSi 18

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112Table 3.8 Available creep data for carbon, low and high alloyed ferritic steels: castings; creep-resistant steels; consumables; generic products

Product form Castings Creep-resistant steels Consumables GenericEN 10295 EN 10302 EN 14532-2 ECCC

1% strain creep Creep rupture 1% strain creep Creep rupture Creep Creep rupture

Material gradestrength strength strength strength rupture strength


GP240GHGP280GHP195 x 400–500 10k 100k

200k, 250k,P235GH x 400–500 10k, 100k,

200k, 250k,P245GHP250GHP265GH x 400–500 10k, 100k,

200k, 250k,P280GHP295GHP305GHP355GH x 400–500 10k, 100k,

200k, 250k,C7–C24C35E15 MnCrMoNiV 5 315 MnMoV 4 518 MnMo 4 520 Mn 520 MnMoNi 4 58 MoB 5 4 400–550 10k, 100k,

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12 MoCrV 6 2 2 x14 MoV 6 316 Mo 3 x 450–550 10k, 100k,

200k, 250k17 Mo 520 MnNb 6G 20 Mo510 CrMo 5 510 CrMo 9 1011 CrMo 9 1011 CrMo 9 10 x 450–600 10k, 100k,

200k, 250k11 CrMo 9 10 x 400–520 10k, 100k12 CrMo 9 1012 CrMoV 12 1012 CrMoV 6 2 2 475–600 10k, 100k,

200k, 250kG 12 MoCrV 5 213 CrMo 4 5 x 450–600 10k, 100k,

200k, 250k13 CrMo 4 513 CrMoSi 5 513 CrMoV 9 10G 17 CrMo 5 5G 17 CrMo 9 10G 17 CrMoV5 1020 CrMoV 13 5 420–550 10k, 100k20 CrMoV13 5 520 CrMo x 450– 600 10k, 30k,VTiB 4 10 xNotch 100k, 200k21 CrMoV 5 7 420–550 10k, 30k,

100k200k, (1%)

25 CrMo 4 420–550 10k, 100k

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40 CrMo 5 6 xxNotch

40 CrMoV 4 6 x40 CrMoV 4 6 x

xNotch42 CrMo 5 6 450–550 10k, 30k,

100k, 200k15 NiCuMo 400–500 10k, 100kNb 5 6 4X 3 CrAlTi 18 2X 8 CrCo 500–600 10k, 100kNiMo 10 6X 10 CrAlSi 13X 10 CrAlSi 18X 10 CrAlSi 25X 10 CrAlSi 7X 10 Cr 470 –650 10k, 470 –650 10k, x 500–670 10k, 30k,

100k 100k,MoVNb 9 1 200k 100k, 200k,X 11 CrMo 5 450–630 10k, 100k,

200k, 250kX 11 CrMo 5 xX 11 CrMo 5 xX 11 CrMo 5 x

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X 11 CrMo 9 1 450–640 10k, 100k,200k

X 11 CrMo 9 1 x 460–600 10k, 100k,X 11 CrMo 9 1 x 450–650 10k, 100k,

200kX 11 CrMo 480–650 10k, 480–650 10k, 520–650 10k, 30k,WVNb 9 1 1 (100k) 100k 100k, (200k)X 11 CrWMo 520–650 10k, 30k,VNb 9 1 100k, (200k)GX 12 CrMo 470–620 10k, 30k,VNbN 9 1 100k, 200k,X 12 CrMo 5 450–600 10k, 100k,

200k 250kX 12 CrNi 500–600 10k,

100kMoV 12 3GX 15 CrMo 5X 15 CrMo 5 1X 16 CrMo 5 1X 16 CrMo 5 1X 18 CrN 28X 19 CrMoN 450–600 10k, 100k 450–600 10k, 100k,bVN 11 1 200kX 20 CrMoV 470–650 10k, 100k 470–650 10k, 100k, x 480–650 10k, 100k,11 1 200k 200kX 20 CrMoWV 470–650 10k, 100k 470–650 10k, 100k,12 1 200kX 20 CrMoWV 450–600 10k, 100k 450–600 10k, 100k1 2 1X 22 CrMoV 450–600 10k, 100k 450–600 10k, 100k1 2 1GX 23 CrMoV 12 1GX 30 CrSi7 600–800 10kGX 40 CrSi 13 600–900 100. 1k 600–800 100, 1k

10k

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GX 40 CrSi 17 600–900 10kGX 40 CrSi 24 600 –900 10kGX 40 CrSi 28 600–900 10k 600–900 100, 1kGX 130 CrSi 29 600–900 10k 800–900 100, 1kGX 160 CrSi 18 600–900 10k

Duration or temperature in brackets ( ) represent strength values after extended extrapolation in stress.Duration or temperature with an asterisk * represent strength values after extended extrapolation in time.For steels assessed by ECCC the symbol (1%) means that a 1% strain creep strength assessment is also available.EN 14532-2 does not report strength data but does report assessed equations relating strength, duration and temperature.In the EN 14532-2 column, the remark ‘x notch’ means that an equation for creep notch rupture strength is also provided.






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Table 3.9 Chemical composition of austenitic steels contained in European Specifications


Concentration in mass% of the following elements

C Cr Mo Ni Ti Other

X 2 CrNi 18 9 <0.03 17.5–19.5 – 8.0–10.0 – N<0.11X 5 CrNi 18 10 <0.07 17.5–19.5 – 8.0–10.5 – N<0.11GX 5 CrNi 19 10 <0.07 18.0–20.0 – 8.0–11.0 –X 6 CrNi 18 10 0.04–0.08 17.0–19.0 – 8.0–11.0 –X 6 CrNi 23 13 0.04–0.08 22.0–24.0 – 12.0–15.0 –X 6 CrNi 25 20 0.04–0.08 24.0–26.0 – 19.0–22.0 –X 8 CrNi 25 21 <0.1 24.0–26.0 – 19.0–22.0 –X 12 CrNi 23 13 <0.15 22.0–24.0 – 12.0–14.0 –X 2 CrNiN 18 10 < 0.03 17.5–19.5 – 8.5–11.5 – N 0.12–0.22X 2 CrNiMo 17 12 2 < 0.03 16.5–18.5 2.0–2.5 10.0–13.0 N <0.11X 3 CrNiMo 17 13 3 <0.05 16.5–18.5 2.5–3.0 10.5–13.0 – N <0.11X 5 CrNiMo 17 12 2 <0.07 16.5–18.5 2.0–2.5 10.0–13.0 – N <0.11GX 5 CrNiNb 19 11 <0.07 18.0–20.0 – 9.0–12.0 – Nb 8*C–1.0GX 5 CrNiMo 19 11 2 <0.07 18.0–20.0 2.0–2.5 9.0–12.0 –X 6 CrNiNb 18 10 A1000 <0.08 17.0–19.0 – 9.0–12.0 – Nb 10*C–1.0X 6 CrNiNb 18 10 A1100 <0.08 17.0–19.0 – 9.0–12.0 – Nb 10*C–1.0X 6 CrNiTi 18 10 A1000 <0.08 17.0–19.0 – 9.0–12.0 5*C–0.7X 6 CrNiTi 18 10 A1100 <0.08 17.0–19.0 – 9.0–12.0 5*C–0.7X 6 CrNiMo 17 13 2 0.04–0.08 16.5–18.5 2.0–2.5 12.0–14.0 – –X 7 CrNiNb 18 10 0.04–0.10 17.0–19.0 – 9.0–12.0 – Nb 10*C–1.2X 7 CrNiTi 18 10 0.04–0.08 17.0–19.0 – 9.0–13.0 5*(C+N)– N <0.11

0.8X 8 CrNiTi 18 10 <0.1 17.0–19.0 – 9.0–12.0 5*C–0.8X 8 CrNiNb 19 11 0.06–0.1 17.0–20.0 – 9. 0– 13 – Si <0.75; Nb 8*C–1X 8 CrNiNb 16 13 0.04–0.1 15.0–17.0 – 12.0–14.0 – Nb 10*C–1.2

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X 12 CrCoNi 21 20 A 0.08–0.16 20–22.5 2.5–3.5 19.0–21 – Nb 0.75–1.25; W 2–3.Co 18.5–21

X 12 CrCoNi 21 20 P 0.08–0.16 20–22.5 2.5–3.5 19.0–21 – Nb 0.75–1.25; W 2–3.Co 18.5–21

X 15 CrNiSi 20 12 <0.2 19–21 – 11.0–13.0 – Si 1.5–2.5;X 15 CrNiSi 25 21 <0.2 24.0–26.0 – 19.0–22.0 – Si 1.5–2.5;X 15 CrNiSi 25 4 A 0.1–0.2 24.5–26.5 – 3.5–5.5 – Si 0.8–1.5GX 25 CrNiSi 18 9 0.15–0.35 17–19 <0.5 8.0–10.0 – Si 0.5–2.5GX 25 CrNiSi 20 14 0.15–0.35 19–21 <0.5 13–15 – Si 0.5–2.5GX 40 CrNiSi 27 4 0.3–0.5 25–28 <0.5 3.0–6.0 – Si 1–2.5GX 40 CrNiSi 22 10 0.3–0.5 21–23 <0.5 9.0–11.0 – Si 1–2.5GX 40 CrNiSi 25 12 0.3–0.5 24–27 <0.5 11.0–14.0 – Si 1–2.5GX 40 CrNiSi 25 20 0.3–0.5 24–27 <0.5 19.0–22.0 – Si 1–2.5X 2 CrNiMoN 17 11 2 <0.03 16.5–18.5 2.0–2.5 10.0–12.0 – N 0.12–0.22X 2 CrNiMoN 17 13 3 <0.03 16.5–18.5 2.5–3.0 11.0–14.0 – N 0.12–0.22X 3 CrNiMoN 17 13 3 <0.04 16.0–18.0 2.0–3.0 12.0–14.0 – N 0.1–0.18; B 0.0015–0.005X 5 CrNiNbN 18 10X 5 CrNiMoB 17 13 3X 6 CrNiTiB 18 10 0.04–0.08 17.0–19.0 – 9.0–12.0 5*C–0.8 B 0.0015–0.005X 6 CrNiMoTi 17 12 2 <0.08 16.5–18.5 2.0–2.5 10.5–13.5 5*C– 0.7X 6 CrNiMoB 17 12 2 0.04–0.08 16.5–18.5 2.0–2.5 10.0–13.0 – B 0.0015–0.005X 8 CrNiMoNb 16 16 0.04–0.1 15.5–17.5 1.6–2.0 15.5–17.5 – Nb 10*C–1.2X 8 CrNiNbN 25 21 <0.1 23.0–27.0 – 17.0–23.0 – Nb 0.2–0.6; N 0.15–0.35X 10 CrNiCuNb 18 10 0.07–0.13 17–19 – 7.5–10.5 – Cu 2.5–3.5; Nb 0.3–0.6;

N 0.05–0.12




C Cr Mo Ni Ti Other

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X 12 CrNiWTiB 16 13 A 0.07–0.15 15.5–17.5 – 12.5–14.5 0.4–0.7 W 2.5–3; B 0.0015–0.006X 25 CrMnNiN 25 9 7 0.2–0.3 24–26 – 6.0–8.0 – Mn 8–10; N 0.2–0.4GX 40 CrNiSiNb 24 24 0.3–0.5 23–25 <0.5 23.0– 25.0 – Si 1–2.5. Nb 0.8–1.8X 3 CrNiMoBN 17 13 3 <0.04 16.0–18.0 2.0–3.0 12.0–14.0 – N 0.1–0.18; B 0.0015–0.005X 6 CrNiWNbN 16 16 0.04–0.1 15.5–17.5 – 15.5–17.5 – Nb 10*C–1.2; W 2.5–3.5;

N 0.06–0.14X 6 CrNiWNbN 0.04–0.1 15.5–17.5 – 15.5–17.5 – Nb 10*C–1.2; W 2.5–3.5;16 16 WW N 0.06–0.14X 6 CrNiMoTiB 17 13 A 0.04–0.08 16–18 2–2.5 12.0–14.0 5*C–0.8 B 0.0015–0.006X 6 CrNiSiNCe 19 10 0.04–0.08 18–20 – 9.0– 11.0 – Si 1–2; N 0.12–0.2;

Ce 0.03–0.08X 7 CrNiMoBNb 16 16 0.04–0.1 15.5–17.5 1.6–2.0 15.5–17.5 – B 0.05–0.1; Nb+Ta 10*C–1.2X 8 CrNiMoVNb 16 13 0.04–0.1 15.5–17.5 1.1–1.5 12.5–14.5 – Nb 10*C–1.2; V 0.6–0.85X 9 CrNiSiNCe 21 11 2 0.05–0.12 20–22 – 10.0–12.0 – Si 1.4–2.5;. N 0.12–0.2;

Ce 0.03–0.08X 10 CrNiMnNbV15 10 6 1X 12 CrNiWTiB 16 13 WW 0.07–0.15 15.5–17.5 – 12.5–14.5 0.4–0.7 W 2.5–3; B 0.0015–0.006X 10 CrNiMoMnNbV 0.07–0.13 14.0–16.0 0.8–1.2 9.0–11.0 – Mn 5.5–7; B 0.003–0.009;B 15 10 1 V 0.15–0.4; Nb

0.75–1.25; N<0.11X 10 NiCrSi 35 19 <0.15 17.0–20.0 – 33–37 – Si 1–2;X 12 NiCrSi 35 16 <0.15 15.0–17.0 – 33–37 – Si 1–2;GX 35 NiCrSi 25 21 0.2–0.5 19–23 <0.5 23.0–27.0 – Si 1–2.0GX 40 NiCrSi 35 17 0.3–0.5 16–18 <0.5 34–36 – Si 1–2.5GX 40 NiCrSi 38 19 0.3–0.5 18–21 <0.5 36–39 – Si 1–2.5GX 40 NiCrSi 35 26 0.3–0.5 24–27 <0.5 33–36 – Si 1–2.5GX 40 NiCrNb 45 35 0.35–0.45 32.5–37.5 – 42–46 – Si 1.5–2. Nb 1.5–2GX 50 NiCrCo 20 20 20 0.35–0.65 19–22 2.5–3.0 18–22 – Si<1:. Nb 0.75–1.25;

Co 18.5–22, W 2–3X 5 NiCrAlTi 31 20 0.03–0.08 19.0–22.0 – 30.0–32.5 0.2–0.5 Al 0.2–0.5; Cu <0.5; Nb<0.1X 5 NiCrAlTi 31 20 (RA) 0.03–0.08 19.0–22.0 – 30.0–32.5 0.2–0.5 Al 0.2–0.5; Cu <0.5; Nb<0.1

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C Cr Mo Ni Ti Other

X 6 NiCrNbCe 32 27 0.04–0.08 26–28 – 31–33 – Ce 0.05–0.1; Nb 0.6–1X 8 NiCrAlTi 32 21 0.05–0.1 19.0–22.0 – 30.0–34.0 0.25–0.65 Al 025–0.65; Cu<0.5;

Co<0.5; Al+Ti<0.5X 8 NiCrAlTi 32 21 (RK) 0.05–0.1 19.0–22.0 – 30.0–34.0 0.25–0.65 Al 0.25–0.65; Cu<0.5;

Co<0.5; Al+Ti 0.85–1.2X 10 NiCrAlTi 32 21 <0.12 19.0–23.0 – 30–34 0.15–0.6 Al 0.15–0.6X 10 NiCrSiNb 35 22 <0.15 20–23 – 33–37 – Si 1–2; Nb 1–1.5GX 10 NiCrSiNb 32 20 0.05–0.15 19–21 <0.5 31–33 – Si 0.5–1.5; Nb 0.5–1.5GX 40 NiCrSiNb 35 18 0.3–0.5 17–20 <0.5 34–36 – Si 1–2.5; Nb 1–1.8GX 40 NiCrSiNb 38 19 0.3–0.5 18–21 <0.5 36–39 – Si 1–2.5; Nb 1.2–1.8GX 40 NiCrSiNb 35 26 0.3–0.5 24–27 <0.5 33–36 – Si 1–2.5; Nb 0.8–1.8GX 50 NiCrCoW 35 0.45–0.55 24–26 – 33–37 – Si 1–2. Co 14–16; W 4–625 15 5X 6 NiCrSiNCe 35 25 0.04–0.08 24–26 – 34–36 – Si 1.2–2; N 0.12–0.2;

Ce 0.03–0.08X 2 NiCrMoNbBN 25 22 <0.04 21.5–23 1.0–2.0 22–28 <0.2 Nb 0.1–0.4; B 0.002–0.01;

N 0.1–0.25X 6 NiCrTiMoVB 25 15 2 0.03–0.08 13.5–16.0 1.0–1.5 24.0–27.0 1.9–2.3 B 0.003–0.01; V 0.1–0.5X 8 NiCrMoNbBN 25 20 0.05–0.1 19–21 1.0–2.0 22–28 <0.2 Nb 0.1–0.4; B 0.002–0.01;

N 0.1–0.25

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Table 3.10 Available creep data for austenitic steels: plates; heat-resistant steels and Ni alloys; castings; seamless tubes/pipes

Product form Plates Heat-resisting steels and Ni-alloys Castings Seamless tubes/pipesEN 10028-7 EN 10095 EN 10213-4 EN 10216-5

1% strain creep Creep rupture 1% strain creep Creep rupture Creep rupture Creep rupturestrength strength strength strength strength strength

Material grade T range Duration T range Duration T range Duration T range Duration T range Duration T range Duration(°C) (h) (°C) (h) (°C) (h) (°C) (h) (°C) (h) (°C) (h)

X 2 CrNi 18 9X 5 CrNi 18 10GX 5 CrNi 19 10 550–700 10k, 100kX 6 CrNi 18 10 500–750 10k 100k 500–750 10k, 30k, 500–700 10k, 100k,

50k, 100k,200k (750) 200k

X 6 CrNi 23 13 550–800 10k 100k 550–800 10k, 100kX 6 CrNi 25 20 600–910 10k, 30k,

50k, 100k,150k, 200k,250k

X 8 CrNi 25 21 600–900 1k 10k 600–900 1k, 10k,100k

X 12 CrNi 600–900 1k 10k 600–900 1k, 10k,23 13 100kX 2 CrNiN18 10X 2 CrNiMo17 12 2X 3 CrNiMo17 13 3

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X 5 CrNiMo17 12 2GX 5 CrNiNb 550–700 10k, 100k 19 11GX 5 CrNiMo 550–700 10k, 100k19 11 2X 6 CrNiNb18 10 A1000X 6 CrNiNb18 10 A1100X 6 CrNiTi18 10 A1000X 6 CrNiTi18 10 A1100X 6 CrNiMo 550–700 10k, 100k17 13 2X 7 CrNiNb 540–700 10k, 100k,18 10 200k*X 7 CrNiTi 550–800 10k, 100k18 10




Material gradeT range Duration T range Duration T range Duration T range Duration T range Duration T range Duration(°C) (h) (°C) (h) (°C) (h) (°C) (h) (°C) (h) (°C) (h)

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X 8 CrNiTi 600–800 1k 10k 600–800 1k,10k,18 10 100k

X 8 CrNiNb 19 11X 8 CrNiNb 580–750 10k, 100k 580–750 10k, 100k, 580–750 10k, 100k,16 13 200k 200kX 12 CrCoNi21 20 AX 12 CrCoNi21 20 PX 15 CrNiSi 600–800 1k 10k 600–900 1k, 10k,20 12 100kX 15 CrNiSi 600–1000 1k 10k 600– 1k, 10k,25 21 1000 100kX 15 CrNiSi 500–900 1k 10k 500–900 1k 10k25 4 AGX 25 CrNiSi18 9GX 25 CrNiSi20 14GX 40 CrNiSi27 4GX 40 CrNiSi22 10GX 40 CrNiSi25 12GX 40 CrNiSi25 20X 2 CrNiMoN17 11 2X 2 CrNiMoN17 13 3

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X 3 CrNiMoN17 13 3X 5 CrNiNbN18 10X 5 CrNiMoB17 13 3X 6 CrNiTiB 550–700 10k, 100k, 550–700 10k, 100k,18 10 200k 200kX 6 CrNiMoTi17 12 2X 6 CrNiMoB17 12 2X 8 CrNiMo 580–750 10k, 100k,Nb 16 16 200kX 8 CrNiNbN25 21X 10 CrNiCuNb18 10X 12 CrNiWTiB16 13 AX 25 CrMnNiN 700–900 1k, 10k 700–900 1k, 10k25 9 7





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GX 40 CrNiSiNb 24 24X 3 CrNiMoBN 550–800 10k, 100k, 550–800 10k, 100k,17 13 3 200k (200k)X 6 CrNiWNbN16 16X 6 CrNiWNbN16 16 WWX 6 CrNiMoTiB17 13 AX 6 CrNiSiNCe 600– 1k, 10k, 600– 1k, 10k,19 10 (1000) 100k (1000) 100kX 7 CrNiMoBNb 16 16X 8 CrNiMoV 580–650 10k, 100kNb 16 13 200kX 9 CrNiSiNCe 600– 1k, 10k, 600– 1k, 10k,21 11 2 (1000) 100k (1000) 100kX 10 CrNiMnNbV 15 10 6 1X 12 CrNiWTiB16 13 WWX 10 CrNiMoMn 600–780 10k, 100kNbVB 15 10 1 200k, 250k*X 10 NiCr 600– 1k 10k 600–900 1k, 10kSi 35 19 (1000)X 12 NiCr 600–900 1k 10k 600– 1k, 10k,Si 35 16 100k 1000 100kGX 35 NiCrSi 25 21GX 40 NiCrSi 35 17

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GX 40 NiCrSi 38 19GX 40 NiCrSi 35 26GX 40 NiCrNb 45 35GX 50 NiCrCo20 20 20X 5 NiCr 600–700 10k 500–700 10k, 100k, 500–700 10k, 100kAlTi 31 20 (100k) (200k) (200k)X 5 NiCrAlTi 550–700 10k 500–700 10k, 100k, 500–700 10k, 100k31 20 (RA) (100k) (200k) (200k)X 6 NiCrNb 800– 10k, 100kCe 32 27 (1000)X 8 NiCr 700– 10k 100k 700– 10k, 30k 700– 10k, 100kAlTi 32 21 1000 1000 100k, 200k 1000 (200k)X 8 NiCrAlTi32 21 (RK)X 10 NiCr 600–900 1k 10k 600 –900 1k, 10k,AlTi 32 21 100kX 10 NiCr 600– 1k 10k 600–900 1k, 10kSiNb 35 22 (1000)





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GX 10 NiCrSiNb 32 20GX 40 NiCrSiNb 35 18GX 40 NiCrSiNb 38 19GX 40 NiCrSiNb 35 26GX 50 NiCrCoW 35 25 15 5X 6 NiCrS 600– 1k, 10k 600– 1k, 10kiNCe 35 25 1000 100k 1000 100kX 2 NiCrMoNbBN 25 22X 6 NiCrTiMoVB 25 15 2X 8 NiCrMoNbBN 25 20

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128Table 3.11 Available creep data for austenitic steels: forgings; bars for fasteners and bolts; castings; creep-resistant steels

Product form Forgings Bars for fasteners and bolting Castings Creep-resistant steelsEN 10222-5 EN 10269 EN 10295 EN 10302

Creep rupture 1% strain creep Creep rupture 1% strain creep Creep rupture 1% strain creep Creep rupture



X 2 CrNi 18 9 550–700 10k, 100kX 5 CrNi 18 10 550–700 10k, 100kGX 5 CrNi 19 1X 6 CrNi 18 10 550–700 10k, 100k, 550–700 10k, 550–700 10k, 100k,

200k 100k 200kX 6 CrNi 23 13X 6 CrNi 25 20X 8 CrNi 25 21X 12 CrNi 23 13X 2 CrNiN 18 10X 2 CrNiMo17 12 2X 3 CrNiMo 540–700 10k, 100k,17 13 3 200kX 5 CrNiMo 540–700 10k, 100k,17 12 2 200kGX 5 CrNiNb19 11GX 5 CrNiMo19 11 2X 6 CrNiNb 540–700 10k, 100k,18 10 A1000 200kX 6 CrNiNb 540–700 10k, 100k,18 10 A1100 200k

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X 6 CrNiTi18 10 A1000X 6 CrNiTi 540–700 10k 100k18 10 A1100 200kX 6 CrNiMo17 13 2X 7 CrNiNb 540–700 10k 100k18 10 200kX 7 CrNiTi 18 10X 8 CrNiTi 18 10X 8 CrNiNb 19 11X 8 CrNiNb 580–750 10k 580– 10k, 100k,16 13 100k 750 200kX 12 CrCoNi 550–850 10k 550– 10k, 100k,21 20 A 100k 900 200kX 12 CrCoNi 550–850 10k 550– 10k, 100k,21 20 P 100k 900 200kX 15 CrNiSi 20 12X 15 CrNiSi 25 21X 15 CrNiSi 25 4 AGX 25 CrNiSi 18 9 600–900 10k 600–900 100, 1kGX 25 CrNiSi 600–900 10k20 14GX 40 CrNiSi 600– 10k 600–900 100, 1k27 4 1000GX 40 CrNiSi 600–900 10k22 10GX 40 CrNiSi 700– 10k 700– 100, 1k25 12 1000 1000GX 40 CrNiSi 700– 10k 700– 100, 1k25 20 1100 1100X 2 CrNiMoN 550–700 10k 100k17 11 2 200kX 2 CrNiMoN 550–700 10k 100k17 13 3 200k

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X 3 CrNiMoN 550–800 10k, 100k,17 13 3 200kX 5 CrNiNbN18 10X 5 CrNiMoB17 13 3X 6 CrNiTiB 550–700 10k, 100k, 550–700 10k, 100k,18 10 200k 200kX 6 CrNiMo 540–700 10k, 100k,Ti 17 12 2 200kX 6 CrNiMo 550–850 10k, 100k,B 17 12 2 200kX 8 CrNiMo 580–750 10k, 580–750 10k, 100k,Nb 16 16 100k 200kX 8 CrNiNbN 25 21X 10 CrNiCuNb 18 10X 12 CrNiW 600–750 10k, 600–750 10k, 100kTiB 16 13 A 100kX 25 CrMnNiN 25 9 7GX 40 Cr 700– 10k 700– 100, 1kNiSiNb 24 24 1000 1100






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X 3 CrNiMo 550–800 10k, 100k, 550–800 10k, 100k,BN 17 13 3 200k 200kX 6 CrNiWN 580–750 10k, 580–750 10k, 100k,bN 16 16 100k 200kX 6 CrNiWNbN (550)– 10k, 100k (550)– 10k, 100k16 16 WW 650 650X 6 CrNiMoTiB 600–700 10k, 100k 600–700 10k, 100k17 13 AX 6 CrNiSiNCe 19 10X 7 CrNiMo 580–670 10k, 580–670 10k, 100kBNb 16 16 100kX 8 CrNiMo 580–650 10k, 100k 580– 10k, 100k,VNb 16 13 650 200kX 9 CrNiSiNCe 21 11 2X 10 CrNiMnNbV 15 10 6 1X 12 CrNiWTiB 500–700 10k, 500– 10k, 100k,16 13 WW 100k, 700 200k

200kX 10 CrNiMo 550–650 10k,

100k,MnNbVB 15 10 1 200kX 10 NiCrSi 35 19X 12 NiCrSi 35 16GX 35 NiCr 700– 10k 800–1100 1kSi 25 21 1000GX 40 NiCr 700– 10k 700–1100 100 1kSi 35 17 1000GX 40 NiCr 700– 10k 700–1100 100 1kSi 38 19 1100GX 40 NiCr 700– 10kSi 35 26 1000GX 40 NiCr 1000 10k 900–1100 100 1kNb 45 35

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GX 50 NiCr 900– 10k 800–1000 100 1kCo 20 20 20 1100X 5 NiCr 600–700 10k, 500–700 10k, 100k,AlTi 31 20 (100k) (200k)X 5 NiCr 550–700 10k, 500–700 10k, 100k,AlTi 31 20 (RA) (100k) (200k)X 6 NiCrNbCe 32 27X 8 NiCr 700– 10k, 100k 700– 10k, 100k,AlTi 32 21 1000 1000 (200k)X 8 NiCrAlTi 32 21 (RK)X 10 NiCrAlTi 32 21X 10 NiCrSiNb 35 22GX 10 NiCr 700– 10k 700– 100 1kSiNb 32 20 1000 1000GX 40 NiCr8 700– 100 1kSiNb 35 1 1100GX 40 NiCr 700– 10k 700–900 100 1kSiNb 38 19 1000GX 40 NiCr 700– 10k 700–1 100 1kSiNb 35 26 1100 1100






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GX 50 NiCrCo 1000– 10kW 35 25 15 5 1100X 6 NiCrSiNCe 35 25X 2 NiCrMoNbBN 25 22X 6 NiCrTiMo 500–650 10k, 100k 500–650 10k, 100k 500–650 10k, 100k 500– 10k, 100kVB 25 15 2 650X 8 NiCrMoNbBN 25 20

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Table 3.12 Available creep data for austenitic steels: consumables; genericproducts

Product form Consumables Generic

Material gradeEN 14532-2 ECCC

Creep rupture Creep rupture strength

T range (°C) Duration (h)

X 2 CrNi 18 9 550–720 (10k)–(100k)X 5 CrNi 18 10GX 5 CrNi 19 10X 6 CrNi 18 10 x 500–(750) 10k, 100k, 200kX 6 CrNi 23 13X 6 CrNi 25 20X 8 CrNi 25 21 600–910 10k, 30k, 50k, 100k,

150k, 200k, 250kX 12 CrNi 23 13X 2 CrNiN 18 10 580–800 10k, 100k, 200kX 2 CrNiMo 17 12 2 500–700 10k, 30k, 100k, 200kX 3 CrNiMo 17 13 3X 5 CrNiMo 17 12 2 x 500–850 10k, 30k, 100k, 200kGX 5 CrNiNb 19 11GX 5 CrNiMo 19 11 2X 6 CrNiNb 18 10 A1000 x 540–710 10k 30k, 50k, 100k,

150k, 200k, 250kX 6 CrNiNb 18 10 A1100 x 540–730 10k 30k, 50k, 100k,

150k, 200k, 250kX 6 CrNiTi 18 10 A1000 550–700 10k 100k 200kX 6 CrNiTi 18 10 A1100 540–720 10k 30k, 50k, 100k,

150k, 200k, 250kX 6 CrNiMo 17 13 2X 7 CrNiNb 18 10X 7 CrNiTi 18 10 x (A1000)

x (A1100)X 8 CrNiTi 18 10X 8 CrNiNb 19 11 600–750 10k, 100kX 8 CrNiNb 16 13 580–750 10k, 100k, 200k (1%)X 12 CrCoNi 21 20 AX 12 CrCoNi 21 20 PX 15 CrNiSi 20 12X 15 CrNiSi 25 21X 15 CrNiSi 25 4 AGX 25 CrNiSi 18 9GX 25 CrNiSi 20 14GX 40 CrNiSi 27 4GX 40 CrNiSi 22 10GX 40 CrNiSi 25 12GX 40 CrNiSi 25 20X 2 CrNiMoN 17 11 2X 2 CrNiMoN 17 13 3 550–750 10k, 100k, 200kX 3 CrNiMoN 17 13 3 xX 5 CrNiNbN 18 10 xX 5 CrNiMoB 17 13 3 xX 6 CrNiTiB 18 10

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X 6 CrNiMoTi 17 12 2 580–820 10k, 30k, 100k, 200kX 6 CrNiMoB 17 12 2X 8 CrNiMoNb 16 16X 8 CrNiNbN 25 21 580–770 10k, 30k, 100kX 10 CrNiCuNb 18 10 600–750 10k, 100kX 12 CrNiWTiB 16 13 AX 25 CrMnNiN 25 9 7GX 40 CrNiSiNb 24 24X 3 CrNiMoBN 17 13 3 550–800 10k, 100k, (200k)X 6 CrNiWNbN 16 16X 6 CrNiWNbN 16 16 WWX 6 CrNiMoTiB 17 13 AX 6 CrNiSiNCe 19 10X 7 CrNiMoBNb 16 16X 8 CrNiMoVNb 16 13X 9 CrNiSiNCe 21 11 2 550–1100 10k, 20k, 100k, 200kX 10 CrNiMnNbV 15 10 6 1 xX 12 CrNiWTiB 16 13 WWX 10 CrNiMoMnNbVB 15 10 1 x 600–790 10k, 30k, 100k, 200kX 10 NiCrSi 35 19X 12 NiCrSi 35 16GX 35 NiCrSi 25 21GX 40 NiCrSi 35 17GX 40 NiCrSi 38 19GX 40 NiCrSi 35 26GX 40 NiCrNb 45 35GX 50 NiCrCo 20 20 20X 5 NiCrAlTi 31 20 x 500–700 10k, 100k, (200k)X 5 NiCrAlTi 31 20 (RA) 500–700 10k, 100k (200k)X 6 NiCrNbCe 32 27 580–950 (10k–100k)X 8 NiCrAlTi 32 21 550–1000 10k, 20k, 30k, 50k,

70k, 100k, 150k, 200k,250k, 300k

X 8 NiCrAlTi 32 21 (RK) 600–1000 10k, 20k, 30k, 50k,70k, 100k, 150k, 200k,250k, 300k (1%)

X 10 NiCrAlTi 32 21X 10 NiCrSiNb 35 22GX 10 NiCrSiNb 32 20GX 40 NiCrSiNb 35 18GX 40 NiCrSiNb 38 19GX 40 NiCrSiNb 35 26GX 50 NiCrCoW 35 25 15 5X 6 NiCrSiNCe 35 25X 2 NiCrMoNbBN 25 22 600–820 10k, 30k, 100kX 6 NiCrTiMoVB 25 15 2X 8 NiCrMoNbBN 25 20 580–770 10k, 30k, 100k

Duration or temperature in brackets ( ) represent strength values after extendedextrapolation in stress.Duration or temperature with an asterisk * represent strength values after extendedextrapolation in time.For steels assessed by ECCC the symbol (1%) means that a 1% strain creep strengthassessment is also available.EN 14532-2 does not report strength data but does report assessed equations relatingstrength, duration and temperature.In the EN 14532-2 column, the remark ‘x notch’, means that an equation for creepnotch rupture strength is also provided.

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136Table 3.13 Chemical analysis of nickel and cobalt base alloys contained in European Standards

Gross chemical composition

Material grade Heat Concentration in mass% of the following elements

treatment C Cr Co Fe Mo Ni Other

G NiCr15 As cast 0.35–0.65 12.0–18.0 – Balance <1 58–66G NiCr28W As cast 0.35–0.55 27–30 – Balance <0.5 47–50 W 4–6G NiCr50Nb As cast <0.1 48–52 – <1 <0.5 Balance Nb 1–1.8. N<0.16NiCo20Cr20MoTi P 0.04–0.08 19–21 19–21 <0.7 5.6–6.1 Balance Al 0.3–0.6; Ti

1.9–2.4; BNiCr15Fe A 0.05–0.1 14–17 <1.5 6.0–10.0 – Balance Al. TiNiCr15Fe7TiAl AT P <0.08 14–17 <1 5.0–9.0 – Balance Al 0.4–1; Ti

2.25–2.75;Nb+Ta 0.7–1.2;

NiCr15Fe7TiAl P 980 <0.08 14–17 <1 5.0–9.0 – Balance Al 0.4–1; Ti2.25–2.75;Nb+Ta 0.7–1.2;

NiCr15Fe7TiAl P 1170 <0.08 14–17 <1 5.0–9.0 – Balance Al 0.4–1; Ti2.25–2.75;Nb+Ta 0.7–1.2;

NiCr19Fe19Nb5Mo3 P 0.02–0.08 17–21 <1 Balance 2.8–3.3 50–55 Al 0.3–0.7; Nb+Ta4.7–5.5; Ti 0.6–1.2;B 0.002–0.006

NiCr20Co13Mo4 P 0.02–0.1 18–21 12.0–15 <2 3.5–5 Balance Al 1.2–1.6; TiTi3Al 2.8–3.3; B; Zr;

NiCr20Co18Ti P <0.13 18–21 15–21 <1.5 – Balance Al 1–2; Ti 2–3; Zr; BNiCr20Ti AT 0.08–0.15 18–21 <5 <5 – Balance Al; Ti 0.2–0.6;NiCr20TiAl AT P 2 stage 0.04–0.1 18–21 <1 <1.5 – Balance Al 1–1.8; B <0.008.

Ti 1.8–2.7

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NiCr20TiAl P 3 stage 0.04–0.1 18–21 <1 <1.5 – Balance Al 1–1.8; B <0.008.Ti 1.8–2.7

NiCr22Fe18Mo AT 0.05–0.15 20.5–23 0.5–2.5 17–20 8.0– 10 Balance Al <0.5; B; W 0.2–1NiCr22Mo9Nb A 0.03–0.1 20–23 <1 <5 8.0–10.0 Balance Al <0.4; Nb+Ta

3.15–4.15NiCr23Co12Mo AT 0.05–0.1 20–23 11.0–14 <2 8.5–10 Balance Al 0.7–1.4; Ti

02–0.6; BNiCr23Fe AT 0.03–0.1 21–25 <1.5 <18 – Balance Al 1–1.7; Ti

<0.4; B<0.006NiCr25Co20TiMo P 1100 0.03–0.07 23–25 19–21 <1 1–2.0 Balance Al 1.2–1.6; Nb+Ta

0.7–1.2; Ti 2.8–3.2;B 0.01–0.015; Zr0.03–0.07; Ta

NiCr25FeAlY AT 0.15–0.25 24–26 – 8.0–11 – Balance Al 1.8–2.4; Ti0.1–0.2; Y0.05–0.12; Zr

NiCr26MoW AT 0.03–0.08 24–26 2.5–4 Balance 2.5–4 44–47 W 2.5–4NiCr28FeSiCe AT 0.05–0.12 26–29 <1.5 21–25 – Balance Ti <0.5; Si 2.5–3;

Ce 0.03–0.09NiCr29Fe AT <0.05 27–31 – 7.0–11 – Balance Al <0.5CoCr20W15Ni AT 0.05–0.15 19–21 Balance <3 – 9.0–11 W 14–16G CoCr28 As cast 0.05–0.25 27–30 48–52 Balance <0.5 <4 Nb <0.5

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138Table 3.14 Available creep data for nickel and cobalt base alloys: heat-resistant steels and Ni alloys; bars for fasteners and bolts; castings

Product form Heat-resisting steels and Ni-alloys Bars for fasteners and bolting CastingsEN 10095 EN 10269 EN 10295

Heat 1% strain creep Creep rupture 1% strain creep Creep rupture Creep rupture Creep rupturetreatment strength strength strength strength strength strength


G NiCr15 As cast 900–1000 100. 1kG NiCr28W As cast 700–1100 10k 800–1100 1kG NiCr50Nb As cast 700–1000 10k 700–1100 100, 1kNiCo20Cr P20MoTiNiCr15Fe A 500–900 10k, 100k 600–1000 1k, 10k,

100kNiCr15Fe7TiAl AT P 500–800 10k 100k 500–800 10k 100kNiCr15Fe7TiAl P 980NiCr15Fe7TiAl P 1170NiCr19Fe19 PNb5Mo3NiCr20Co P13Mo4Ti3AlNiCr20Co18Ti PNiCr20Ti AT 600–1000 1k, 10k,

100kNiCr20TiAl AT P 2 500–800 10k 100k 500–800 10k 100k

stageNiCr20TiAl P 3 stage

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NiCr22Fe18Mo ATNiCr22Mo9Nb A 700–900 1k, 10kNiCr23Co12Mo ATNiCr23Fe AT 600–1000 1k, 10k,

100kNiCr25Co P 110020TiMoNiCr25FeAlY ATNiCr26MoW ATNiCr28FeSiCe AT 700–1000 10k, 100k 700–1000 10k, 100kNiCr29Fe ATCoCr20W15Ni ATG CoCr28 As cast 700–1100 10k 900–1000 100, 1k

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140Table 3.15 Available creep data for nickel and cobalt base alloys: creep-resistant steels; generic products

Product form Creep-resistant steels GenericEN 10302 ECCC

Material gradeHeat 1% strain creep strength Creep rupture strength Creep rupture strength

treatment T range (°C) Duration (h) T range (°C) Duration (h) T range (°C) Duration (h)

G NiCr15 As castG NiCr28W As castG NiCr50Nb As castNiCo20Cr20MoTi P 500–900 10k, 100k 500–900 10k, 100kNiCr15Fe ANiCr15Fe7TiAl AT PNiCr15Fe7TiAl P 980 650–800 10k, 100k 650–800 10k, 100kNiCr15Fe7TiAl P 1170 500–800 10k, 100k 500–800 10k, 100kNiCr19Fe19Nb5Mo3 P 500–800 10k, 100k 500–800 10k, 100kNiCr20Co13Mo4Ti3Al P 650–800 10k, 100kNiCr20Co18Ti P 700–800 10k, 100k 550–800 10k, 100kNiCr20Ti ATNiCr20TiAl AT P 2 500–800 10k, 100k 500–800 10k, 100k 600–900 10k, 30k, 100k, 200k

stageNiCr20TiAl P 3 stage 450–670 10k, 30k, 100k, 200kNiCr22Fe18Mo AT 600–1000 10k, 100k 550–1000 10k, 100kNiCr22Mo9Nb ANiCr23Co12Mo AT 650–1000 10k, 100k 580–1000 10k, 100k 580–950 10k 30k 100kNiCr23Fe ATNiCr25Co20TiMo P 1100 600–(900) 10k, 100k 550–900 10k, 100kNiCr25FeAlY AT 650–1200 10k, 100k 650–1200 10k, 100kNiCr26MoW AT 600–1050 10k, (100k) 600–1000 10k, (100k)NiCr28FeSiCe AT

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NiCr29Fe AT 700–1050 10k, 100k 700–1050 10k, 100kCoCr20W15Ni AT 700–1000 10k, (100k) 700–1000 10k, (100k)G CoCr28 As cast

Duration or temperature in brackets ( ) represent strength values after extended extrapolation in stress.Duration or temperature with an asterisk * represent strength values after extended extrapolation in time.For steels assessed by ECCC the symbol (1%) means that a 1% strain creep strength assessment is also available.EN 14532-2 does not report strength data but does report assessed equations relating strength, duration and temperature.In the EN 14532-2 column, the remark ‘x notch’, means that an equation for creep notch rupture strength is also provided.

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Although these marks provide obviously significant information, it is notclear how design should consider them with regard to safety factors, or howand if the marked strengths can be further extrapolated. The challenge tostate guidelines on the handling of ‘different reliability’ strength values is atarget of new design standards currently under development.

In some cases, the standard also provides details about the scatter band ofthe data set used and the assessment method. If the standard data wereassessed by ECCC, details of the data set, the collation specification (mechanicaldata, chemical composition, quantity and duration of tests, etc) and a straightequation for the chosen assessment (relating rupture time, temperature andstrength) can be found in the ECCC Data Sheets.23,24

For bolting materials, stress relaxation may also become a design issue aswell as the creep properties themselves. Therefore some relaxation propertieshave been collated and published, in EN 10269 and the ECCC Data Sheets.23,24

Table 3.16 gives an overview of available materials and the scope of relaxationstrengths included.

3.5.3 Welding consumables and qualification

Welding consumables for creep-resistant steels are now classified, like allothers, according to:

• EN 12070 (Welding consumables – Wire electrodes and wires and rodsfor arc welding of creep resisting steels – classification);

• EN 12072 (Welding consumables – Wire electrodes and wires and rodsfor arc welding of stainless and heat resisting steels – classification);

• EN 12536 (Welding consumables – Rods for gas welding of non-alloyand creep resisting steels – classification);

• EN 1599 (Welding consumables – Covered electrodes for manual arcwelding of creep resisting steels – classification); or

• EN 1600 (Welding consumables – Covered electrodes for manual arcwelding of stainless and heat resisting steels – classification).

The general welding procedure qualification comes under EN ISO 15613or EN ISO 15614. The qualification of the consumables themselves is regulatedby EN 14532 (Welding Consumables – Test methods and quality requirements)and a particular, creep-related ‘supplementary prescription’ is included in:EN 14532-2 (Welding Consumables – Test methods and quality requirements– Part 2: Supplementary methods and conformity assessment of consumablesfor steel, nickel and nickel alloys). Consumables intended for use in thecreep regime, defined numerically by a temperature range for ferritic–martensitic, austenitic and nickel-base materials, shall be qualified for themaximum service temperature by comparison tests with a ‘sufficiently similar’parent material: a stress rupture all weld material creep test series of at least

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urope143

Table 3.16 Relaxation strength in European Standards

Gross chemical composition EN 10269 ECCC

HeatConcentration in mass% of the following elements Relaxation strength Creep rupture strength

Material gradetreat- C Cr Mo Ni V Other T range Strain range Duration T range Strain Durationment (°C) (%) (h) (°C) range (%) (h)

42 CrMo 5 6 QT 0.39–0.45 1.2–1.5 0.5–0.7 – – 350–500 0.15 1k, 10k, 30k40 CrMoV 4 6 QT 0.26–0.44 0.9–1.2 0.5–0.65 – 0.25–0.35 400–500 0.15 1k, 10k, 30k21CrMoV 5 7 QT 0.17–0.25 1.2–1.5 0.55–0.8 <0.6 0.2–0.35 300–540 0.2 1k, 10k, 30k20 CrMoVTiB QT 0.17–0.23 0.9–1.2 0.9–1.1 <0.2 0.6–0.8 B 0.001– 400–600 0.15 1k, 10k, 30k4 10 0.010; Ti

0.007–0.15

X 22 CrMoV QT 0.18–0.24 11–12.5 0.8–1.2 0.3–0.8 0.25–0.35 400–580 0.2 1k, 10k12 1 30k,X 19 CrMo QT 0.17–0.23 10–11.5 0.5–0.8 0.2–0.6 0.1–0.3 Nb 0.25– 400–600 0.2 1k, 10k, 430–590 0.15 10k, 30k.,NbVN 11 1 0.55; N 30k 100k,

0.05–0.1 200kX 10 CrNiMoMn 0.07–0.13 14.0–16.0 0.8–1.2 9.0–11.0 – Mn 5.5–7; 550–700 0.15 10kNbVB 15 10 1 B 0.003–

0.009;V 0.15–0.4;Nb 0.75–1.25;N<0.11

NiCr20TiAl AT P 2 0.04–0.1 18–21 <1 <1.5 – Al 1–1.8; 450–640 0.15 1k, 10k, 30kstage B <0.008. 650–750 0.15– 30k

Ti 1.8–2.7 0.2 1k, 10k30k

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four tests has to be produced with the longest test not shorter than 10 000 h,one test above 1000 h and one between 50 and 250 h; tests shorter than50 h are not considered. The regression line of this test series is then comparedwith the ±20% scatter band around the mean rupture strength lines of thereference parent material, for which EN 14532-2 Appendix E gives theequations based on Manson–Brown-parameter approaches following thediscontinued ISO 6303 Appendix. The source of the coefficients to the equationsis the British Guideline PD 6525.

All weld material specimens – according to EN 14532-2 – are expected tofall within the scatter band of the matching parent material. If this is not thecase, as is likely for longer testing times, detailed warnings about the longeror shorter duration have to be stated on the consumables qualification certificate.To date, EN 14532-2 is the only standard that includes creep strength equations.Some materials are marked with an ‘x’ in Tables 3.5 to 3.15 becausethe equation applicability is not just related to the temperature–duration plane.

3.5.4 Testing and testing standards

A side product of the ECCC WG1 Creep Data Generation and AssessmentProcedures and of the EC-funded Standard Measurement and Testing (SMT)Projects conducted in the 1990s16,17 produced a group of testing procedures,which:

• harmonised the rules of the main National Standards (BS 3500, DIN50118, NF A03-355, ASTM E139/E292, UNI 5111, etc);

• collated the results from the relevant SMT projects; and• critically overviewed the creep testing procedures of 15 leading European

labs, as laid down in Reference 19 and Reference 20: Vol 3 part I.

These recommendations were taken up by CEN/ECISS TC1 and became,with slight modifications, the standards:

• EN 10291:2000 (Metallic materials – Uniaxial creep testing in tension),which included stress-rupture, creep-rupture and creep testing with bothinterrupted and non-interrupted strain measurement;

• EN 10319:2003 (part 1) – 2006 (part 2) (Metallic materials – Tensilestress relaxation testing – Part 1: Procedure for testing machines andPart 2: Procedure for bolted joint models);

• CWA 15261-3:2005 (Measurement uncertainties in mechanical tests onmetallic materials – Part 3: The evaluation of uncertainties in creeptesting). This standard has the support and is based on the ground rules asestablished by the EC funded projects UNCERT and UNCERT-AM.26

EN 10291 is currently part of the discussions regarding revision ofinternational standard ISO 204 (Metallic materials – Uninterrupted uniaxial

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creep testing in tension – method of test). Creep testing on notched specimenshas been proposed for inclusion according to the new ECCC-recommendednotch geometry (Reference 19 and Reference 20: Vol. 3 part V; References27, 28) and methods following the High Temperature Mechanical TestingCommittee HTMTC Working Group on ‘Notch Creep Behaviour’.29,30

A particular testing technique (miniature specimen testing with the smallpunch method) is also under discussion in CEN WorkShop Agreement 21(Small Punch Test Method for Metallic Materials) with the support of EPERC;the issuance of CWA 15627 concluded WS21’s work).

Testing technique

Compared to standards prepared elsewhere, the new European creep testingnorms are based on the ECCC experience in refining testing requirementsand some additional testing options. EN 10291 and EN 10319 were designedto best suit long duration tests, while realising that although most relevant todesign and dimensioning, quality control laboratories would also use themfor the more common short-duration stress rupture testing.

• Test methods and strain measurement: in Europe a large amount of theavailable long-term creep data was obtained in stress rupture tests, i.e.tests in which the specimen is held at constant temperature and loaduntil fracture occurs. Such tests are often performed in multi-specimentesting machines, which may include up to ten strings each including tenor more specimens. Creep rupture testing, where creep strain iscontinuously measured during the test, is considered more expensiveowing to the need to maintain thermal and electrical stability of theextensometer, or because ambient temperature must be controlled toreduce thermal drift in the transducers. It must be taken into account thatduring the initial stages of long- and medium-term creep rupture tests,elongations of a few micrometres must be measured, to keep uncertaintyto a reasonable level. As a less expensive alternative,16 interrupted creeprupture tests are performed using techniques originally developed in the1940s31 and improved in Germany in the 1980s.32 During these tests, thespecimen is cooled, unloaded, its elongation measured, re-assembled,reloaded and re-heated at regular, generally logarithmically equi-spaced,time intervals. Although the ‘off line’ strain measurement, requiresparticular care and should be performed in a suitable metrologicallaboratory, the practical comparison of results achieved in a round robintest on BCR reference materials ‘Alloy 75’ (CRM 425, NiCr20Ti EN10095) and ‘Durhete 1055’ (20 CrMoVTiB 4 10 EN 10269, 1%CrMoVTiB)showed,16 that interrupted and non-interrupted tests provide creep curvescontained in the same scatter-band of approximately ±25% in strain

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around the mean line. Owing to these premises, EN 10291 and probablythe future ISO 204 will regulate this alternative technique.

• Temperature: temperature stability is an essential parameter during long-term tests, and the main consideration here is distinction and considerationof the various factors contributing to thermal stability: EN 10291 definestemperature stability as the minimisation of the difference between truetest temperature and specified test temperature, where these definitionsaccount for all errors contained in the temperature measurement andregulation chain, i.e. regulation deviation plus random errors. Becauseof this, the allowable temperature differences are somewhat higher than,for instance, those in ASTM E139, where thermal stability has beenconsidered equivalent to indicated temperature minus specifiedtemperature, that is, to the regulation-induced deviations only.

Another topic concerning thermal stability is the calibration ofthermocouples, generally expensive PtRh–Pt thermocouples of type R,S or B, which unavoidably deteriorate during high temperature exposureowing to metallurgical instability.33 For long-duration tests, severalcalibration methods were compared and discussed34 and recommendationsfor both off-line and in situ calibration were stated (Reference 20: Vol 3part I). Both of these consider the temperature range and gradient towhich the thermocouple is exposed during its use, i.e. the depth to whichthe thermocouple is inserted into the testing furnace, which could bevery high in multi-specimen machines where discarding the exposedthermocouple part may be very expensive.

• Data assessment is subdivided into two distinct steps:(i) Single test raw data assessment consists of transforming elongation

time data into strain–time data by dividing elongation by the referencelength, which, as a function of the test piece geometry, may bedifferent from the cylindrical length because collars used for theextensometer fixation will contribute to the measured deformation.The reference length is therefore computed according to expectedcreep behaviour and the Norton stress exponent and may have to bere-assessed when the true results are obtained (Reference 20: Vol 3part I Appendix 1).

From the creep curve (strain versus time), initial plastic strain (ifpresent) and times to specified plastic strains are collated, for whichlinear interpolation in a bi-logarithmic strain-time diagram appearsto be best [20 Vol 3 Part I Appendix 1]. Small differences betweeninterrupted and non-interrupted testing due to inelastic recoverymay be negligible in most technical applications.

(ii) Data assessment of an entire test programme is described in Chapter6 by Holdsworth.

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Testing programmes

Setting up a testing programme for parent materials depends on what thedata are needed for. In material quality assurance, single specimen tests withsmooth or combined smooth-notched specimens are required. Productionqualification tests generally require an isothermal approach including four tosix specimens per temperature and a duration not shorter than 10 000 h,sometimes 30 000 h to meet the requirement to have an extrapolation factorbelow 3 when computing 100 000 h strengths. The ECCC approach tocommon testing programmes, intended to fill gaps identified in the currentlyavailable creep data population of a given material, also follows isothermalprogramme set-ups, for which durations of 1000, 3000, 10 000, 30 000 and70 000 h are requested. To fully meet significance criteria suitable forstandardisation, the data additionally have to meet the requirements listed inTable 3.3.

In situations where creep properties must be verified in a relatively shortperiod of time, such as during residual life assessment or, in some cases,particular material appraisal according to PED for non-harmonised materials,alternative methods are adopted, although these are generally not standardisedand strongly based on the individual assessor’s experience. In such cases‘iso-stress’, ‘MPC-Ω’ approaches35,36 or reduced test programmes whichproduce an estimate of a Larson–Miller master curve are used. The effectivenessand credibility of these tests, although in some cases clearly demonstrated,has proven to be strongly dependent on the assessor’s skill and experiencewith the testing procedure and the material tested (Reference 20: Vol 5 partIII).

Component creep tests

In the past, a considerable number of near- and full-scale tests for componentshave been developed in Europe. As a consequence, procedures to test close-to-component specimens, generally tubes and pipes, have also been developed.Owing to the complexity of the testing devices and procedures, few testingfacilities are available, although an impressive number of tests has beencollated (Reference 20: Vol. 9 part III). In addition, codes of practice fortubular specimens under internal pressure and combined internal pressureand tension or torsion have been published.37,38

Relaxation testing

The use of bolting at high temperature has to take into account the loss inpre-loading with time (relaxation) owing to creep effects under constanttotal strain. Various testing techniques have been proposed:

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• Uniaxial relaxation testing, keeping a standard specimen under totalstrain control, i.e. keeping its elasto–plastic strain constant over time(generally to a value comparable to the pre-load of a bolt, ca. 0.1–0.5%strain). Two testing approaches are used:

– Isothermal relaxation: the specimen is heated to test temperature,and loaded to the foreseen strain, which is then kept constant, whilethe loss in stress over time is recorded. Severe technical testingissues need to be addressed during such tests,17 for instance closed-loop testing systems proved to be very difficult to keep stable overlong periods of time, when the stress decrease becomes small.

– Non-isothermal relaxation: in this case the specimen is loaded atroom temperature with the required total strain and is then heated tothe testing temperature.

• Model bolt relaxation testing, developed and already standardised inGermany,39 consists of preloading a series of bolt-flange assemblies atroom temperature and then tightening the screws to a known and controlledpre-load (generally measuring the bolt shaft strain using strain gauges).The assemblies are then put into a furnace and heated. One by one thebolt-flange assemblies are extracted at specified time intervals, cooledand loosened, using strain gauges to measure the remnant bolt pre-load.

A comparison of the two methods was conducted on a typical bolting steel(“X19” = X19CrMoNbVN 11 1 EN 10269, 11%CrMoNbNV) under theSMT project.17 Results showed that for long periods of time (ca. 10 000 h),the remnant stress measured using the two different methods are comparable.For both methods, detailed guidelines are found in ECCC Recommendations2005, Vol 320 and in EN 10319, parts 1 and 2.

Design standards

At present there are no generally available European standards for designthat directly address components which operate under European creep regimes.Nevertheless, considerable work is underway for the new standards EN 13445part 3 (Unfired pressure vessels – Part 3: Design) and EN 13480 part 3(Metallic industrial piping – Part 3 Design) and their various appendices. Tobe in line with PED, the new standards will, besides traditional creep designmethods, also have to cover the dimensioning of welds, the use of finiteelement simulation techniques, the applicability of fracture mechanics andthe consideration of statistics and/or risk-based methods (CEN Workshop 22on FITNET and CEN Workshop 24 on Risk Based Inspection andMaintenance). The related CEN-TCs, EPERC-TP and ECCC are workingtowards compiling collations of reliable creep strength equations, weld creepstrength reduction factors and other material related design features.

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3.5.5 Specifications of creep-resistant steel users

The market for plant building in refineries, chemical plants and power unitshas changed since the publication of the PED with the breakup and privatisationof monolithic electricity boards and mergers/closures of big boilermakers.Plant building now tends to be in the hands of general engineering companies,which have significantly less skill and experience in handling creep problemsand the materials suitable for such operating conditions. In addition, plantowners are asking for longer and longer service periods, so that creep strengthvalues of 200 000 h or 250 000 h are no longer the exception. As a consequence,the plant builders’ specifications for creep-resistant steels tend to refer as faras possible to any available standards and they ask material, componentfitting or spool suppliers to guarantee the values in accordance with PED andunder all possible configurations during assembly and operation.

Since long-term creep test results have up to now been collected andstored by specialist companies (boiler makers, turbine manufacturers, largefurnace and reactor builders) rather than the material producers, negotiationsto find a PED-conforming compromise are often very difficult. Even thespecialist companies are often unable to provide experimental results formaterials after complex forming and heat-treating operations (shapes, fittings,repaired castings, etc) or on welds after unusual heat treatments (for instancewelds subjected to more than three post-weld heat treatments owing to repairor close assembly to manufacturing welds).

As a consequence, the creep strength values of EN standards tend tobecome mandatory values in engineering companies’ specifications, sometimeslowered by a percentage apparently accounting for the usual ±20% scatterband in stress on rupture strength values. In most typical situations, wherethe material, component, fitting, spool or vessel manufacturer cannot complywith pre-existing data, ‘qualification’ creep test rows are requested, which,owing to the short time frame and the belief that ‘time–temperature parameterbased’ extrapolations are safe, are generally reduced to a few tests below1000 h, then grandly extrapolated to the required 200 000 h strengths. Validationprocedures, like ECCC’s PATs (Reference 20: Vol 5), are generally unknownand not applied.

3.5.6 Residual life assessment

Residual life assessment, that is, activities to decide whether a given plantcan be used beyond its original design life, is generally led by in-serviceinspection procedures, dictated by national rules and/or by more advancedsystems like RBI (risk based inspection). Nevertheless, some nationalregulations require that design reviews are also conducted, during whichpast as well as future service is checked against the most credible available

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creep strength values. This requires on the one side sound assessmentprocedures, because residual life assessment tends to ask for very long durations,longer than available and possible tests, and on the other side the refinedstrength values published in the various documentation for the standards.

Some guidelines on how to handle these particular needs were laid downin Reference 20, Vol 5 and are assessed in the CEN WorkShop Agreements22 (FITNET – European Procedure for Fitness for Service) and 24 (RiskBased Inspection and Maintenance Procedures – RIMAP) and in the newItalian standard for residual life assessment, UNI 11096: 2004.

3.6 Future trends

The new concept of CEN standards, including creep strength values, and thenew ECCC strength assessment procedure were a first step towards improvingthe reliability of creep strength and design values. Nevertheless, a hugeamount of work is still to be done:

• New materials are continuously offered by leading companies withmetallurgical skills, for instance the new 9–12%Cr steels and the newlow-alloyed steels of 2.25%Cr type, as well as many new austenitic andnickel base grades. These materials need big duration tests and thoroughassessment, but, since industrial users never have the time to wait forthese results to come through, credible creep strength reference values atleast are needed as soon as possible. The prescriptions of EN 12952-2and ECCC (Reference 20: Vol 5 part I), along with the data requirementsaccording to Table 3.3, make it clear that testing has to be taken further.Reliable assessments are only possible with very large, or at leaststrategically planned, Europe-wide coordinated test campaigns.

• ECCC and other bodies have also identified large gaps in the understandingof creep behaviour and data availability in consolidated materials, whichurgently need to be filled.

• Some materials, obsolete or no longer used for their original applications,may suddenly find new users or applications (for instance alloy 617 forpiping and forgings), so that data previously disregarded become relevantagain and may need to be up-dated.

• Cast steel grades in general, with the exception of the 9–12%Cr steels,have not been a priority for assessment.

• A very particular area, mainly the concern of refineries and petrochemicalplants, are the cast austenitic high-alloyed steels (i.e. G-X40CrNi25 20or G-X40NiCr35 25, partially included in EN 10295) and theirmodifications based on mixtures of Co, W and Nb. These have not yetbeen considered in detail.

• CEN will review all standards on a five-to-ten year basis, so that creepstrength values will be updated and upgraded in the light of on-going

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tests and new high-priority materials (for instance grades 23 (10CrWMoNbB 9 6), 24 (10 CrMoVTiB 10 10), 92 (X 10 CrWMoVNb 92), 911 (X 19 CrWMoVNb9 1 1), 709 (X 6 NiCrMoTiBN 25 20), etc).Some changes and refinements, particularly for those grades assessedwithout fully representative data sets or under reduced testing times (forinstance grade 91, X 9 CrMoVNNb 9 1) will need to be upgraded in thepublished strength values.

Besides the specifically material-related issues, there are other topics thatwill require strong collaboration between material specialists and designers,such as how to handle strength values of extended extrapolation, how toinclude time-dependent weld creep strength reduction factors, and so on.There is also an urgent need to harmonise the PED interpretation of creephandling in pressure vessel design, which, owing partly to a somewhat vagueformulation, has led to radically different interpretations among the notifiedbodies. Agreement and reciprocal understanding is also desirable betweenthe mandatory European PED and the forthcoming very advanced voluntarydesign codes (EN 13445, etc) and the well-established ASME, API andJapanese codes.

Finally a few basic issues should also be addressed, including how toextrapolate from data sets of limited size or duration in the most reliable wayand how to consider scatter or probabilistic effects for creep strength dataassessments. These should be addressed by more basic research projects.

In the past, Europe has shown that, despite the inherent difficulties,collaboration between different technical schools of thinking is possible andcan lead to outstanding results. The next step is to include non-Europeansand to enlarge the multi-disciplinarity of the interacting specialists.

3.7 References

1 French, Tucker ‘Flow in a low-carbon steel at various temperatures’ NIST reportT296, 1925; (re-published by Geil and Carwile, National Institute of Standards andTechnology (NIST) RP2329, 1952).

2 Pomp A. and Dahmen A., ‘Entwicklung eines abgekürzten Prüfverfahrens zurErmittlung der Dauerstandfestigkeit von Stahl bei erhöhten Temperaturen’, Mitteilunga. d. Kaiser-Wilhelm-Institut für Eisenforschung zu Düsseldorf, Abhandlung 72–95IX Band (1927), 33–52.

3 Parravano N. and Guzzoni G., ‘Prove statiche delle leghe ultraleggere’, La MetallurgiaItaliana, 1930, XXII, 367 ff.

4 DIN 50 112 (DVM-Creep Test), ca. 1930.5 German Iron and Steel Society (VDEh), Ergebnisse deutscher Standzeitversuche

langer Dauer, in collaboration with the German Creep Resistant Steel Committee(AGW) and the Committee for High Temperature Engineering Materials, VerlagEisen und Stahl, Düsseldorf, 1969.

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6 Former CECA/ECSC (European Carbon and Steel Collaboration): http://cordis.europa.eu/ecsc-coal/home.html, and actual http://cordis.europa.eu/coal-steel-rtd

7 COST website: www.cost.esf.org8 Information on EC programme Joule–Thermie: http://erg.ucd.ie/ttp.html9 Official PED-site http://ec.europa.eu/enterprise/pressure_equipment/ped/index_en.html

10 European Standardisation Body: www.cenorm.be11 European Committee for Iron and Steel Standardization ECISS: http://www.cenorm.be/

cenorm/aboutus/structure+/relations/workprogrammeeciss.asp12 European Collaborative Creep Committee website: www.etd1.co.uk/eccc/

advancedcreep13 EC funded Brite Euram Concerted Action, Creep –BE-5524, 1992–1996.14 EC funded Thematic Network, BET2-0509 - Weldcreep, 1997–2001.15 EC funded Framework V Thematic Network GTC2-2000-33051 Advanced Creep,

2001–2005.16 EC funded Standard–Measurement and Testing Project MAT1-CT-940065,

Development and Validation of a Code of Practice for a Low Cost Method of StrainMeasurement from Interrupted Creep Testing, 1994–1998.

17 EC funded Standard–Measurement and Testing Project MAT1-CT-940078,Development of Standard European Methodology for Stress Relaxation Testing ofMetals, 1994–1998.

18 International Conference on Creep & Fracture in High Temperature Components –Design and Life Assessment Issues, organized by ECCC, September 12–14, London,2005.

19 ECCC Recommendations 1994, ERA Technology Ltd, Leatherhead. Volume 1 (issue1) 1994: Creep Data Validation and Assessment Procedures – Overview, HoldsworthS.R. (ed.). Volume 2 (issue 2) 1994: Terms and Terminology for use with StressRupture, Creep and Stress Relaxation: Testing, Data Collation and Assessment, OrrJ. (ed.). Volume 3 (issue 2) 1994: Recommendations for Data Acceptability Criteriafor Creep, Creep Rupture, Stress Rupture and Stress Relaxation Data, HoldsworthS.R. and Granacher J. (eds). Volume 4 (issue 1) 1994: Guidance for the Exchangeand Collation of Creep Rupture, Creep Strain–Time and Stress Relaxation Data forAssessment Purposes, Merckling G. & Bullough C.K. (eds). Volume 5 (issue 1)1995: Guidance for the Assessment of Creep Rupture, Creep Strain snd RelaxationData, Holdsworth S.R. (ed.).

20 ECCC Recommendations 2005, ETD Ltd, Leatherhead. Volume 1 (issue 6) 2005:Creep Data Validation and Assessment Procedures – Overview, Holdsworth S.R.(ed.). Volume 2, 2005: Terms and Terminology for use with Stress Rupture, Creep,Stress Relaxation, Creep Crack Initiation and Multi-Axial Creep: Testing, DataCollation and Assessment, Morris P., Orr J., Servetto C., Seliger P., Holdsworth S.R.and Brown T.B. (eds). (a) part I (issue 8): Parent Material, (b) part Iia (issue 2):Welding Processes and Weld Configurations, (c) part Iib (issue 2): Weld CreepTesting, (d) part III (issue 4): Post Exposure Material, (e) part IV (issue 2): Generationand assessment of creep crack initiation data, (f) part V (issue 1): Generation andassessment of multi-axial feature specimen and component data. Volume 3, 2005:Recommendations for Data Acceptability Criteria and Generation for Creep, CreepRupture, Stress Rupture, Stress Relaxation, Creep Crack Initiation and Multi-AxialCreep Data, Holdsworth S.R., Granacher J., Klenk A., Buchmayr B., Gariboldi E.,Brett S., Merckling G., Müller F., Gengenbach T., Dean D. and Brown T.B. (eds). (a)

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part I (Issue 5): Generic recommendations , (b) part II (issue 3): Creep data forwelds, (c) part III (issue 4): Creep testing of post exposure (ex-service) material, (d)part IV (issue 2): Creep crack initiation, (e) part V (issue 1): Multi-axial featurespecimen and component test data. Volume 4, 2005: Guidance for the Exchange andCollation of Creep Rupture, Creep Strain–Time, Stress Relaxation, Creep CrackInitiation and Multi-Axial Creep Test Data for Assessment Purposes, Merckling G.,Calvano F., Bullough C.K., Holdsworth S.R. and Tonti A. (eds). (a) part I (issue 6):Creep rupture, strain–time and relaxation data; (b) part II (issue 1) Creep crackinitiation, (c) part III (tba); (d) part IV (issue 1): Creesty user manual. Volume 5,2005, Guidance for the Assessment of Creep Rupture, Creep Strain and StressRelaxation Data, Holdsworth S.R. and Merckling G., (eds). (a) Part Ia (issue 5):Full-size creep rupture data sets datasets, (b) Part Ib (issue 2): Creep strain and creepstrength data, (c) Part Ic (issue 2): Full-size stress relaxation datasets, (d) Part Iia(issue 1): Sub-size datasets, (e) Part Iib (issue 1): Weld creep rupture datasets, (f)Part III (issue 2): Post exposure (ex-service) creep data. Volume 6, 2005 (issue 1):Residual Life Assessment and Microstructure, Concari S. (ed.). Volume 9, 2005:High Temperature Component Analysis, Patel R. (ed.). (a) part II (issue 1) Overviewof assessment & design procedures, (b) part III (issue 1): Database of componenttests and assessments.

21 Holdsworth S.R., Orr J., Granacher J., Merckling G. and Bullogh C.K. on behalf ofthe ECCC-WG1, ‘European Creep Collaborative Committee activities on creep datageneration and assessment methodologies’ in Materials for Advanced PowerEngineering, D. Coutsouradis, F. Schubert, D.V. Thornton and J.H. Davidson (eds),Liege, 3–6 October 1994, Kluwer Academic Publishers, 1994, 591–600.

22 Merckling G. and Holdsworth S.R., ‘Long term creep rupture strength assessment:the development of the European Collaborative Creep Committee Post assessmenttests’ , International Conference on Creep & Fracture in High Temperature Components– Design and Life Assessment Issues, September 12–14, London, 2005.

23 ECCC Data Sheets, issue 1996. ERA Technology Publishers, 1996.24 ECCC Data Sheets, issue 2005. European Technology Development Publishers,

200525 European Pressure Equipment Research Council Web Site: www.mpa-lifetech.de/

eperctp/26 EC-funded Thematic Network and CEN Workshop 11 “UNCERT-AM”: http://

www.mpa-lifetech.de/UNCERT-AM/HTML_Files/Main/UNCERTDefault.htm27 Morris P. and Granacher J., ECCC-Information Day, European Technology

Development Ltd, Prague, 2001.28 Scholz A., Schwienheer M. and Morris P.F., ‘European notched testpiece for creep

rupture testing’, in: Proceedings of the 21st Symposium of the German Iron & SteelInstitute (VDEh) & German Society for Material Testing (DVM), Herausforderungdurch den Industriellen Fortschritt, 4–5 December 2003, Buchholz W.O. and GeislerS. (eds), Bad Neuenahr, Verlag Stahleisen, Düsseldorf, 2003, 308–314.

29 Webster G.A., Holdsworth S.R., Loveday M.S., Perrin I.J. and Purper H., A Code ofPractice for Conducting Notched Bar Creep Rupture Test and the Interpretation ofthe Data, ESIS TC11, 1999.

30 Webster G.A., Holdsworth S.R., Loveday M.S., Nikbin K., Perrin I.J., Purper H.,Skelton R.P. and Spindler M.W., ‘A code of practice for conducting notched barcreep tests and for interpreting the data, issue 3’, Fatigue and Fracture of EngineeringMaterials & Structures, 2004, 27 (4), 319–342.

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31 Matteoli L. and Andreini, B., ‘Le prove di scorimento interrotte. Influenza delleinterruzioni di sollecitazione e di riscaldamento sulle proprietà di scorrimento’, LaMetallurgia Italiana, 1947, XXXIX, 41.

32 Granacher J., Oehl M. and Preußler T., ‘Comparison of interrupted and uninteruptedcreep rupture test’, Steel Research, 1992, 63, 39–45.

33 Granacher J. and Scholz J., ‘Über die langzeitige Temperaturgenauigkeit vonZeitstandsprüfanlagen’, Materialprüf., 1973, 15, 116–123.

34 McCarthy P. and Loveday M.S. (eds), Proceeding of the Seminar on the Practicalitiesof Thermocouple Calibration and the Usage for Materials Testing, High TemperatureMechanical Testing Committee HTMTC, 16th December 1996.

35 Regis V., Livraghi M. and Di Pasquantonio F. Invecchiamento dei materiali perscorrimento viscoso, ENEL-CRTN, 1972.

36 Prager M., ‘Development of the MPC omega method for life assessment in the creeprange’, J. Pressure Vessel. Technology, 1995, 117, 95.

37 How I.M., Browne R.J., Coleman M.C., Craig I.H., Ham M.W., Hurst R.C. andMeecham P.C., ‘A code of practice for internal pressure testing of tubular componentsat elevated temperatures’, in Proceedings of the HTMTC Symposium on HarmonisedTesting Practice for High Temperature Materials, Loveday M.S. and Gibbons T.B.(eds), Ispra, Italy 18–19 October 1990, 363–400.

38 Rees D.W.A., Brown M.W., Hyde T., Lohn R.D., Morrison C.J. and Shammas M.,‘A code of practice for torsional creep testing of tubular testpieces at elevatedtemperatures’, in Proceedings of the HTMTC Symposium on Harmonised TestingPractice for High Temperature Materials, Loveday M.S. and Gibbons T.B. (eds),Ispra, Italy, 18–19 Octobers 1990, 331–361.

39 SEP 1260, Relaxationsversuch bei erhöhter Temperatur mitSchraubenverbindungsmodellen, VDEh, 1996.

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4.1 Introduction

Creep-resistant steels have been widely used in structural components andmachinery operated under high temperature and high pressure. There aremany different types of heat-resistant steels, as the properties required varyaccording to application and these properties can generally be categorizedinto physical, chemical, mechanical, manufacturability and economical aspects.Chemical composition influences physical and chemical properties such asthermal conductivity, thermal expansion and high temperature corrosion/oxidation resistance. Mechanical properties, including hardness, tensile, impact,creep and fatigue properties, depend on microstructures formed by heat-treatment and chemical composition. The general specifications for heat-resistant steels in Japan established JIS (Japanese Industrial Standards)1 andin METI (Ministry of Economy, Trade and Industry)2 Codes and JSME (JapanSociety of Mechanical Engineers)3 Codes for power applications. This sectionreviews the status of specifications for heat-resistant steels in Japanese codesand standards.

4.2 Types of heat-resistant steels in Japan

Table 4.1 lists heat-resistant steels (SUH) and stainless steels (SUS) in JISG4311 and G4312 for use under high temperature conditions, indicatingtheir applications as well. In the case of ferritic steels, cooling in air from theannealing temperature of 780–880°C is necessary, since slow cooling ataround 600°C causes material embrittlement owing to sigma phaseprecipitation. These steels can be used for heater boxes, burners and heaterequipment operated up to 1100°C in air and gas containing sulphur. Martensiticsteels are annealed in the temperature range of about 800–900°C, slow cooledat about 1000–1100°C, oil quenched and tempered at about 650–800°C, andthen subjected to rapid cooling or air cooling. These steels are used for hightemperature intake valves or turbine blades, applications that utilize their

4Specifications for creep-resistant steels:

Japan

F . M A S U YA M A, Kyushu Institute ofTechnology, Japan

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156Table 4.1 Heat-resistant steels in JIS

Designation Nominal composition (equivalent) Applications

Ferritic SUH 21 19Cr–3Al (18SR, JIS FCH2) Heater, automobile exhaust gas cleanerSUH 409 11Cr–Ti Automobile exhaust gas cleaner, mufflerSUH 446 25Cr–N–0.4C Combustion chamberSUS 405 13Cr–0.2Al–0.06C Combustion turbine compressor bladeSUS 430 18Cr-0.1C Furnace, bumer parts up to 900°C

Martensitic SUH 1 9Cr–3Si–0.4C Diesel engine intake valveSUH 3 11Cr–2Si–1Mo–0.4C Valve, pre-combustion charierSUH 4 20Cr–1.5Ni–2Si–0.8C Inhale/exhaust valve, abrasion resistant partsSUH 11 9Cr–1.5Si–0.5 Gas/diesel engine intake valve up to 750°CSUH 600 12Cr–0.6Mo–0.3V–0.5Nb–N–0.15C (H46) Steam turbine blade, disk, rotor, boltSUH 616 12Cr–1Mo–1W–0.25C AlSl 422) Steam turbine blade, disk, rotor, boltSUS 403 13Cr–low Si–0.1C Steam turbine blade, nozzleSUS 410 13Cr–0.5Mo–0.1C High temperature oxidation resistant parts up to 800°CSUS 410J1 13Cr–Mo–0.15C Steam turbine blade, high temp./press. componentsSUS 420J1 13Cr–0.2C Steam turbine blade, valve stem, nozzle, pump shaftSUS 420J2 13Cr–0.4C Piston ring, fuel injection nozzle, seat ringSUS 431 16Cr–2Ni–0.15C Shaft, bolt, springSUS 440A 18Cr–0.5Mo–0.7C Valve parts, ball valveSUS 440B 18Cr–0.5Mo–0.85C Bearing, roller, valve partsSUS 440C 18Cr–0.5Mo–1C Bearing, abrasion resistant parts

Precipitation SUS 630 17Cr–4Ni–4Cu–Nb–0.05 (17–4PH) Steam/combustion turbine bladehardening SUS 631 17Cr–7Ni–1Al–0.07C (17–7PH) High temp. spring, bellowsAustenitic SUH 31 15Cr–14Ni–2Si–2.5W–0.4C Gas/diesel engine exhaust valve up to 1150°C

SUH 35 21Cr–4Ni–9Mn–0.45N–0.5C (21–4N) Gas/diesel engine exhaust valve with high strengthSUH 36 21Cr–4Ni–9Mn–0.45N–0.5C–high S (21–4N) Gas/diesel engine exhaust valve with high strengthSUH 37 21Cr–11Ni–0.25N–0.2N (21–11N) Gas/diesel engine exhaust valve with oxidation

resistanceSUH 38 20Cr–11Ni–2.3Mo–0.2P–0.3C–B (20–11P) Gas/diesel engine exhaust valve, bolt

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157

Austenitic SUH 309 22Cr–12Ni–0.2C Furnace, burner cyclically operated up to 980°CSUH 310 25Cr–20Ni–0.2C Furnace, burner cyclically operated up to 1035°CSUH 330 15Cr–35Ni–0.1C Furnace, oil cracking equipmentSUH 660 15Cr–25Ni–1.5Mo–0.5V–2Ti–Al–B (A286) Turbine rotor, bolt, shaft up to 700°CSUH 661 22Cr–20Ni–20Co–3Mo–2.5W–Nb–0.15N Turbine rotor, bolt, blade, shaft up to 750°C

(LCN155)SUS 304 18Cr–8Ni Cyclic heating equipment with oxidation resistance up

to 870°CSUS 309S 22Cr–12Ni–0.06C Furnace parts cyclically operated up to 980°CSUS 310S 25Cr–20Ni–0.06C Furnace, automobile exhaust gas cleaner up to 1035°CSUS 316 18Cr–12Ni–2.5Mo–0.06C Heat exchanger with high creep strengthSUS 317 18Cr–12Ni–3.5Mo–0.06C heat exchanger with high creep strengthSUS 321 18Cr–8Ni–Ti–0.06C Welded parts for corrosion-resistant use at 400–900°CSUS 347 18Cr–8Ni–Nb–0.06C Parts for corrosion resistant use at 400–900°CSUSXM15J1 18Cr–13Ni–4Si (ASTM XM15) Automobile exhaust gas cleaner

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superior oxidation-resistant and/or creep-resistant properties. Austenitic heat-resistant steels listed in JIS G4311 and G4312 are for furnace equipment,heat exchangers and chemical plants taking advantage of their superior hightemperature corrosion/oxidation resistance and creep resistance. Except forSUH35 to 38 and SUH660/661, which need to be age-hardened, austeniticsteels are used as-solution annealed at about 1030–1200°C, followed byrapid cooling. For the aforementioned exceptional steels, aging treatmentsare applied at about 700–800°C after solution annealing.

4.3 Specifications for high temperature

tubing and piping steels

Table 4.2 shows JIS and METI Code Specifications for high temperaturetubing and piping steels. These tables include steel designation in the JIS andMETI Code, minimum tensile and yield strengths, heat treatment conditionsand application components in power boilers. However, several of thesesteels are not actually used in power boilers despite being listed in the codesand standards. The maximum usage temperature is determined by the designer,but the allowable stresses are listed in the METI Code up to the temperatureindicated in the code for individual steels. The criteria for allowable stressesin the creep temperature range in the METI Code are stated as follows.

The allowable tensile stress in the creep temperature range shall not exceedthe minimum values given below:

1 average value of the stress by which creep rate of 0.01%/1000 h occursat the relevant temperature;

2 67% of the average value of the stress by which rupture takes place in100 000 h at the relevant temperature;

3 80% of the minimum value of the stress by which rupture takes place in100 000 h at the relevant temperature.

On the other hand, the JSME Code recently changed the above criteria (2),substituting 67% with 100Favg% in accordance with the ASME criteria,where Favg is a multiplier applied to average stress for rupture in 100 000 h.At 815°C and below, Favg is 0.67. Above 815°C, it is determined from theslope of the log time-to-rupture versus log stress plot at 100 000 h, such thatlog Favg = 1/n, but it may not exceed 0.67. n is a negative number equal toδ log time-to-rupture divided by δ log stress at 100 000 h. METI Codematerials are defined as KA-XXXX, with XXXX, taking on a JIS-likedesignation that has not yet been integrated into JIS and indicating powerapplications only. However, new steels recently developed in Japan for powerboiler applications, such as high strength ferritic steels and austenitic steels,are included in the METI Code with the symbol ‘J’ followed by the developmentnumber in the designation. Tables 4.3 and 4.4 show the chemical compositions

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Part IC-steel tube for STB340 C-Steel 340 175 No specification

boiler & heat STB410 C-Steel 410 255

exchanger STB510 C-Steel 510 295

JIS G 3461(1988)

C-steel tube KA-STB480 C-Steel 480 275 Ann. or nor.

for powerMETI Code

Alloy steel STBA12 0.5Mo 380 205 Low temp. ann. iso-

tube for thermal ann., full ann.boiler & heat STBA13 0.5Mo 410 205 Nor. or nor. + temper

exchanger STBA20 0.5Cr, 0.5Mo 410 205 Low Temp. ann.

Iso-thermal ann.JIS G 3462 STBA22 1Cr, 0.5Mo 410 205 Full ann. or nor. +

(1988) temperSTBA23 1.25Cr, 0.5Mo 410 205 iso-thermal ann.

full ann. or nor. +temper

STBA24 2.25Cr, 1Mo 410 205

STBA25 5Cr, 0.5Mo 410 205STBA26 9Cr, 1Mo 410 205

Table 4.2 JIS and METI Code specification for high temperature tubing and piping materials

Designation Nominal composi- Min. Min. Heat treatment Applicationstion (mass%) tensile yield

(MPa) (MPa)

Eco

no

miz

er

Wat

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all

Hea

der

Co

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Alloy steel tube KA-STBA10 1.25Cr, 0.3Cu 410 255 Nor.for power boiler KA-STBA21 1Cr, 0.3Mo 410 205 Ann. or nor. + temper

METI Code KA-STBA24EG 2.25Cr, 1Mo 410 205 Ann. or nor. + temper

at min. 650°CKA-STBA24J1 2.25Cr, 1.6W 510 400 Nor. + temper

KA-STBA27 9Cr, 2Mo 510 295 Nor. at min. 900°C+

temper at min. 700°CKA-STBA28 9Cr, 1Mo, Nb, V 590 410 Nor. at min. 1040°C+

temper at min. 730°CKA-STBA29 9Cr, 1.8W 620 440 Nor. at min. 1040°C+

temper et min. 730°C

Part IIStainless steel SUS304TB 18Cr, 8Ni 520 205 1010°Ctube for SUS304HTB 18Cr, 8Ni 520 205 1040°C

boiler & heat SUS304LTB 18Cr, 8Ni, Low C 480 175 1010°Cexchanger SUS309TB 23Cr, 12Ni 520 205 1030°CJIS G 3463 SUS310TB 25Cr, 20Ni 520 205 1030°C(1994) SUS316TB 16Cr, 12Ni, 2Mo 520 205 1010°C

SUS316HTB 16Cr, 12Ni, 2Mo 520 205 1040°C

SUS316LTB 16Cr, 12Ni, 2Mo, Low C 480 175 1010°CSUS321TB 18Cr, 10Ni, Ti 520 205 920°C

Table 4.2 Cont’d


(MPa) (MPa)

Eco

no

miz

er

Wat

er w

all

Hea

der

Co

nn

ect

pip

e

Su

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hea

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Ste

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SUS321HTB 18Cr, 10Ni, Ti 520 205 1095°C for cold

finished, 1050°Cfor hot finished

SUS347TB 18Cr, 10Ni, Nb 520 205 980°CSUS347HTB 18Cr, 10Ni, Nb 520 205 1095°C for cold


SUS410TB 13Cr 410 205 700°C AC or FCSUS430TB 16Cr 410 245 700°C AC or FC

Part IIIStainless steel KA-SUS304J1HTB 18Cr, 9Ni, 3Cu, Nb, N 590 235 1040°C

tube for KA-SUS309J1TB 24Cr, 15Ni, 1Mo, N 690 345 1050°C

power boiler KA-SUS309J2TB 22Cr, 14Ni, 1.5Mo, N 590 245 1050°C

METI Code KA-SUS309J3LTB 25Cr, 14Ni, 0.8Mo, 690 345 1050°C

N, 0.2SiKA-SUS309J4HTB 22Cr, 15Ni, Nb 590 235 1120°C

KA-SUS310J1TB 25Cr, 20Ni, Nb, V 660 295 1030°C

KA-SUS310J2TB 20Cr, 25Ni, 1.5Mo 640 270 1100°C

KA-SUS310J3TB 22.5Cr, 18.5Ni, 1.8W, 3Cu, 650 295 1030°C

0.45Nb, 0.2NKA-SUS321J1HTB 18Cr, 10Ni, Ti, Nb 520 205 1100°C

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KA-SUS321J2HTB 18Cr, 10Ni, 3Cu, Nb, 500 205 1160°C

Ti, B, NKA-SUSTP347HTB 18Cr, 10Ni, Nb 520 205 1150°C

KA-SUS347J1TB 18Cr, 9Ni, V, W, Nb, N 650 270 1100°C

KA-SUS410J2TB 12Cr, 1Mo, W, V, Nb 590 390 Nor. + temper

KA-SUS410J3TB 11Cr, 2W, 0.4Mo, V, Nb 620 400 Nor. + temper

KA-SUS410J3DTB 12Cr, 2W, 0.4Mo, V, Nb 620 400 Nor. + temper

C-steel for high STPT370 C-steel 370 215 As rolled for hot

temp. pipe JIS STPT410 C-steel 410 245 finished low

G3456 (1988) STPT480 C-steel 480 275 temp. ann. or nor.

for cold finished

Part IVAlloy steel STPA12 0.5Mo 380 205 Low temp. ann. iso-

piping thermal ann. full ann.nor. or nor. + temper

JIS G 3458 STPA20 0.5Cr, 0.5Mo 410 205 low temp. ann. iso-

(1988) thermal ann. full ann.or nor. + temper

STPA22 1Cr, 0.5Mo 410 205 Iso-thermal ann. full

STPA23 1.25Cr, 0.5Mo 410 205 ann. or nor. + temper

STPA24 2.25Cr, 1Mo 410 205

Table 4.2 Continued


(MPa) (MPa)

Eco

no

miz

er

Wat

er w

all

Hea

der

Co

nn

ect

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Su

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STPA25 5Cr, 0.5Mo 410 205STPA26 9Cr, 1Mo 410 205

Alloy steel KA-STPA21 1Cr, 0.3Mo 410 205 Ann. or Nor. + Temper

power piping KA-STPA24J1 2.25Cr, 1.6W 510 400 Nor. + Temper

METI code KA-STPA27 9Cr, 2Mo 510 295 Nor. at min. 900°C +

temper at min. 700°CKA-STPA28 9Cr, 1Mo, Nb, V 590 410 Nor. at min. 1040°C +

temper at min. 730°CKA-STPA29 9Cr, 1.8W 620 410 Nor. at min. 1040°C +

temper at min. 730°C

Part VStainless steel SUS304TP 18Cr, 8Ni 520 205 1010°Cpiping JIS G SUS304HTP 18Cr, 8Ni 520 205 1040°C3459(1997) SUS304LTP 18Cr, 8Ni, Low C 480 175 1010°C

SUS309TP 23Cr, 12Ni 520 205 1030°CSUS310TP 25Cr, 20Ni 520 205 1030°CSUS316TP 16Cr, 12Ni, 2Mo 520 205 1010°CSUS316HTP 16Cr, 12Ni, 2Mo 520 205 1040°CSUS316LTP 16Cr, 12Ni, 2Mo, Low C 480 175 1010°CSUS321TP 18Cr, 10Ni, Ti 520 205 920°CSUS321HTP 18Cr, 10Ni, Nb 520 205 1095°C for cold

finished, 1050°C forhot finished

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SUS347TP 18Cr, 10Ni, Nb 520 205 980°CSUS347HTP 18Cr, 10Ni, Nb 520 205 1095°C for cold


Stainless steel KA-SUS410J3TP 11Cr, 2W, 0.4Mo, V, Nb 620 400 Nor. + temper

power piping

Ann = annealing, Nor = normalizing, Temper = tempering, AC = air cooling, FC = furnace cooling.

Table 4.2 Continued


(MPa) (MPa)

Eco

no

miz

er

Wat

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all

Hea

der

Co

nn

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Table 4.3 Chemical compositions for ferritic boiler tubing steels


C Si Mn P S Ni Cr Mo V Nb Al N W B

C–Steel STB410 ≤ 0.32 ≤ 0.35 0.30 ≤ 0.035 ≤ 0.035 – – – – – – – – –~0.80

KA-STB510 ≤ 0.25 ≤ 0.35 1.00 ≤ 0.035 ≤ 0.035 – – – – – – – – –~1.50

KA-STB480 ≤ 0.30 ≤ 0.10 0.29 ≤ 0.048 ≤ 0.058 – – – – – – – – –~1.06

Low KA-STBA10 ≤ 0.10 0.20 ≤ 0.80 ≤ 0.025 0.015 – 1.00 – – – – – – –alloy ~0.80 ~0.030 ~1.50steel STBA12 0.10 0.10 0.30 ≤ 0.035 ≤ 0.035 – – 0.45 – – – – – –

~0.20 ~0.50 ~0.80 ~0.65KA-STBA21 0.10 ≤ 0.50 0.30 ≤ 0.035 ≤ 0.035 – 1.90 0.87 ~ – – – – – –

~0.20 ~0.60 ~2.60 1.13STBA23 ≤ 0.15 0.50 0.30 ≤ 0.030 ≤ 0.030 – 1.00 0.45 – – – – – –

~1.00 ~0.60 ~1.50 ~0.65STBA24 ≤ 0.15 ≤ 0.50 0.30 ≤ 0.030 ≤ 0.030 – 1.90 0.87 – – – – – –

~0.60 ~2.60 ~1.18KA-STBA24E-G ≤ 0.15 ≤ 0.50 0.30 ≤ 0.030 ≤ 0.030 – 1.90 0.87 – – – – – –

~0.60 ~2.60 ~1.18KA-STBA24J1 0.04 ≤ 0.50 0.10 ≤ 0.030 ≤ 0.010 – 1.90 0.05 0.20 0.02 ≤0.03 ≤0.03 1.45 0.0006

~0.10 ~0.60 ~2.60 ~0.30 ~0.30 ~0.08 ~ 1.75

9Cr STBA26 ≤ 0.15 0.25 0.30 ≤ 0.030 ≤ 0.030 – 8.00 0.90 – – – – – –~ 1.00 ~0.60 ~10.00 ~1.10

KA-STBA27 ≤ 0.08 ≤ 0.50 0.30 ≤ 0.030 ≤ 0.030 – 8.00 1.80 ~ – – – – – –~0.70 ~10.00 2.20

KA-STBA28 0.08 0.20~ 0.30 ≤ 0.020 ≤ 0.010 ≤ 0.40 8.00 0.85 ~ 0.18 0.06 ≤0.04 0.030 – –~0.12 0.50 ~0.60 ~9.50 1.05 ~0.25 ~0.10 ~0.070

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Table 4.3 Continued


C Si Mn P S Ni Cr Mo V Nb Al N W B

KA-STBA29 0.07 ≤ 0.50 0.30 ≤ 0.020 ≤ 0.010 ≤ 0.40 8.00 0.03 0.15 0.04 ≤0.04 0.030 1.50 0.001~0.13 ~0.60 ~9.50 ~ 0.60 ~0.25 ~0.09 ~0.070 ~2.00 ~0.005

12Cr KA-SUS410J2TB ≤ 0.14 ≤ 0.50 0.30 ≤ 0.030 ≤ 0.030 – 11.00 0.80 0.20 ≤0.20 ≤0.04 – 0.80 –~0.70 ~13.00 ~1.20 ~0.30 ~1.20

KA-SUS410J3TB 0.07 ≤ 0.50 ≤ 0.70 ≤ 0.020 ≤ 0.010 ≤ 0.50 10.00 0.25 0.15 0.04 ≤0.04 0.040 1.50 0.0005~0.14 ~11.50 ~0.60 ~0.30 ~0.10 ~0.100 ~2.50 ~0.005

KA- 0.07 ≤ 0.50 ≤ 0.70 ≤ 0.020 ≤ 0.010 ≤ 0.50 11.51 ~ 0.25 0.15 0.04 ≤0.04 0.040 ~1.50 0.0005SUS410J3DTB ~0.14 12.50 ~0.60 ~0.30 ~0.10 0.100 ~2.50 ~0.005

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Table 4.4 Chemical compositions for austenitic boiler tubing steels


C Si Mn P S Ni Cr Mo V Nb Al N Cu W B Others

SUS304HTB 0.04 ≤0.75 ≤ 2.00 ≤0.040 ≤ 0.030 8.00 18.00 – – – – – – – – –~0.10 ~11.00 ~20.00

SUS321HTB 0.04 ≤0.75 ≤ 2.00 ≤0.030 ≤ 0.030 9.00 17.00 – – – – – – – – Ta: 4 × C~ 0.60

~0.10 ~13.00 ~20.00SUS316HTB 0.04 ≤0.75 ≤ 2.00 ≤0.030 ≤ 0.030 11.00 16.00 2.00 – – – – – – – –

~0.10 ~14.00 ~18.00 ~3.00SUS347HTB 0.04 ≤1.00 ≤ 2.00 ≤0.030 ≤ 0.030 9.00 17.00 – – – – – – – – Nb+Ta:

~0.10 ~13.00 ~20.00 8 × C ~1.00KA–SUS304J1HTB 0.07 ≤0.30 ≤ 1.00 ≤0.040 ≤ 0.010 7.50 17.00 – – – 0.30 0.05 2.50 – – –

~0.13 ~10.50 ~19.00 ~0.60 ~0.12 ~3.50KA–SUS309J1TB ≤ 0.06 ≤1.50 ≤ 2.00 ≤0.040 ≤ 0.030 12.00 23.00 0.50 – – – 0.25 – – – –

~16.00 ~26.00 ~1.20 ~0.40KA–SUS309J2TB ≤ 0.04 ≤1.00 2.50 ≤0.030 ≤ 0.030 12.50 21.00 1.00 – – – 0.10 – – – –

~3.50 ~15.50 ~23.00 ~2.00 ~0.25KA–SUS309J3LTB ≤ 0.025 ≤0.70 ≤ 2.00 ≤0.040 ≤ 0.030 13.00 23.00 0.50 – – – 0.25 – – – –

~16.00 ~26.00 ~1.20 ~0.40KA–SUS309J4HTB 0.03 ≤1.00 ≤ 2.00 ≤0.040 ≤ 0.030 14.50 21.00 – – – 0.50 0.10 – – ≤ 0.006 –

~0.10 ~16.50 ~23.00 ~0.80 ~0.20KA–SUS310J1TB ≤ 0.10 ≤1.50 ≤ 2.00 ≤0.030 ≤ 0.030 17.00 23.00 – – – 0.20 0.15 – – – –

~23.00 ~27.00 ~0.60 ~0.35KA–SUS310J2TB ≤ 0.10 ≤1.00 ≤ 1.50 ≤0.030 ≤ 0.010 22.00 19.00 1.00 ≤ 0.20 – 0.10 0.10 – – 0.002 –

~28.00 ~23.00 ~2.00 ~0.40 ~0.25 ~0.010KA–SUS310J3TB 0.05 ≤1.50 ≤ 2.00 ≤0.030 ≤ 0.010 15.00 21.00 – – – 0.30 0.15 2.00 0.80 – –

~0.12 ~22.00 ~24.00 ~0.60 ~0.30 ~4.00 ~2.80KA–SUS321J1HTB 0.07 ≤1.00 ≤ 2.00 ≤0.040 ≤ 0.030 9.00 17.50 – ≤ 0.20 – ≤ 0.40 – – – – (Ti+Nb/2)C

~0.14 ~12.00 ~19.50 2.0~4.0

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KA–SUS321J2HTB 0.07 ≤ 1.00 ≤ 2.00 ≤0.040 ≤ 0.010 9.00 17.50 – 0.10 – 0.10 – 2.50 – 0.0010 (Ti+Nb/2)C~0.14 ~12.00 ~19.50 ~0.25 ~0.45 ~3.50 ~0.0040 2.0~4.0

KA–SUS3TP347HTB 0.04 ≤ 0.75 ≤ 2.00 ≤0.030 ≤ 0.030 9.00 17.00 – – – 8 X C% – – – – –~0.10 ~13.00 ~20.00 ~1.00

KA–SUS347J1TB ≤ 0.05 ≤ 1.00 ≤ 2.00 ≤0.040 ≤ 0.030 8.00 17.00 – – 0.20 0.25 0.10 – 1.50 – –~11.00 ~20.00 ~0.50 ~0.50 ~0.25 – ~2.60

Table 4.4 Continued


C Si Mn P S Ni Cr Mo V Nb Al N Cu W B Others

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Specifications for creep-resistant steels: Japan 169

of JIS and METI Code steel tubes used extensively in high temperaturecomponents such as economizers, water walls, superheaters and reheaters inJapanese power boilers. The heat treatment conditions and tensile requirementsfor these steels are indicated in Tables 4.2–4.4. The specific designations forproduct forms other than boiler tubing and piping are shown in comparisonwith ASME Code in Table 4.5 for recently developed high strength ferriticsteels.

4.4 Specifications for steam turbine steels

Table 4.64 shows the chemical compositions for steam turbine steels used inJapanese power plants. Steels are classified for specific components such asblades and discs, rotors and casings. The rotor steels shown here are recentlydeveloped 12%Cr high strength materials, although CrMo/CrMoV rotor steeland casings are still generally used for conventional steam turbines.Specifications for these turbine steels are controlled by turbine manufacturesindividually and not by the regulatory specifications. The technical controlof melting and refining is also very important for turbine steels, as well asheat treatment control in order to maintain appropriate mechanical (strength/toughness) and creep/fatigue properties. Thus, the listed nominal chemicalcompositions and heat treatment conditions in the table are suggested asguidelines.

4.5 Heat-resistant super alloys

Table 4.7 shows the chemical compositions of heat-resistant alloys (knownas super alloys) listed in JIS G4901 for bars and G4902 for plates. Thesealloys are designated similarly to NCF XXX, with the lower three columnsof the unified numbering system (UNS) number or ASTM designation. Thesealloys are annealed or solution annealed, except several grades that are agedafter solution annealing. JIS does not specify the recommended annealing orsolution temperature and aging temperature, although tensile/yield strength,rupture elongation and hardness are specified. These alloys exhibit extremelyhigh creep strength and high temperature corrosion/oxidation resistancecompared with other steels.

4.6 Summary

Specifications for creep-resistant steels in Japan are well established in JISfor ordinary use including high temperature applications, with METI Codesand JSME Codes only for boiler and nuclear applications. These specificationsshow general recommendations and regulations for chemical compositions,heat treatment conditions, mechanical properties, and so on. Mechanical

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Table 4.5 Comparative code designations in METI Code and ASME for new ferritic boiler steels

Steels Code Tube Pipe Plate Forgings(trade name)

2.25Cr HCM2S ASME SA-213 T23 SA-335 P23 SA-1017 Gr.23 SA-182 F23METI KA-STBA24J1 KA-STPA24J1 KA-SCMV4J1 KA-SFVAF22AJ1

HCM9M ASME – – – –METI KA-STBA27 KA-STPA27 – KA-SFVAF27

9Cr Mod.9Cr ASME SA-213T91 SA-335P91 SA-387Gr.91 SA-182F91METI KA-STBA28 KA-STPA28 KA-SCMV28 KA-SFVAF28

NF616 ASME SA-213T92 SA-336P92 – SA-182F98METI KA-STBA29 KA-STPA29 – KA-SFVAF29

HCM12 ASME – – – –METI KA-SUS410J2TB – – –

12Cr HCM12A ASME SA-213T122 SA-335P122 SA-1017Gr.122 SA-182F122METI KA-SUS410J3TB KA-SUS410J3DTP KA-SUS410J3 KA-SUS410J3

KA-SUSF410J3DTB – – –

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171Table 4.6 Chemical compositions and heat treatment for turbine steels

Steels Nominal chemical composition (mass%) Heat—————————————————————————————————————— treatment

Nor.-temperC Si Mn Ni Cr Mo W Co V Nb B N (°C)

Small H46 (12Cr–0.5MoVNbN) 0.15 0.40 0.60 – 12.0 0.5 – – 0.30 0.25 – 0.050 1,150 –650components TAF (10.5Cr–1.5MoVNbB) 0.18 0.30 0.50 – 10.5 1.5 – – 0.20 0.15 0.030 0.015 1,150 –700

TAF650 0.11 0.01 0.50 – 11.0 0.2 2.6 3.0 0.20 0.08 0.015 0.020 1,100 –750(11Cr –2.6W–3CoVNbB)TOS203 0.11 0.05 0.50 0.6 10.5 0.1 2.5 1.0 0.20 0.10 0.010 0.030 1,120 –680(10.5Cr –2.5W–1CoVNbBRe)

Rotor GE (10.5Cr–1MoVNbN) 0.18 0.30 0.60 0.6 10.5 1.0 – – 0.20 0.06 – 0.060 1,050 –620TMK1 (10.3Cr–1.5MoVNbN) 0.14 0.05 0.50 0.6 10.3 1.5 – – 0.17 0.06 – 0.040 1,100 –680HR1100 0.15 0.03 0.60 0.6 10.3 1.2 0.3 – 0.20 0.05 – 0.050 1,075 –660(10.3Cr–1.2Mo–0.3WNbN)TOS107 0.14 0.05 0.60 0.7 10.0 1.0 1.0 – 0.20 0.07 – 0.050 1,050 –660(10Cr–1Mo–1MVNbN)TMK2 0.14 0.05 0.50 0.5 10.2 0.5 1.8 – 0.17 0.06 – 0.040 1,050 –700(10.2Cr–0.5Mo 1.8WVNbN)TR1200 0.13 0.05 0.50 0.8 11.0 0.2 2.5 – 0.20 0.08 – 0.050 1,050 –710(11Cr–0.2W–2.5WVNbB)HT1200 0.11 0.05 0.50 0.5 11.0 0.2 2.6 3.0 0.20 0.08 0.015 0.025 1,050 –720(11Cr–2.6W–3CoNiVNbB)TOS110 (10Cr–0.7Mo– 0.11 0.08 0.10 0.2 10.0 0.7 1.8 3.0 0.20 0.05 0.010 0.020 1,070 –6801.8W –3CoVNbB)

Casing MJC12 (9.5Cr –1MoVNbN) 0.10 0.70 0.70 0.5 9.5 1.0 – – 0.15 0.06 – 0.040 1,050 –710TOS302 0.12 0.25 0.50 1.0 10.0 1.0 0.8 – 0.20 0.10 – 0.050 1,050 –710(10Cr –1Mo–0.8WVNbN)TOS303 0.12 0.15 0.50 0.2 10.0 0.7 1.8 3.0 0.20 0.05 0.006 0.020 1,100 –730(10Cr–1.8W–3CoVNbB)

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172Table 4.7 Chemical compositions of heat-resistant alloys (JIS G4901 and G 4902)

Designation Chemical composition (mass%)

C Si Mn P S Ni Cr Fe Mo Cu Al Ti Nb + BTa

NCF600 ≤ 0.15 ≤ 0.50 ≤ 1.00 ≤ 0.030 ≤ 0.015 72.00~ 14.00 6.00 – ≤ 0.50 – – – –~17.00 ~10.00

NCF601 ≤ 0.10 ≤ 0.50 ≤ 1.00 ≤ 0.030 ≤ 0.015 58.00 21.00 Bal. – ≤ 1.00 1.00 – – –~63.00 ~25.00 ~1.70

NCF625 ≤ 0.10 ≤ 0.50 ≤ 0.50 ≤ 0.015 ≤ 0.015 58.00~ 20.00 ≤ 5.00 8.00 – ≤ 0.40 ≤ 0.40 3.15 –~23.00 –10.00 ~4.15

NCF690 ≤ 0.05 ≤ 0.50 ≤ 0.50 ≤ 0.030 ≤ 0.015 58.00~ 27.00 7.00 – ≤ 0.50 – – – –~31.00 ~11.00

NCF718 ≤ 0.08 ≤ 0.35 ≤ 0.35 ≤ 0.015 ≤ 0.015 50.00 17.00 Bal. 2.80 ≤ 0.30 0.20 0.65 4.75 ≤0.006~55.00 ~21.00 ~3.30 ~0.80 ~1.15 ~5.50

NCF750 ≤ 0.08 ≤ 0.50 ≤ 1.00 ≤ 0.030 ≤ 0.015 70.00~ 14.00 5.00 – ≤ 0.50 0.40 2.25 0.70 –~17.00 ~9.00 ~1.00 ~2.75 ~1.20

NCF751 ≤ 0.10 ≤ 0.50 ≤ 1.00 ≤ 0.030 ≤ 0.015 70.00~ 14.00 5.00 – ≤ 0.50 0.90 2.00 0.70 –~17.00 ~9.00 ~1.50 ~2.60 ~1.20

NCF800 ≤ 0.10 ≤ 1.00 ≤ 1.50 ≤ 0.030 ≤ 0.015 30.00 19.00 Bal. – ≤ 0.75 0.15 0.15 – –~35.00 ~23.00 ~0.60 ~0.60

NCF800H 0.05 ≤ 1.00 ≤ 1.50 ≤ 0.030 ≤ 0.015 30.00 19.00 Bal. – ≤ 0.75 0.15 0.15 – –~0.10 ~35.00 ~23.00 ~0.60 ~0.60

NCF825 ≤ 0.05 ≤ 0.50 ≤ 1.00 ≤ 0.030 ≤ 0.015 38.00 19.50 Bal. 2.50 1.50 ≤ 0.20 0.60 – –~46.00 ~23.50 ~3.50 ~3.00 ~1.20

NCF80A 0.04 ≤ 1.00 ≤ 1.00 ≤ 0.030 ≤ 0.015 Bal. 19.50 ≤ 1.50 – ≤ 0.20 1.00 1.80 – –~0.10 ~23.50 ~1.80 ~2.70

Bal. = balance.

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Specifications for creep-resistant steels: Japan 173

properties, including hardness, tensile, impact, creep and fatigue properties,are strongly influenced by microstructures resulting from heat treatment andchemical composition. In the case of turbine steels, the steel manufacturingprocess is very important to achieve the desired properties and performance.In the case of boiler/piping and pressure vessel applications, allowable stressesare provided in the METI Codes and JSME Codes, which feature criteriasimilar to ASME for determination of the allowable stress values fortemperatures. In the future, specifications for creep-resistant steels will beupdated in accordance with codes and standards developed in other areasworldwide.

4.7 References

1 http://www.jsa.or.jp/default_english.asp2 http://www.meti.go.jp/english/index.html3 http://www.jsme.or.jp/English/4 Masuyama F. ‘History of power plants and progress in heat resistant steels’, ISIJ

International, 2001 41, 612–625.

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174

5.1 Introduction

Increases in steam temperature and pressure have made a large contributionto improvement in the efficiency of fossil power plants. The importance ofthe increase in efficiency is a stringent problem for CO2 reduction andconservation of resources. Realisation of advanced fossil power plants thatincorporate advanced steam conditions depends on the development ofadvanced heat-resistant steels and of production technology for large turbinerotor forgings. So far, many advanced type high pressure (HP) turbines,intermediate pressure (IP) turbines and high pressure low pressure combinationturbines (HLP) have been developed successfully and contribute to an increasein plant efficiency. These advanced turbines are realised through the applicationof state-of-the-art production technology developed together with cultivatedwidely used and conventional technologies over the years. Points for theproduction of high-performance and reliable turbine rotor forging can besummarized as follows.

Production of high purity steels with minimized residual elements andfreedom from non-metallic inclusions is important in the steelmaking process.Homogeneous ingots with minimal segregation, delta ferrite, non-metallicinclusions and porosities should be made with a homogeneous distributionof chemistry throughout the casting process. For each forging process, asufficient forging effect at the centre of large diameter in the ingots andforging blocks needs to be attained to consolidate the porosities in the ingotsand sufficient forging strain should be given to the ingots to eliminate thesolidification structure (e.g. dendrite) and promote formation of equiaxedgrain through dynamic recrystallization. In the heat treatment process, heattreatment effects need to be exerted to develop the required properties at thecentre of the forging. A fine grain microstructure needs to be obtained toassure sufficient detectability of defects.

In this chapter, the development of production technologies for turbinerotor forgings is reviewed and several key processes are introduced. In

5Production of creep-resistant steels

for turbines

Y. T A N A K A, Japan Steel Works, Japan

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Production of creep-resistant steels for turbines 175

addition, the properties of forgings made from heat-resistant steels areintroduced.

5.2 Overview of production technology of rotor

shaft forgings for high temperature steam

turbines

Figure 5.1 shows a schematic of the production process for large turbinerotor forgings. The detail of each process is essentially different dependingon the designation of steels and the required properties of the rotor shaft. Inthis section, the facilities for the production and typical technologies areoverviewed.

5.2.1 Steelmaking and casting process

Figure 5.2 shows the historical change in steelmaking methods and equipmentin one company since 1950, as an example.1 Similar changes in facilities andproduction technologies have taken place in the majority of the forgemastersin the world.

Turbine material used to be refined by open hearth furnaces and was castin air. At that time, the absorption of hydrogen in steel was one of the mostserious problems of the process since hydrogen causes defects such as flaking.Installation of vacuum degassing equipment like the Bochumer–Verein typewas a solution for the degassing of hydrogen during casting. By use ofvacuum degassing equipment, the basic open hearth furnace and the basicelectric arc furnace (EAF), which tend to absorb hydrogen during refiningbut are superior in refining ability, were able to be used, leading to animprovement in the material properties. The efficiency of the vacuum degassingequipment has been improved to attain a higher vacuum. With the introductionof vacuum casting equipment, vacuum carbon deoxidization (VCD) technologyhas been successfully applied to steam turbine materials.2

With the increase in capacity of power plants, larger turbine rotor forgingswere needed and new production technologies were developed. A pouringmethod using multiple furnaces was developed as the casting technology forsuch large forging. Ladle refining furnaces (LRF) were installed to keepsteel molten after refining in an electric furnace. After that, several ladlerefining furnaces were additionally installed and these make the productionof ingots up to 600 tonnes possible using a fully ladle refined melt.3 On theother hand, an electroslag remelting (ESR) furnace was used for refining andcasting of high quality steels. The capacity was enlarged to meet the demandsof melting large high temperature turbines made from low alloy and highalloy steels. Vacuum induction melting (VIM) was also installed. The use ofVIM for the production of rotor forging, however, is uncommon. Through

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176

5.1 Typical production process for large steam turbine forging.

1 2 3 4

8765

Electric arcfurnace Vacuum casting Press

Horizontal furnace

Steelmaking casting Forging Preliminary heat treatment

VacuumVacuumLadle

refiningfurnace

Vertical furnace

Quenching Tempering

FCWQOilFC

Quality heat treatment NDE, machining Stress relieving

Stressrelieving

Vertical furnace

Rough machining

Shipping

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Production of creep-resistant steels for turbines

177

Acid and basic openhearth furnace

Year1950 55 60 65 70 75 80 85 90 95

Meltingand

refining

Ingotmaking

Maximumingot

weight

5 ton VIM

100 ton ESR20 ton ESR

150 tonx2

30 tonx2

150 ton130 ton

Holding furnaceLadle refining furnace

Electric furnaceBasic open hearth furnace

Electric furnace

Air castingMechanical pump Steam ejector (high vacuum)

Bochumer–Verein type mould stream degassing

Multi-pouring Low Si-VCD

Prediction ofsegregation

Doubledegassing

600 ton570 ton500 ton400ton

250ton220 ton140 ton

5.2 The history of the steelmaking and casting process: an example.

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the development of these facilities and production technologies, highperformance and high reliability rotor forgings have been manufactured.

Current steelmaking processes used in the production of high temperaturesteam turbine rotor forging are as follows.

Basic electric arc furnace and ladle refining furnace

Formerly, acid and/or basic open hearth furnaces were used for refiningsteels. Since the introduction of vacuum degassing facilities, basic electricarc furnaces have become the major equipment used in the melting andrefining of turbine rotor steels. In this process, refining is performed by thedouble slag method. After melting the raw materials, oxidizing refining iseffected by adding a basic oxidizing slag to reduce the C, Mn, Si and Pcontent. Then the oxidizing slag is removed to avoid the oxidized elementsreturning to the molten steel. When removing oxidizing slag, it is especiallyimportant to remove the P, which is harmful to the material properties. Thenreducing slag is added to decrease the S content. After adjustment of thechemistry, the molten steel is poured into a mould in the vacuum degassingchamber. In order to reduce the residual elements as low as possible to makehigh purity steels, the double slag process was further improved by usingEAF and subsequently LRF. Figure 5.3 shows an example of a typical currentsteelmaking process using EAF and LRF.3 The raw materials are melted inan EAF where oxidizing refining is performed. Then the molten steel ispoured into a ladle and the oxidizing slag in the EAF is completely removed.After reladling, the reduction refining and degassing processes take place in

Ar gas Ar gas

Vacuum

Vacuum

Ladlefurnace

Electricfurnace

Ingot makingLadle refiningReladleMelting/refining

De-sulphurisationDe-phosphorisation

5.3 Typical steelmaking process with EAF and LRF for high puritysteel.

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a vacuum after agitation with Ar gas. Subsequently, the molten steel is castin vacuum through mould stream degassing.

Figure 5.4(a), (b) and (c) shows the appearance of tapping from the EAFto the ladle, ladle refining and casting from ladle to ingot mould in thevacuum chamber, respectively. The use of multiple ladles make it possible tocast a large ingot.

Figure 5.5 shows the history of P and S content in materials for LP (lowpressure steam turbine) rotor forging steel. The significant effect of advances

(a) (b) (c)

5.4 The steelmaking and casting process. (a) Tapping from an electricarc furnace, (b) ladle refining, (c) vacuum pouring.

: S%: P%

Year1960 1970 1980 1990 2000

P a

nd

S c

on

ten

ts (

wt%

)

0.024

0.022

0.020

0.018

0.016

0.014

0.012

0.010

0.008

0.006

0.004

0.002

0

5.5 Historical change of P and S content in steels for LP rotor forging.

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in refining technology is demonstrated by the reduction of P and S. Table 5.1summarizes the methods of reducing the impurity elements in currentsteelmaking processes and currently superclean steels has been developedby reducing not only residual elements such as P, S, As, Sn, Sb but also Siand Mn which are usually added to effect deoxidization.3–5

Electroslag remelting (ESR)

The ESR equipment consists of a large capacity power supply and a watercooled crucible. Figure 5.6(a) shows a schematic of the ESR process andFig. 5.7 shows an ESR facility. In the ESR process, an electrode is preparedby casting or forging after the conventional melting, refining and castingprocess. The melting of the electrodes occurs in the mould by heating causedby the electric resistance of the slag. As the droplets of electrode material

5.6 Equipment for (a) ESR process and (b) VAR process.

Table 5.1 Elimination of tramp elements for clean steel

Element Technology and process to be applied

Mn Oxidizing refining in EAF → reladleSi Oxidizing refining in EAF and VCDP Oxidizing refining in EAF → reladleSi Ladle refiningSn Selection of raw materialAs Selection of raw materialSb Selection of raw materialAl VCDH Vacuum treatmentO Ladle refining and VCDN Ladle refining

Vacuum

Consumable electrode

Cooling water

CrucibleSlag pool

Molten steel pool

IngotCooling water

Consumable electrode

Furnace

Cooling water

Molten steel poolCrucibleIngot

Cooling water

(b)(a)

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fall, refining of the material proceeds. The material solidifies at the bottomof the molten material pool. By application of the ESR process, the purity,cleanliness of the material and homogeneity of the ingot can be improved.Formation of macro segregation is suppressed and the distribution of chemicalelements becomes more uniform compared to ingots made by conventionalcasting. Slag composition is especially important in the ESR process in orderto attain the expected properties. Presently, the ESR process is frequentlyused for the production of rotor forgings made from heat-resistant steelssuch as CrMoV steels, advanced CrMoV steels and various 12Cr steels.

Hot topping process

Several applications of ESR technology combined with a conventional cas-ting method such as the Bohler electroslag topping process (BEST)6 and

5.7 An ESR facility.

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electroslag hot topping (ESHT-J)7 have been developed. In both methods,refining is performed by a conventional process in an EAF and the productis poured into an ingot mould. The slag is fed on to the surface of the moltensteel in the mould and heating by the consumable electrode (together with agraphite electrode) takes place. By using these processes, the formation ofporosity, shrinkages and segregation can be suppressed.

Vacuum arc remelting (VAR)

Figure 5.6(b) shows a schematic of VAR equipment. The electrode is arcmelted in a water-cooled crucible inside a vacuum chamber. In contrast to theESR process, molten steel refining does not proceed unless there is degassingthrough this process. Therefore, the electrode for the VAR process should beproduced through a conventional refining and casting process or by vacuuminduction heating process so as to attain the target chemistry and purity. Onthe other hand, cooling of the molten steel takes place at a faster rate thanduring the ESR process for a given mould dimension, resulting in a superiorsolidification structure with less segregation. The capacity of VAR iscomparatively small owing to installation of a vacuum chamber; experienceof application to the production of rotor steel forging is therefore limited.

Vacuum induction melting (VIM)

VIM uses a melting furnace for raw materials by induction heating in avacuum. Since refining does not result from the VIM process, the raw materialshould be high purity ferroalloys and high purity metals, depending on therequirement of the products. Although application to large steel forging productsis quite limited, the process coupled with VAR/ESR processes is indispensableto the production of super alloys coupled with VAR and/or ESR process.

5.2.2 Ingot making process

After the conventional melting and refining process, the molten steel is castinto an ingot mould from top or bottom of a mould made from cast iron. Inthe case of ingots for large steam turbine forging, top pouring into the ingotplaced in vacuum tank is performed, as shown in Fig. 5.3. During pouring,a further reduction in the gas elements proceeds in the vacuum tank. Forsmaller rotor forgings, casting of ingots weighing several tens of tons ofsteel is done by bottom pouring. In this process, the refined molten steel ispoured into the ingot mould from bottom of the mould set on a steelplate through refractory tube until the mould fills up. In the case of ESR andVAR, the re-melted steel is again solidified into ingots in a water-cooledcrucible.

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In order to reduce segregation in the ingot and to improve the performanceof the material, a low Si-VCD process is applied.2 In the VCD process,molten steel with a low Si content is cast in a high vacuum; C and O in thesteel react and are exhausted as CO. During application of the VCD process,it is possible to achieve much lower levels of Si than in the Si deoxidationprocess. Decrease in Si content not only develops a fine solidification structurewhich suppresses formation of macro-segregation in large ingots but alsodecreases the susceptibility to temper embrittlement which is critical problemfor materials used in the temperature range between 350 and 550 °C.8,9 Adecrease in Si is also preferable for increasing creep and creep rupture strength.

Figure 5.8 shows sulphur prints of the axial sections of rotor forging,although the material is NiCrMoV steel in an LP turbine, manufactured by aconventional Si-deoxidization and VCD method. The segregation streak inthe forging made during the VCD process has almost disappeared while thatin Si deoxidized forging is clearly observed. Low Si-VCD technology hasbeen applied in the production of rotor forging not only for LP turbines butalso for high temperature turbines.

The design of the dimensions of the ingot is important, to reduce defectsand inhomogeneities such as porosities and macro-segregation. The importanceis increased for large ingots. The ratio of height to diameter (H/D), taper ofthe ingot and hot top design are major affecting factors. It is considered thatan H/D of less than 1 is preferable to reduce the porosity size.10 A quality canbe obtained equivalent to ingots prepared by ESR by reducing the H/D to0.75 for the case of 12Cr steel. Nowadays, solidification behaviour can bepredicted by the development of numerical simulation technology contributingto the manufacture of high quality ingots.

5.2.3 Forging process

The consolidation of porosities formed during solidification andhomogenization of the material are the major aim of the initial stage of theforging. Then the material is forged to form the shape of products. Specificforging processes have been developed and applied to exert the optimumforging effects. These lead to sound and desirable properties from forging.Figure 5.9 shows an open die forging press. To consolidate porosities in thelarge ingot, a strong forging effect in the centre of a large ingot is requiredand many processes which optimize the forging temperature, shape anddimension of the dies, pressing sequence, and so on, have been developedand applied. For example, in the warm forging process, by cooling a uniformlyheated material to a surface temperature of around 800°C, forging developsa temperature difference between the interior and surface of the material. Asthe result of the difference in flow stress between the interior and outerportions of the material, the process gives strong consolidation effect. Recently

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Bottom

0 500mm

Bottom

0 500mm

(a)

C Si Mn P S Ni Cr Mo V

0.22 0.02 0.31 0.006 0.008 3.47 1.79 0.41 0.12

(b)

C Si Mn P S Ni Cr Mo V

0.34 0.38 0.51 0.021 0.008 3.60 0.19 0.43 0.15

Top

Dec. 1975

Top

Sep. 1952

5.8 Cross-sections of rotor forgings cast in air and cast through a VCD process. (a) NiCrMoV LP turbine rotor shaft from140 ton VCD ingot; (b) acid open hearth furnace air cast 75 ton ingot.

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the application of the finite element method (FEM) analysis to the problemsof plastic deformation have become general and applied widely to simulatethe effect of hot working and to determine the forging process.

An example of a forging process for CrMoV rotor forging is shown in Fig.5.10. Generally, the forging process consists of several steps. The optimumheating temperature for each hot working step is determined by consideringthe dynamic recrystallization behaviour of the material, its resistance to hotworking, the grain growth behaviour, the diffusion effect of inhomogeneitiessuch as segregation, and so on. Before the forging operation, generally, thehot top and bottom side of the ingot are discarded to remove the portion thatcontains heavy segregation and non-metallic inclusions. In the early stage ofthe forging process, upsetting is performed. The ingot height is reducedby upsetting and the diameter is increased to improve homogeneity and toincrease the forging ratio. The forging ratio generally required to develop ahomogeneous microstructure is around 3 for a conventional ingot. In case ofan ESR ingot, a lower forging ratio is acceptable owing to its inherently

5.9 A free forging press.

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good solidification structure. Dies and hot working steps are carefully designedto exert the largest forging effect. Figure 5.11(a) and (b) shows upsetting andfinish forging.

5.2.4 Heat treatment

The role of heat treatment is not only in the development of target mechanicalproperties such as strength, toughness and creep strength in forging but also

Process Sketch of process

Discard

Discard Discard

Ingot

Makinghandling stem

Upsettingrounding

Finishing thebody

Finishing thejournal

5.10 An example of the forging process for a CrMoV rotor shaft.

5.11 (a) Upsetting and (b) finish forging.

(a) (b)

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the formation of microstructure with good inspectability and heat stability.The level of performance of the forging that is achieved is defined by thefine and uniform microstructure. The overall process of heat treatment involvesseveral heating steps and largely depends on the designation of the material.The difference in heating rate, cooling rate and holding time from the surfaceto the centre of the large diameter forging need to be considered in order toattain the target properties. Heat treatment of turbine rotor forgings consistsof preliminary heat treatment which is first performed after forging andquality heat treatment which is performed subsequently. Stress relief heattreatment is also conducted after quality heat treatment.

Preliminary heat treatment

After the forging process, preliminary heat treatment is performed aimed atthe relaxation of strain introduced by hot working and refining the coarsegrain formed during the forging process. Since, generally, it is difficult todevelop a small grain structure in large forgings through dynamicrecrystallization by hot working, preliminary heat treatment is important indeveloping a fine grain microstructure that exhibits toughness and inspectabilityby an ultrasonic test (UT). A typical heat treatment that aims to refine coarsegrain is a normalizing treatment and an alternative is heat treatment applyingpearlite transformation.

Figure 5.12 shows a schematic of normalizing heat treatment and pearlitetransformation heat treatment. In the normalizing process, the material is

(a)

T3 : Temperature for pearlite transformation(b)

T1 T3

AC1AC3

AC1AC3

T1T2

5.12 Two typical heat cycles for preliminary heat treatment afterforging. (a) Normalizing and tempering heat treatment; (b) pearlitetransformation heat treatment.

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cooled once after finish forging to develop the ferritic microstructure andthen heated again to greater than the temperature at which austenitetransformation (AC3) is completely finished. When the surface portion ofthe forging reaches the austenitizing temperature, the material is kept forenough time for the temperature at the centre of the forging to reach theaustenitizing temperature. Setting the austenitizing temperature for normalizingis important in grain refining since certain types of material require acomparatively higher austenitizing temperature than AC3 to completerecrystallization as small austenite grains.

Grain refining behaviour during austenite transformation may also beaffected by the heating rate. Higher heating rates tend to develop finer austenitegrain. However, in the large forging, the heating rate that is attainable at thecentre is small. For forgings where there is difficulty in grain refining, dependingon the designation of material and dimensions of the forging, normalizingheat treatment is repeated, reducing the grain in each treatment. Pearlitetransformation is also a measure for grain refining. Pearlite transformationproceeds during cooling from the austenitizing temperature and the temperatureand time for completion of pearlitic transformation depend on the chemistryof the material. The material needs to be kept at around the temperature ofpearlite transformation nose of the time-temperature transformation (TTT)diagram. Since materials for rotor forging require good hardenability makingit difficult to proceed with pearlite transformation, a considerable holdingtime may often be required to complete the pearlite transformation. After thecompletion of pearlite transformation, the material is again austenitized todevelop a fine austenite grain.

Quality heat treatment

In the case of turbine forging materials, the target properties are developedby quenching heat treatment followed by tempering heat treatment. Quenchingis heat treatment by cooling from the austenitizing temperature which iscommonly selected at a temperature that dissolves the carbides in steels anddevelops the desired material properties such as creep strength. Attentionshould be given, however, to the avoidance of excessive grain coarsening byheating.

In order to attain high strength and toughness for rotor forging,microstructures transformed at low temperatures like martensite and lowerbainite are preferable. The formation of ferrite and pearlite, however, shouldbe avoided in order to develop good toughness and creep strength. In largeforgings, the cooling rate deep inside the forging is significantly slower thanthat of the surface region. Attention should be focused on developing thedesired microstructures even at the centre of the forging. The microstructureof a large forging after quenching can be estimated by referring to the continuous

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cooling transformation (CCT) diagram of the material corresponding to thecooling rate during quenching in each portion of the forging. Figure 5.13shows examples of CCT diagrams of quenching for a CrMoV steel and a12CrMoV steel.

Water quenching and water spray quenching are commonly applied inorder to achieve sufficient cooling to develop a martensitic or bainiticmicrostructure throughout the whole volume of the large diameter forging. A

488 489 433 450 370 363 362 269228

Ferrite

Martensite

Time (s)(b)

102 103 104 105

Tem

per

atu

re (

°C)

1000

800

600

400

200

418 381 377 344 340 335

Time (s)(a)

Tem

per

atu

re (

°C)

1000

800

600

400

200

Ferrite

Bainite

105104103102

HV

5.13 Examples of continuous cooling transformation (CCT) diagramsfor (a) CrMoV steel at an austenizing temperature of 970°C and (b)12CrMoVNbN steel at an austenizing temperature of 1050°C. HV isthe Vickers hardness valve after cooling to room temperature.

HV

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rather mild cooling effect can be obtained by immersion into an oil medium.Oil quenching reduces the risk of quenching crack but may also reduce thestrength and toughness of the material. Forced air cooling and air cooling areother methods for effecting a mild cooling rate. In the case of high temperatureturbine forging material like CrMoV steels, forced air cooling is sometimesapplied in order to attain higher creep strength.11

After quenching heat treatment, tempering heat treatment is performedconsecutively to develop the target mechanical properties below the initiationtemperature of austenite transformation (AC1) of the material. The propertiesare controlled by the tempering temperature and holding time. Highertemperature and longer holding time generally reduces strength and increasesthe toughness. Since ultimate safety is required for turbines, the forgingmust have homogeneous microstructures and inhomogeneity leading tovibration must be avoided. In order to attain a uniform heat treatment effect,quality heat treatment of rotor forging is commonly performed in a verticalfurnace under rotation of the forging. Figure 5.14 shows a rotor forgingheated in the vertical furnace.

Stress relief heat treatment is performed after machining to reduce theresidual stresses generated in the forging by the heat treatment and machining.The heating temperature is set at more than 30°C lower than the temperingtemperature so that no change in mechanical properties occurs.

5.2.5 Machining

Machining is performed in several stages of the production of rotor shaftforging as shown in Fig. 5.1. After the preliminary heat treatment, generally,machining using a lath is performed to make a smooth surface for ultrasonicexamination. The smooth surface also contributes a homogeneous heat treatingeffect in quality heat treatment. Scale from oxides formed during the forgingoperation, surface defects and decarburized surface are removed by machining.After quality heat treatment, machining for gashing and finish machining areperformed. Sometimes the central bore is machined following requests fromturbine builders/users for non-destructive examination (NDE) of the centreof the forging. Turbine builders often conduct further machining for bladeattachment, and so on.

5.2.6 Metallurgical and mechanical tests and non-destructive examination (NDE)

After heat treatment, specimens are removed from the forging and subjectedto metallurgical and mechanical tests. These specimens are typically severalportions taken from the surface and both ends of the forging. Metallurgicaland mechanical tests on the centre core are also performed for forgings with

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a central bore. The followings are typical features of metallurgical andmechanical test, although all of these are not necessarily performed for allrotor forgings. Each test is performed in accordance with standards such asISO, ASTM, DIN, JIS, and so on.

• chemical analysis• macrostructure and sulphurprint• microstructure• cleanliness• hardness test• tensile tests (at room temperature and elevated temperature)• Charpy impact test (absorbed energy, shear fracture)• creep test• creep rupture test• low cycle and high cycle fatigue tests

5.14 A rotor forging heat treated in a vertical furnace.

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• fracture toughness tests• fatigue crack growth tests.

In order to assure the quality of the forging, NDE is also performed.Surface examination includes visual examination, liquid penetrant examinationand magnetic particle examination. Ultrasonic examination is performed asa volumetric examination to detect defects in the forging. In order to havegood detectability of the defects, as already mentioned, a fine and homogeneousgrain microstructure in the forging should be formed in order to reduce thenoise level and to be able detect small indications of effects in the forging.

After stress relief heat treatment, a heat indication test (HIT) is performedto confirm the heat stability of the forging.12 Rotor forging placed in a testfurnace is heated from room temperature up to approximately operatingtemperature under rotation and then is cooled down to room temperature.The movement of the central axis of the forging during the heat cycle ismeasured precisely. Inhomogenity, such as asymmetry of the microstructureand the existence of a large residual stress, may cause large abnormal movementof the central axis which may lead to vibration in the turbine. The amount ofmovement is limited to the specified value.

5.3 Production and properties of turbine rotor

forgings for high temperature applications

5.3.1 Production and properties of CrMoV steelrotor forging

CrMoV steel designated ASTM A470 Class 8 and DIN 30 CrMoNiV 5 11are typical low alloy steels for high temperature rotor forging for HP/IPturbines. The material bears V and a fine precipitation of vanadium carbidedevelops good creep strength. The CrMoV rotor is generally used at steamtemperatures up to 566°C. Although the high temperature creep strength ofthe material is appreciable, CrMoV steel is poor in fracture toughness.Considerable discussions took place between turbine builders to determinethe optimum balance of toughness and creep strength and heat treatments inthe development of the target properties.11 In the early 1950s in the USA, anotch sensitivity problem caused by extremely high austenitizing temperaturesover 1000°C was disclosed. This sensitivity disappeared when the austenitizingtemperature was decreased to 954°C. Air cooling from the austenitizingtemperature has been adopted by turbine builders preferring a higher creepstrength. Oil quenching is also applied to improve toughness by turbinebuilders for whom toughness is important. In order to evaluate the effect ofadvanced steelmaking technology on the properties of CrMoV rotor forging,the Electric Power Research Institute (EPRI) carried out a project to evaluatethe performance of three rotor forgings made by three different advanced

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steelmaking and casting processes.13,14 The processes applied are electroslagremelting (ESR), low S and vacuum carbon deoxidization (VCD). Theseevaluation tests on the forgings suggest that improvement in productiontechnology significantly contributes to the performance of the rotor forgings.

On the other hand, improvement in the toughness by alloy modificationwas successfully attained.15,16 2CrMoNiWV steel was proved to have superiortoughness and equivalent creep strength to that of conventional CrMoVsteel. Table 5.2 shows the chemistry of CrMoV steel and 2CrMoNiWV steelfor rotor forging.

Manufacturing processes for CrMoV rotor forgings are not necessarilythe same for all the forgemasters. The process and detailed conditions aredesignated depending on their facilities and technologies. Figure 5.15 showsan example of the production process for CrMoV steel rotor forging. In thiscase, the steel is refined in EAF and LRF. Then the ingot is cast and forgingtakes place followed by preliminary heat treatment. After preliminary heattreatment, the surface of the forging is machined for a uniform heat treatmenteffect in quality heat treatment. Then quality heat treatment, quenching andtempering are performed followed by more machining. Test specimens forthe mechanical tests are removed at this stage. After machining and NDE,stress relieved heat treatment is performed followed by the heat indicationtest.

In general, EAF and subsequent LRF refining followed by vacuum castinginto a mould is a typical process widely applied for steelmaking. The ESRprocess is also applied as an alternative refining and ingot-making process.

An example of the forging process of a conventional CrMoV rotor shaftis shown in Fig. 5.10. In this case, four hot working steps are performed tocomplete the forging process.

Figure 5.16 shows an example of heat treatment of CrMoV rotor forging.After finish forging, normalizing and subsequent tempering are performedas the preliminary heat treatment. Quenching is done by forced air coolingfrom the austenitizing temperature of 950°C. Then tempering at 670°C andstress relief heat treatment at 640°C are performed.

The microstructure of the forgings quenched by forced air cooling and oilquenching is essentially an upper bainitic microstructure from the surface tothe centre of the forging. Increase of cooling rate may enhance thetransformation to lower bainite which gives better toughness. Table 5.3 showsan example of mechanical and impact properties at the surface and centreregion of the CrMoV rotor forgings with a maximum body diameter ofapproximately 1.2 m investigated in an EPRI advanced CrMoV rotorproject.13,14

The tensile properties are homogeneous from the surface to the centreportion. The fracture appearance transition temperature (FATT) of the centralregion is slightly higher compared to that of the surface region. The

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194

Table 5.2 Typical chemistry of CrMoV rotor steel, in wt%

Steel C Si Mn P S Ni Cr Mo W V

ASTM A470 Cl.8 0.25–0.36 <0.10 <1.0 <0.015 <0.018 <0.75 0.9–1.5 1.0–1.5 0.2–0.3

DIN 30 Cr 0.28–0.34 <0.10 0.30–0.80 <0.007 <0.007 0.70–0.80 1.1–1.4 1.0–1.2 0.25–0.35MoNiV 5 11*

Alloy 88 0.21–0.23 <0.10 0.65–0.75 <0.007 <0.007 0.50–0.75 2.05–2.15 0.80–0.90 0.60–0.70 0.25–0.32(22CrMoNiWV8–8)*

* VCD

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2CrMoNiWV steel is reported to develop better toughness from surface tocentre compared to the conventional CrMoV steel forging. Addition of Niand increase of Cr content in this steel is beneficial to improve the hardenabilityleading to better toughness. Figure 5.17 shows an example of creep rupturetest of CrMoV steels from oil quenched and forced air cooled forgingsshowing similar rupture strength for both quenching methods. The creeprupture strength of the 2CrMoNiWV steel forging is equivalent to that ofCrMoV steel.16

5.3.2 Production and properties of 12Cr steels

12Cr steel was introduced into service in 1960 for high temperature rotorforging operating at 566°C but its use was infrequent up until the early1970s. The material contains V, Nb and N, in addition to Mo, V and Cr. Fineprecipitates of vanadium carbide and niobium carbonitride in this type ofsteel suppress the recovery of microstructure during creep and develop highcreep strength. The material also exhibits better toughness and higher creepand creep rupture strength than CrMoV steels. Table 5.4 summarizes thetypical chemical composition of 12Cr steels for rotor forgings developed sofar.17–23 Since the application of 12CrMoVNbN steel by the General Electric

Steelmaking (EAF, LRF), casting

NDE, mechanical andmetallurgical tests

HIT

NDE (Ultrasonic test)

Forging


Rough machining

Quenching (oil)

Tempering

Machining

Stress relief

Shipping

5.15 Production sequence for a CrMoV rotor shaft.

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(c)

640°C

FC

Quenching Tempering(b)

FC

Forcedair cool

670°C

950°C

Normalizing Tempering(a)

720°C1020°C

FCFC

5.16 An example of heat treatment for CrMoV rotor shaft forging: (a)Preliminary heat treatment; (b) quality heat treatment; (c) stressrelieving heat treatment.

Table 5.3 Example of mechanical properties of EPRI advanced rotor forging project

Low S forging VCD forging ESR forging

Quenching 955°C × 25 h 950°C × 23 h 850°C × 28.5 hForced air cool Forced air cool Forced air cool

Tempering 680°C × 35 h 670°C × 52 h 670°C × 35 h

Surface Centre Surface Centre Surface Centre

0.2% yield 619 627 632 635 666 663strength (MPa)

Tensile strength 775 796 779 787 817 813(MPa)

Elongation (%) 20.5 20.1 23 22.2 19.1 20.5

Reduction in 59.3 60.3 68.6 67.3 60.7 60.3area (%)

FATT (°C) 73 94 46 67 86 97

Uppershelf 157 134 160 136 125 110energy (J)

vE24°C (J) 14 11 49 18 22 16

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Company, experience of the application of 12Cr rotor forging for hightemperature, rotor has increased. Furthermore, with the development ofadvanced power plants such as ultra super critical (USC) plants, improvedtype 12Cr steels with higher creep and creep rupture strength such as TOS107,HR1100, TMK1, COST E and COST B rotors have been developed.

Trial rotor forgings were evaluated in the EPDC project and COST501 inEurope. Improvement in the creep strength to withstand an increased steamtemperature of around 593°C, and pressure condition was made by adding orincreasing the contents of alloying elements such as W and Mo while reducingthe C content. This enhances the solid solution strengthing and stabilizationof carbides. The amount of W and Mo are generally controlled by referencingthe Mo equivalent, Mo + W/2 (wt%), of 1.5. An increase in the Mo contenthas a beneficial effect on toughness whereas an increase in W is effective indeveloping better creep strength.

A rotation test was performed on these advanced 12Cr steels and noproblems were disclosed for operation at 593°C in the advanced type 12Crsteels that have been developed. However, a ferritic superalloy A286 forgingtested at 650°C disclosed difficulty in application owing to progressive thermalfatigue damage caused by a large coefficient of thermal expansion and asmall thermal conductivity.24 In order to realize a 12Cr steel that is serviceableup to 650°C, advanced 12Cr steels have been further improved as a substitutefor superalloy A286. In these new 12CrMoV steels such as TOS110, MTR10Aand HR1200, W content is increased to enhance solid solution strengthening;Co and B are also added. Addition of Co is effective in suppressing the

Forced air coolOil quench

Larson Miller parameter, T(20 + log t) 10–3 (K, h)21.020.520.019.519.018.518.017.517.0

400

350

300

250

200

150

100

Str

ess

(MP

a)

CrMoV scatterband

5.17 Example of creep rupture strength of CrMoV rotors.

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198

Table 5.4 Typical chemistry of developed 12Cr rotor steels in wt%

Material C Si Mn P S Ni Cr Mo W V Ta Nb N Co B Ref

12CrMoVNbN 0.19 0.30 0.65 0.6 10.5 1.0 – 0.20 – 0.085 0.060 – –12CrMoVTaN 0.18 0.27 0.62 0.016 0.017 0.30 10.3 0.94 – 0.25 0.089 – 0.0412 – –TOS107 0.14 0.03 0.52 0.009 0.0024 0.73 10.36 1.05 1.06 0.21 – 0.07 0.0414 – – 17HR1100 0.13 0.28 0.58 0.58 10.23 1.13 0.23 0.22 – 0.06 0.045 – – 18TMK1 0.14 0.07 0.51 0.008 0.001 0.60 10.28 1.46 – 0.17 – 0.056 0.046 – – 19TMK2 0.12 0.06 0.49 0.52 10.38 0.28 1.98 0.19 – 0.047 0.051 – – 20COST B 0.17 0.07 0.06 0.007 0.001 0.12 9.34 1.58 – 0.27 – 0.059 0.015 – 0.0080 21COST E 0.12 0.10 0.45 0.008 0.002 0.74 10.39 1.06 0.81 0.18 – 0.045 0.052 – 0.0002 21TOS110 0.11 0.08 0.1 0.2 10.0 0.65 1.8 0.2 – 0.05 0.02 3.0 0.01 17MTR10A 0.12 0.05 0.05 <0.05 10.2 0.65 1.75 0.2 – 0.06 0.02 3.3 0.002 22HR1200 0.10 0.06 0.46 0.25 10.2 0.14 2.51 0.21 – 0.07 0.017 2.44 0.013 23

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formation of delta ferrite during solidification of the ingot by reducing theCr-equivalent without reducing the creep strength. Addition of B significantlycontributes to stabilizing carbides and suppressing the progress of recovery.25

In Europe, COST522 was performed after COST50126 and the COST536project is running with the aim of developing improved creep strength 12Crsteels for turbine rotor forging operating up to 650°C.27

Figure 5.18 shows an example of the production sequence for a 12Crrotor forging. The production process for 12Cr steel rotor forging is almostthe same as that for CrMoV steel with regard to machining after quality heattreatment, but differs for the journal overlay welding which is followed bystress relieving heat treatment.

In order to develop desirable properties for 12Cr steels, production of asound ingot is particularly important. In casting large 12Cr steel ingots,possible problems are the formation of delta ferrite and the macro-segregationin the ingot. Formation of delta ferrite basically depends on the shape of the

5.18 Production sequence for 12Cr steel rotor shaft.

Steelmaking and ingot making

NDE, Mechanical andmetallurgical test

NDE for Journal, HIT

UT

Forging


Rough machining

Quenching

Tempering

Machining

Journal overlay

Stress relieving

Machining of journal

Shipping

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gamma-loop in the phase diagram of the material. Alloy modification byadding or increasing alloying elements such as Cr, Mo, W for better creepstrength tends to increase the potential for formation of delta ferrite. Theexistence of delta ferrite reduces the strength, ductility and toughness of thematerial. Since delta ferrite is hardly eliminated by heating and hot workingduring the forging process, formation of delta ferrite during solidification ofan ingot must be avoided. The Cr equivalent value is mostly used as ameasure of the susceptibility to form delta ferrite. The value depends on thechemistry of the material and is given as follows:

Cr-eq = Cr + 6Si + 4Mo + 1.5W + 11V

+ 5Nb – 40C – 30N – 4Ni – 2Mn – 2Co (%)

It is generally said that delta ferrite will not present at a Cr-eq of less than 10.The potential for delta ferrite formation, however, strongly depends not onlyon the chemistry but also solidification rate which is significantly affectedby the size of the ingot. Therefore, the Cr-eq value should be controlledbased on experience of particular compositions and casting processes. Themajor aim of Co addition in new 12Cr steels is to decrease the Cr equivalentto avoid the formation of delta ferrite. The formation of segregation andprecipitation in eutectic carbonitrides like NbCN is another problem.Enrichment of alloying elements in the segregation zone promotes formationof eutectic carbonitrides and causes significant deterioration in the toughnessand ductility of the material. In order to avoid the formation of NbCN, thecontent of C and Nb need to be controlled.

As shown in Fig. 5.18, the raw materials are melted and subjected tooxidizing refining in an EAF. Then reduction refining and degassing isperformed in LRF after which the material is poured into an ingot mouldduring a VCD process. An ESR process can be also applied as an alternativemethod of refining and ingot making and several turbine builders have askedfor 12CrMoV steel ingots to be made by ESR in order to eliminate problemsrelated to segregation. In addition, elctroslag hot topping is also applied inthe manufacture of 12Cr steel ingots. New 12Cr rotors with alloying elementslike Mo, W, Co and B are being produced by the ESR process to avoidsegregation and to develop homogeneous material properties.

The forging process for 12Cr steels is basically the same as that forCrMoV steels, as shown in Fig. 5.10. In case of the 12Cr steels, the resistanceto deformation in hot working is larger than that of low alloy steels andforging of 12Cr steel requires a larger load and sometimes needs additionalsteps to finishing compared with CrMoV steels. Forging at high temperatureis preferable from the standpoint of hot workability. However, increasing theheating temperature encourages formation of delta-ferrite and grain growth.In addition, careful attention is needed to avoid loss of ductility caused bythe formation of low melting boride for boron bearing 12Cr steels.23

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An example of a heat treatment diagram for conventional 12Cr steel isshown in Fig. 5.19. After the forging process, normalizing heat treatmentfollowed by tempering is performed before quality heat treatment. The pearlitetransformation process can be applied as an alternative preliminary heattreatment. The major purpose of these preliminary heat treatments is grainrefining to provide good toughness and enhanced detectability indicated byNDE.

It is essential to attain a rapid cooling rate in the quenching process of12Cr steels, to avoid precipitation of carbides in the austenite phase and tocomplete martensitic transformation. Since the material has good hardenability,cooling is performed by oil quenching or water spray quenching. Theaustenitizing temperature for quenching is set at the resolving temperature ofNb carbides. In 12Cr steels, some of the retained austenite may exist in thematerial after quenching. In order to eliminate retained austenite, doubletempering heat treatment is commonly performed after quenching. The firsttempering is performed at around 600°C for complete transformation of theretained austenite. The second tempering is to develop the target mechanicalproperties. After quality heat treatment, stress relieving heat treatment isusually conducted by heating at more than 30°C below the second temperingtemperature.

Overlay welding on the journal of rotor forging is a peculiar process in themanufacture of 12Cr steel forging. Since the thermal conductivity of 12Crsteel is small, seizure is likely to occur in the bearing area.28 In order to

1050°C

Normalizing Tempering

(a)

FCFC

1100°C700°C

Quenching 1st Tempering 2nd Tempering

FCFC

Oilquench

570°C660°C

(b)

630°C

FC

(c)

5.19 An example of heat treatment for 12Cr steel rotor shaft forging.(a) Preliminary, (b) quality and (c) stress relieving heat treatments.

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avoid seizure, a CrMo steel sleeve, platings and overlay weldings have beenapplied to the journal of the 12Cr rotor forging. The reliability of the overlayis superior to other steel sleeve and platings and is currently applied to the12Cr rotor steel forging. Overlay welding is often applied to CrMo typematerials. Tensile residual stress occurs owing to the difference in thermalexpansion coefficients between the forging material and the overlay material.On the standpoint of fatigue strength, the overlay surface is sometimes rolledto attain compressive residual stress in the bearing area. Figure 5.20shows a journal overlay just after welding. The applied welding processand overlay material are selected according to the properties of the forgingmaterial.

Since 12Cr steels develops good hardenability, the microstructure of the12Cr rotor forging shows a tempered martensite structure from the centre tothe surface of the forging. Table 5.5 shows the typical mechanical propertiesof developed 12Cr steels. The balance of strength versus toughness is muchbetter than that of CrMoV steels. Generally, alloying to increase the creepstrength increases the martensitic transformation temperature and reducesthe toughness. Figure 5.21 shows the creep strength of conventional, advancedand new 12Cr steels. The 105 h rupture temperature under a stress of 100MPa is around 580°C for conventional 12CrMoVNbN steels and around600°C for advanced type steels. Although the new 12Cr steels are aimed atapplication at 650°C, the long term properties tend to deteriorate, so thatpresently 630°C is assumed to be the optimum. Efforts have been made tocharacterize premature fracture and to develop a material serviceable at650°C.

5.20 A 12CrMoV rotor forging with overlay welding on the journal.

Overlay

Overlay

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203

Table 5.5 Typical mechanical properties of the 12Cr rotor steels

Material Body Austenitizing Tempering 0.2%YS (MPa) TS El (%) RA (%) Centre Centrediameter (°C) (°C) (MPa) FATT (°C) vE (J)(mmϕ)

12CrMoVNbN 1050 570 + 62012CrMoVTaN 1262 1050 570 + 640 (0.02%YS = 709) 935 16.7 43.1 61 24TOS107 1200 1050 570 + 660 (0.02%YS = 700) 920 20.9 55.7 47 35HR1100 1200 1050 665 (0.02%YS = 695) 918 20.8 52.6 43 35TMK1 1220 1090 550 + 665 765 900 20 60 20 70TMK2 1050 550 + 680 (0.02%YS = 759) 883 19.0 59.0COST B 840 1100 590 + 700 642 801 60 33COST E 1150 1070 570 + 690 801 914 5 86TOS110 1296 1070 570 + 690 (0.02%YS = 615)MTR10A 1200 730 830 20 65 62

YS, yield strength; TS, tensile strength; EL, elongation; RA, reduction in area.

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5.3.3 High pressure–low pressure combination rotor

A combined cycle power plant, which acomplishes high efficiency, consistsof gas turbines and steam turbines. A high pressure–low pressure combinationrotor (HLP) has been often used in combined cycle power plants since thecombination type turbine has advantages, such as smaller building space,lower cost, easy maintenance, and so on, compared with the separate typesteam turbines. Figure 5.22 shows a schematic of HLP turbine rotor forgingand its required properties. Generally, both the high and low temperaturesections of the HLP turbine are made from the same material. Optimizationof chemistry of the material and manufacturing process is important indeveloping target creep strength in the HP section and good centre toughnessin the LP section simultaneously. With the increase in the capacity of thecombined cycle power plant, a larger HLP turbine has been required.

Extensive research on the development of material has been carriedout.29–34 The material needs to have good hardenability to develop a lowtemperature transformation microstructure, which leads to good toughness,even at the centre of a large diameter LP section. A creep strength equivalentto that of CrMoV steel is also required in the HP section. Table 5.6 summarizesthe typical chemistry of the low alloy materials developed for large HLPturbine rotor forging. The chemistry of these materials is almost midwaybetween the 3.5NiCrMoV steel and CrMoV steel with a small addition of theother elements. In order to attain the good toughness at the centre of the LPsection, the Cr and Ni content are optimized taking account of the balancewith creep strength. Addition or increase in the content of elements such as

Temperature (°C)650600550500

Temperature (F)950 1000 1050 1100 1150 1200

100

000

h r

up

ture

str

eng

th (

MP

a)

400

300

200

100908070

6050

New 12Cr steel

CrMoV steel

12Cr steel

Advanced 12Cr steel

5.21 Representative creep rupture strengths of CrMoV and 12Crsteels.

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Nb, W is effective in increasing the creep strength. The chemistry of thematerials was determined taking account of the optimum balance of toughnessand creep strength which is achieved by the application of the special heattreatment process.

The steelmaking process for materials for HLP turbine rotor forging isalmost the same as those applied to CrMoV steels. Conventional refiningprocess by EAF and LRF and subsequent casting into moulds can be applied.The ESR process is also used to avoid the segregation and increase thehomogeneity of the ingots. The forging process for an HLP turbine is similarto those for HP/IP turbine rotor shaft forging. After the upsetting, coggingand finish forging are performed.

Normalizing and tempering or pearlite transformation process can be appliedas the preliminary heat treatment. Differential quality heat treatment is apeculiar process in developing the creep strength in the HP portion and highcentre toughness in the LP portion simultaneously. Figure 5.23 shows aschematic of special equipment for differential heat treatment. Figure 5.24shows an example of a differential heat treatment diagram. The HP sectionis heated to a higher temperature and subjected to forced air cooling. On theother hand, the LP section is heated at a comparatively low austenitizingtemperature to suppress grain growth and develop the fine grain microstructure.Then the LP portion is water spray cooled to form a low temperaturetransformation microstructure such as lower bainite, which develops goodtoughness in all the LP portion. Differential heat treatment is performed in avertical furnace which consists of two separate temperature sections underrotation of the forging. In another differential quenching technology, controlof the cooling rate has been developed to simulate that of oil quenching byadjusting the intensity of air and water spray.29 Subsequently, each section istempered, usually differentially, to reach the target strength level.

HLP turbine (single cylinder type)

LP and HIP turbine (separate type)

LP portionHP portion

Creepstrength

High strengthHigh toughness

LP turbineHIP turbine

Creep strengthHigh strength

High toughness

5.22 Requirements for materials in a HLP turbine.

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206

Table 5.6 Chemical composition of the HLP rotor steels, in wt%

C Si Mn P S Ni Cr Mo W V Nb Ref.

EPRI-Europe 0.22 0.07 0.67 0.006 0.002 0.74 2.10 0.84 0.66 0.30 – 29

2CrMoNiWV

EPRI-Japan 0.22 0.03 0.02 0.003 0.001 2.49 1.58 1.19 – 0.23 – 30

2.5NiCrMoV

2.25CrNiMoVWNb 0.24 0.02 0.45 0.004 0.0009 1.69 2.22 1.08 0.19 0.19 0.015 31

2Cr1.8NiMoV 0.23 0.01 0.20 0.004 0.002 1.74 2.03 1.17 – 0.26 – 32

2CrMo1.5NiV 0.24 0.02 0.20 0.004 0.004 1.48 1.99 1.69 – 0.23 – 33

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Owing to the differential heat treatment, a transition area of mechanicaland impact properties unavoidably exists in the forging although the forginghas a bainitic microstructure. Several trail rotors of various chemistries,shown in Table 5.6, were investigated in detail and it was confirmed that agood balance between strength and toughness is attained at the centre of thelow pressure portion, sufficient to cope with a larger body diameter and acreep rupture strength exceeding that of CrMoV steel was attained in thehigh pressure portion.29–34 Table 5.7 summarizes the mechanical and impactproperties of the HLP rotor forgings developed.

Figure 5.25 shows the relationship between strength and 50% fractureappearance transition temperature (FATT) at the centre of the low pressureportion in the developed low alloy steel rotor forging. The balance betweenstrength versus toughness is remarkably improved in these advanced steelscompared with that of CrMoV steel. Figure 5.26 also shows an example ofcreep rupture strength.30 Including these results, the HP section of all thetrial rotor forging listed in Table 5.6 is confirmed to have the equivalentcreep rupture strength to that of CrMoV steel. Figure 5.27 shows an HLProtor forging.

5.4 Future trends

The production technology for rotor shaft forgings made of low alloy steelsand 12Cr steels has been established and high quality forgings are being

Differential coolingDifferential heating

Austenitizing temp.for LP portion

Austenitizing temp.for HP portion

Electric furnace

HP

LP

HP

LP

Fan

Waterspray

5.23 Schematic of differential heat treatment.

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manufactured and operated in power plants. Many forgemasters have beenequipped with advanced and large capacity facilities for the production ofheat-resistant steel forgings. With regard to ferritic heat-resisitant steels, asalready mentioned, new 12Cr steel developed for application up to 650°Cstill presents the problem of premature fracture at around 650°C. Therefore,efforts to eliminate the premature fracture and manufacture ferritic steel at650°C have been continuing.27

The importance of higher efficiency fossil power plant is still increasingas a solution for the problem of global warming. Several projects are attemptingto realise advanced ultra super critical (A-USC) power plants using steamtemperatures over 700°C, which is around 100°C higher than the currenttypical USC power plants. The THERMIE project was first commenced in

5.24 An example of a heat treatment process for HLP turbine rotorforging. (a) Preliminary, (b) quality and (c) stress, relieving heattreatments. (*) represents differential heat treatment.

630°C

FC

(c)

Quenching 1st tempering 2nd tempering

(∗) (∗)

(b)

FCFCHP

640°C580°C

970°C

Forcedair cool

FC

680°C

1100°C

Pearlite transformation(a)

Quenching 1st tempering 2nd tempering(∗) (∗)

FCFCLP

580°C 625°C

900°C

Waterspray

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209

Table 5.7 Mechanical properties of the HLP rotor forgings

Material Body Quenching Tempering 0.2%YS TS EL RA Centre Centre vEdiameter (°C) (°C) (MPa) (MPa) (%) (%) FATT (°C) at RT (J)(mmϕ)

EPRI/Japan 2.5MoCrMoV LP 1750 935WSQ 650 651 805 22 68 22 165HP 1720 950FAN 650 653 972 22 67 3 157

EPRI/Europe 2CrMoNiWV LP 1750 954WSQ 655 716 925 20 70 57 52HP 1250 954WSQ/Air 655 726 840 19 66 55 41

2.25CrNiMoVWNb LP 1750 900WAQ 625 682* 839 22 61 23 71HP 1000 970FAN 640 655* 781 23 63 40 46

2Cr1.8NiMoV LP 1720 950WSQ 666* 864 21 59 46 48HP 1320 970FAN

2CrMo1.5NiV LP 1356 910WQ 650 620 620 18 60 –5 72HP 968 940FAN 658

*0.02%YS; Data is not available for the blank rows.

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Tensile strength (MPa)700 800 900 1000 1100

120

80

40

0

–40

LP c

entr

e FA

TT

(°C

)

2CrMoNiWV

CrMoV

Advanced CrMoV

2Cr1.8NiMoV

2.5NiCrMoV

2.25CrNiMoVWNb2CrMoNiWV2CrMoNiWV2.5NiCrMoV

5.25 Balance between toughness of LP centre and tensile strength.

5.26 An example of creep rupture strength of material for HLP rotor.

Tangential

Longitudinal

550°C 600°C 630°C

Larson Miller parameter, T (20 + log t) × 10–3

538°C, 105 h

1716 18 19 20 21 22 23

Centre materialof HP portion

Mean creep curve ofconventional CrMoV steel

700

600

500

400

300

200

100

50

Str

ess

(MP

a)

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1998 to develop such a A-USC plant and was superceded by the AD700 andCOMTES projects.35,36 In Japan and USA, activities to realize A-USC powerplant are also continuing.

One of the most critical points for enabling the provision of A-USC plantswill be the material used for high temperature application over 700°C. Forhigh temperature turbines, several candidate materials, Alloy 617, Alloy625, Alloy 263 and Alloy 718 were selected for forgings in the THERMIEproject. Large super alloy ingots of Alloy 718 and Alloy 706 have been madeby ESR and/or VAR and much experience of production exists for land-based gas turbine disks. However the weight of ingots produced so far is atmost 20 ton. Formation of segregation depends significantly on the chemistryof the alloys and the ingot diameter. Production of a sound ingot from superalloys become more difficult with an increase in the ingot diameter. Productionwas demonstrated to be feasible for the application of superalloys to turbinerotor forging, and a 1-m diameter ESR ingot was produced.36 Weld typeconstruction type is proposed in the THERMIE project for the design of ahigh temperature rotor shaft in a A-USC plant.35,36 The turbine is constructedby welding a superalloy and 12Cr steel. This makes it possible to constructa large turbine using smaller pieces of superalloy forging(s)

Many high efficiency plants such as A-USC, combined cycle, IGCC,IGFC, and so on, have been designed. Current production technology ofheat-resistant steel turbine forgings supports the realization of these power

5.27 An HLP rotor forging.

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plants and the development of advanced production technologies undoubtedlycontributes to a further increase of plant efficiency.

5.5 References

1 Tanaka Y and Ishiguro T, ‘Development of high-purity large-scale forgings for energyservice’, Phisica Stati Solidi(a), 1996, 160 305–320.

2 Sawada S and Kawaguchi S, ‘The beneficial effect of vacuum carbon deoxidation onrotor forging properties’, Workshop on Rotor Forgings for Turbines and Generators,Palo Alto, EPRI, 1980, 5–1.

3 Ikeda Y, Yoshida H, Tanaka Y and Fukuda T, ‘Production and properties of supercleanmonoblock LP turbine rotor forging’, in Clean Steel: Superclean Steel, Nutting J andViswanathan R (eds), The Institute of Metals, 1996, 71–87.

4 Mayer W, Bauer R and Zeiler G, Development of production technology andmanufacturing experiences with superclean 3.5NiCrMoV steels’, in Clean Steel:Superclean Steel, Nutting J and Viswanathan R (eds), The Institute of Metals, 1996,89–100.

5 Jaffee R I, ‘Development of superclean rotor steels’, in Superclean Steels, Jaffee RI (ed.), Pergamon Press, 1991, 3–27.

6 Fielder H, Richter G and Scharf G, ‘Application of special metallurgical processesfor the production of highly stressed forgings’, The 8th International ForgemastersMeeting, Paper No. 15, Kyoto, Japan, 1977.

7 Morinaka K, Futamura Y, Kitagawa I and Watanabe S, ‘The manufacture of the largeESHT-J ingot’, I&SM, 1989, April, 9–15.

8 Gould G C, Long Time Isothermal Embrittlement in 3.5Ni, 1.75Cr, 0.50Mo, 0.20CSteel, ASTM STP407, ASTM, 1968, 90–105.

9 Tanaka Y, Azuma T, Yaegashi N and Ikeda Y, ‘10000H isothermal ageing test resultsof NiCrMoV Steels for low pressure steam turbines’, in Clean Steel: SupercleanSteel, Nutting J and Viswanathan R (eds), The Institute of Metals, 1996, 71–87.

10 Takenouchi T, Ikeda Y and Tanaka T, ‘Production of 12CR rotor forgings for steamturbines using advaced VCD process’, Recent Developments in Rotor Forging Steels,Iron & Steel Society, Warrendale, PA, 1990.

11 Kolar M, ‘Discussion on Heat Treatment Practices’, Workshop on Rotor Forgingsfor Turbines and Generators, Palo Alto, EPRI, 1980 5–105.

12 E-472/472M : Standard Test Method for Heat Stability of Steam Turbine Shaft Forgings,in ASTM Standard, ASTM 01.05, 2006.

13 Swaminathan V R, Steiner J E and Mitchel A, Advanced RotorForgings for High-Temperature Steam Turbine-Vol 1, EPRI Report CS-4516, Palo Alto, EPRI, 1986.

14 Swaminathan V R, Steiner J E and Mitchel A, Advanced Rotor Forgings for High-Temperature Steam Turbine – Vol 2, EPRI Report, CS-4516, Palo Alto, EPRI, 1986.

15 Finkler H and Potthast E, New 2Cr-Mo-Ni-V Steel for High-Pressure Rotors, ASTMSTP 903, ASTM, 1986, 107–123.

16 Potthast R, Viswanathan R and Wieman W, ‘Advanced 2%CrMoNiWV Steel forCombination Rotors, Proceedings of 12th International Forgemasters Meeting,Chicago, Forging Industry Education and Research Foundation, Cleveland, OH, 1994.

17 Tsuda Y, Yamada M, Ishii R, Tanaka Y, Azuma T and Ikeda Y, ‘Development of highstrength 12% Cr Ferritic steel for turbine rotor operated above 600°C’, Proceedingsof 13th International Forgemasters Meeting, Pusan, Korea, October 1997, 417–428.

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18 Iijima K, Siga M, Yoshioka T, Fukui Y, Kaneko R, Simomura K and Sasaki R,‘Steam turbine materials for improved coal fired plants’, The First InternationalConference on Improved Coal-Fired Power Plants, Palo Alto, CA, USA, EPRI,1986.

19 Tsuchiyama T, Suzuki K, Kohno M, Arihara H, Okamura M and Ohizumi H,‘Manufacturing and quality of large ESR 12%Cr rotor forging’ , Proceedings of 11thInternational Forgemasters Meeting, Terni, Italy, 1991 IX-4, 1–10.

20 Hizume A, Takeda Y, Yokota H, Takano Y, Suzuki A, Kinoshita S, Kohno M andTushichiyama T, ‘The probability of a new 12% Cr rotor steel applicable for steamtemperatures above 593°C’, Proceedings of International Conference on Advancesin Material Technology for Fossil Power Plants, Chicago, ASM International, 1987,143.

21 Berger C, Beech S M, Mayer K H, Scarlin R B and Thorton D V, ‘High temperaturerotor forgings of high strength 10%CrMoV steel’, Proceedings of 12th InternationalForgemasters Meeting, Chicago, 1994, Section 8 – 1, 1–17.

22 Kagawa Y, Tamura F, Ishiyama O, Matsumoto O, Honjo T, Tsuchiyama T, ManabeY, Kadoya Y, Magoshi R and Kawai H, ‘Development and manufacturing of the nextgeneration of advanced 12Cr steel rotor for 630C steam temperatures’ , Proceedingsof 14th International Forgemasters Meeting, Wiesbaden, Germany, 2000, 301–308.

23 Arai M, Doi H, Fukui Y, Kaneko R, Azuma T and Fujita T, ‘Improvement of longtime creep rupture properties of High WcoB containing 12Cr rotor steels for use of650°C in USC power plant’, Proceedings of the 3rd Conference on Advances inMaterial Technology for Fossil Power Plants, Wales, UK, 2001, 415–423.

24 Muramatsu K, ‘Development of ultra-super critical plant in Japan’, in AdvancedHeat Resistant Steel, Viswanathan R and Nutting J (eds), IOM Communications,London, UK, 349–364.

25 Azuma T, Kazuhiro M and Tanaka Y, ‘Effect of boron on creep strengthening in12% Cr heat resistant steel’, 14th International Fogemasters Meeting, Wiesbaden,Germany, September 2000, 283–289.

26 Thornton D V and Mayer K H, ‘European high temperature materials developmentfor advanced steam turbines’, in Advanced Heat Resistant Steel, Viswanathan R andNutting J (eds), IOM Communications, London, UK, 349–364.

27 Scarlin B, Kern T-U and Staubli M, ‘The European efforts in material developmentfor 650C USC Power Plants-COST 522’, in Advances in Material Technologies forFossil Power Plant, Viswanathan R, Gandy D and Coleman K (eds), ASM International,2004, 80–99.

28 Haas H, Simmermann A and Termuehlen H, ‘Turbines for advanced steam conditions– operational experience and development’ The First International Conference onImproved Coal-Fired Power Plants, Palo Alto, USA, EPRI, 1986.

29 Potthast E, Poppenhager J, Wiemann W and Mayer K H, ‘Advanced 2%CrMoWVsteels for combination rotors’, Proceedings of 11th International Forgemasters Meeting,Terni, Italy, Societa Delle Fucine, 1991, IX–8.

30 Tanaka Y, Ikeda Y, Ohnishi K, Kawaguchi S, Watanabe O, Kaplan A, Schwant R C,Jaffee R I and Poe G, ‘Development of a superclean 2.5NiCrMoV rotor steel for HPand LP application’, Proceedings of 11th International Forgemasters Meeting, Terni,Italy, Societa Delle Fucine, 1991, IX–7.

31 Yamada M, Tsuda Y, Watanabe O, Miyazaki M, Tabaja Y, Takenouchi T and IkedaY, ‘HLP single cylinder steam turbine rotor forgings for combined cycle power

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plants’, Proceedings of the Robert I. Jaffee Memorial Symposium on Clean MaterialsTechnology, Chicago, USA, ASM International, 1992, 161.

32 Fukui Y, Shiga M, Hidaka K, Kaneko R and Tan T, ‘Development of superclean0.2Mn-1.8Ni-Cr-Mo-V steel rotor for Hp-Lp turbine’, Proceedings of the Robert I.Jaffee Memorial Symposium on Clean Materials Technology, ASM Materials Week,Chicago, USA, ASM International, 1992, 249.

33 Tsuchiyama T, Miyakawa M, Okamura M, Matsumura K, Morita M, Yamamoto Tand Nishida S, ‘Development and production of a new HP-LP combined turbinerotor of 2CrMoV steel’, Proceedings of the Robert I. Jaffee Memorial Symposium onClean Materials Technology, ASM Materials Week, Chicago, USA, ASM International,1992, 181.

34 Kitagawa, K. Soeda, Tsuji I and Kadoya Y, ‘Manufacturing of 2 1/4CrMoV steelHP-LP merged type steam turbine rotor forgings’, Proceedings of 11th InternationalForgemasters Meeting, Terni, Italy, Societa Delle Fucine, 1991, IX–10.

35 Kern T-U, Wieghardt K and Kirchner H, ‘Material and design solutions for advancedsteam power plants’, in Advances in Material Technologies for Fossil Power Plant,Viswanathan R, Gandy D and Coleman K (eds), ASM International, 2004, 20–34.

36 Scarlin B, ‘Material developments for ultrasupercritical steam turbines’, in Advancesin Material Technologies for Fossil Power Plant, Viswanathan R, Gandy D andColeman K (eds), ASM International, 2004, 51–67.

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Part II

Behaviour of creep-resistant steels

215

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217

6.1 Introduction

The general physical and elastic properties of creep-resistant steels aresometimes the first to be considered in design calculations. For example,ferritic/martensitic steels are sometimes more favoured than austenitic steelsin power plant owing to their merit of lower thermal expansion and higherthermal conductivity. The former provides more structural stability and thelatter reduces temperature gradient within a component and therefore yieldslower thermal stress levels and greater heat transfer rates when the temperatureof a component changes during heat up and cool down.

The historical problem with superheaters, reheaters and associatedcomponents has been that of thermal fatigue, in which problems are exacerbatedby weak, heavy section components, complex geometries and bending stresses(Starr, 2002). A low coefficient of expansion and low elastic modulus isobviously advantageous, because the thermal stress, σthermal caused by atemperature change of ∆T is directly proportional to the product of thecoefficient of linear thermal expansion, α and Young’s modulus, E, i.e.

σthermal = α∆TE [6.1]

Fatigue stresses can result from pipework movement in the plant, duringheat up and cool down when load changes occur (secondary stresses). Here,the advantage is with strong thin wall members having innate flexibility andwhose deadweight does not overwhelm pipe support systems. However,during start up, rapid changes in temperature in the plant can lead to significantthrough wall temperature differences. This situation occurs more frequentlynowadays as two-shifting operations of power plants are more common. Inthis case, as well as good high strength properties, a high thermal conductivityis also of advantage. An estimate of likely susceptibility to thermal fatigue ofsteel is given by the thermal stress parameter (TSP), which is defined as:

TSP Ek

= α [6.2]

6Physical and elastic behaviour of


Y. F. Y I N and R. G. F A U L K N E R,Loughborough University, UK

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where α is the coefficient of linear thermal expansion with units of10–6K–1, E is Young’s modulus in GPa and k is the thermal conductivity inWm–1K–1. The lower the thermal stress parameter, the more resistant is thealloy to thermal stress. Thus, low thermal expansion, low elastic modulusand high thermal conductivity are desirable. More recent studies show thesusceptibility to thermal fatigue is also related to the yield strength of thematerial at a particular temperature and a modified parameter, R-value, hasbeen proposed to describe the resistance-to-crack of a material (Skelton andBeckett, 1987):

RS k

EY = 0.2

α [6.3]

where SY0.2 is the 0.2% yield strength. Opposite to TSP, higher R-valuesindicate lower probability for thermal induced crack initiation. R-values forsome common steels and a Ni-base alloy at 650°C (except those specified)are shown in Fig. 6.1. It is clear from Fig. 6.1 that ferritic/martensiticsteels and Ni-base alloys are much more resistant to thermal induced crackinitiation.

In this chapter, the general physical behaviour, mainly the thermal propertiesof creep resistant steels will be discussed. Where available, data concerningthe measured properties will be given. The implications of these propertiesfor industrial applications of creep-resistant steels are also discussed.

R-v

alu

e

7000

6000

5000

4000

3000

2000

1000

0

Austenites1Cr

(550C)NiCr20TiA1

New10Cr

(600C)

Ni-basealloy Martensites

X8CrNiMoBNb16–16

X8CrNiMoVNb16–13

X8CrNiMoNb16–16

NF 709

A286

Type304

Type316 Type

321

Type347

Esshete1250

Alloy800

6.1 R-values (resistance-to-crack) of some ferritic/martensitic,austenitic and Ni-base alloys at 650°C (except those individuallylabelled in the figure).

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Physical and elastic behaviour of creep-resistant steels 219

6.2 Elastic behaviour

The response of any material to externally applied forces is deformation, thatis a change of their size, shape and/or volume. These changes in dimensionand volume will be reversed, remain permanent or be a combination of thetwo when the externally applied load is withdrawn. When the changes arereversed, this is called elastic deformation and when the changes are permanent,this is called plastic deformation. This chapter only deals with elastic behaviourof creep-resistant steels. Plastic deformation will be discussed in later chapters.To understand the behaviour of a material under load, it is necessary todefine the terms stress and strain.

6.2.1 Stress and strain

When a pair of balanced forces, F, acting on the opposite side of a bulkmaterial with cross-sectional area, A, the force can be resolved to twocomponents, one perpendicular to the surface, Fn, and the other within theplane of the surface, Fτ, as illustrated in Fig. 6.2. The action of such a pairof applied forces can be represented by two stresses, the tensile stress, σ, andthe shear stress, τ. They are defined as follows:

σ = FAn [6.4]

A

Fn F

Fτ

Fn

Fτ

F

6.2 Tensile and shear stresses caused by a pair of externally appliedforces.

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τ τ = FA

[6.5]

The metric unit of stress is N m–2. In engineering applications, this unit is toosmall and is often replaced by MPa (1 MPa = 106 N m–2).

The basic response of a block of material to applied load is its dimensional,and/or volumetric and/or shape change depending on the nature of the appliedload. In the simplest case, the applied stress is a pure tensile stress (Fig.6.3(a)) and the block would elongate from the original length, l0, to a newlength, l0 + u. This dimensional change is described by the nominal tensilestrain, εn, which is defined as:

ε nul

= 0

[6.6]

In general, the block would shrink sidewise when it is stretched. This can berepresented by a nominal lateral strain:

ε lw

w = –

0[6.7]

These two strains are related through Poisson’s ratio:

ν εε = – l

n[6.8]

A shear stress will cause a shear strain. The engineering shear strain isdefined as (see Fig. 6.3(b)):

V0 – ∆Vp

p

p

pV0

ww τ

τ

l

σ

σw/2w/2

u/2

u/2

w0 l0

6.3 Illustration of various forms of deformation and thecorresponding strains. (a) Tensile, (b) shear, (c) dilatation.

(a) (b) (c)

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γ = wl

[6.9]

The strain caused by hydrostatic pressure is called dilatation and is definedas (see Fig. 6.3(c)):

∆ ∆ = 0

VV

[6.10]

6.2.2 Modulus of elasticity

At small strains, the relation between strain and the applied stress obeysHooke’s Law, that is the deformation of the material is linear–elastic. Therefore,in the case of simple tension, the nominal tensile strain εn is proportional tothe tensile stress:

σ = Eεn [6.11]

where E is the Young’s modulus. Similarly, the following relations hold

τ = Gγ [6.12]

p = – K∆ [6.13]

where G and K are called the shear modulus and the bulk modulus, respectively.The four elastic constants, namely E, G, K and ν, are related to each other bythe following equations (Cottrell, 1995):

K E = 2(1 – 2 )ν [6.14]

G E = 2(1 + )ν [6.15]

E KGK G

= 93 +

[6.16]

It is sometimes useful to know that the above equations can be approximatelyrewritten as (Ashby and Jones, 1996):

K ≈ E, G ≈ 3/8E and ν ≈ 0.33 [6.17]

The values of Young’s modulus of some common creep-resistant steels(chemical composition as shown in Table 6.1) are listed in Table 6.2, togetherwith values for pure iron for comparison. From Table 6.2, it is clear that thevalues of Young’s modulus do not change very much from one alloy toanother within a specific type of alloy. For example, austenitic steels listedin Table 6.2 have Young’s modulus varying from 193–200 GPa at 500°C,although their compositions differ considerably.

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222

Table 6.1 Chemical compositions of some creep-resistant steels (weight percent, Fe balance)

Steel C Si Mn Cr Ni Mo W V Nb B N Cu Ta Ti P Ref.

T22 0.12 0.3 0.45 2.25 0 1 0 0 0 0 0 0 0 0 0 1P91 0.1 0.4 0.45 9 0 1 0 0.2 0.08 0 0.05 0 0 0 0 1P92 0.07 0.03 0.45 9 0 0.5 1.8 0.2 0.05 0.004 0.06 0 0 0 0 1E911 0.1 0.17 0.47 9 0 1 1 0.2 0.07 0 0.07 0 0 0 0.007 1HCM12A 0.11 0.1 0.6 12 0 0.4 2 0.2 0.05 0.003 0.03 1 0 0 0 1410 0.14 1 1 12.5 0.5 0 0 0 0 0 0 0 0 0 0.04 2

A286 0.08 1 2 13.5–16 24–27 1–1.75 0 0.1–0.5 0 0.003–0.01 0 0 0 1.9–2.3 0 3304L 0.03 0.6 2 18 8 0 0 0 0 0 0 0 0 0 0.04 4304 0.08 0.6 2 18 8 0 0 0 0 0 0 0 0 0 0.04 5304H 0.08 0.6 1.6 18 8 0 0 0 0 0 0 0 0 0 0.04 6309 0.2 0.75 2 22–24 12–15 0 0 0 0 0 0 0 0 0 0.045 7310 0.08 0.6 1.6 25 20 0 0 0 0 0 0 0 0 0 0 5316 0.06 1 2 17 12 2.5 0 0 0 0 0 0 0 0 0.045 5321 0.08 0.75 2 17–19 9–13 0 0 0 0 0 0.1 0 0 (a) 0.045 8347 0.08 0.75 2 17–19 9–13 0 0 0 10 × 0 0 0 (b) 0 0.045 9

C–1

(a) 5 × (C + N) – 0.7 (b) 10 × C-1 minus Nb1, Klueh (2005); 2, Tsai et al, (2002); 3, De Cicco et al, (2005); 4, Ravi Kumar et al. (2006); 5, Alyousifa and Nishimura (2006); 6, Zelada-Lambri et al. (1999); 7, Ye and Wang (2006); 8, Chênea et al. (2007); 9. Laha et al. (2007).

Au

sten

itic

Ferr

itic

/m

arte

nsi

tic

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Physical and elastic behaviour of creep-resistant steels

223

Table 6.2 Some physical properties of the steels listed in Table 6.1, together with those for pure iron

Steel Tempera- Thermal Thermal Young’s Heat capacity Electrical Referencesture T (°C) expansion conductivity modulus C (J kg–1 K–1) resistivity

coefficient k (Wm–1K–1) E (GPa) ρ(10–6Ωm)α (10–6 K–1) at 20°C

Iron Fe 600 14.5 38.9 196 699 10.1 Brandes and Brook (1992)

T22 600 14.6 33 167 Starr (2002P91 600 12.6 30 168 770 Haarmann et al. (2002)P92 600 13.1 29.8 170 630 0.992 Richardot et al. (2000)E911 600 ~12 ~27 ~180 Starr (2002)HCM12A 550 12 29.5 179 744 Yoo (2004)410 600 11.6 24.9 57 Sandmeyersteel (2007)9Cr–0.12C–1Mo 20 11.15 26 402 49.9 Brandes and Brook (1992)

304L 500 18.7 21.5 193 500 72 Matweb (2007)304 500 18.7 21.5 193 500 72 Matweb (2007)304H 500 19.8 21.4 200 500 72 Hightempmetals (2007)309 500 17.6 18.7 200 502 78 Hightempmetals (2007)310 500 17.1 18.7 200 502 78 Hightempmetals (2007)316 Various 17.5 19.5 193 500 74 Matweb (2007)A286 538 17.6 23.8 163 461 Hightempmetals (2007)321 500 18.5 21.4 193 500 72 Matweb (2007)347 500 18.4 21.4 193 500 72 Matweb (2007)

Au

sten

itic

Ferr

itic

/m

arte

nsi

tic

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The values of Young’s moduli of austenitic steels are also very close tothat of pure iron at 600°C. However, the difference between these values foraustenitic steels and the Young’s modulus of ferritic steels is marked; about10–15%. This indicates that it might be difficult to achieve a desired Young’smodulus value by simple addition of alloying elements. This property ismainly determined by the type of the steel, that is ferritic/martensitic oraustenitic. Indeed, from an atomic point of view, Young’s modulus is determinedby (1) the bonding between atoms; (2) the number of atoms contributing toholding an atom in place in the structure; and (3) the distance between theatoms. The bonding of atoms in steels is mainly metallic, whether it isferritic/martensitic (BCC) or austenitic (FCC). However, the FCC austeniticsteel is generally more densely packed when compared with the BCCmartensitic steel. In another words, the average distance between atoms inaustenitic steels is less than that in martensitic steels. Therefore, the attractingforces between atoms which hold the atoms together are stronger in austeniticsteels than those in martensitic steels. This means that higher stressesmust be applied to cause the same deformation, that is, a higher Young’smodulus.

At higher temperatures, the atoms are more active because they havemore kinetic energy. Therefore, it is easier to pull them apart and this resultsin a decrease in Young’s modulus. An example of decreasing Young’s modulusas temperature increases is shown in Fig. 6.4 for T/P91 and T/P92 steels.Clearly, the dependence of Young’s modulus on temperature is very significant.The slope of the curves increases with increasing temperature. It is also

P91P92

Temperature, T(°C)700600500400300200

210

200

190

180

170

160

You

ng

’s m

od

ulu

s, E

(G

Pa)

6.4 Dependence of Young’s modulus of T/P91 (Haarmann et al., 2002)and T/P92 (Richardot et al., 2000) on testing temperature.

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worth noting that the Young’s modulus of T/P92 in Fig. 6.4 is it is slightlyhigher than that of T/P91 at the same testing temperature. At normal servicetemperatures, Young’s modulus of creep-resistant steels can be as low as halfthat at room temperature. Therefore, it is necessary to give the testingtemperature when reporting Young’s modulus values, otherwise confusionmay occur.

6.3 Thermal properties of creep-resistant steels

6.3.1 Thermal expansion

Thermal expansion describes the response of a material to heat input. Theextent of expansion or contraction depends on the material and temperaturechange and is commonly described by the linear coefficient of thermalexpansion or CTE in short. Assume a steel bar of initial length l0 undergoesa temperature increase from T to T + ∆T and its length changes to l = l0 + ∆l.The CTE, α, is defined as:

α ∆∆ =

0

ll T

[6.18]

Therefore, CTE is the relative increase in length of material when temperatureincreases by one degree Kelvin. It is clear that the metric unit of CTE ism · m–1 · K–1 or K–1. In practical applications, the expansion or contractionof a material is very small, this unit is too big and 10–6K–1 is commonly usedinstead. The values of some common power plant creep resistant ferritic/martensitic and austenitic steels are listed in Table 6.2.

Design of the structures or components must take into consideration thechange of dimension of the material as temperature changes during normalservice. This means that in the design process, the structure must be allowedto accommodate the dimensional change of the material from room temperatureto normal operational temperature. This produces the requirement to putexpansion joints or expansion loops in the structure. If the component isrestrained, thermal expansion may cause buckling or bending of the component.Buckling and bending can also occur when two steels with markedly differentthermal expansion coefficients are fabricated together and subsequently heatedor cooled.

Another effect is thermal stress. When a component expands, tensile forcesare created. On the other hand, compressive forces are created when acomponent contracts. If we recall that strain in a pure tensile situation isdefined as the relative increase of the length of the material, (Equation [6.6],we can define a thermal strain here as well,

ε ∆thermal

0 = l

l[6.19]

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From Equation [6.18], we obtain

εthermal = α∆T [6.20]

Therefore, thermal strain is directly proportional to the change intemperature. However, it should be noted that the CTE varies with temperature,thus Equation [6.20] only can be applied in a small temperature range or ina sense of average behaviour. In fact, CTE as defined in Equation [6.18] isthe mean CTE of the material over the temperature range T to T + ∆T.Outside this range, the CTE might have a different value. To obtain a CTEat temperature T, the range ∆T must be acceptably small. Therefore, in reportingthermal expansion coefficients of steels, it is common for the mean values ofCTE from room temperature to a test temperature to be cited. Of course,these mean values are also dependent on the test temperature. Figure 6.5shows the mean thermal expansion of three creep-resistant steels, namelyT/P91, T/P22 and TP304H, as a function of test temperature. In the temperaturerange shown, mean CTE for all materials increases approximately linearlywith increasing temperature and the increase in CTE from room temperatureto 600°C is about 25%. Of course, in design of power plant components, themean CTE in the range from room temperature to the operational temperatureis more relevant.

The difference in thermal expansion coefficient between different creep-resistant steels is significant. The CTEs for austenitic steels are about 50%higher than those for ferritic and martensitic steels, as shown in Fig. 6.5 andTable 6.2. CTE generally increases with increasing bond energy (Incroperaet al., 2006). Bond energy depends on the nature of the interaction betweenatoms forming the solid and the bond length. The stronger the interaction,the higher is the bond energy. On the other hand, the shorter the bond length,

Temperature (°C)

Temperature (°F)

6005004003002001000

11129327525723922123220

15

10

Co

effi

cien

t o

f lin

ear

exp

ansi

on

(10

–6 K

)

Co

effi

cien

t o

f lin

ear

exp

ansi

on

(10

–6 /°

F)

11.2

8.4

5.6

6.5 The temperature dependence of mean linear thermal expansioncoefficients of T/P91, T/P22 and TP304H (Haarmann et al., 2002).

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the higher is the bond energy. In the case of steels, the main bonding isbetween iron atoms and the nature of interaction is the same in all the casesof ferritic, martensitic and austenitic steels. Therefore the bond energy ismainly determined by the bond length. The bond length can be approximatedusing the equilibrium distance between the centres of atoms. Thus, the bondenergy increases with decreasing interatomic distance. As mentioned earlier,the FCC austenitic steel is denser than the BCC ferritic/martensitic steel,hence it has a higher thermal expansion coefficient.

However, the bond strength may be altered by adding different alloyingatoms to the material. This provides the possibility for developing low thermalexpansion steels via alloying. The significant variation of CTE within theferritic/martensitic group is good evidence of such an alteration. The variationsof CTE as a function of the main alloying elements, namely carbon andchromium for ferritic/martensitic steels and carbon, chromium and nickelfor austenitic steels are shown in Figs. 6.6 and 6.7. Figure 6.6 shows cleartrends of the decrease of CTE with increasing concentration of both carbonand chromium in the ferritic/martensitic steels. The only exception in Fig.6.6(a) is T/P22 which has exceptionally low chromium content (2.25 wt%).This exception is believed to be due to the effect of chromium which outweighsthe effect of carbon. However, the effects of both carbon and chromium onCTE of austenitic steels are not so clear, as shown in Fig. 6.7(a) and (b). Onthe other hand, Fig. 6.7(c) shows a clear decrease in CTE with increasingnickel content of the steels. On the whole, it is reasonable to conclude thatCTE decreases with increasing concentration of C, Cr and Ni in both ferritic/martensitic and austenitic steels.

Yamamoto et al. (2003) have carried out some regression analysis of CTEdata for Ni-base superalloys and found the contribution of different alloyingelements to CTE from room temperature to 700°C could be described by theformula below:

α700°C = 13.8732 + 7.2764 × 10–2 × [Cr] + 3.751 × 10–2

× [Ta + 1.95Nb + 1.9774 × 10–2[Co] + 7.3 × 10–5 × [Co]

× [Co] – 1.835 × 10–2 × [Al] – 7.9532 × 10–2

× [W] – 8.2385 × 10–2[Mo] – 1.63381 × 10–1 × [Ti]

[6.21]

where α700°C is the mean thermal expansion coefficient from room temperatureto 700°C in 10–6 K–1 and [element name] is the concentration of the elementin weight percent. This kind of equation is very useful in developing lowCTE alloys.

Thermal stress caused by thermal expansion is described by Equation[6.1], that is, σthermal = Eεthermal = α∆TE. Localized stresses from thermal

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expansion during heating and cooling can contribute to the problem of stresscorrosion cracking in an environment which would not normally attack thesteel. These applications require design to minimize the adverse effects oftemperature variations such as the use of expansion joints to permit movementwithout distortion of the component and the avoidance of notches and abruptchanges of section. Thermal stress is also a contributor to fatigue as discussedin the introduction.

Carbon content, C (wt%)0.160.140.120.100.080.06

15

14

13

12

11

CT

E, α

(10

–6 K

–1)

(a)

AISI410

T/P22

T/P92

T/P91

E911 HCM12A

Cr content, C (wt%)

(b)

1614121086420

CT

E, α

(10

–6 K

–1)

15

14

13

12

11

AISI410

HCM12A

T/P92

T/P91

E911

T/P22

6.6 Linear thermal expansion coefficient of some ferritic/martensiticcreep-resistant steels as a function of the content of (a) carbon and(b) chromium at 600°C.

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Ni content, C (wt%)(c)

CT

E, α

(10

–6 K

–1)

28242016128

20

19

18

17

16

310

A286309316

347321

304H

304/304L

Cr content, C (wt%)(b)

CT

E, α

(10

–6 K

–1)

272421181516

17

18

19

20

21

310

309316A286

304H

304/304L321347

Carbon content, C (wt%)(a)

CT

E, α

(10

–6 K

–1)

0.250.200.150.100.050.0016

17

18

19

20

309

310

A286316

347304

321304L

304H

6.7 CTE as a function of the content of (a) carbon, (b) chromium and(c) nickel in some austenitic creep-resistant steels.

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6.3.2 Thermal conductivity

In physics, thermal conductivity, k, is the property of a material that measuresits capability to conduct heat. Heat transfer by conduction involves transferof energy within a material without any motion of the material as a whole.The rate of heat transfer depends on the temperature gradient and thermalconductivity of the material. Higher thermal conductivity indicates higherability for transferring heat. This implies that if the same volume of twodifferent materials is heated, the one with higher thermal conductivity willhave lower thermal gradient throughout the bulk. Therefore, higher thermalconductivity is desirable in at least two ways. First, higher thermal conductivityreduces the temperature gradient within a component and therefore reducesthermal stress caused by temperature change. This has a direct influence onthe thermal stress parameter, as indicated in Equation [6.2]. Second, higherthermal conductivity allows thicker sections of material to be used, hencedecreasing the demand on the strength of the material.

The definition of thermal conductivity is straightforward. Assume a thinslab of a material with cross-sectional area A and thickness ∆x; one side ofthe slab is kept at temperature T + ∆T and the other kept at T, as shown in

Fig. 6.8, then the temperature gradient across the slab is ∆∆

Tx

. The heat Q

transferred through area A from the high temperature to low temperature side

is directly proportional to the temperature gradient ∆∆

Tx

, area A and the time

∆t during which the heat transfer takes place, that is:

Q Tx

A t = – κ ∆∆ ∆ [6.22]

Q A

∆xT T + ∆T

6.8 Definition of thermal conductivity.

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The minus sign here signifies that the heat transfer is from the highertemperature to lower temperature side, that is, opposite to the direction oftemperature gradient. The proportionality constant is the thermal conductivityof the material. Rearranging Equation [6.22],

κ ∆∆ ∆

= –

QTx

A t[6.23]

Therefore, thermal conductivity is the quantity of heat, Q, transmitted in unittime through a unit normal area caused by a unit temperature gradient understeady state conditions.

Different materials transfer heat in different ways and therefore havevastly different thermal conductivities. Gases transfer heat by direct collisionsbetween molecules and their thermal conductivities are low compared tomost solids since they are dilute media. Non-metallic solids transfer heat bylattice vibrations. In metals, heat transfer is accomplished by mobile electronswhich also participate in electrical conduction. Therefore, metals are muchbetter thermal conductors than non-metals. Because the carriers of heat inmetals are the mobile electrons, the thermal conductivity of metals is determinedby the number density and mobility of the mobile electrons in the metal.Higher mobile electron density and higher electron mobility result in higherthermal and electrical conductivity. Thermal conductivities of some creep-resistant steels are listed in Table 6.2.

Thermal conductivity of ferritic/martensitic steels is about two-thirds ofthat of pure iron, but is about 50% higher than that of austenitic steels (Table6.2). As mentioned before, austenitic steels have denser structures than ferritic/martensitic steels; therefore the number density of mobile electrons is higherthan that in ferritic/martensitic steels. This should lead a higher thermalconductivity for austenitic steels. However, owing to the smaller inter-atomicdistance (this is also a measure of the atom size), the attraction between thenuclei and the electrons is stronger and the mobility of mobile electrons inaustenitic steels is much lower than that in ferritic/martensitic steels. Thiseffect outweighs the effect of the number density of mobile electrons andtherefore the thermal conductivity of austenitic steels is lower than that offerritic/martensitic steels.

As thermal conductivity is a kind of measure of the effectiveness ofenergy transfer by the collision of mobile electrons, anything that affects themobility of electrons may have an influence on thermal conductivity. Forexample, grain boundaries, second phase particles and non-metallic inclusionsin steels all affect the thermal conductivity of the material. Figure 6.9 showsthermal conductivity as a function of steel composition for (a) carbon steelsas a function of carbon content; (b) chromium steels as a function of chromiumcontent; (c) nickel steels as a function of nickel concentration; and (d) tungsten

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steels as a function of tungsten concentration. Figure 6.9 is plotted accordingto data from www.EngineersEdge.com.

As can be seen from Fig. 6.9, there is a general trend of decreasingthermal conductivity with increasing concentration of alloying element. Theeffect is significant. For example, the thermal conductivity of chromiumsteel decreases to less than one-third of its original value when the concentrationof chromium increases from 0 to about 20 wt%. Interestingly, there is anupturn in the relation between thermal conductivity of nickel steel and theconcentration of nickel at 40 wt% (Fig. 6.9(c)). Therefore, the minimum ofthermal conductivity appears to occur at a nickel concentration of ~50 wt%.Further increase of nickel content results in the material becoming a nickel-base alloy rather than a steel. Nickel-base alloys have tighter compositionspecifications than steel, which means that they contain fewer impurities.Therefore, thermal conductivity increases.

W content, C (wt%)(d)

1086420

Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1)

130

120

110

100

90

80

Ni content, C (wt%)(c)

806040200Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1)

140

120

100

80

60

40

20

0

Cr content, C (wt%)(b)

20151050Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1)

140

120

100

80

60

40


1.61.41.21.00.80.60.4Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1)

95

90

85

80

75

70

65

60

6.9 Thermal conductivity of some steels as a function of content(a) carbon, (b) chromium, (c) nickel and (d) tungsten. Data fromwww.EngineersEdge.com.

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Creep-resistant steels are more complicated, often containing up to around15 or more alloying elements. Therefore, no straightforward relationshipsbetween conductivity and composition are expected. However, it still holdsthat there is a general trend of decreasing thermal conductivity with increasingconcentration of alloying element, as shown in Fig. 6.10 for some commonferritic/martensitic creep-resistant steels. Thermal conductivity of ferritic/martensitic steels is much lower than that of pure iron and decreases withincreasing concentration of carbon or chromium. The trend in Fig. 6.10(a) isnot so clear owing to the exceptionally high thermal conductivity of T/P22.


Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1)

25

30

35

40

0.160.120.080.040.00

Pure Fe

AISI410

E911T/P92

T/P91

T/P22

HCM12A

Chromium content, C (wt%)(b)

14121086420

40

35

30

25

Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1) Pure Fe

T/P22

T/P92E911

T/P91

AISI410

HCM12A

6.10 Thermal conductivity of some common ferritic/martensiticcreep-resistant steels as a function of (a) carbon and (b) chromiumcontent.

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This is because T/P22 has very low chromium content and the effect oflower chromium concentration outweighs that of higher carbon concentration.Similar trends are found for austenitic steels as shown in Fig. 6.11. However,thermal conductivity values for austenitic steels are much lower than thosefor ferritic/martensitic steels and its variation as a function of composition issmaller than that in the case of ferritic/martensitic steels.

Thermal conductivity of creep-resistant steels is also a function oftemperature. As mentioned earlier in this section, metallic solids conductheat via the collision of mobile electrons and thermal conductivity is ameasure of the effectiveness of the energy transfer during the collision.Generally speaking, electrons move faster at higher temperatures and theprobability of collision with others is much higher. Therefore, energy can bemore effectively transferred from locations with higher temperature to locationswhere temperature is lower and thermal conductivity generally increaseswith increasing temperature. Figure 6.12 shows that this is true for T/P91,TP304H and T/P22 at low temperatures.

6.4 Electrical resistivity and conductivity of creep-

resistant steels

Metals, including steels, are good electrical conductors. However, they doresist the flow of electrical current. The ability of a metal to resist electricalflow is represented by its electrical resistivity. If a sample of metal withuniform cross-sectional area A and length l has an electrical resistance R,then its resistivity, ρ is defined by

R lA

= ρ [6.24]

or

ρ = Rl/A

[6.25]

Therefore, the resistance of a sample of steel is inversely proportional to itsconducting area. This has implications for the monitoring of crack initiationand propagation in steels. If a crack grows or propagates in a steel component,the conducting area will be reduced and the resistance to electrical currentwill be increased. The simplest way of measuring resistance is to use Ohm’slaw, that is:

I VR

R VI

= or = [6.26]

where I is the electrical current and V the applied electrical potential orvoltage. In SI units, the unit of electrical resistivity is ohm-meter or Ωm. Forsteels, the resistivity is in the range 1~10 × 10–7 Ωm and some values arelisted in Table 6.2.

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0.200.150.100.050.00

40

35

30

25

20

15

Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1)

309

Pure Fe

304L316 310

A286304/304H321/347

Nickel content, C (wt%)(c)

Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1) 40

35

30

25

20

15302520151050

Pure Fe

A286

310309

316

321/347304L/304304H

Chromium content, C (wt%)(b)

Th

erm

al c

on

du

ctiv

ity,

k (

W m

–1 K

–1) 40

35

30

25

20

152520151050

Pure Fe

309310316

A286304L/304/304H321/347

6.11 Thermal conductivity of some austenitic creep-resistant steels asa function of (a) carbon, (b) chromium and (c) nickel content.

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Electrical conductivity, often represented using the Greek letter σ, is a measureof the ability of a conductor to conduct electrical current. It is defined as thereciprocal of electrical resistivity, i.e.

σ ρ = 1 [6.27]

Metals conduct electricity by the movement of free electrons, which alsoconduct heat as discussed earlier. Therefore, electrical conductivity of steelsis related to their thermal conductivity. Generally speaking, a good thermalconductor is also good at conducting electricity. This relationship is describedby the Wiedemann–Franz Law:

κσ = LT [6.28]

where κ is thermal conductivity, σ is electrical conductivity and T is theabsolute temperature. The proportional constant L is called Lorenz numberand it has the value:

L ke

= 3

= 2.45 10 W K2 2

2–8 –2π × Ω

where k is the Boltzmann constant and e is the charge of electrons.Electrical resistivity of creep-resistant steels is dependent on various factors,

such as the purity of the material and the temperature. From Table 6.2,electrical resistivity of both ferritic/martensitic and austenitic steels is muchhigher than that of pure iron, indicating that electrical resistivity increaseswith the degree of alloying. For austenitic steels, electrical conductivity does

TP304H

T/P91

T/P22

Temperature (°C)

Temperature (°F)

6005004003002001000

40

30

20

10

Th

erm

al c

on

du

ctiv

ity

(Wm

K–1

)

22.4

16.8

11.2

5.6 Th

erm

al c

on

du

ctiv

ity

((B

tuft

h–1

°F–1

)

932 111275257239221232

6.12 Examples of thermal conductivity as a function of temperaturefor some creep-resistant steels.

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not differ very much from alloy to alloy, but higher than the usual values, asfor ferritic/martensitic steels, indicates that the structure of the material playsa vital role in conducting electrical current. Electrical resistivity is alsotemperature dependent. Examples of the dependence of electrical resistivityon temperature for ferritic/martensitic and austenitic steels are shown in Fig.6.13. Clearly, electrical resistivity of both ferritic/martensitic and austeniticsteels increases with increasing temperature.

Austenitic steels have much higher electrical resistivity than ferritic/martensitic steels, that is, ferritic/martensitic steels are much better electricalconductors than austenitic steels, which is as expected because ferritic/martensitic steels have much higher thermal conductivity, as discussed inSection 6.3. It is interesting to note that the temperature dependence ofelectrical conductivity of steels is opposite to that of thermal conductivity. InSection 6.3 we have shown that thermal conductivity generally increaseswith increasing temperature. Figure 6.13 shows that electrical resistivity ofsteels rises as temperature increases. Therefore, electrical conductivity decreaseswith increasing temperature. This is due to the different mechanisms ofconducting heat and electricity. Although metals conduct both heat andelectricity by the movement of mobile electrons, the nature of the movementis different. Heat is transferred from a location of higher temperature to thatof lower temperature by collision of electrons which depends on the randommovement of mobile electrons, no directional motion of the electrons as awhole is involved. Temperature in fact is a measure of the intensity of suchrandom movement and the higher the temperature, the higher is the frequencyof collision. As a result, mobile electrons conduct heat more effectively athigher temperatures. It can be shown that thermal conductivity is directly

6.13 Examples of temperature dependence of electrical resistivity offerritic/martensitic and austenitic creep-resistant steels.

9Cr–0.12C–1MoAISI 321/347

Temperature, T(°C)800 10006004002000

120

100

80

60

40

Ele

ctri

cal r

esis

tivi

ty, ρ

(Ω

m)

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proportional to the average speed of electrons, which in turn is directlyproportional to the square root of absolute temperature. Therefore, thermalconductivity increases with increasing temperature.

In the case of conducting electricity, the mobile electrons move in thedirection opposite to the electrical current as a whole, in addition to therandom movement. Electrical conductivity is dependent on the mass directionalmovement (or sometimes called drift) of the electrons rather than the randommotion. In fact, the collision between the electrons owing to random movement(and lattice vibration) scatters the electrons and prevents them from drift.Therefore, electrical conductivity is inversely proportional to the averagespeed of Brownian motion of electrons. At higher temperatures, the randommovement is more intense and it is more difficult for electrons to drift in acertain direction. Therefore electrical conductivity decreases with increasing

temperature. Combining κ ∝ T and σ 1∝T

, we obtain the Wiedemann–

Franz Law, Equation [6.28].

6.5 Implications for industries using


Although general physical properties of creep-resistant steels are not discussedso much as creep strength in the literature, they have important implicationsfor the design and service of the component. Generally speaking, lowermodulus of elasticity, lower thermal expansion and higher thermal conductivityare desirable if secondary stresses caused by heat input and temperaturechanges are considered. As two-shifting and weekend close down becomecommon, thermal fatigue caused by such operations at power plants must betaken into consideration. Thermal stress caused during cool down and heatup of power plant components is directly linked to all three parametersdiscussed in this chapter. Higher thermal expansion coefficient leads to higherthermal strain. Higher Young’s modulus results in higher thermal stress atthe same thermal strain. Material with lower thermal conductivity cannotconduct heat effectively and gives rise to a higher temperature gradient andtherefore a higher thermal strain. If the strength of the material is acceptableunder operational conditions, the material with lower thermal expansion,higher thermal conductivity and lower Young’s modulus should be used. Ifthis material cannot meet the strength requirements and the use of materialswith higher thermal expansion, lower thermal conductivity and higher Young’smodulus are inevitable and acceptable ways of reducing or avoiding relatedproblems must be found. For example, based on calculations for thick sectionaustenitic pipework, Fleming et al. (1997) pointed out that to avoid unacceptablethermal stresses in a fully austenitic plant it would take several days to heatup and cool down.

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The combination of high thermal expansion and low thermal conductivitymeans that precautions must be taken to avoid adverse effects. For example,during the welding of austenitic steels, measures such as low heat input anddissipation of heat using other media may be required. When heated fromroom temperature (20°C) to 600°C, an austenitic pipe of length 1 m mayexpand by 1 cm. Expansion loops/joints may need to be used to accommodatesuch a dimensional change. Precautions must also be taken to reduce stresscorrosion cracking caused by or assisted by thermal stress caused by lowthermal conductivity and high thermal expansion. On the other hand, higherthermal conductivity allows thicker section components to be used and thisreduces the demand on the strength of the material.

6.6 Future trends

Owing to their low thermal expansion and high thermal conductivity,tremendous efforts have been made to develop better ferritic/martensiticsteels. In recent years, material scientists and engineers seem to have concludedthat with less than 11–12 wt% chromium content, ferritic/martensitic steelscannot overcome the problem of steam oxidation. Therefore, much attentionhas been paid to developing higher chromium ferritic/martensitic steels.However, the attempt has not been successful. These higher chromium ferritic/martensitic steels seem to promote the formation of the so-called Z phasewhich has detrimental effects on the long term creep strength of the alloy.For both environmental and economic reasons, power plant designers arekeen to operate plant at higher temperatures and pressures to improve theefficiency. Under such conditions, the problem of oxidation of ferritic/martensitic steels becomes more serious. Thus, more recently, materialsscientists and engineers have suggested abandonment of ferritic/martensiticsteels and development of new austenitic steels for higher temperatureoperation. Although thermal expansion may be reduced via the route ofalloying, it is difficult to obtain low thermal stress during heat up or cooldown periods via alloying because alloying inevitably decreases thermalconductivity of the alloy. Under such considerations, means of avoidingunacceptable thermal stress must be found in the future at the design stageto accommodate higher thermal expansion and lower thermal conductivityof austenitic steels.

6.7 References

Alyousifa O M and Nishimura R (2006), ‘The effect of test temperature on SCC behaviourof austenitic stainless steels in boiling saturated magnesium chloride solution’, CorrosionScience, 48, 4283–4293.

Ashby M F and Jones R H (1996), Engineering Materials 1 – An Introduction to theirProperties and Applications, Butterworth Heinemann, Oxford.

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Brandes E A and Brook G B (eds) (1992), The Smithells Metal Reference Book, ButterworthHeinemann, Oxford.

Chênea J, Brassa A-M, Trabucb P and Gastaldi O (2007), ‘Role of microstructure andheat treatments on the desorption kinetics of tritium from austenitic stainless steels’,Journal of Nuclear Materials, 360, 177–185.

Cottrell A (1995), An Introduction to Metallurgy, 2nd edition, The Institute of Materials,London.

De Cicco H, Luppo M I, Raffaeli H, Di Gaetano J, Gribaudo L M and Ovejero-García(2005), ‘Creep behavior of an AISI286 type stainless steel’, Materials Characterization,55 97–105.

Fleming A, Buchanan L W and Maskell R V (1997), ‘Overview of material developmentrequirements for boiler plant’, Materials Issues in Heat Exchangers in Boilers Conference,Starr F and Meadowcroft B (eds), IOM London, 109–117.

Haarmann K, Vaillant J C, Vandenberghe B, Bendick W and Arbab A (2002), The T91/P91 Book Vallourec and Mannesmann tubes, Boulogne.

Hightempmetals (2007), www.hightempmetals.comIncropera F P, DeWitt D P, Bergman T L and Lavine A S (2006), Fundamentals of Heat

and Mass Transfer, 6th edition, Wiley, New York.Klueh R L (2005), ‘Elevated temperature ferritic and martensitic steels and their applications

to future nuclear reactors’, International Materials Review, 50 (5), 287–310.Laha K, Kyono J and Shinya N (2007), ‘An advanced creep cavitation resistance Cu-

containing 18Cr–12Ni–Nb austenitic stainless steel’, Scripta Materialia, 56, 915–918.

Matweb (2007), www.matweb.comRavi Kumar B, Das S K, Mahato B, Arpan Das and Ghosh Chowdhury S (2006), ‘Effect

of large strains on grain boundary character distribution in AISI 304L austenitic stainlesssteel’, Materials Science and Engineering A, 454–455, 239–244.

Richardot D, Vaillant J C, Arbab A and Bendick W (2000), The T92/P92 book, Vallourec& Mannesmann Tubes, Boulogne.

Sandmeyersteel (2007), www.sandmeyersteel.comSkelton P and Beckett B E (1987), ‘Thermal Fatigue Properties of Candidate Materials

for Advanced Steam Plant’, Conference on Advances in Material Technology forFossil Power Plants, ASM, Ohio, 359–366.

Starr F (2002), ‘Potential issues in the cycling of advanced power plants’, OMMI, 1 (1),1–19.

Tsai M C, Chiou C S, Du J S and Yang J R (2002), ‘Phase transformation in AISI 410stainless steel’, Materials Science and Engineering A, 332, 1–10.

Yamamoto R, Kadoya Y, Kawai H, Magoshi R, Noda T, Hamano S, Ueta S and Isobe S(2003), ‘New wrought Ni-based superalloys with low thermal expansion for 700Csteam turbines’, Energy Technology, 21, 1351–1360.

Ye W, Li Y and Wang F (2006), ‘Effects of nanocrystallization on the corrosion behaviourof 309 stainless steel’, Electrochimica Acta, 51, 4426–4432.

Yoo, Y-S (2004), ‘Study of LBB assessment methodology applied to a 12Cr series ferritesteel piping structure for FBRs’, 24, 27–36.

Zelada-Lambri G I, Lambria O A and Rubiolob G H (1999), ‘Amplitude dependentdamping study in austenitic stainless steels 316H and 304H. Its relation with themicrostructure’, Journal of Nuclear Materials, 273, 248–256.

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241

7.1 Introduction

In the early 1950s, Sherby and coworkers (e.g., Sherby et al., 1954) analysedcreep data and diffusion data of several metals and claimed that thetemperature dependence of both phenomena is similar to each other. Thisfinding first gave us a sound physical base for discussing creep mechanisms.In those days, however, creep data were very limited and diffusion datawere also a few and their values were scattered over a wide range. Therefore,they had to base their discussion on average values of these limiteddata.

Today, the view is well established (Sherby and Burke, 1967; Mukherjee,et al., 1969) that the temperature dependence of creep at high temperature isclose to that of diffusion in pure metals (see Fig. 7.1). A similar correlationhas also been found in solid solution alloys (Monma et al., 1964).

Diffusion is one of most fundamental processes governing creepdeformation. In this chapter diffusion behaviour in metals and alloys will beoutlined and then the role of diffusion in creep deformation will be discussed.Finally, some fundamental diffusion data which are deemed to be useful indiscussing creep of steels will be cited.

7.2 Diffusion and creep

7.2.1 Activation energies

It is worthwhile noting that there is an essential difference in the physicalmeaning of temperature dependence of creep and of diffusion, although thetemperature dependences of these two phenomena are close to each otherunder some conditions.

Diffusion in simple metals can be recognized as a thermally activatedprocess, where the activation energy is the sum of the formation energy andthe migration energy of vacancies. The pre-exponential (or frequency) term

7Diffusion behaviour of creep-resistant

steels

H. O I K A WA and Y. I I J I M A,Tohoku University, Japan

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is essentially the number of lattice sites in the system and not a function oftemperature.

Creep is a time-dependent deformation process which is a result of complexdislocation behaviour. Its temperature dependence is influenced by manyfactors which may have their own temperature dependence. The so-calledactivation energy of creep is a temperature dependence of the steady-state(or minimum) creep rate under a given stress, derived on the assumption thatone thermally activated process occurs in the phenomenon. Therefore, theactivation energy of creep is an apparent one in the sense of rigorous thermalactivation rate theory. It must be kept in mind during discussion of creepmechanisms that the pre-exponential term in the form of a thermally activatedrate equation contains many factors which might depend on temperatureeven under constant stress.

7.2.2 Time-dependent deformation and diffusion

Creep deformation is time-dependent straining under a constant applied stress,or under a given load in many cases. At very high temperature under verylow stress, ‘diffusional creep’ occurs in (pure) metals. Under these conditions,creep strain arises directly from the movement of atoms. The temperaturedependence (the activation energy) of creep is the same as that of vacancydiffusion.

At temperatures higher than about a half of the melting temperature,0.5Tm, and under ordinal creep conditions, the temperature dependence of

Qc/(kJ mol–1)100050020010050

Qd/(

kJ m

ol–1

)

1000

500

200

100

MoTa

Nb

γ Feβ Co

Ni

α Fe

PtCu

Ag Auα Ti

β Ti

AlMg

Pb

Zn

Cd

7.1 Correlation between the activation energies of high-temperaturecreep, Qc, and of lattice self-diffusion, QD, in pure metals.

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Diffusion behaviour of creep-resistant steels 243

creep (strain rate) is also similar to that of diffusion. This similarity, however,does not mean that creep strain arises directly from the movement of atoms.Rather, the similarity indicates that the rate-controlling step of ‘so-called’high-temperature creep is a kind of restoration process relating intimately todiffusion.

When steady-state creep-rates and/or minimum creep rates are taken asthe parameter of a creep process, a similarity in the temperature dependencebetween creep and diffusion can be observed, at least in (pure) metals andmany simple alloys. Correlation between diffusion and creep behaviour ofpractical alloys under practical creep conditions is not simple, as in thecase of pure metals and solid solution alloys at high temperatures. In theseconditions temperature is usually less than 0.5Tm, significant structurechanges occur in the matrix and the influence of surrounding atmospherebecomes obvious with time. These factors strongly affect the creep behaviourof material and any simple relation between creep and diffusion is difficultto observe.

7.2.3 Influence of short-circuit diffusion

In creep of metals and alloys, the essential rate-determining stage is a kindof restoration process which is governed by diffusion of vacancies and/oralloying elements. At temperature higher than about 0.5Tm, diffusion ofvacancies/atoms through the nearly perfect crystal lattice governs creep processand the temperature dependence of creep rate is similar to that of the latticediffusion of vacancies/atoms.

At temperatures lower than about 0.5Tm, however, the effects of short-circuit diffusion (see Section 7.3.4) become obvious in many cases. Theabsolute values of short-circuit diffusion coefficients are always larger thanthe lattice diffusion coefficients, but the net effect of these types of diffusionis not always large, because their cross-sections are small. In some cases, achange in the activation energy of creep is observed at temperatures below0.5Tm. This change in the activation energy can be explained not by a changein creep mechanism, but simply by the effect of short-circuit diffusion(Lüthy et al., 1980).

7.3 Diffusion characteristics

7.3.1 Moving species: vacancies

In most metallic materials, direct exchange of an atom with an adjacent atomhardly ever occurs without vacant lattice sites. In pure metals, therefore,(long-range) atom movement has the same meaning as vacancy travel but inthe opposite direction. The (self-)diffusion coefficient of constituent atoms

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using radio isotopes, D*, is essentially the same as the diffusion coefficientof vacancies.

In (binary) solid solutions consisting of the solute atoms B with the solventatoms A, the measured diffusion coefficients using isotopes, D*B or D*A, ata fixed concentration of B, are again an indication of the movement ofvacancies replacing the lattice site of atom B or atom A. These coefficientsdo not relate to the concentration gradient in a material. This type of diffusioncoefficient in alloys have been measured in limited cases only.

When the site exchange between atom A and vacancies occurs independentlyof the presence of other kinds of atoms (B), diffusion of atom A and atom Boccurs as parallel reactions independent with each other. The resultant vacancydiffusion is simply the sum of vacancies exchanging with atoms A and withatoms B. Diffusion of vacancies under this condition, D , can be expressedby:

D = NA D*A + NB D*B [7.1]

Here, NA and NB are the mole fraction of element A and B, and D*A and D*B

are self-(tracer) diffusion coefficients in the alloy, respectively. Diffusion ofthis type is affected preferentially by the faster atoms. This type of diffusionresults in a segregation of atoms, because of the difference between D*A andD*B .

When the condition does not allow segregation (long-range concentrationfluctuation), then vacancy diffusion must proceed as a series of reactions ofthe movement of atoms A and atoms B, ′D :

′D = D*AD*B/(NAD*B + NBD*A) [7.2]

This equation can be more easily understood from another form of expression:

1/ ′D = NA/D*A + NB/D*B [7.3]

This formulation is analogous to series-combined resistance which is a simplesum of resistance of each component. Diffusion of this type is affectedpreferentially by the diffusion of slower atoms.

The climb motion of an edge dislocation is controlled by vacancy formation(or annihilation) at the elementary jog and the moving away of this vacancyfrom the jog. The diffusion coefficients suitable for use in the analyses arenot D (parallel vacancy diffusion coefficients), but ′D (combined vacancydiffusion coefficients).

7.3.2 Moving species: constituent atoms

In diffusion experiments with alloys, usually two pieces of alloys with differentlevels of the solute (B) concentration are welded and the concentration changeof B with time (homogenization process) is measured. The available diffusion

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coefficients are called interdiffusion coefficients, or chemical diffusioncoefficients, and expressed customarily as D .

In interdiffusion reactions, the concentration profile changes with time;usually a kind of homogenization process proceeds. When diffusion of atomA occurs independently of that of atom B, then the resultant diffusion is asimple sum of movement of A and B (in the inverse direction from eachother under ordinal conditions) and the so-called Kirkendall effect can beseen. Interdiffusion coefficients, D , can be estimated from the equation(Darken equation):

D = NAD*B + NBD*A [7.4]

When a change of concentration profile is not allowed, but the change in theposition only is allowed, then the movement of atoms (solute and solvent) isexpressed by:

˜ ′D = D*AD*B/(NAD*A + NBD*B) [7.5]

This equation can be more easily understood from another form of expression:

1/ ˜ ′D = NA/D*B + NB/D*A [7.6]

This formulation is analogous to series-combined resistance which is a sumof resistance of each component.

In analyses of the glide motion of an edge dislocation, which is controlledby diffusion of the solute atmosphere, the diffusion coefficient suitable foruse in the analyses is not D (interdiffusion, a parallel sum of independentdiffusion of constituent components), but ˜ ′D (combined diffusion coefficientsof constituent components). When a solute atmosphere is being formed arounda fresh dislocation, the diffusion coefficients that should be used in theanalyses are ordinary interdiffusion coefficients.

An example of concentration dependence of several types of diffusioncoefficient is shown schematically in Fig. 7.2 for an A–B binary solid solutionalloy system. Diffusion coefficients suitable for employment in analyses ofcreep phenomena are those under the condition of a no-concentration gradient,that is, ′D and/or ˜ ′D , not D and/or D. Unfortunately most diffusion datareported are D and/or D . Values of ′D and/or ˜ ′D are available in verylimited cases.

7.3.3 Diffusion paths: lattice diffusion

In common metals and alloys, atom movement occurs through interchangeof an atom in a crystal lattice site with a vacant site. Diffusion in nearlyperfect crystals is called lattice diffusion or volume diffusion. Lattice diffusionin metals depends on Tm and the crystal lattice system. Among metals with

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the same crystal lattice system, (self-) diffusion coefficients depend closelyon their Tm value. In metals with a higher Tm, diffusion coefficients aresmaller than those in metals with lower Tm values. In metals with close-packed lattice systems (e.g. face centred cubic (fcc) and hexagonal closepacked (hcp)), diffusion is slower than that in those with less close-packedlattice systems (e.g. body centred cubic (bcc)). In pure iron (see Fig. 7.3)self-diffusion coefficients in the γ-phase are more than two orders of magnitudesmaller than those in the δ- and α-phases.

7.3.4 Diffusion paths: short-circuit diffusion

There are some diffusion paths that go through imperfect crystal lattice sitesand these are termed short-circuit diffusion. Typical examples are thosethrough grain boundaries, Dgb, and over the surface, Ds. Diffusion alongdislocation cores, Dd, is also a typical example of short-circuit diffusion.

Short-circuit diffusion coefficients are larger than lattice diffusioncoefficients even at the melting temperature of the crystal. The differencebetween the lattice diffusion coefficient and other short-circuit diffusioncoefficients increases with decreasing temperature (see Fig. 7.4). The differencecan be many orders of magnitude. The practical effect of short-circuit diffusion,however, is not very large, at least above 0.5Tm, because the effective cross-sections of short-circuit diffusion are usually very small. At lower temperaturesthe effect of short-circuit diffusion becomes obvious in some cases. In these

Mole fraction of B

1.0(B)

0(A)

0.5

DB (self)*

DA (imp)*

DB (imp)*

DA (self)*

Dif

fusi

on

co

effi

cien

t(l

og

. sca

le)

DA*

DB*

D ′~

D~

D ′D

7.2 Schematic diagram of concentration dependence of diffusioncoefficients in an A–B binary solid solution alloy.

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T –1 (10–3 K–1)

1.21.00.80.6

DFe

(m

2 s–1)

*

10–10

10–12

10–14

10–16

10–18

10–20

10–22

1800 1400 1000 800T (K)

αp

αf

γ

δ

7.3 Lattice self-diffusion coefficients in pure iron as a function oftemperature (Oikawa,1982).

T –1 (10–3 K–1)

1.81.61.41.21.00.8

T (K)1100 900 800 700 600

Dv,

Dd a

nd

Dg

b (

m2 s

–1)

10–23

10–22

10–21

10–20

10–19

10–18

10–17

10–16

10–15

10–14

10–13

10–12

10–11

10–10

10–9

10–8

10–7

10–6

Dgb

Dd

DvT α–γ

= 1

184

K

T c =

104

3 K

7.4 Short-circuit diffusion coefficients in high purity α-iron (Iijima etal., 2005).

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cases, effective (total) diffusion coefficients, Deff, is estimated by the followingequation:

Deff = flDl + fgbDgb + fdDd [7.7]

Here, fl , fgb and fd are the cross-sections of the lattice (volume), the grainboundary and the dislocation core diffusion, respectively.

Practical creep conditions of common heat-resistant steels are located inthe temperature range where the effects of short-circuit diffusion cannot beneglected. It is unfortunate that diffusion data in this temperature range arevery limited in number and also the reliability of most existing data is nothigh.

Short-circuit diffusion coefficients are greatly affected by the segregationof minor elements and also the character of grain boundaries. Practical materialsusually contain many minor alloying elements and also some impurities. Thecharacter of grain boundaries in practical materials depends greatly on thehistory of their heat treatments. Hence, actual estimation of the short-circuitdiffusion coefficients is not easy in heat-resistant steels.

7.4 Roles of atom/vacancy movement in creep

7.4.1 Deformation of matrix: dislocation climb controlled

In pure metals belonging to fcc, bcc and hcp lattice systems, dislocations canmove rather easily on their glide planes, but find it difficult to get out ofthese planes. To continue straining under a given stress, dislocations must beeliminated from the glide plane by cross glide in the case of screw dislocationsor climbing out of the glide plane in the case of edge dislocations. Crossgliding can occur with increasing stress level, but in climbing in edgedislocations (motion of jogs) a net change in number of atoms around thedislocation is necessary. To move jogs while keeping mass balance, vacanciesmust be created (atoms must go away) or disappear (atoms must come in) atthe jogs. Therefore, formation and migration of vacancies (atom flow in theinverted direction) are an essential stage in order to continue creep deformation.

In creep of most pure metals, the thermally activated rate-controlling stepis believed to be diffusion of vacancies (Sherby and Weertman, 1979). Insolid solution alloys, usually easiness of dislocation glide becomes less thanthat in pure metals. However, in some alloys, dislocation climbing becomessignificantly more difficult and diffusion of vacancies (jog movement) remainsas the rate-controlling step. This type of behaviour can be seen in somealloys of low stacking-fault energy. A typical example has been reported increep of α-Cu–Al solid solution alloys (Hasegawa et al., 1972). A similarsituation may be seen in some austenitic steels, but not in ferritic steels inwhich the stacking-fault energy is reasonably high.

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7.4.2 Deformation of matrix: dislocation glide controlled

In some solid solution alloys, solute atoms gather around edge dislocationsand make a solute atmosphere. A typical cause of such segregation (clustering)is the elastic strain around edge dislocations and this type of segregation iscalled a Cottrell(-type) atmosphere. Once this type of atmosphere is generated,edge dislocations receive extra resistance to move even on the original glideplane.

Sometimes the glide velocity is retarded to the level of the climbingvelocity. In this case the motion of edge dislocations is controlled by themovement of solute atmosphere being formed around the dislocations.Therefore, in many alloys the creep rate is controlled by the diffusion ofsolute atoms rather than vacancies. This situation can easily be observed inalloys with solutes that have a large size misfit. An example can be seen inα-Fe–Mo alloys (Oikawa et al., 1980). In practical heat-resistant steels,especially in ferritic steels, most alloying elements have this type ofstrengthening effect. Chromium is an exceptional solute, because the sizedifference of Cr (versus α-Fe) is very small.

7.4.3 Influence through microstructure change

During creep deformation of complex alloys such as steels, significant changesin microstructure occur. At the very first stage of creep, a significant changein dislocation configuration occurs, followed by slow changes, not only indislocation structures, but also in the kind, shape, size and amount ofprecipitates.

Diffusion of the solute plays an important role in the formation of a newphase (precipitates). In this case the solute atoms gather to make precipitates.This process is a de-homogenization process and D (parallel process) issuitable for analysing the process. If a strengthening solute element is consumedfrom the matrix to make a different phase as precipitates, the concentrationof the solute causing drag resistance decreases and the material is weakenedby the precipitate formation. When the third (minor) element affects thediffusion of main strengthening alloying elements, this minor element mayact as a decelerator (or accelerator) of the weakening of the material. Atypical example can be seen in the addition of Re in α-Fe–15Cr–5W alloy(Kunieda et al., 2006).

In the early stage of precipitation, the size of the newly precipitated phaseis very small but the number of precipitates is very large. Precipitatestrengthening may occur at this stage, and some of the weakening by soluteconsumption may be cancelled.

In a later stage of precipitation, the size of the precipitates increases andthe number of precipitates decrease. In this precipitate coarsening stage, the

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diffusion of precipitate constituent atoms, D , must be taken into account. Inthis case, the average concentration of the solute is unchanged, but still thestrength of the material decreases because of weakening of the pinning forceof precipitates on dislocations.

7.4.4 Failure

Correlation between creep rupture (time to failure) and diffusion is not wellestablished. Failure is a more large-scale and complicated phenomenon thanthe straining of matrix and many factors other than simple diffusion must betaken into account in the analyses. At the very beginning stage of failure increep of a specimen, the movement of vacancies may have a role in clusteringvacancies and formation of embryos or nuclei of microcracks. The role ofdiffusion, however, may not be significant in the overall weakening of thematrix itself, because creep failure is a complicated phenomenon. Oxidation/cracking, resulting stress concentration and so on, come to play more essentialroles. The diffusion of oxygen into the matrix, especially through grainboundaries, diffusion of oxygen and solute atoms into/within oxide layersand diffusion behaviour in the coating layers are important processes to beconsidered.

7.5 Influence of some factors on creep through

their effects on diffusion

7.5.1 Magnetic transformation

The ferromagnetic state retards diffusion, because of its spin ordering. TheCurie temperature, TC, of ferritic iron is 1043 K so that most practical creepconditions of ferritic steels lie in the ferromagnetic state and diffusion isaffected by this ferromagnetism (see Fig. 7.5).

In pure iron, the effect of the ferromagnetic state on (self-)diffusion hasbeen well studied. The effect can be quantitatively expressed as

Dmag = Dnon-mag × exp( –αQ s 2/RT) [7.8]

Here, s is the spontaneous magnetization relative to the value at 0 K and αa factor for the energy increase in the ferromagnetic state. The value s wasmeasured by Potter (1934) and Crangle and Goodman (1971) for iron and s2

is shown as a function of temperature in Fig. 7.6.The value of α in α-iron has been reported for several elements (see

Section 7.6). The magnitude of the effect depends greatly on the soluteelement and α is well correlated with the change in magnetization of the firstand second nearest neighbours of the diffusing atom in iron (Nitta and Iijima,2005). Cobalt diffusion in α-iron is an exceptional case, in which theferromagnetic effect can be recognized up to about 100 K over TC.

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FeCoCrNbMo

T –1/(10–3 K)1.41.31.21.11.00.90.8

T (K)1173 1073 973 873 773

D (

m2 s

–1)

10–14

10–15

10–16

10–17

10–18

10–19

10–20

10–21

10–22

10–23

Mo

Nb

T c

T α–γ

7.5 Self-diffusion and impurity diffusion coefficients in α-iron aroundTC (Iijima et al., 1988; Nitta and Iijima, 2005).

T (K)11001000900800700600500

s2

1.0

0.8

0.6

0.4

0.2

0.0

7.6 Square of the spontaneous magnetization (relative to the value at0K) of iron as a function of temperature.

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Creep rate is also affected by the ferromagnetic transformation through itseffect on diffusion. An obvious change in the temperature dependence ofcreep rate is observed near TC in α-iron (e.g. Karashima et al., 1972) andmany α-iron base alloys (e.g. Karashima et al., 1968b). The temperaturedependence of creep rate becomes larger in the lower temperature range thanthat in the higher temperature range, that is, the (apparent) activation energyof creep is larger in the temperature range below TC, similar to the case ofdiffusion.

In creep of α-Fe–Co alloys, a discrepancy between the effect offerromagnetism on creep and TC of α-iron has been observed (Karashima etal., 1968b) as in the case of Co diffusion in α-iron. In some alloys such asthe α-Fe–Mo system, the effect of ferromagnetism on creep can be observedup to 50–100 K higher than TC, although no such a discrepancy is reportedin diffusion. A similar discrepancy between diffusion and creep in the effectof ferromagnetic transformation is also reported in γ-Fe–Ni high alloys(Karashima et al., 1968a).

7.5.2 Grain boundaries

It is well known that grain boundaries act as a strengthening factor in lowtemperature deformation, because of their resistance (barriers) to dislocationmotion. In high temperature deformation, in contrast, grain boundaries act assources and sinks of vacancies and also render a path of rapid diffusion foratoms. As a result, grain boundaries act as a weakening factor in hightemperature deformation.

Grain boundary diffusion coefficients depend greatly on the character ofgrain boundaries. They are largest in random boundaries, decreasing with adecrease in the randomness of boundaries, to being smallest in twin boundaries(∑3 coincidence boundary). The effects of grain boundaries on creep dependgreatly on the character of grain boundaries as in the case of grain boundarydiffusion. A significant effect can be recognized in random boundaries, butthe effect becomes smaller as the degree of the randomness decreases. Thesmallest (almost nil) effect is expected in twin boundaries.

In most experiments on grain boundary diffusion, Dgb was derivedfrom penetration a small distance from the surface. Under this condition,diffusion through grain boundaries of relatively less random types can bedetermined. Recently, careful experiments were done on diffusion throughgrain boundaries of more random types. Values of Dgb determined in theseexperiments are significantly higher and the temperature dependence islower than those reported hitherto. An example is shown in Fig. 7.7 forDgb

* in high purity α-iron. Similar results have been reported for grainboundary diffusion of Cr in austenitic stainless steels (Mizouchi et al., 2004).Therefore, a great care must be exercised in experiments on short-circuit

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diffusion to realize the types of short-circuit diffusion actually observed inthe experiments.

Speaking generally, the order of relative values in three kinds of diffusioncoefficients is Dl << Dd < Dgb (see Fig. 7.4 for α-iron). The order of thesediffusion coefficients is not likely to change with temperature, but theirratios depend greatly on temperature, increasing with decreasing temperature.This effect of temperature is significant near TC.

7.5.3 Segregation of minor elements

Care must be exercised because short-circuit diffusion coefficients aresignificantly sensitive to the segregation of impurity atoms and/or minorelements at the dislocation core and/or grain boundaries. Grain boundarydiffusion coefficients depend greatly on the segregation of other elements atgrain boundaries. The segregation usually retards the grain boundary diffusion,although the absolute value of Dgb is still quite high when comparing diffusionin the matrix.

Iijima (2005)Bokstein (1959)Leymonie (1960)Borisov (1964)James (1965)Bernardini (1981)Hänsel (1985)

T α–γ

= 1

184

K

T c =

104

3 K

T –1 (10–3 K–1)

2.01.81.61.41.21.00.8

T (k)1100 900 800 700 600 500

δDg

b (

m3 s

–1)

10–26

10–27

10–25

10–24

10–23

10–22

10–21

10–20

10–19

10–18

10–17

10–16

7.7 Grain boundary diffusion coefficients in α-iron showing the effectof penetration depth (Iijima, 2005).

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A similar situation is expected in the diffusion along dislocation cores, Dd.

In Fig. 7.8 the effect of carbon content in impurity levels on diffusion alongdislocation cores in α-iron is shown as a typical example. The presence of 16ppm C causes a decrease in Dd of an order of magnitude or more.

Practical creep conditions of heat-resistant steels may be located in anintermediate temperature range, where grain boundaries can act as astrengthening factor and also as a weakening factor. Heat-resistant steelscontain many impurity (minor) elements other than the (intended) alloyingelements. Therefore, actual grain boundaries may have quite complicatedeffects on diffusion and so also on creep strength.

7.5.4 Stacking faults

In metals and alloys with a low stacking-fault energy, γsf, a perfect dislocationmay separate into two partial (imperfect) dislocations, and the separationdistance between these two partial dislocations depends inversely on γsf.

Shima et al.(0.5 mass ppm C)

Mehrer and Lübbehuesen(16 mass ppm C)

T –1 (10–4 K–1)1817161514131211109

T (K)1073 973 873 773 673 573

T α–γ

= 1

183

K

T c =

104

3 K

Dv

Dd/(

m2 s

–1)

10–22

10–20

10–18

10–16

10–14

10–12

10–10

10–8

7.8 Dislocation (core) diffusion in α-iron showing the effect of carboncontent (Shima et al., 2002; Mehrer and Lübbehusen, 1989).

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Between these two partial dislocations, a stacking-fault layer is usually formed.In solid solution alloys, the activity coefficient of the solute element in thestacking fault may be different from that of the (perfect) matrix and thesolute element segregates in the stacking fault. When the partial dislocationsmove (on the glide plane), segregated atoms must move with the partialdislocations. Therefore, stacking faults may cause extra resistance to glidingof the (perfect) dislocation. The glide velocity of dissociated dislocationscan thus be retarded.

In some alloy systems, γsf decreases significantly with increasing soluteconcentration. In this type of alloy, the solute segregates easily in the stackingfaults. The segregation of the solute may accelerate the decrease of γsf, whichmay result in more increase in creep strength. Even in a perfect dislocationof a pure screw type, at least one partial dislocation has some degree of edgecomponent, so that the segregation of the solute (Cottrell atmosphere type)occurs and this solute atmosphere may raise resistance to the glide of thedislocation.

The effect, however, of the solute segregation in the stacking fault andaround a partial dislocation (having some edge component) on the glidevelocity is removed by a pronounced retardation of the climb process inthese low γsf materials. Therefore, in low γsf alloys dissociated dislocationsglide viscously with the solute atmosphere, but the rate-determining step isstill the climb motion, similar to that in pure metals.

Some austenitic steels have intermediate/low γsf values and the glide ofdislocations in these steels may be of a viscous type, but the rate-controllingstep is likely to be the climbing (jog movement). Hence, the diffusion thatneeds to be considered in creep of austenitic steels is not that of solute atomsbut of the vacancies.

7.6 Diffusion data in iron and in some

iron-base alloys

7.6.1 Diffusion in α-iron and in bcc alloys

Diffusion parameters for self-diffusion and impurity diffusion in α-iron arelisted in Table 7.1 and the temperature dependence of diffusion coefficientsis shown in Fig. 7.9. References for data in the table are presented in aconcise format for simplicity, that is, the second and subsequent authors areomitted and only the initial page number is cited. Most data were obtainedusing radioactive tracers. In cases where radioisotopes are difficult to use,diffusion data at very dilute concentrations of the solute are listed, which areindicated by (ID).

No influence of the magnetization can be seen in the diffusion of interstitialelements, such as carbon, nitrogen and oxygen. The effect of the magnetization

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is reported for many substitutional elements. For some elements the value ofα in Equation [7.8] is determined and Dimp

* can be estimated for a widerange of temperature. In cases of manganese and phosphorus, diffusionparameters, D0 and Q, are reported as two sets for the paramagnetic (highertemperature) and the ferromagnetic (lower temperature) ranges, respectively.

Tracer diffusion of constituent elements in iron-base bcc alloys are listedin Table 7.2. Most data are obtained in the paramagnetic α-phase. In some

Table 7.1 Self-diffusion and impurity diffusion in α-iron

Diffusant Temperature (K) D0 (m2 s–1) Q (kJ mol–1) Reference

C 238–1168 2.0 × 10–6 83.9 1N 244–1146 1.4 × 10–6 79.1 1O 1023–1123 3.8 × 10–7 92.1 2Al (ID) 1173–1373 5.2 × 10–4 246 37Be 1073–1773 1.7 × 10–3 228 457Co 859–1173 2.8 × 10–4 251 (α = 0.23) 551Cr 885–1174 3.7 × 10–3 267 (α = 0.133) 664Cu 1063–1175 4.2 × 10–4 244 759Fe 766–1148 2.8 × 10–4 251 (α = 0.156) 854Mn (ferro) 973–1033 1.5 × 10–4 234 954Mn (para) 1073–1173 3.5 × 10–5 220 999Mo 833–1163 1.5 × 10–2 283 (α = 0.074) 1095Nb 823–1163 1.4 × 10–1 300 (α = 0.061) 1163Ni 788–1160 4.2 × 10–3 268 (α = 0.060) 1232P (ferro) 932–1017 1.4 × 10 332 1332P (para) 1078–1153 2.9 × 10–2 271 13Si (ID) 1100–1173 1.7 × 10–4 229 1444Ti 948–1174 2.1 × 10–1 293 (α = 0.079) 1548V 1058–1172 1.2 × 10–2 274 16181W 833–1173 1.5 × 10–2 287 (α = 0.086) 17

1. Weller M (1996), Nichtmetalle in Metallen ’96, DGM Informations-gesellschaftmbH, Verlag, Oberursel, Germany

2. Takagi J (1986), Z Metallkd, 77, 6.3. Nishida K (1971), Trans Jpn Inst Metals, 12, 310.4. Grigorev GV (1968), Fiz Metallov Metalloved, 26, 946.5. Iijima Y (1993), Mater Trans JIM, 34, 20.6. Lee, CG (1990), Mater Trans JIM, 31, 255.7. Oikawa H (1983), Tech Reports Tohoku Univ, 48, 27.8. Iijima Y (1988), Acta Metall, 36, 2811.9. Nohara K (1971), Trans Iron Steel Inst Jpn, 11, 1267.

10. Nitta H (2002), Acta Mater, 50, 4117.11. Oono N (2003), Mater Trans, 44, 2078.12. Cermák J (1989), Z Metallkd, 80, 213.13. Matsuyama T (1983), Trans Jpn Inst Metals, 24, 587.14. Bergner D (1989), Defect Diffusion Forum, 66, 1407.15. Klugist P (1995), Phys Stat Sol (a), 148, 413.16. Geise J (1987), Z Metallkd, 78, 291.17. Takemoto S (2006), Phil Mag, 87, 1619.

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alloys diffusion parameters have been determined over a wide range oftemperature. Some are represented by the factor α in Equation [7.8]. Othersare represented by two sets of parameters for the para- and ferro-magneticranges, respectively.

Interdiffusion data in iron-base bcc alloys are listed in Table 7.3. Only afew systematic studies have been reported on interdiffusion and all data arethose in the paramagnetic region.

7.6.2 Diffusion in γ-iron and fcc alloys

Diffusion parameters for self-diffusion and impurity diffusion data in γ-ironare listed in Table 7.4 and the temperature dependence of diffusion coefficientis shown in Fig. 7.10. The mark (ID) has the same meaning as in Table 7.1.In the case of oxygen, diffusion data were obtained from the observation ofmicrostructure changes and are indicated by (RD).

Reliable diffusion data in the γ-phase alloys have been reported for limitedsystems only. Tracer diffusion values for constituent elements in binary andternary alloys are listed in Table 7.5. There are some data for alloys of a Fe–Cr–Ni ternary system, which is the basic system of austenitic heat-resistantsteels.

T –1 (10–4 K–1)1312111098

D (

m2 s

–1)

10–12

10–13

10–14

10–15

10–16

10–17

10–18

10–19

10–20

10–21

10–22

10–23

Fe

Co

W

Ni

NbMo

Cr

P

Mn

Ti

P V

BeSi

AlCu

Mn

7.9 Self-diffusion and impurity(tracer) dffusion coefficients in α-iron.

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Interdiffusion values in fcc binary alloys are listed in Table 7.6. Verylimited data have been reported on interdiffusion in the γ-phase. Concentrationdependences of interdiffusion parameters are reported in Fe–Co, Fe–Mn andFe–Ni systems.

Table 7.2 Tracer diffusion in bcc alloys

Composition Diffusant Tempera- D0 (m2 s–1) Q (kJ mol–1) Reference

(atom%) ture (K)

Fe–6 Al 59Fe 1088–1478 4.2 × 10–5 198 1Fe–10Al (para) 59Fe 1156–1450 2 × 10–6 184 1Fe–10Al (ferro) 59Fe 765–942 6 × 10–6 196 1Fe–6.8Co(para) 60Co 1153–1193 5.7 × 10–9 146 2Fe–6.8Co(ferro) 60Co 903–1073 4.7 × 10–5 187 2Fe–0.87Cr 59Fe 1040–1173 1.2 × 10–4 241 3Fe–1.43Cr 59Fe 1040–1173 2.8 × 10–4 249 3Fe–3.09Cr 59Fe 1040–1173 6.7 × 10–4 256 3Fe–5.05Cr 59Fe 1073–1173 8.5 × 10–5 237 3Fe–13Cr 51Cr 1073–1673 6.4 × 10–5 232 4Fe–16Cr 51Cr 1073–1673 1.9 × 10–5 218 4Fe–19Cr 51Cr 1073–1673 1.8 × 10–5 217 4Fe–9.13Cr(ferro) 59Fe 848–999 9.3 × 10–4 231 5Fe–15.2Cr(ferro) 59Fe 868–950 1.3 × 10–4 227 5Fe–15.2Cr(para) 59Fe 999–1050 2.7 × 10–5 216 5Fe–19.7Cr(ferro) 59Fe 848–919 6.5 × 10–5 217 5Fe–19.7Cr(para) 59Fe 963–1098 1.8 × 10–5 208 5Fe–0.4Mo 59Fe 823–1173 4.6 × 10–4 253(α = 0.153) 6Fe–0.4Mo 99Mo 834–1143 2.0 × 10–2 285(α = 0.090) 6Fe–1.5Mo 59Fe 823–1173 5.7 × 10–4 253(α = 0.151) 6Fe–1.5Mo 99Mo 838–1120 2.7 × 10–2 290(α = 0.076) 6Fe–1.48Si 59Fe 1063–1373 1.0 × 10–4 276 7Fe–1.87Si 59Fe 1063–1373 7.7 × 10–3 276 7Fe–2Ti 59Fe 1173–1473 2.8 × 10–4 276 8Fe–2Ti 59Fe 1273–1673 5.6 × 10–5 216 4Fe–4Ti 59Fe 1273–1673 2.7 × 10–5 205 4Fe–6Ti 59Fe 1273–1673 4.0 × 10–5 209 4Fe–1.8V 59Fe 1173–1773 1.4 × 10–4 237 8Fe–2V 48V 1273–1723 3.9 × 10–4 241 9Fe–5V 48V 1273–1723 3.0 × 10–4 239 9Fe–5.3V 59Fe 1173–1466 1.9 × 10–4 240 8

1. Raghunathan VS (1981), Phil Mag A, 43, 427.2. Hirano K (1970), Trans Jpn Inst Metals, 13, 96.3. Ku c

∨era J (1974), Acta Metall, 22, 135.4. Bowen AW (1970), Met Trans, 1, 2705.5. Ray SP (1968), Acta Metall, 16, 981.6. Nitta H (2006), Acta Mater, 54, 2833.7. Treheux D (1981), Acta Metall, 29, 931.8. Lai DYF (1967), US Atomic Energy Commission Report–50314.7. Bowen AW (1970), Met Trans, 1, 2767.

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7.10 Self-diffusion and impurity(tracer) diffusion coefficients in γ-iron.

T –1/(10–4 K–1)9876

10–17

10–16

10–15

10–14

10–13

10–12

10–11

10–10

10–9

D (

m2 s

–1)

N

CO

P

WBe

TiMo

CuCo

FeNi

NbV

MnCr

Table 7.3 Interdiffusion in bcc alloys

Composition Temperature (K) D0 (m2 s–1) Q (kJ mol–1) Reference(at %)

Fe–9Al 1193–1483 2.7 × 10–5 188 1Fe–17Al 1193–1483 3.6 × 10–7 142 1Fe–3Co 1123–1168 3.0 × 10–1 318 2Fe–5Co 1123–1168 4.1 343 2Fe–2.5Mo 1073–1573 3.6 × 10–4 257 3Fe–4.1Mo 1073–1573 4.2 × 10–4 259 3Fe–5.0Mo 1073–1573 4.0 × 10–4 262 3Fe–0.47Si 1175–1708 7.8 × 10–5 221 4Fe–0.94Si 1175–1708 7.7 × 10–5 220 4Fe–1.87Si 1175–1708 9.6 × 10–5 220 4Fe–2.81Si 1175–1708 1.1 × 10–4 219 4Fe–3.74Si 1175–1708 1.3 × 10–4 220 4

1. Hishinuma A (1968), J Jpn Inst Metals, 32, 516.2. Hirano K (1977), Trans Iron Steel Inst Jpn, 17, 194.3. Nohara K (1977), Tetsu-to-Hagane, 63, 926.4. Borg RJ (1970), J Appl Phys, 41, 5193.

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7.7 Concluding remarks

7.7.1 Practical application of diffusion data

When the thermally activated step controlling creep phenomena is the climbprocess as in the case of pure metals, vacancy diffusion coefficients shouldbe used in (theoretical) analyses. This will also be the case in some solidsolution alloys of low/intermediate γsf , where dislocations glide viscouslybut the climb velocity is affected more significantly with the addition ofsolute elements, because of a wide stacking fault which is not easy to climb.In heat-resistant steels with an austenite base, the diffusion of vacanciesshould be noted.

Table 7.4 Self-diffusion and impurity diffusion in γ-iron

Diffusant Temperature (K) D0 (m2 s–1) Q (kJ mol–1) Reference

C (ID) 1123–1578 2.3 × 10–5 148 1N 1173–1623 9.1 × 10–5 169 2O (RD) 1223–1373 1.3 × 10–4 166 37Be 1373–1623 3.3 × 10–5 256 460Co 1411–1613 1.3 × 10–4 305 551Cr 1173–1618 1.7 × 10–5 264 664Cu 1378–1641 4.3 × 10–5 280 759Fe 1444–1634 4.9 × 10–5 284 854Mn 1193–1553 1.6 × 10–5 262 9Mo (ID) 1323–1633 3.6 × 10–6 240 1095Nb 1210–1604 8.3 × 10–5 267 1163Ni 1426–1560 1.1 × 10–4 297 1232P 1553–1623 2.8 × 10–3 292 13Ti (ID) 1348–1498 1.5 × 10–5 251 1448V 1210–1607 6.2 × 10–5 274 15W (ID) 1258–1578 5.1 × 10–5 272 16

1. Agren J (1986), Scr Metall, 20, 1507.2. Grieveson P (1964), Trans Met Soc AIME, 230, 407.3. Takagi J (1986), Met Trans, 17A, 221.4. Grigorev GV (1986), Fiz Met Metalloved, 25, 836.5. Suzuoka T (1961), Trans Jpn Inst Metals, 2, 176.6. Oikawa H (1983), Tech Reports Tohoku Univ, 48, 22.7. Oikawa H (1983), Tech Reports Tohoku Univ, 48, 27.8. Heumann Th (1968), J Phys Chem Solids, 29, 1613.7. Nohara K (1971), Trans Iron Steel Inst Jpn, 11, 1267.

10. Oikawa H (1983), Tech Reports Tohoku Univ, 48, 36.11. Geise J (1985), Z Metallkd, 76, 622.12. Hanatake Y (1978), Trans Jpn Inst Metals, 19, 667.13. Seibel G (1963), Compt Rend Acad Sci, C256, 4661.14. Moll SH (1959), Trans TMS-AIME, 215, 613.15. Geise J (1987), Z Metallkd, 78, 291.16. Ruzickova J (1981), Kov Mater, 19, 3.

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Table 7.5 Tracer diffusion in fcc alloys

Composition (at%) Diffusant Temperature D0 (m2 s–1) Q (kJ Reference

(K) mol–1)

Fe–6.8Co 60Co 1283–1583 1.1 × 10–3 326 1Fe–2Cr 51Cr 1073–1673 3.2 × 10–4 245 2Fe–6Cr 51Cr 1073–1673 1.2 × 10–4 237 2Fe–9.13Cr 59Fe 1173–1313 1.2 × 10–5 237 3Fe–1.04Mn 59Fe 1263–1513 9.0 × 10–6 265 4Fe–1.04Mn 54Mn 983–1573 5.5 × 10–6 250 5Fe–2.03Mn 59Fe 1263–1513 1.1 × 10–5 263 4Fe–2.03Mn 54Mn 983–1573 2.0 × 10–6 235 5Fe–2.97Mn 59Fe 1263–1513 5.8 × 10–6 256 4Fe–2.97Mn 54Mn 983–1573 9.6 × 10–7 222 5Fe–4.90Mn 59Fe 1263–1513 6.6 × 10–6 255 4Fe–4.90Mn 54Mn 983–1573 1.7 × 10–6 229 5Fe–14.9Ni 59Fe 1258–1578 2.1 × 10–4 286 6Fe–14.9Ni 63Ni 1258–1578 1.9 × 10–4 289 6Fe–29.7Ni 59Fe 1258–1578 1.0 × 10–3 306 6Fe–29.7Ni 63Ni 1258–1578 2.4 × 10–4 292 6Fe–15Cr–20Ni 59Fe 1236–1673 5.3 × 10–4 308 7Fe–15Cr–20Ni 51Cr 1236–1673 8.3 × 10–4 309 7Fe–15Cr–20Ni 57Ni 1236 –1673 1.5 × 10–4 300 7Fe–15Cr–20Ni–1.4Si 59Fe 1236–1673 5.1 × 10–4 303 7Fe–15Cr–20Ni–1.4Si 51Cr 1236–1673 7.1 × 10–4 303 7Fe–15Cr–20Ni–1.4Si 57Ni 1236–1673 4.8 × 10–4 310 7Fe–15Cr–45Ni 59Fe 1236–1673 2.1 × 10–4 288 7Fe–15Cr–45Ni 51Cr 1236–1673 4.0 × 10–4 293 7Fe–15Cr–45Ni 57Ni 1236–1673 1.8 × 10–4 293 7Fe–17Cr–12Ni 59Fe 873–1570 3.6 × 10–5 279 8Fe–17Cr–12Ni 51Cr 849–1568 1.3 × 10–5 264 8Fe–18Cr–8Ni 59Fe 1173–1473 5.8 × 10–5 281 9Fe–18Cr–8Ni 51Cr 923–1123 8.0 × 10–6 245 10Fe–20Cr–25Ni/Nb 59Fe 1029–1560 1.7 × 10–4 284 11Fe–20Cr–25Ni/Nb 54Mn 823–1523 2.1 × 10–5 248 12Fe–22Cr–45Ni 59Fe 1236–1673 1.5 × 10–4 286 7Fe–22Cr–45Ni 51Cr 1236–1673 4.1 × 10–4 295 7Fe–22Cr–45Ni 57Ni 1236–1673 1.1 × 10–4 291 7

In alloys of high γsf, viscous gliding of (edge) dislocations controls thecreep straining. In this case, the diffusion coefficients that should be used in(theoretical) analyses of creep are those of the solute elements. In heat-resistant steels with a ferrite base, diffusion of the solute elements should benoted.

It is unfortunate that not all suitable data are available and we have to useour second choice, estimated values in practice. When self-(tracer) diffusioncoefficients of the constituent elements are already available as a function ofsolute concentration ′D of vacancies, ˜ ′D of the solutes can be derived fora given composition. This situation, however, occurs in very limited alloy

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systems. When interdiffusion coefficients D at/around the composition inquestion have already been determined, then this D can be used as a roughestimation of ˜ ′D for the solute. In dilute solid solution alloys, this alternativeusage may give a reasonable value.

Table 7.5 Continued

Composition (at%) Diffusant Temperature D0 (m2 s–1) Q (kJ Reference

(K) mol–1)

SUS316 stainless steel 59Fe 1178–1483 1.2 × 10–6 229 13SUS316 stainless steel 51Cr 1023–1473 6.3 × 10–6 243 14SUS316 stainless steel 54Mn 1023–1473 4.1 × 10–5 260 15SUS316 stainless steel 60Co 1242–1423 1.8 × 10–3 296 16SUS316 stainless steel 95Zr 1200–1490 1.0 × 10–3 184 16SUS316 stainless steel 99Mo 1178–1497 1.7 × 10–4 143 16

1. Hirano K (1972), Trans Jpn Inst Metals, 13, 96.2. Bowen AW (1970), Met Trans, 1, 2705.3. Ray SP (1970), Trans Ind Inst Metals, 23, 77.4. Nohara K (1973), J Jpn Inst Metals, 37, 51.5. Nohara K (1971), Trans Iron Steel Inst Jpn, 11, 1267.6. Milliom B (1981), Mater Sci Eng, 50, 43.7. Rothman SJ (1980), J Phys F, 10, 383.8. Perkins RA (1973), Met Trans, 4, 2535.9. Linnenbom V (1955), J Appl Phys, 26, 932.

10. Stawstorm C (1969), J Iron Steel Inst, 207, 77.11. Smith AF (1969), Met Sci, 3, 93.12. Smith AF (1975), Z Metallkd, 66, 692.13. Patil RV (1982), Met Sci, 16, 387.14. Smith AF (1975), Met Sci, 9, 375.15. Smith AF (1975), Met Sci, 9, 181.16. Patil RV (1980), Met Sci, 14, 525.

Table 7.6 Interdiffusion in fcc alloys

Composition Temperature (K) D0 (m2 s–1) Q (kJ mol–1) Reference(atom%)

Fe–5Co 1323–1573 4.5 × 10–3 329 1Fe–10Co 1323–1573 2.1 × 10–3 319 1Fe–15Co 1323–1573 1.1 × 10–3 314 1Fe–3Cr 1173–1473 1.2 × 10–7 219 2Fe–5Mn 1123–1573 7.2 × 10–6 393 3Fe–10Mn 1123–1573 1.8 × 10–6 264 3Fe–15Mn 1123–1573 3.0 × 10–5 269 3Fe–10Ni 1273–1561 5.3 × 10–4 318 4Fe–20Ni 1273–1561 8.9 × 10–4 317 4

1. Hirano K (1977), Trans Iron Steel Inst Jpn, 17, 194.2. Davin A (1963), Rev Metall, 60, 275.3. Nohara K (1973), J Jpn Inst Metals, 37, 51.4. Goldstein JI (1965), Trans Met Soc AIME, 233, 691.

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Without having the self-(tracer) diffusion coefficients of the solute andsolvent at the composition in question, it is not easy to estimate diffusioncoefficients of vacancies. In many cases only D of the impurity, Dimp, andself-(tracer) diffusion coefficients, D*, in the pure (base) metal are knownexperimentally, but no other diffusion data. In these cases we inevitably useD* and Dimp as rough alternatives for ′D and ˜ ′D in the alloy.

7.7.2 Diffusion data searching

Diffusion data up to the end of the 1980s were well compiled in the chaptersof a Landort–Börnstein book (Mehrer H (ed), 1990). It contains data forimpurity diffusion (LeClaire A D and Neumann G 1990), self-diffusion inbinary alloys (Bakker, 1990), chemical diffusion in binary alloys (Murchand Bruff 1990), diffusion in ternary alloys (Dayananda, 1990) and boundarydiffusion (Kaur and Gust, 1990). Diffusion data reported thereafter can befound in the review journal Defect and Diffusion Forum published in Zürich-Uetikon by Trans Tech Publications.

7.8 References

Bakker H (1990), ‘Self-diffusion in homogeneous binary alloys and intermetallic phases’,in Diffusion in Solid Metals and Alloys Mehrer H (ed,), Landort-Börnstein, NewSeries, Group 3, Springer Verlag, Berlin, Volume 26, 213–278.

Crangle J and Goodman G M (1971), ‘The magnetization of pure iron and nickel’, ProcRoy Soc London, Ser A, 321, 477–491.

Dayananda M A (1990), ‘Diffusion in ternary alloys’, in Diffusion in Solid Metals andAlloys, Mehrer H (ed.), Landort-Börnstein, New Series, Group 3, Springer Verlag,Berlin, Volume 26, 372–436.

Hasegawa T, Ikeuchi Y and Karashima S (1972), ‘Internal stress and dislocation structureduring sigmoidal transient creep of a copper-16at.%aluminium alloy’, Met Sci J, 6,78–82.

Iijima Y (2005), ‘Diffusion in high purity iron: influence of magnetic transformation ondiffusion’, J Phase Equil Diff, 26, 466–471.

Iijima Y, Kimura K and Hirano K (1988), ‘Self-diffusion and isotope effect in α-iron’,Acta Metall, 36, 2811–2820.

Iijima Y, Nitta H, Nakamura R, Takasawa K, Inoue A, Takemoto S and Yamazaki Y(2005), ‘Precise measurement of low diffusion coefficients using radioactive tracers’,J Jpn Inst Met, 69, 321–331.

Karashima S, Motomiya T and Oikawa H (1968a), ‘High temperature creep of nickel–iron alloys”, Technol Rept Tohoku Univ, 33, 193–206.

Karashima S, Oikawa H and Watanabe T (1968b), ‘The effect of ferromagnetism uponcreep deformation of alpha-iron and its solid solution alloys’, Trans Metall Soc AIME,242, 1703–1708.

Karashima S, Iikubo T and Oikawa H (1972), ‘On the high-temperature creep behaviourand substructures in alpha-iron single crystals’, Trans Jpn Inst Met, 13, 176–181.

Kaur I and Gust W (1990), ‘Grain and interphase boundary diffusion’, in Diffusion in

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Solid Metals and Alloys, Mehrer H (ed.), Landort-Börnstein, New Series, Group 3,Springer Verlag, Berlin, Volume 26, 630–716.

Kunieda T, Yamashita K, Murata Y, Koyama T and Morinaga M (2006), ‘Effect of Readdition on W diffusivity in Fe–Cr alloys’, Mater Trans, 47, 2106–2108; erratum 47,2888.

LeClaire A D and Neumann G (1990), ‘Diffusion of impurities in solid metallic elements’,in Diffusion in Solid Metals and Alloys, Mehrer H (ed.), Landort-Börnstein, NewSeries, Group 3, Springer Verlag, Berlin, Volume 26, 85–212.

Lüthy H, Miller A K and Sherby O D (1980), ‘The stress and temperature dependencesof steady-state flow at intermediate temperature for pure polycrystalline aluminum’,Acta Metall, 28, 169–178.

Mehrer H (ed) (1990), ‘Diffusion in Solid Metals and Alloys’, Landort-Börnstein, NewSeries, Group 3, Springer Verlag. Berlin, Vol. 26, p. 747.

Mehrer H and Lübbehusen M (1989), ‘Self-diffusion along dislocations and in the latticeof alpha-iron’, Defect Diffus Forum, 66–69, 591–604.

Mizouchi M, Yamazaki Y, Iijima Y and Arioka K (2004), ‘Low temperature grain boundarydiffusion of chromium in SUS316 and 316L stainless steels’, Mater Trans, 45, 2945–2950.

Monma K, Suto H and Oikawa H (1964), ‘Relation between high-temperature creep anddiffusion in alloys’, J Jpn Inst Met, 28, 387–400.

Mukherjee A K, Bird J E and Dorn J E (1969), ‘Experimental correlations for high-temperature creep’, Trans Amer Soc Met, 62, 155–177.

Murch G E and Bruff C M (1990), ‘Chemical diffusion in inhomogeneous binary alloys’,in Diffusion in Solid Metals and Alloys, Mehrer H (ed.), Landort-Börnstein, NewSeries, Group 3, Springer Verlag, Berlin, Volume 26, 279–371.

Nitta H and Iijima Y (2005), ‘Influence of magnetization change on solute diffusion iniron’, Philos Mag Lett, 85, 543–548.

Oikawa H (1982), ‘Lattice self-diffusion in solid iron: a critical review’, Technol ReptTohoku Univ, 47, 67–77.

Oikawa H, Saeki M and Karashima S (1980), ‘Steady-state creep of Fe-4.1 at%Mo alloyat high temperatures’, Trans Jpn Inst Met, 21, 309–318.

Potter H (1934), ‘The magneto-caloric effect and other magnetic phenomena in iron’,Proc Roy Soc London, Ser A, 146, 362–387.

Sherby O D and Burke P M (1967), ‘Mechanical behaviour of crystalline solids at elevatedtemperature’, Progress in Materials Science, 13, 325.

Sherby O D and Weertman J (1979), ‘Diffusion-controlled dislocation creep: A defense’,Acta Metall, 27, 387–400.

Sherby O D, Orr R L and Dorn J E (1954), ‘Creep correlation of metals at elevatedtemperatures’, Trans Amer Inst Min Metall Engr, 200, 71–80.

Shima Y, Ishikawa Y, Nitta H, Yamazaki Y, Miura K, Isshiki M and Iijima Y (2002), ‘Self-diffusion along dislocations in ultra high purity iron’, Mater Trans JIM, 43, 173–177.

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265

8.1 Introduction

This chapter deals with the phenomenology of creep deformation. Mechanicalproperties such as yield stress and ultimate tensile strength are determinedby tensile tests under constant strain rate. Deformation modes are differentbelow (elastic) and above (plastic) the yield stress. On the other hand, creeptests under constant stress (or load) provide the creep properties necessaryfor structural materials for high temperature use. Creep deformation ratevaries during a test from a high speed just after loading to a low speedaround the minimum creep rate. It should be understood what is going on inthe course of creep deformation. This chapter provides suggestions on howto understand creep behavior of engineering materials based on the stress–strain response of materials. Special attention is also paid to stress-acceleratedcreep tests for evaluating long term creep behavior. The athermal yield stressis a key point for appropriate evaluation of creep properties under serviceconditions. The athermal yield stress is the critical stress above whichinstantaneous plastic deformation takes place during loading. The differencein creep deformation behavior above and below the athermal yield stresswill be examined.

Creep tests are carried out over wide ranges of stress and temperature, andcreep deformation rate varies widely from 10–8 to 10 h–1. Several creepdeformation modes appear depending on the creep test conditions. Thedeformation mechanism map is another subject of this chapter. It can provideuseful information, for example which creep mode is operative under agiven creep condition.

8.2 Stress–strain response of materials

Suppose a specimen is subjected to a tensile test at a strain rate, ε , at roomtemperature. The specimen first deforms elastically and stress, σ, increaseslinearly with increasing strain, ε, as shown in Fig. 8.1. Then plastic deformation

8Fundamental aspects of creep deformation

and deformation mechanism map

K. M A R U YA M A , Tohoku University, Japan

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starts at the yield stress, σy, and the stress–strain response deviates from theelastic line. Strain hardening proceeds in the course of plastic deformationand the flow stress, σ, increases with increasing strain usually due to increasein dislocation density and formation of dislocation substructures.

In scientific considerations the true stress σΤ should be used. The truestress is defined by:

σT = L/A [8.1]

where L is the applied load and A is the current cross-section area at strain ofε. Engineering stress, σE, defined by the following equation is often morepractical and we often use it in engineering expressions:

σE = L/A0 = σT/(1 + ε) [8.2]

where A0 is the initial cross-section area and ε in this equation is the engineeringstrain (elongation/initial length of specimen). As Fig. 8.1 shows, the engineeringstress is lower than the true stress in tensile tests and their difference increaseswith increasing strain. The engineering stress–strain curve has a maximumcaused by the reduction of cross-section area and the subsequent necking,despite the absence of such a peak on the true stress–strain curve. Themaximum stress is called ultimate tensile strength (UTS). In the ASMEcode, allowable stress is given by two-thirds of yield stress when the two-thirds of yield stress is lower than one-quarter of UTS, but given by one-quarter of UTS when the one-quarter of UTS is lower. UTS together withyield stress play an important role when determining the allowable stress. Inthe case of brittle materials, UTS corresponds to fracture stress. However,

Strain, ε

Str

ess,

σ

σu

σyPlastic

Elastic

True stress

Engineeringstress

8.1 Difference between a true stress–strain curve and an engineeringstress–strain curve at room temperature.

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Fundamental aspects of creep deformation 267

the UTS of a ductile material gives the maximum load attainable during atensile test and may not have sufficient scientific importance.

Stress–strain curves tested at elevated temperatures show a slightly differentshape from those tested at room temperature, at which diffusion of atoms canbe neglected. At elevated temperature dislocations can climb and annihilatethemselves with other dislocations. Therefore, after a sufficient amount ofdeformation, dislocation density reaches a stationary value determined bythe dynamic balance between multiplication and annihilation of dislocations.The steady state dislocation substructure results in a steady state flow undera constant true stress as depicted in Fig. 8.2. In the case of pure metals andsolid solution alloys with a low density of initial dislocations, dislocationdensity increases to a steady state value. The increase results in the truestress–strain curve of strain hardening type drawn in Fig. 8.2. High Cr ferriticsteels often contain a high density of dislocations introduced during martensitictransformation. Plastic deformation assists in annihilation of dislocations insuch steels. The recovery of dislocation substructures results in a true stress–strain curve of a strain softening type. We sometimes use the steady stateflow stress–strain rate relation as a substitute for the steady state creep rate–stress relation in creep.

8.3 Temperature and strain rate dependence of

yield stress

The yield stress of steel increases with decreasing temperature at roomtemperature and below (see Fig. 8.3). In plastic deformation at low temperature

Strain, ε

σy

σy

Tru

e st

ress

, σ

Ela

stic

lin

e

Strainsoftening

Strainhardening

Steadystate

8.2 Strain hardening type and strain softening type stress–straincurves at elevated temperature.

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the Peierls barrier (a short range obstacle) is the main obstacle to dislocationmotion. Dislocations overcome the obstacle with the assistance of thermalvibration of atoms and applied stress. Since the thermal energy decreaseswith decreasing temperature, the yield stress increases at low temperature.

At intermediate temperatures ranging from 100–450°C for steel, the thermalvibration energy becomes large enough for dislocations to overcome thePeierls barrier and the main obstacle to dislocation motion changes to anotherlong range obstacle, such as other dislocations and particles. Since theseobstacles are too large for dislocations to overcome by thermal energy,dislocations can overcome the obstacles with the aid of applied stress only.Therefore, the yield stress at intermediate temperature is essentially independentof temperature and strain rate. Such a yield stress is called the athermal yieldstress σa. The yield stress is related to dislocation density ρ and interparticlespacing λ by the following equations:

σa = α M G b ρ [8.3]

σa = β M G b/λ [8.4]

where α is a constant of about 0.4, M is the Taylor factor (= 3), G is the shearmodulus, b is the length of Burgers vector and β is a constant of about 0.8.The yield stresses show weak dependence on testing temperature owing tothe temperature dependence of G. They are truly independent of temperaturewhen normalized by an elastic constant such as G or Young’s modulus E.Dynamic strain aging caused by carbon and nitrogen atoms may introduce apeak on the yield stress–temperature curve at intermediate temperature, butthe peak is insignificant when 0.02 or 0.2% proof stress is used for the yieldstress.

Temperature, T0 K

High ε.

Low ε.

Athermalyield stress

No

rmal

ized

yie

ld s

tres

s, σ

y/E

σa

8.3 Temperature dependence of yield stress σy normalized withYoung’s modulus E, and effects of strain rate on the curve. σa is theathermal yield stress independent of temperature and strain rate.

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At elevated temperature, diffusion of atoms and vacancies assists dislocationsto pass through the athermal obstacles, resulting in another temperature andstrain rate dependent plastic deformation, namely creep. This temperaturerange is called the creep regime. Yield stress at high temperature is expressedas:

σy = E ( / )01/˙ ˙ε ε n exp(QD/nRT) [8.5]

where ε is the strain rate, ε 0 is a material constant, n is the stress exponentwhich is usually greater than 3, QD is the activation energy for lattice diffusion,R is the universal gas constant and T is the absolute temperature. Therefore,the yield stress of high temperature deformation decreases with increasingtemperature or decreasing strain rate.

In low temperature deformation, thermal energy assists dislocations toovercome the Peierls barrier. The assistance of thermal energy increaseswith decreasing strain rate ε , resulting in a decrease in yield stress. However,σa is the lower limit of yield stress at low temperature, since applied stressbelow σa does not allow dislocations to pass through the long range obstacles.At high temperature, yield stress increases with increasing strain rate, asexpected from Equation [8.5]. However the yield stress cannot exceed theathermal yield stress σa, since above σa dislocations can pass through theathermal obstacles without the aid of diffusion. The athermal yield stress isan important value even in high temperature deformation, since it is theupper limit of yield stress.

8.4 Deformation upon loading of creep test

During loading of creep tests at elevated temperature specimens deform at ahigh strain rate, suggesting athermal deformation during the loading. Toexamine this expectation the amount of deformation εi upon loading of 2.25Cr–1Mo steel is plotted in Fig. 8.41 as a function of creep stress σ normalized byYoung’s modulus E. The steel was normalized at 930°C for 20 min and thentempered at 720°C for 2 h. The εi versus σ/E relationships tested at varioustemperatures overlap with each other, confirming the athermal nature ofthe deformation during the loading. The εi versus σ/E curve resembles atypical stress–strain curve of tensile test. Linear response, namely elasticdeformation occurs at low stress. Above a critical stress (σ/E = 6–7 × 10–4)the εi versus σ/E curve deviates upwards from the elastic line. One can easilyobtain 0.02% proof stress (micro yield stress) and 0.2% proof stress (macroyield stress) from the εi versus σ/E curve. The proof stresses normalized byE are independent of testing temperature, confirming the prediction of Fig.8.3. The micro yield stress is two-thirds of the macro yield stress in thisexample. The micro yield stress is used as the athermal yield stress σa in thischapter.

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8.5 Creep behavior below and above

athermal yield stress

Suppose a creep stress of σ1 (σ1 < σa) is applied to a specimen, then thespecimen first deforms elastically to the strain corresponding to point A asshown in Fig. 8.5 (a). The deformation stops at this point at room temperature.Atoms and vacancies can move a sufficient distance above 40% of the meltingpoint of the material; 450°C in the case of iron. At these elevated temperatures,dislocations can climb over particles on their slip planes with the aid ofdiffusion. Recovery of dislocation substructure proceeds during hightemperature exposure, resulting in a decrease in σa given by Equation [8.3].Diffusional flow of atoms themselves can bring about creep deformation(diffusion creep). Because of these reasons, time-dependent plastic deformation,namely creep, occurs below the athermal yield stress as depicted in Fig. 8.5(b). The strain–time curve is called the creep curve.

Slope of the creep curve (dε/dt) is called creep rate. On the creep curve,creep rate first decreases owing to strain hardening (primary creep) andreaches a stationary value (secondary creep). Degradation of microstructuresand accumulation of creep damage, such as cavities and micro cracks, proceedduring creep, resulting in creep acceleration to rupture. This stage is calledtertiary creep. Since the stationary creep rate in the secondary creep stage isminimum during the whole creep process, the creep rate is often called theminimum creep rate and used as a measure of creep deformation resistance.The minimum (or steady state) creep rate ε m is expressed as:

ε m = ε 0(σ/E)n exp(–Qc/RT) [8.6]

450°C475°C500°C525°C550°C600°C650°C

Normalized stress, σ /10–3 E0.1 0.2 0.4 0.6 1 2 4

2/3Elastic

0.2%

σa

0.02%

Elo

ng

atio

n u

po

n lo

adin

g, ε

i

10–1

10–2

10–3

10–4

8.4 Elongation upon loading in creep test for 2.25Cr–1Mo steel as afunction of stress normalized with Young’s modulus E.

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where ε 0 is a material constant, σ is the creep stress and Qc is the activationenergy for creep. The values of n and Qc are closely related to creep mechanismas will be explained in Section 8.7. Practical heat-resistant steels arestrengthened by various obstacles, such as dislocation substructures, precipitatesand so on. Degradation of these obstacles starts at a very early stage of creep.In such materials the minimum creep rate is attained by dynamic balancebetween strain hardening and microstructural degradation.2,3

At applied stress σ2 higher than σa, the specimen deforms elastically topoint B and then plastically to point C within a short period of time. Thespecimen is strain hardened during the plastic deformation by introducingdislocations. The athermal plastic deformation stops at this point and nofurther deformation continues at room temperature. Since recovery ofdislocation substructure proceeds at elevated temperature, another time-dependent plastic deformation (creep) continues at elevated temperature.Creep behavior above σa is similar to that below σa. However, it should benoted that above σa, dislocations are introduced during fast plastic deformationupon loading and the dislocation substructure at the beginning of creep testsabove σa is different from that below σa.

8.6 Change in creep behavior at athermal yield

stress σa

Let us take a look at how creep characteristics change at σa in a particle-strengthened material. Orowan stress, σOr, of a material with average particlespacing λ is given by Equation [8.4]. As depicted in Fig. 8.6, the particles

Stress, σσ2σaσ1

Time, ttr

Str

ain

, ε

(a) (b)

σ2

σ1

Tertiary

Secondary

Primary

εm.

BA

C

8.5 (a) Athermal stress–strain curve and (b) creep curves at stressesσ1 (below athermal yield stress σa) and σ2 (above σa).

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alone can sustain the applied stress σ below σOr without introducing additionaldislocations. Above σOr dislocations pass through the particles leaving anOrowan loop at each particle. The Orowan loops bring about strain hardeningwithout appreciable increase in dislocation density. This type of strain hardeningoccurs below the macro-yield stress σmy. Above σmy the particles withdislocation loops cannot sustain the applied stress and dislocations travellong distance upon loading, introducing additional dislocations into the material.These drawings explain how the dislocation substructure just after loadingabove σmy is different from that below σOr.

4

Experimental results for an aluminum alloy containing particles (2.2 vol%Al6Mn) are given in Fig. 8.7;4 minimum creep rates ε m are plotted againstcreep stress σ normalized by shear modulus G. The positions of σΟr (0.02%proof stress) and σmy (at 0.08% plastic strain) measured experimentally bytensile test are indicated in the figure. There are three regions with a differentstress exponent n: the value of n is low (n = 9) in the high stress region H (σ> σmy), takes a medium value (n = 14) in the intermediate stress region M(σOr < σ < σmy), and is high (n = 26) in the low stress region L below σOr.

Additional dislocation

Orowan loop

(b) Region M: σmy > σ > σOr

(c) Region H: σ > σmy

Particle

(a) Region L: σOr > σ

8.6 Changes in dislocation substructures in particle-strengthenedmaterial with decreasing creep stress. (a) Above macro yield stressσmy, (b) between σmy and Orowan stress σOr and (c) below σOr.

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The ε m versus σ curves exhibit threshold like behavior typical of particlestrength materials.

Figure 8.84 shows the evolution processes of dislocation substructures inthe alloy during loading and the subsequent creep deformation. In Fig. 8.8

350°C300°C250°C

Normalized stress, σ /10–3 G

0.6 0.8 1 2 3

Min

imu

m c

reep

rat

e (s

–1)

10–9

10–7

10–5

10–3

10–1

σmy

σOr

H

M

L

Al–2.2vol% Al6 Mn

8.7 Minimum creep rate of a particle-strengthened alloy as a functionof creep stress σ normalized with shear modulus G. σOr and σmy areOrowan (micro-yield) stress and macro-yield stress, respectively.

2.0 (H)1.3 (M/H)1.0 (L)

1.0 (M)0.98 (M)0.95 (L)0.81 (L)

σ /10–3 G

σ /10–3 G

(a) 250°C (b) 350°C

Al–2.2vol% Al6Mn

Strain0.20.10.20.10 0

Dis

loca

tio

n d

ensi

ty (

1013

m–2

)

0.8

0.6

1

2

4

6

8

8.8 Changes in dislocation density during creep tests. The solidsymbol is the density before testing. Stress levels H, M and Lcorrespond to the three regions indicated in Fig. 8.7.

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the stress levels L, M and H correspond to the three regions indicated in Fig.8.7. The dislocation density versus strain relationships accord very well withthe prediction from Fig. 8.6. In region L (σ < σOr) the dislocation density iskept close to the original value during the whole creep test, since no additionaldislocation is needed to sustain the applied stress. Bowing out of dislocationsbetween particles can explain the slight increase in dislocation density fromthe value before loading. At σ > σmy (Region H) the presence of particles isnot sufficient to sustain the applied stress and a high density of dislocationsis introduced into the alloy during loading. The recovery of dislocationsproceeds towards a steady state value during the primary creep. A similarrecovery process of dislocation substructures has often been reported increep of pure metals. In region M (σOr < σ < σmy) an obvious increase indislocation density is absent upon loading since the alloy can sustain appliedstress only by introducing Orowan loops around particles. The dislocationdensity increases up to a steady state level during the primary creep, suggestingthat the hardening by Orowan loops is replaced by hardening by dislocationsubstructures.

The stationary densities of dislocations in the Al–Mn alloy are plottedagainst applied stress in Fig. 8.9.4 The boundaries between regions L and Mand between M and H are indicated in the figure. The boundaries L/M shiftslightly among the three testing temperatures. The following well knownrelationship holds between the dislocation density ρ and creep stress σnormalized by shear modulus G:

ρ = (σ/α MGb)2 [8.7]

250°C300°C350°C

Normalized stress, σ /10–3 G

0.80.6 1 2 3

M/H

Al – 2.2vol% Al6Mn

350

300

250

L/M

0.6

0.8

1

2

4

6

8

Sta

tio

nar

y d

islo

cati

on

den

sity

(10

13 m

–2)

8.9 Stationary dislocation density as a function of creep stress σnormalized with shear modulus G. L/M and M/H indicate theboundaries between regions L and M and between regions M and H,respectively.

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The equation is essentially the same as Equation [8.3]. Below the boundaryL/M, on the other hand, the dislocation density is independent of creep stressand kept at a value close to the original dislocation density before creeploading.

As demonstrated in this section creep behavior above macro-yield stressshould be different from that below micro-yield stress. The minimum creeprates for the 2.25Cr–1Mo steel used in Fig. 8.4 are plotted in Fig. 8.10 as afunction of creep stress normalized by Young’s modulus E. 0.02% (σ0.02)and 0.2% (σ0.2) proof stresses are indicated in the figure. The stress exponentdecreases with decreasing stress and then discontinuously jumps up to 15below σ0.02. This threshold like-behavior of the 2.25Cr–1Mo steel resemblesFig. 8.7 which was plotted for a particle-strengthened aluminum alloy.Engineering materials are used below the micro-yield stress and creep behaviorchanges at the micro-yield stress. Their creep properties under service conditionsshould be evaluated from short-term creep data obtained below the micro-yield stress.

8.7 Deformation mechanism maps

A material deforms by several mechanisms at elevated temperatures dependingon its creep conditions, namely stress and temperature. Representatives ofthe mechanisms are diffusion creep controlled by volume diffusion (Nabarro–Herring creep) or grain boundary diffusion (Coble creep) and dislocationcreep controlled by volume diffusion (high temperature power law creep) or

500°C475°C450°C

650°C600°C550°C525°C

2.25Cr–1Mo steel

σ0.02 σ0.2

Normalized stress, σ /10–3 E

0.20.1 0.4 0.6 1 2 4

Min

imu

m c

reep

rat

e (h

–1)

10–8

10–6

10–4

10–2

8.10 Stress dependence of minimum creep rate of normalized andtempered 2.25Cr–1Mo steel.

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by pipe diffusion (low temperature power law creep). The creep rates of allthe mechanisms are represented by the following general equation:

ε = ε 0(σ/G)nd pD [8.8]

where ε 0 is a material constant characteristic of the mechanism and material,d is the grain size, p is the grain size exponent and D is the diffusion coefficientrelevant to the mechanism. The values of n, p and D are typical of each creepmechanism and are listed in Table 8.1. Dl, Dp and Dgb are the diffusioncoefficient of lattice, dislocation pipe and grain boundary diffusion, respectively.The four representative creep mechanisms are independent of each other andthe creep strain produced by each mechanism contributes additively to thetotal creep strain. Therefore, the mechanism giving the highest value of εdominates creep deformation at a given stress and temperature. This assumptionpredicts that a mechanism with a high stress exponent n operates at highstress. The diffusion creep mechanisms take the stress exponent of n = 1 andare operative at low stress, while a low temperature power law creep of n =nd + 2 dominates at high stress, where nd is the stress exponent of hightemperature power law creep and nd = 3 – 5. A high temperature power lowcreep of n = nd is operative at intermediate stress. The Coble creep (controlledby grain boundary diffusion) and low temperature power law creep (bydislocation pipe diffusion) are rate controlled by short circuit diffusion.Activation energies for such short circuit diffusion paths are lower and almosthalf that of lattice diffusion. Therefore, the Coble creep and the low temperaturespower law creep with low activation energy appear at lower temperaturesthan the Nabarro–Herring creep and the high temperature power law creep.

The aforementioned consideration provides the deformation mechanismmap drawn in Fig. 8.11.5 The diffusion creep mechanisms (Coble creep andNabarro–Herring creep) appear in the lowest stress range and the dislocationcreep mechanisms (high temperature and low temperature power law creep)in the intermediate stress range. A dislocation glide mechanism without theaid of diffusion takes over the role of plastic deformation above the athermalyield stress. The critical stress for the onset of the dislocation glide mechanism

Table 8.1 Stress exponent n, grain size exponent p and diffusion coefficient D inEquation [8.8] for each creep mechanism

Deformation mechanism n p D

Dislocation creepLow temperature power law creep 5–7 0 Dp

High temperature power law creep 3–5 0 Dl

Diffusion creepCoble creep 1 3 Dgb

Nabarro–Herring creep 1 2 Dl

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corresponds to the athermal yield stress σa, and is independent of temperatureand strain rate as mentioned in Sections 8.3 and 8.4. Nabarro–Herring creepand high temperatures power law creep operate at higher temperatures owingto their higher activation energy for creep. Creep rates of both dislocationcreep mechanisms are independent of grain size, but those of the diffusioncreep mechanisms increase with decreasing grain size. Therefore, the diffusioncreep fields, especially the Coble creep field, expand, while the dislocationcreep fields shrink with decreasing grain size. Deformation mechanism mapsof various materials are available in a book by Frost and Ashby.5 By usingthis kind of map we can predict a relevant deformation mechanism of amaterial under any creep conditions. Correct evaluation of creep propertiesrequires accelerated creep tests conducted in the same deformation mechanismfield as the service conditions.

Below the athermal yield stress, creep deformation by one of the fourmechanisms starts after elastic deformation upon loading, whereas plasticdeformation by the dislocation glide mechanism proceeds during loadingabove the athermal yield stress. The deformation mechanism map is correctin this sense. However, above the athermal yield stress dislocation creepbegins after the athermal plastic deformations upon loading (see Fig. 8.10).1

The main obstacles for dislocation creep below the athermal yield stress arethe inherent obstacles present in the material, such as solute atoms in solution-treated materials, precipitates in usual engineering steels and dislocationsubstructure in tempered martensite steels. On the other hand, the mainobstacles above the athermal yield stress are dislocation substructuresintroduced into specimens during loading, as demonstrated in Section 8.6.One should keep this fact in mind when using deformation mechanism maps.

8.11 Schematic drawing of deformation mechanism map. Tm is themelting temperature of the material under consideration.

Ideal strength

Dislocation glide

σa

High temperaturepower law creep

Low temperaturepower law creep

Coble creep Nabarro–Herringcreep

Normalized temperature, T/Tm

0.8 10 0.60.40.2

No

rmal

ized

str

ess,

σ/G

10–7

10–5

10–3

10–1

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Plastic deformation at room temperature is essentially time independent andfinishes within a short period of time. At elevated temperatures, thermallyactivated migration of atoms and vacancies, namely diffusion, occursextensively and time-dependent plastic deformation (creep) continues for along period of time until fracture. Several deformation mechanisms, such asdiffusion creep and dislocation creep appear, depending on testing stress andtemperature. Deformation mechanism maps are useful in predicting thedeformation mechanism operative under the creep condition of interest. Inthe case of heat-resistant steel, the stress and temperature ranges of interestare 1 × 10–4E – 3 × 10–3E (E: Young’s modulus) and 0.4Tm – 0.55Tm (Tm:melting temperature). Under such conditions, the stress exponent for creeprate is usually greater than 3, suggesting that dislocation creep is the relevantdeformation mechanism of engineered steels.1

Creep deformation occurs both above and below athermal yield stress.The athermal yield stress is the yield stress of a material at a high strain rateat a elevated temperature, for example during loading of the creep test.Below the athermal yield stress, creep (time-dependent plastic deformation)starts after elastic deformation upon loading. Above the stress, time independentplastic deformation (dislocation glide) occurs upon loading before time-dependent plastic deformation (creep). Since plastic deformation upon loadingalters the major obstacle to creep deformation from the inherent obstacles todislocation substructures, creep deformation behavior is different above andbelow the athermal yield stress. Engineering materials are used below theathermal yield stress. Their creep properties should be evaluated from short-term creep data obtained below their athermal yield stress.

8.9 References

1 Maruyama K, Sawada K, Koike J, Sato H and Yagi K, ‘Examination of deformationmechanism maps in 2.25Cr–1Mo steel by creep tests at strain rates of 10–11s–1 to 10–

6s–1’, Mater Sci Eng, 1997, A224, 166–172.2 Evans R W, Parker J D and Wilshire B, ‘An extrapolation procedure for long-term creep

strain and creep life prediction with special reference to 1/2Cr1/2Mo1/4V ferriticsteels’, in Recent Advances in Creep and Fracture of Engineering Materials andStructures, Wilshire B and Evans R W (eds), Pineridge Press, Swansea, 1982, 135–184.

3 Maruyama K, Tanaka C and Oikawa H, ‘Long-term creep curve prediction based onthe modified θ projection concept’, Trans ASME, J Press Vess Technol, 1990, 112, 92–97.

4 Maruyama K and Nakashima H, Materials Science for High Temperature Strength –Creep Theories and their Application to Engineering Materials, Uchida-Rokakuho,Tokyo, 1997.

5 Frost H J and Ashby M F, Deformation Mechanism Maps, Pergamon Press, Oxford,1982.

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279

9Strengthening mechanisms in steel for

creep and creep rupture

F. A B E, National Institute for Materials Science (NIMS), Japan

9.1 Introduction

Design stress of elevated-temperature components used under creep conditionsis usually determined on the basis of a 100 000 h creep rupture strength at theoperating temperature, and sometimes also a 200000–300 000 h creep rupturestrength. Therefore, creep-resistant steels must be reliable over long periodsexceeding 100000 h at elevated temperature and the characteristic parametersfor the strength of creep-resistant steels are long-term creep and creep rupturestrength. In this chapter, the strengthening mechanisms in creep-resistantsteels are described with emphasis on long-term creep strength, mainly fortempered martensitic 9–12Cr steels.

9.2 Basic ways of strengthening steels at elevated

temperature

The basic ways in which creep-resistant steels can be strengthened are bysolute hardening, precipitation or dispersion hardening, dislocation hardeningand boundary hardening.1–4 These should be helpful in examining the behaviorof engineering creep-resistant steels at elevated temperature. It is possible tocombine several strengthening mechanisms but it is often difficult to quantifyeach contribution to the overall creep strength.

9.2.1 Solid solution hardening

Substitutional solute atoms such as Mo and W, which have much largeratomic sizes than those of solvent iron, have been favored as effective solidsolution strengtheners for both ferritic and austenitic creep-resistant steels. Itshould be noted that the contribution of solid solution hardening by Mo andW to the overall creep strength of engineering creep-resistant steels is practicallysuperimposed on other strengthening mechanisms, for example precipitationhardening. As will be described later, the addition of Mo and W sometimes

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causes precipitation of the Fe2(Mo,W) Laves phase and enhances finedistributions of M23C6 carbides during exposure at elevated temperature.Taking the Hume–Rothery size effect and large solid solubility in iron intoaccount, Ir5 and Re,6 which are located in the lower part of periodic table, arealso promising for effective solid solution hardening. In terms of interstitialsolute atoms, it is well known that nitrogen is beneficial for long-term creepstrength of engineering ferritic and austenitic creep-resistant steels throughsolid solution hardening as well as precipitation hardening by fine nitrides.

Figure 9.1 shows the minimum creep rate of α-Fe and solid-solution α-Fe–Mo–W alloys (FE, MH, MWH, WH) with a ferrite matrix and precipitationhardened α-Fe–C–Mo–W–V–Nb alloys (PO, PSM, PSW) at 600°C.7 Theprecipitation hardened α-Fe–C–Mo–W–V–Nb alloys contained MXcarbonitrides of (Nb, V)(C, N) which were 40 nm in size in the matrix. Theminimum creep rate of solid solution α-Fe–Mo–W alloys is three orders ofmagnitude lower than that of α-Fe, indicating an effective solid solutionhardening by Mo and W. On the other hand, the minimum creep rate ofprecipitation hardened α-Fe–C–Mo–W–V–Nb alloys (PSM, PSW) is onlyone-half lower than that of precipitation hardened α-Fe–C–V–Nb alloys(PO). This indicates that the strengthening mechanisms in precipitationhardened α-Fe–C–Mo–W–V–Nb alloys (PSM, PSW) come mainly fromprecipitation strengthening by fine MX. It should be noted that the contribution

: FE: MH: MWH: WH: PO: PSM: PSW

600°Cα – Fe

α – Fe–Mo–WSolid solution

7

911

α – Fe–Mo–W–V–Nb

20 40 60 100 200Stress (MPa)

Min

imu

m c

reep

rate

(h

–1)

1

10–1

10–2

10–3

10–4

10–5

10–6

α – Fe–V–Nb

9.1 Minimum creep rate of α-Fe and solid-solution α-Fe–Mo–W alloyswith a ferrite matrix and precipitation hardened α-Fe–C–Mo–W–V–Nballoys at 600°C. MH: Fe–0.120Mo (mass%), MWH: Fe–0.34Mo–1.63W,WH: Fe–2.27W, PO: Fe–0.06C–0.2V–0.08Nb, PSM: Fe–0.06C–0.51Mo–0.2V–0.08Nb, PSW: Fe–0.06C–0.15Mo–0.70W–0.2V–0.08Nb. FE is pureiron.

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Strengthening mechanisms in steel for creep and creep rupture 281

of solid solution hardening by Mo and W to the overall creep strength ofprecipitation strengthened α-Fe–C–Mo–W–V–Nb alloys (PSM, PSW) ismuch smaller than that of solid solution Fe–Mo–W alloys (MH, MWH,WH). This suggests that an additive rule for solid solution hardening andprecipitation hardening does not hold. Transmission electron micrograph(TEM) observations show that the creep deformation of the alloys shownin Fig. 9.1 is controlled by the mobility of dislocations within sub-grainsand at sub-grain boundaries.

9.2.2 Precipitation or dispersion hardening

Precipitation or dispersion hardening is one of the important strengtheningmechanisms in creep-resistant steels at elevated temperature. To achieveenough strengthening using this effect, engineering creep-resistant steelsusually contain several kinds of precipitate particles in the matrix and atgrain boundaries: carbonitrides such as M23C6, M6C, M7C3, MX and M2X,where M denotes the metallic elements, C are the carbon atoms and X are thecarbon and nitrogen atoms, intermetallic compounds such as the Fe2(Mo,W)Laves phase, Fe7W6 µ-phase, χ-phase and so on, and a metallic phase suchas Cu. In a special case of oxide dispersion strengthened (ODS) steels, fineparticles of alloy oxides such as Y2O3 are dispersed in the matrix by mechanicalalloying. A dispersion of fine precipitates stabilizes free dislocations in thematrix and sub-grain structure, which enhances dislocation hardening andsub-boundary hardening.

Several mechanisms have been proposed for the threshold stress,corresponding to the stress needed for the dislocation to pass through precipitateparticles, for example, in the Orowan mechanism, local climb mechanism,general climb mechanism and Srolovitz mechanism, see Fig. 9.2.8 The Orowanstress σor is given by:

σor = 0.8MGb/λ [9.1]

where M is the Taylor factor (= 3), G is the shear modulus, b is the Burgersvector and λ is the mean interparticle spacing.3 Typical values of the volumefraction, diameter and spacing of the major particles contained in temperedmartensitic high Cr steels after tempering are listed in Table 9.1, togetherwith the Orowan stress estimated from the values of interparticle spacing.3

The coarsening of fine precipitates of M23C6, MX and Fe2(W, Mo) Lavesphase and the dissolution of fine MX to form massive precipitates of Zphase, which have been observed in 9–12Cr steels during creep, cause anincrease in λ in Equation (9.1) and hence a decrease in Orowan stress overlong periods of time.3,4 The coarsening and dissolution of fine precipitatessometimes takes place preferentially in the vicinity of grain boundaries duringcreep, which promotes the formation of localized weak zone and promotes

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localized creep deformation near grain boundaries.9,10 This results in prematurecreep rupture and is time and temperature dependent.

The strengthening mechanisms caused by a dispersion of oxide particleswere examined for a 13Cr–3W–0.5Ti–0.4Y2O3 ODS steel with ferrite matrixat 650°C, by comparing the threshold stress measured by a stress abruptlyloading test (SAL test) with the calculated Orowan and void-hardening stresses,σor and σV.11 Figure 9.3(a) shows the relationship between the creep stressand strain upon loading for the ODS steel at 650°C, using the time elapsedafter applying the stress as a parameter. The Orowan and void-hardeningstresses are calculated to be 135–192 and 114–163 MPa, respectively, fromthe histogram for size distribution of Y2O3 particles in the steel (Fig. 9.3(b)).These values are also shown in Fig. 9.3(a). The threshold stress, caused by

Dislocation Dislocation

Particle Particle

(a) (b)

1 2 3 1 2 43

Dislocation

Dislocation

x x

(c) (d)

9.2 Schematic drawings of a dislocation passing through particles.(a) Orowan mechanisms, (b) Srolovitz mechanism, (c) general climbmechanism and (d) local climb mechanism.

Table 9.1 Volume fraction, diameter and spacing of each kind of precipitate inhigh Cr ferritic steel, together with Orowan stress estimated from the valuesof interparticle spacing

Particle Volume fraction Diameter Spacing Orowan stressV(%) dp(nm) σp(nm) σor (MPa)

Fe2(W, Mo) 1.5 70 410 95M23C6 2 50 260 150MX 0.2 20 320 120

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the dislocation passing through the oxide particles, just after loading (at t =100 ms) is measured to be 175 MPa as shown by arrow A in Fig. 9.3(a),which agrees with the calculated Orowan stress. As the time elapses to 4 s,the threshold stress decreases to 150 MPa as shown by arrow B, whichagrees with the calculated void-hardening stress. This suggests that theoriginating mechanisms of the threshold stress come from the Srolovitzmechanism in this steel. According to the Srolovitz mechanism,12 when thematrix-particle interface is incoherent, the normal traction of dislocationstress field on the particle surface is relaxed by interface sliding and volumediffusion and the dislocation is attracted to the particle, see Fig. 9.2(b). After

t = 100 mst = 1 st = 2 st = 3 st = 4 s

Str

ain

(×

10–3

)

16

14

12

10

8

6

4

2

00 100 200 300 400

Stress (MPa)

(a)

Elastic line

AB

σvσv

Rel

ativ

e fr

equ

ency

(%

)

50

40

30

20

10

0

Total number: 1243

0 1 2 3 4 5 6 7 8 9 10Radius (nm)

(b)

9.3 (a) Relationship between creep stress and instantaneous strain at650°C and (b) size distribution of Y2O3 particles in a 13Cr–3W–0.5Ti–0.4Y2O3 ODS steel.

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the relaxation is completed, the particles are felt by dislocations as voids andthe threshold stress should be equal to the void-hardening stresses. In Fig.9.3(a), the time of 1–4 s required for the change from the Orowan stress tothe void-hardening stress corresponds for the time necessary to the fullrelaxation. The Srolovitz mechanism was also confirmed for the thresholdstress in high-temperature deformation of Al–1.5Be and Al–0.7Mn alloyscontaining incoherent precipitate particles.13

9.2.3 Dislocation hardening

Dislocation hardening given by:

σρ = 0.5 MGb (ρf)1/2 [9.2]

where ρf is the free dislocation density in the matrix, is an important meansof strengthening steel at ambient temperature. Tempered martensitic 9–12Crsteels usually contain a high density of dislocations even after tempering,usually in the range 1–10 × 1014 m–2 in the matrix.3 The density of dislocationsproduced by martensitic transformation during cooling after austenitizationcan be controlled by changing the tempering temperature. Tempering isusually carried out at low temperatures of 700°C or lower for turbine steelsto ensure enough tensile strength at ambient temperature by the dislocationhardening, whereas it is as high as 750–800°C for boiler applications.

At elevated temperature, on the other hand, cold working enhances softeningby promoting the recovery of excess dislocations and the recrystallization ofdeformed microstructure. A comparison of creep rupture strength of 12Cr–1Mo–1W–VNb steel at 600°C and 650°C between the two tempering treatmentsat 750°C and 800°C shows that the low temperature tempering gives highercreep rupture strength than high temperature tempering for short times belowabout 15 000 h and 6000 h at 600°C and 650°C, respectively.14 However, thestress versus time to rupture curve is crossed over during long peiods of timeby that of the steel subjected to high temperature tempering at 800°C. Thisis because excess dislocations accelerate recovery and recrystallization duringcreep with the aid of stress. The dislocation density after tempering at ahigher temperature of 800°C is too low to promote recovery andrecrystallization during creep. This effectively suppresses a rapid decreasein creep rupture strength. Also in austenitic steels, the creep strength of coldworked materials is typically higher for short times than that of solutionannealed materials but it is reversed for long periods of time.15,16 Theseresults indicate that the dislocation hardening is useful in the creep strengthonly for short times but that it is not useful for long-term creep strength atelevated temperature.

Figure 9.4 shows the effect of cold rolling on the creep rate versus timecurves and creep rate versus strain curves of a tempered martensitic 9Cr–

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1W–0.1C steel at 600°C and 78 MPa.17 V, Nb and nitrogen were not addedto this steel and hence only M23C6 carbides were distributed along grainboundaries and lath boundaries after tempering. The decrease in creep ratewith time and strain in the transient creep region is less pronounced withincreasing cold rolling level. The onset of acceleration creep significantlyshifts to shorter times with increasing cold rolling level. This results in amuch higher minimum creep rate and hence a shorter time to rupture with

Standard QTQT + 20% CRQT + 40% CR

Standard QTQT = 20% CRQT + 40% CR

Cre

ep r

ate

(h–1

)

10–2

10–3

10–4

10–1 100 101 102 103

Time (h)(a)

Cre

ep r

ate

(h–1

)

10–2

10–3

10–4

10–3 10–2 10–1

True strain(b)

9.4 Effect of cold rolling on (a) creep rate versus time and (b) creeprate versus strain curves for a tempered martensitic 9Cr–1W–0.1Csteel at 600°C and 78 MPa.

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increasing cold rolling level. The strain to reach a minimum creep rateincreases with increasing cold rolling level. The cold rolling provides a highdensity of free dislocations, which accelerates the creep deformation at elevatedtemperature. The movement and annihilation of excess dislocations withinthe lath is considered to be the major process in the transient creep region.

In austenitic steels, cold working is sometimes employed after solutionannealing to accelerate the nucleation of fine carbides such as NbC atdislocations, resulting in a homogeneous distribution of a high density offine NbC carbides in the matrix.18 This enhances the precipitation hardening.

9.2.4 Sub-boundary hardening

Tempered martensitic high Cr steels subjected to normalizing and temperingare usually observed to have a lath martensitic microstructure consisting oflath and block with a high density of dislocations and a dispersion of finecarbonitrides along the lath and block boundaries and in the matrix. The lathand block can be regarded as elongated sub-grains. The lath and blockboundaries provide the sub-boundary hardening given by:

σsg = 10Gb/λsg [9.3]

where λsg is the short width of elongated subgrains.3 The subgrain width λsg,corresponding to the width of lath and block, is in the range 0.3–0.5 µm inmartensitic high Cr steels after tempering. Using the values of G = 64 GPaat 650°C, b = 0.25 nm, λsg = 0.3–0.5 µm, we obtain σsg = 530–320 MPa,which are much larger than the Orowan stress in Table 9.1 for Fe2(W, Mo),M23C6 and MX. Therefore, the sub-boundary hardening gives an importantmeans of strengthening the creep strength of tempered martensitic high Crferritic steels.

Many people have considered that M23C6 carbides precipitate preferentiallyalong the lath, block and prior austenite grain boundaries in temperedmartensitic high Cr steels, while MX carbonitrides are distributed in thematrix. However, modern energy-filtered transmission electron microscopy(FE-TEM) studies have revealed that a large number of fine MX carbonitridesare distributed along the lath, block and prior austenite grain boundaries, aswell as in the matrix after tempering,19 as shown in Fig. 9.5. The finedistributions of M23C6 carbides and MX carbonitrides along the lath andblock boundaries stabilize these boundaries and exert a pinning force againstthe migration of lath and block boundaries and the coarsening of lath andblock during creep. This suggests that the strengthening achieved by sub-boundary hardening is further enhanced by a dispersion of fine precipitatesalong sub-grain boundaries.

The coarsening of lath and block with creep strain, which takes placemainly in the tertiary or acceleration creep region20,21 and causes an increase

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in λsg in Equation (9.3), indicates the mobile nature of lath and block boundariesunder stress. It is well known that polygon and sub-grain boundaries freefrom precipitates in pure metals and solid solution alloys are highly mobileunder applied stress.1 The movement of lath and block boundaries can absorbor scavenge excess dislocations inside the lath and block. This correspondsto a dynamic recovery process, resulting in softening. Thus lath and blockboundaries act as sink sites for the recovery of excess dislocations as well asact as barrier sites exerting back stress given by Equation (9.3) for themovement of dislocations inside the lath and block during creep. The role oflath and block boundaries in the strengthening mechanisms for creep andcreep rupture is not yet fully understood.

9.3 Strengthening mechanisms in modern creep-

resistant steels

9.3.1 Bainitic low-Cr steels

The most recent advancement in bainitic low Cr steels is the improvement inthe creep rupture strength of 2.25Cr and 3Cr steels.4 Figure 9.6 shows theextrapolated 100 000 h creep rupture strength for 2.25Cr–1Mo (T22) and2.25Cr–1.6W–VNb (T23), 2.25Cr–1Mo–VTi (T24) and 3Cr–1.5W–0.75Mo–0.25V without Ta (grade A) and with 0.1Ta (grade B) as a function oftemperature, compared with martensitic 9Cr–1Mo–VNb steel (T91).22 Thecreep rupture strength of the grade B steel is higher than T23 for the entiretest temperature range and also higher than T91 up to 615°C. The higheroxidation rate in air during creep testing of the grade B steel, owing to its3%Cr at temperatures exceeding 615°C, as opposed to T91 which contains9%Cr, is the reason for its lower creep rupture strength at high temperature.The grade A steel has higher creep rupture strength than T23 and T91 up to600°C.

The strengthening mechanisms in 2.25Cr–1.6W–VNb steel (T23) are found

Bright field image Cr map V map

500nm

9.5 Bright field image and elemental mapping of Cr and V for P92 inas tempered steel, by energy-filtered TEM.

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to be the combination of solid solution hardening due to W and precipitationhardening due to fine (V, Nb)C carbides in a fully bainitic matrix.23 Figure9.7 shows the creep rupture data for 2.25Cr–1Mo–0.25V–0.05Nb steel (Mo-steel) and 2.25Cr–1.6W–0.25V–0.05Nb steel (W-steel) at 650°C, comparingwith those for a conventional 2.25Cr–1Mo steel. The creep rupture strengthis increased by the addition of V and Nb and also by the substitution of Wfor Mo. M23C6 forms along prior austenite grain boundaries, while MX (M= V and Nb, X = C and N) forms along lath boundaries and in the bainiticmatrix after tempering. M7C3 is occasionally observed along lath boundaries.In the crept specimens, no M23C6 and M7C3 are observed and instead blockyM6C is formed along prior austenite grain boundaries and fine MX remainsalong lath boundaries and in the bainitic matrix. The pronounced lath structureis kept in 2.25Cr–1Mo–0.25V–0.05Nb steel even after long term creep,suggesting the availability of sub-boundary hardening in addition toprecipitation hardening for long periods of time.

The substitution of W for Mo also constrains the precipitation of M6Cduring creep as shown in Fig. 9.8, which constrains the evolution of thebainitic microstructure.23 The precipitation of M6C loosens W and or Mo insolution, which is the key factor for the solid solution strengthening in thissteel. The growth rate of M6C is 10 to 100 times slower in the W steel thanin the Mo steel. It is also found that the MC carbides in Mo steel havealready lost coherency with the matrix after heat treatment, while those in Wsteel have kept coherency even after long-term exposure. This suggests thatW is superior to Mo in terms of precipitation hardening. In 2.25Cr–1Mo–

100

000

h c

reep

ru

ptu

re s

tren

gth

(M

Pa)

350

300

250

200

150

100

50

0520 540 560 580 600 620 640 660

Temperature (°C)

T22 (2.25Cr–1Mo)T23 (2.25Cr–1.6W–0.25V–0.05Nb)T24 (2.25Cr–1Mo–0.25V–0.07Ti)T91 (9Cr–1Mo–0.2V–0.05Nb)Grade A (3Cr–0.75Mo–1.5W–0.25V)Grade B (3Cr–0.75Mo–1.5W–0.25V–0.1Ta)

9.6 100000 h creep rupture strength as a function of temperature ofT23, T24, Grade A and B of 3Cr–1.5W–0.75Mo–0.25V steel, comparedwith T22 and T91.

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VTi (T24), the precipitation of fine TiC carbides is one of the strengtheningmechanisms.

Both 3Cr–3W–0.25V and 3Cr–3W–0.25V–0.1Ta steels are observed tohave an acicular bainitic structure.24 The addition of 0.1Ta to 3Cr–3W–0.25V steel substantially decreases the prior austenite grain size. Fine TaCparticles are precipitated and dispersed within the grains and also along prioraustenite grain boundaries. The carbides along prior austenite grain boundariesin the Ta-containing steel are smaller than those in the steel without Ta. Thisis one of the major reasons for the improvement of creep rupture strength of3Cr–3W–0.25V steel by the addition of 0.1Ta.

Str

ess

(MP

a)100

908070

60

50

40102 103 104

Time to rupture (h)

650°C

2.25Cr–1Mo

2.25Cr–1Mo–0.25V–0.05Nb

2.25Cr–1.6W–0.25V–0.05Nb

9.7 Creep rupture data for 2.25Cr–1Mo–0.25V–0.05Nb and 2.25Cr–1.6W–0.25V–0.05Nb steels at 650°C, compared with those forconventional 2.25Cr–1Mo steel.

Frac

tio

n o

f M

o p

reci

pit

atio

n (

–)

1

0.8

0.6

0.4

0.2

01 10 100 1000 104 105

Aging time (h)(a)

823K

873K923K

Frac

tio

n o

f M

W p

reci

pit

atio

n (

–)

1

0.8

0.6

0.4

0.2

01 10 100 1000 104 105

Aging time (h)(b)

823K

873K923K

9.8 Concentration of Mo and W in extracted residues of (a) 2.25Cr–1Mo–0.25V–0.05Nb and (b) 2.25Cr–1.6W–0.25V–0.05Nb steels,showing the precipitation of Mo and W as M6C. The curves representJohnson–Mehl–Avrami precipitation curves.

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9.3.2 Tempered martensitic 9–12Cr steels

Figure 9.9 shows the creep rupture data for high strength martensitic 9–12Crsteels at 650°C.4 In this figure, two 9Cr–3W–3Co–0.2V–0.05Nb steels with0.05N–0.002C (MARN: Martensitic 9Cr steel strengthened by MX nitrides)25,26

and with 0.08C–0.008N–0.014B (MARBN: Martensitic 9Cr steel strengthenedby boron and MX nitrides)10,27 were developed by the National Institute forMaterials Science (NIMS), 12Cr–2.6W–2.5Co–0.5Ni–0.2V–0.05Ni steel(NF12)28 and 12Cr–3W–3Co–0.2V–0.05Nb–0.1Ta–0.1Nd–0.05N steel(SAVE12)29 were developed by Japanese steelmaking companies as upgradeversions of P92 and P122, respectively, and oxide dispersion strengthened(ODS) 9Cr steel, 0.13C–9Cr–2W–0.2Ti-0.35Y2O3, with tempered martensiticmicrostructure30 was developed for fast breeder reactor cladding materials.

The strengthening mechanisms in MARN steel come mainly fromprecipitation hardening by fine MX nitrides, which enhances the sub-boundaryhardening. In order to achieve a dispersion of fine MX nitrides alone, whichare thermally stable particles for prolonged periods of exposure at elevated

9Cr–3W–3CO–VNb–0.05N–0.002C (MARN)9Cr–3W–3Co–VNb–0.08C–0.008N–0.014B (MARBN)P92 (9Cr–0.5Mo–1.8W–VNb)T91 (9Cr–1Mo–VNb)ODS–9Cr (Ukai et al)NF 12 (12Cr–2.6W2.5Co–NVNb)SAVE 12 (12Cr–3W–3Co–VNbTaNdN)

Str

ess

(MP

a)

300

200

100

80

60101 102 103 104 105

Time to rupture (h)

9.9 Creep rupture data for 9Cr steels of MARN, MARBN, P92, T91,ODS-9Cr and 12Cr steels of NF12 and SAVE12 at 650°C.

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temperatures, it is crucial for 9Cr steels to reduce carbon content to very lowamounts of less than 50 ppm, because the addition of carbon to a 9Cr steelcauses the formation of a large amount of M23C6 carbides rich in Cr. Thetime to rupture significantly increases in the lower carbon region below0.02%, where fine MX nitrides are dominant, Fig. 9.10. A large number offine precipitates which have a size less than 10 nm are distributed not onlyin the matrix but also along lath, block, packet and prior austenite grainboundaries in the 0.002C steel after tempering, stabilizing the lath–blockstructure for long periods of time. The major component of the MX nitrideswas identified as vanadium nitride.

The strengthening mechanisms in MARBN steel come mainly from thecombination of the boron effect and of precipitation hardening by fine MXnitrides, which enhances the sub-boundary hardening. The nitrogen contentin the MARBN steel (0.008% N) is much lower than that in the MARN steel(0.05% N), because excess addition of boron and nitrogen causes the formationof massive boron nitrides during normalizing at high temperature.31 Theformation of boron nitrides offsets the benefit given by boron and nitrogen.The carbon content for MARBN is increased in comparison with that forMARN. Addition of boron reduces the rate of Ostwald ripening of M23C6

Nu

mb

er o

f m

ole

s

0.04

0.02

0

M23C6

MX

0.00 0.05 0.10 0.15 0.20(a)

9Cr–3W–3Co–0.2V–0.05Nb–0.05N

0 0.05 0.1 0.15 0.2Carbon concentration (wt%)

(b)

12 000

8000

4000

0

Tim

e to

ru

ptu

re (

h)

9.10 (a) Amount of M23C6 and MX after tempering and (b) time torupture of 9Cr–3W–3Co–VNbN steel at 650°C and 140 MPa, as afunction of carbon concentration.

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carbides in the vicinity of prior austenite grain boundaries during creep at650°C. The M23C6 carbides are mainly distributed along prior austenitegrain boundaries and along lath, block and packet boundaries inside thegrain. Therefore, the fine distribution of M23C6 carbides and fine lath-blockstructure is maintained in the vicinity of prior austenite grain boundaries forlong periods of time in the steel with about 100 ppm boron during creep. Afine distribution of M23C6 carbides is also observed in the steel withoutboron after tempering, but extensive coarsening of M23C6 and lath-blocktakes place during high-temperature exposure. With increasing boron content,the transient or primary creep region continues for longer periods of timeand hence the onset of acceleration or tertiary creep retards, which results ina lower minimum creep rate and longer rupture time, as shown in Figure9.11. The effect of boron relates to the stabilization of the fine lath-blockstructure in the vicinity of prior austenite grain boundaries for long periodsof time through the stabilization of fine M23C6 carbides.

Most of early studies on ODS steels concentrated on ODS ferritic steelswith a ferrite matrix.30 The main problem in ODS ferritic steels that keptthem from being used was the anisotropy in creep strength between thecircumferential hoop and the longitudinal direction in tubes. One improvementin this anisotropy has been achieved by using martensitic transformationwhich produces an equi-axed grain structure. The ODS martensitic steelexhibits lower creep strength than the ODS steel with a ferrite matrix in thelongitudinal direction but without anisotropy. TEM observations show finedistributions of Y–Ti oxide particles with a size of about 3 nm in the matrixof tempered martensite, together with Ti oxides of hundreds of nanometersand M23C6 carbides of several hundred nanometers in size. Oxide particlesare introduced by complicated mechanical alloying.

New attempts have also been demonstrated in using fine precipitates ofFePd–L10 intermetallic compound in tempered martensitic 9Cr steel32,33 andfine intermetallic compounds in 15Cr steel with a ferrite matrix.34

9.3.3 Austenitic steels

Austenitic steels are usually solid solution hardened prior to service and arecapable of significant precipitation hardening by fine carbonitrides andsometimes by fine intermetallic compounds during service at elevatedtemperature. The creep strength of austenitic steels for super heater boilertubes has been enhanced from conventional 18Cr–8Ni steel to 20Cr–25Nisteel and then to high Cr and high Ni steels combined with the addition of W,which can be used for a long time at a steam temperature of 700°C. Forexample, the creep rupture strength at 700°C and 100 000 h is estimated tobe 86 MPa and 90 MPa for 20Cr–25Ni–1.5Mo–NbTiN steel (KA-SUS310J2TB)35 and 23Cr–43Ni–6W–NbTiB steel (HR6W),36,37 respectively,

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which are higher than those for high-strength stainless steels such as Super304Hand TP347HFG. A cast 19Cr-12.5Ni–4Mn–MoNbN steel (CF8C-Plus),18

recently developed by Oak Ridge National Laboratory, shows a creep rupturestrength similar to KA-SUS310J2TB.

The strengthening mechanisms in 23Cr–43Ni–6W–NbTiB steel (HR6W)come from the combination of solid solution hardening by W and precipitationhardening by fine M23C6 carbides, fine MX carbonitrides and fine Fe2WLaves phase.36,37 Figure 9.12 shows creep rupture data for 0.08C–23Cr–43Ni–(5–7)W–0.2Nb–0.1Ti–0.003B steel (W steel) at 700, 750 and 800°C,compared with data for 0.08C–23Cr–43Ni–(3–5)Mo–0.2Nb–0.1Ti–0.003B

0 ppmB48 ppmB92 ppmB139ppmB

Cre

ep r

ate

(h–1

)

10–2

10–3

10–4

10–5

10–6

10–7

10–8

10–1 100 101 102 103 104 105

Time (h)

log(time)

log

(cre

ep r

ate)

Basesteel

Stabilization ofM23C6 and lath near

grain boundary Decrease inminimumcreep rate

Boron-steel

Increase increep life

Onset of accelerationcreep by microstructure

recovery near G.B.

9.11 Effect of boron on creep rate versus time curves of 9Cr–3W–3Co–VNb steel at 650°C and 80 MPa and schematics of a creepdeformation mechanism.

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steel (Mo steel). These steels were solution annealed at 1200°C for 1 hfollowed by air cooling. The creep rupture strength is improved by thesubstitution of W for Mo for long periods of time. The main precipitates inthe Mo steel are a coarse σ phase at grain boundaries and M23C6 carbidesand a coarse Fe2Mo Laves phase inside the grain, while those in the W steelare fine M23C6 carbides but no σ phase at grain boundaries and fine Fe2WLaves phase, fine M23C6 and fine MX inside the grain, as shown schematicallyin Fig. 9.13. TEM observations show that the fine precipitates of M23C6, MXand Fe2W Laves phase in the W steel serve as an effective dislocation barrier,Fig. 9.14.

In the 20Cr–25Ni–1.5Mo–NbTiN steel (KA–SUS310J2TB) with 0.3Nband 0.05Ti, three types of precipitate are observed in the crept specimens:massive particles of 0.4–0.5 µm, string-like Cr–Nb nitrides of less than 0.03µm and fine granular and needle-like M23C6 carbides of less than 0.2 µm.35

The addition of a small amount of Nb and Ti is effective in depressing ofprecipitate coarsening.

The cast austenitic steel 19Cr–12.5Ni–4Mn–MoNbN (CF8C-Plus) with0.8Nb was developed as an upgraded version of 19Cr–10Ni–MoNbSi steel(CF8C).18 The alloy-design philosophy for CF8C-Plus is to eliminate all ofthe detrimental phases such as the σ phase, leaving only carbides forstrengthening. After creep rupture testing of CF8C-Plus for 23 000 h at 850°Cand 35 MPa, fine NbC carbides of much less than 50 nm in diameter areclosely spaced and uniformly dispersed throughout the matrix. These fineNbC carbides nucleate on dislocations very early during creep and are one ofthe important strengthening mechanisms.

3% Mo5% Mo5% W7% W

700°C × 105 h

21.5 22.0 22.5 23.0 23.5 24.0 24.5 25.0T(20 + logt) × 10–3

Str

ess

(MP

a)

200

150

100

80

9.12 Creep rupture data for austenitic 0.08C–23Cr–43Ni (5–7)W–0.2Nb–0.1Ti–0.003B steel (W steel) and 0.08C–23Cr–43Ni–(3–5) Mo–0.2Nb–0.1Ti–0.003B steel (Mo steel) at 700, 750 and 800°C.

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9.4 Loss of strengthening mechanisms in 9–12Cr

steels during long time periods

In recent years, efforts have been made to clarify the mechanisms of creepstrength loss in tempered martensitic 9–12Cr steels at 550°C and aboveduring creep exposure. The loss of creep strength often takes the form of asigmoidal inflection over long periods of time in creep rupture data. Theproposed mechanisms relate mainly to a loss of precipitation hardening byfine carbonitrides and also to a loss of dislocation hardening during creepexposure. These accelerate microstructure evolution such as the coarseningof lath and block, resulting in a loss of sub-boundary hardening.4

Fine laves (much)M23C6 Coarse laves (little)

M23C6 σ0.08C-23Cr-43Ni-5/7W–0.2Nb-0.1Ti-0.003B

(a)

0.08C-23Cr-43Ni-3/5Mo–0.2Nb-0.1Ti-0.003B

(b)

9.13 Schematic drawings of microstructure in (a) 0.08C–23Cr–43Ni–(5–7)W–0.2Nb–0.1Ti–0.003B and (b) 0.08C–23Cr–43Ni(3–5)Mo–0.2Nb–0.1Ti-0.003B steels at 700, 750 and 800°C.

9.14 Fine precipitation of the Fe2W Laves phase as well as M23C6 andMX, showing effective obstacles to dislocation motion in 23Cr–43Ni–7W–0.2Nb–0.1Ti–0.003B steel (HR6W), after creep rupture testing for58798.4 h at 700°C.

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9.4.1 Precipitation of new phases and dissolution of finecarbonitrides

The precipitation of Z phase, M6X carbonitrides and Fe2(W, Mo) Lavesphase during creep causes a loss of creep strength over long periods of time,because they consume existing fine precipitates. The Z phase is a complexnitride of the form Cr(Nb,V)N. Figure 9.15(a) shows the sigmoidal behaviorof creep rupture data for 0.1C–10.84Cr–0.14Mo–2.63W–2.86Co–0.55Ni–0.19V–0.06Nb–0.016N–0.019B steel (TAF 650) at 650°C, compared withthe data for original TAF and 9Cr–1Mo–0.2V–0.05Nb steel (Mod.9Cr–1Mo,T91).38 The synergetic effect of Z phase precipitation and tungsten depletionof the solid solution by Fe2W Laves phase formation could be the reason forthe sigmoidal shape of the creep strength curve of TAF 650 steel. Precipitationof the Z phase takes place after a long time at the service temperature and itforms large particles at the expense of the previously formed fine vanadiumnitrides, which leaves a vanadium nitride free zone around the Z phase. Thiscauses a drastic loss of creep strength. Figure 9.15(b) shows an example ofa Z phase formed in a 12Cr steel NF12.39 It has been suggested that theprecipitation of the Z phase may have the most detrimental effect on creepstrength for tempered martensitic high Cr steels. 12Cr steels are moresusceptible to precipitation of the Z phase than 9Cr steels because of the highcontent of Cr.39 A higher content of nitrogen also accelerates the precipitationof the Z phase.40

The coarsening of M23C6 carbides in 12CrMo(W)VNbN steels isaccompanied by dissolution of fine MX carbonitrides owing to the precipitationof coarse M6X and or the Z phase.41 Increase in Ni content in 12CrMoV steelresults in accelerated microstructure degradation with more rapid coarseningof M23C6, dissolution of MX, and precipitation of coarse M6X and Fe2Mo.42

The impurities Al and Ti cause the formation of AlN and TiN in creep-

(b)

Z phaseMX

0.2µm

TAF 650TAFTAF 650Mod. 9Cr-1Mo

100 1000 10 000 100000Time to fracture (h)

(a)

Z6

Z5Z4

Str

ess

(MP

a)

400

300

200

100

80

60

40

9.15 Comparison of stress versus time to rupture curves for varioustempered martensitic 9–12Cr steels at 650°C.

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resistant steels during creep which degrades the creep strength at the expenseof dissolved nitrogen and fine vanadium nitrides.43 Al and Ti are strongnitride-forming elements. Figure 9.16 shows the creep rupture data for thenine heats of tempered martensitic 12Cr–1Mo–1W–0.3V steel at 600°C,where the variation in impurities Al and N was 0.007–0.040Al and 0.015–0.032N, respectively, between the different heats. The time to rupture simplydecreases with decreasing available nitrogen concentration under low stress

69 MPa98 MPa157 MPa216 MPa

RAARABRACRADRAERAFRAGRAHRAJ

Str

ess

(MP

a)

300

100

80

60

40101 102 103 104 105

Time to rupture (h)(a)

Tim

e to

ru

ptu

re (

h)

105

104

103

102

0 0.04 0.08 0.12∆ = N – Al – Ti (atom%)

(b)

9.16 (a) Creep rupture data for nine heats of tempered martensitic12Cr–1Mo–1W–0.3V steel with different Al, Ti and nitrogen contentsat 600°C and (b) time to rupture as a function of available nitrogenconcentration defined as nitrogen – Al – Ti (atom%).

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and long time conditions. The available nitrogen concentration is defined asthe concentration of nitrogen free from AlN and TiN, and is given by thedifference ∆ = nitrogen – Al – Ti (atom%). Higher Al content results in loweravailable nitrogen concentration.

9.4.2 Preferential recovery of microstructure near prioraustenite grain boundaries

The loss of creep rupture strength in T91 steel is due to preferential recoveryof the lath martensitic microstructure in the vicinity of prior–austenite grainboundaries, as shown in Fig. 9.17.44 The preferential recovery promotespreferential and localized creep deformation in the vicinity of prior-austenitegrain boundaries, resulting in a premature creep rupture. The dissolution ofMX and the precipitation of the Z phase promote preferential recovery.9

9.4.3 Progressive coarsening of M23C6

Under the condition of no Z phase formation, a rapid loss of creep rupturestrength was observed for a 9Cr–3W–3Co–0.2V–0.05Nb steel with 0.08%carbon but no addition of nitrogen, for long periods more than about 1000 hat 650°C as shown in Fig. 9.18.25 The residual nitrogen content of the steelwas only 0.0019% (19 ppm), suggesting an extremely low content of MXcarbonitrides and hence an extremely low driving force for Z phase formationduring creep. Therefore, the dissolution of fine M2X and MX carbonitridesand the precipitation of the Z phase can be excluded from the main explanationfor the loss of creep rupture strength. The proposed mechanism is the coarseningof M23C6 carbides and the recovery of lath martensitic microstructure in thevicinity of prior-austenite grain boundaries. The addition of a small amountof boron, about 100 ppm, effectively suppresses the coarsening of M23C6carbides in the vicinity of prior-austenite grain boundaries and hence suppressesthe rapid loss of creep rupture strength.

9.4.4 Loss of creep ductility

Maruyama et al.3 pointed out that loss of ductility is the origin of the loss ofcreep rupture strength in 11Cr 2W–0.3Mo–CuVNb steel at 650°C. Thereduction in area measured after creep rupture was only 11% at 650°C and100 MPa, while it was 86% at 700°C and 100 MPa. The following is theirscenario of the loss of creep rupture strength. Enhanced recovery of sub-grain structure takes place along grain boundaries. Strain concentration alongthe boundary regions forms grain boundary cracks. This results in low ductility,premature failure and loss of rupture strength.

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9.17 (a) Stress versus time to rupture curves for Mod.9Cr–1Mo steel(T91) and (b) TEM micrograph after creep rupture testing for 34141 hat 600°C and 100 MPa.

550 °C

600 °C

650 °C

700 °C

725 °CPredicted(εr = 30%)

101 102 103 104 105

Time to rupture (h)

(a)

Str

ess

(MP

a)

500

300

10080

60

40

20

1 µm

(b)

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9.4.5 Recovery of excess dislocations resulting fromlow-temperature tempering

Excess dislocations resulting from low-temperature tempering cause a rapiddecrease in creep rupture strength for 12Cr–1Mo–1W–VNb steel for longperiods of time at 600–650°C, as described in Section 9.2.3.14 This is becauseexcess dislocations accelerate recovery and recrystallization during creepwith the aid of stress, which also promotes microstructure evolution duringcreep. 12Cr turbine steels are more susceptible to a rapid loss of creepstrength for long periods of time than 9–12Cr boiler steels, because thedislocation density after tempering is much higher in 12Cr turbine steelsthan in boiler steels.

9.4.6 Effect of δ-ferrite

Dual phase 12Cr steel (12Cr–0.4Mo–2W–CuVNb, specified as KA-SUS410J3DTB) exhibits a rapid decrease in creep strength at 600–650°C forlong periods of time. Kimura et al.45 proposed that the degradation is causedby inhomogeneous creep deformation around δ-ferrite. Because diffusion ofcarbon and nitrogen is promoted by the large concentration gap across theinterface between martensite and δ-ferrite, enhanced diffusion promotes thecoarsening of precipitates and microstructure recovery. Igarashi et al.46 reportedthat heterogeneous creep deformation at low stresses is a main factor in thedegradation in long-term creep strength. They pointed out that a decrease in

9.18 Loss of creep rupture strength in 9Cr–3W–3Co–VNb steelwithout boron at 650°C and for long periods of time. The addition ofsmall amount of boron suppresses the degradation.

0 ppmB48 ppmB92 ppmB139ppmB

101 102 103 104 105

Time to rupture (h)

Str

ess

(MP

a)

200

150

100

80

60

50

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the density of fine MX and non-uniform distribution of MX in δ-ferritepromotes heterogeneous creep deformation at low stresses.

9.5 Future trends

The experimental results described in this chapter suggest that a homogeneousdispersion of thermally stable and fine precipitates not only along grainboundaries but also inside the grain provides us with an ideal microstructurefor creep-resistant ferritic and austenitic steels in terms of long-term stability.The inhomogeneous nature of microstructure evolution during creep at lowstresses, such as preferential recovery in the vicinity of grain boundaries,should also be taken into account for the minimization of creep strength lossfor long time periods. The precipitation of the Z phase and the σ phaseenhances inhomogeneous microstructure evolution in tempered martensiticand austenitic steels, respectively.

In tempered martensitic 9–12Cr steels, the key issues for the improvementof long-term creep strength seem to be the long-term stabilization of sub-boundary hardening in the vicinity of prior-austenite grain boundaries. Thestabilization of precipitation or dispersion hardening enhances the sub-boundaryhardening. In future, therefore, much effort should be paid to clarify themechanisms of precipitation and coarsening of carbonitrides and intermetalliccompounds at the lath-block and prior-austenite grain boundaries, the movementof lath and block boundaries and the development of lath and block in thevicinity of prior-austenite grain boundaries.

9.6 References

1 Takeuchi S and Argon A S, ‘Steady-state creep of single-phase crystalline matter athigh temperature’, J Mater Sci, 1976, 11, 1542–1566.

2 Meier M and Blum W, ‘Modelling high temperature creep if academic and industrialmaterials using the composite model’, Mater Sci Eng, 1993, A164, 290–294.

3 Maruyama K, Sawada K and Koike J, ‘Strengthening mechanisms of creep resistanttempered martensitic steel’, ISIJ Int, 2001, 41, 641–53.

4 Abe F, ‘Bainitic and martensitic creep-resistant steels’, Curr Opinion Solid StateMater Sci, 2004, 8, 305–39.

5 Abe F, Igarashi M, Fujitsuna N, Kimura K and Muneki S, ‘Research and developmentof advanced ferritic steels for 650°C USC boilers’, Proceedings of the 6th LiegeConference on Materials for Advanced Power Engineering 1998, Liege, Belgium,1998, 259–268.

6 Murata Y, Morinaga M and Hashizume R, ‘Development of ferritic steels for steamturbine rotors with the aid of a molecular orbital method’, Proceedings of the 4thInternational Charles Parsons Turbine Conference, Newcastle, UK, 1997, 270–282.

7 Kadoya Y and Shimizu E, ‘Effect of solute Mo, W and dispersoid carbonitride onhigh-temperature creep of ferritic steels’, Tetsu-to-Hagane, 1999, 85, 827–834 (inJapanese)

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8 Maruyama K and Nakashima H, Materials Science for High Temperature Strength,Uchida-Rokakuho, Tokyo, 1997 (in Japanese).

9 Kimura K, Kushima H, Abe F, Suzuki K, Kumai S, Satoh A, ‘Microstructural changeand degradation behaviour of 9Cr–1MoVNb steel in the long term’, Proceedings ofthe 5th International Charles Parsons Turbine Conference, Cambridge, UK, 2000,590–602.

10 Abe F, ‘Metallurgy for long-term stabilization of ferritic steels for thick section boilercomponents in USC power plant at 650°C’, Proceedings of the 8th Liege Conferenceon Materials for Advanced Power Engineering 2006, Liege, Belgium, 2006, 965–980.

11 Yoshizawa A, Fujita T, Yoshida F and Nakashima H, ‘Dispersion hardening mechanismof Y2O3 dispersed ferritic steel at high temperature’, Tetsu-to-Hagane 1996, 82, 865–869 (in Japanese).

12 Srolovitz D J, Perkovic-Luton R A and Luton M, J. Phil Mag A 1983, 48, 795–809.13 Yoshida F and Nakashima H, ‘Threshold stress for high-temperature deformation of

dispersion-strengthened alloys with incoherent dispersoids’, Key Eng Mater, 2000,171–174, 261–268.

14 Iseda A, Teranihsi H and Masuyama F, ‘Effects of chemical compositions and heattreatments on creep rupture strength of 12%Cr heat resistant steels for boiler’, Tetsu-to-Hagane, 1990, 76, 1076–1083 (in Japanese).

15 Grant N J, Bucklin A G and Rowland W, ‘Creep-rupture properties of cold-workedtype 347 stainless steel’, Trans ASM, 1955, 48, 446–455.

16 Abe F, Kimura K, Baba E, Kanemaru O and Yagi K, ‘Creep curve analysis and creeplife evaluation of 10Cr–30Mn austenitic steels’, Proceedings International Symposiumon Materials Aging and Component Life Extension, Bicego V, Nitta A and ViswanathanR (eds), Milan, Italy, 1995, 1075–1084.

17 Abe F, ‘Effect of quenching, tempering and cold rolling on creep deformation behaviorof a tempered martensitic 9Cr–1W steel’, Metall Mater Trans A, 2003, 34A, 913–925.

18 Shingledecker J P, Maziasz P, Evans N D and Pollard M J, ‘Creep behavior of a newcast austenitic alloy’, Proceedings of ECCC Creep Conference, London UK, 2005,99–109.

19 Sawada K, Kubo K, Hara T and Abe F, ‘Distribution of MX carbonitrides and itseffect on creep deformation in 9Cr–0.5Mo–1.8W–VNb steel’, Proceedings of the 7thLiege Conference on Materials for Advanced Power Engineering 2002, Liege, Belgium,2002, 1181–1188.

20 Abe F, Nakazawa S, Araki H and Noda T, ‘The role of microstructural instability oncreep behavior of a low radio-activation martensitic 9Cr–2W steel’, Metall Trans1992, 23A, 469–477.

21 Straub S, Meier M, Ostermann J and Blum W, ‘Development of microstructure andstrengthening in ferritic steel X20 CrMoV 12 1 at 823K during long-term creep testsand during annealing’, VGB Kraftwerkstechnik, 1993, 73, 646–653.

22 Sikka V K, Klueh R L, Maziasz P J, Babu S, Santella M L, Jawad M H, Paules J Rand Orie K E, ‘Mechanical properties of new grades of Fe–3Cr–W alloys’, Amer SocMech Eng Pressure Vessel and Piping, 2004, 476, 97–106.

23 Miyata K and Sawaragi Y, ‘Effect of Mo and W on the phase stability of precipitatesin low Cr heat resistant steels’, ISIJ Int, 2001, 41, 281–289.

24 Chen Z, Shan Z-W, Wu N Q, Sikka V K, Hua M H and Mao S X, ‘Fine carbide-strengthened 3Cr–3WVTa bainitic steel’, Metall Mater Trans, 2004, 35A, 1281–1288.

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25 Taneike M, Abe F and Sawada K, ‘Creep-strengthening of steels at high temperaturesusing nano-sized carbonitride dispersions’, Nature 2003, 424, 294–296.

26 Abe F, Horiuchi T, Taneike M and Sawada K, ‘Improvement of creep strength byboron and nano-size nitrides for tempered-martensitic 9Cr–3W–3Co–VNb steel at650°C’, Proceedings of the 6th International Charles Parsons Turbine Conference,Dublin, Ireland, 2003, 379–396.

27 Horiuchi T, Igarashi M and Abe F, ‘Improved utilization of added B in 9Cr heat–resistant steels containing W’, ISIJ Int, 2002, 42, S67–71.

28 Masuyama F, ‘12Cr–2.6W–2.5Co–0.5Ni–V–Nb steel’, in: Landolt-Börnstein NumericalData and Functional Relationships in Science and Technology, Group VIII: AdvancedMaterials and Technologies, Volume 2 Springer-Verlag, Berlin, 2004, 200–203.

29 Masuyama F, ‘12Cr–3W–3Co–V–Nb–Ta–Nd–N steel’, in: Landolt-Börnstein NumericalData and Functional Relationships in Science and Technology, Group VIII: AdvancedMaterials and Technologies, Volume 2 Springer-Verlag, Berlin, 2004, 204–205.

30 Ukai S and Fujiwara M, ‘Perspective of ODS alloys application in nuclear environments’,J Nucl Mater, 2002, 307–311, 749–57.

31 Sakuraya K, Okada H and Abe F, ‘Influence of heat treatment on formation behaviorof boron nitride inclusions in P122 heat resistant steel’, ISIJ Int, 2006, 46, 1712–1719.

32 Igarashi M, Muneki S, Hasegawa H, Yamada K and Abe F, ‘Creep deformation andthe corresponding microstructural evolution in high-Cr ferritic steels’, ISIJ Int, 2001,41, S101–105.

33 Okada H, Muneki S, Yamada K, Okubo H, Igarashi M and Abe F, ‘Effects of alloyingelements on creep properties of 9Cr–3.3W–0.5Pd–V, Nb, N, B steels’, ISIJ Int, 2002,42, 1169–1174.

34 Toda Y, Iijima M, Kushima H, Kimura K and Abe F, ‘Effect of Ni and heat treatmenton long-term creep strength of precipitation strengthened 15Cr ferritic heat resistantsteels’, ISIJ Int, 2005, 45, 1747–1753.

35 Takahashi T, Sakakibara M, Kikuchi M, Ogawa T, Araki S and Fujita T, ‘Elevated-temperature strength and hot corrosion resistance of 20Cr–25Ni steel for tubes inultra-supercritical power boilers’, Tetsu-to-Hagane, 1990, 76, 1131–1138 (in Japanese).

36 Semba H, Igarashi M, Yamadera Y, Iseda A and Sawaragi Y, ‘High temperaturestrength and microstructure of the 23Cr–43Ni–6W steel for USC boilers’, Report ofthe 123rd Committee on Heat Resisting Materials and Alloys, Japan Society for thePromotion of Science, 2003, Volume 44, 119–127.

37 Igarashi M, Semba H and Okada H, ‘Development of high strength austenitic steelsfor 700°C USC plant’, Proceedings of the 8th Ultra-Steel Workshop, National Institutefor Materials Science (NIMS), Tsukuba, Japan, 2004, 194–199.

38 Sklenicka V, Kucharova K, Svoboda M, Kloc L, Bursik J and Kroupa A, ‘Long-termcreep behavior of 9–12%Cr power plant steels’, Mater Character, 2003, 51, 35–48.

39 Danielsen H K and Hald J, ‘Behaviour of Z phase in 9–12%Cr steels’, Energy Mater,2006, 1, 49–57.

40 Sawada K, Taneike M, Kimura K and Abe F, ‘Effect of nitrogen content onmicrostructural aspects and Creep behavior in extremely low carbon 9Cr heat-resistantsteel’, ISIJ Int, 2004, 44, 1243–1249.

41 Vodarek V and Strang A, ‘Effect of nickel on the precipitation processes in 12CrMoVsteel during creep at 550°C’, Scripta Mater, 1998, 38, 101–6.

42 Vodarek V and Strang A, ‘Compositional changes in minor phases present in

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12CrMoVNb steels during thermal exposure at 550°C and 600°C’, Mater Sci Technol,2000, 16, 1207–13.

43 Abe F. ‘12Cr–1Mo–1W–0.3V steel’, in: Landolt-Börnstein Numerical Data andFunctional Relationships in Science and Technology, Group VIII: Advanced Materialsand Technologies, Volume 2, Springer-Verlag, Berlin, 2004, 161–169.

44 Kushima H, Kimura K and Abe F, ‘Degradation of Mod.9Cr–1Mo steel during long-term creep deformation’, Tetsu-to-Hagane, 1999, 85, 841–847 (in Japanese).

45 Kimura K, Sawada K, Kushima H and Toda Y, ‘Degradation behaviour and long-termcreep strength of 12Cr ferritic creep resistant steels’, Proceedings of the 8th LiegeConference on Materials for Advanced Power Engineering 2006, Liege, Belgium,2006, 1105–1116.

46 Igarashi M, Yoshizawa M, Iseda A, Matsuo H and Kan T, ‘Long-term creep strengthdegradation in T122/P122 steels for USC power plants’, Proceedings of the 8th LiegeConference on Materials for Advanced Power Engineering 2006, Liege, Belgium,2006, 1095–1104.

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305

10Precipitation during heat treatment and

service: characterization, simulation andstrength contribution

E. K O Z E S C H N I K and I. H O L Z E RGraz University of Technology, Austria

10.1 Introduction

Precipitation hardening is one of the most prominent ways of strengtheningmaterials. Precipitates can effectively hinder dislocation and subgrainmovement and thus increase the resistance of the material microstructureagainst plastic deformation. In industrial processes, size and number densityof precipitates are controlled by the chemical composition of the alloy aswell as the thermomechanical processing route. Owing to the significantinfluence of precipitates on the mechanical properties of the material, efficientcharacterization, modelling and simulation of precipitation processes in multi-component alloys are of considerable relevance for industry as well as foracademics. The goal of these activities is to be able to produce materials withan optimized spectrum of mechanical properties based on a fundamentalunderstanding of the complex interactions between precipitates andmicrostructure.

Materials with superior creep properties are characterized by amicrostructure, which exhibits a superior long-term resistance against plasticdeformation. This can be achieved by strong pinning forces upon dislocationsand subgrain boundaries. The two major effects of precipitates on the creepproperties of a material are:

• Increase of the creep strength by direct interaction between precipitatesand dislocations. Precipitates effectively hinder dislocations in their abilityto move through the material as a consequence of an external load. Thus,the creep process is considerably slowed down and the creep rate isminimized.

• Stabilization of the initial microstructure by pinning of grain and sub-grain boundaries. The high strength of the materials in the as-receivedcondition, i.e. the conditions in the delivery state before service, isconserved, because grain and sub-grain coarsening is minimized.

The interactions between microstructure and creep properties have been

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investigated in numerous experimental and theoretical studies.1–8 They aredescribed in other chapters of this book in detail, e.g. Chapter 9 by Abe andChapter 13 by Blum.

Figure 10.1 shows a schematic picture of the evolution of a typicalmicrostructure as observed in ferritic/martensitic 9–12%Cr steels. Theschematic emphasizes the role of precipitates in conserving the favourablefine-grained microstructure. After solidification and cooling to roomtemperature, the microstructure of typical 9–12%Cr steels consists primarilyof martensite, usually with coarse and elongated precipitates along the prioraustenite grain boundaries (see Fig. 10.1 left). In the course of the austenitizationand quality heat treatments, a dense distribution of fine precipitates is produced(see Fig. 10.1 centre). If the material is exposed to long-term thermal andmechanical loading during service, coarsening of the precipitate microstructureoccurs (see Fig. 10.1 right). Simultaneously, the material softens because themean distance between individual precipitates increases, which leads to adecrease of the effective pinning force.

In the following sections, the evolution of precipitates throughout theentire manufacturing process is investigated. First, the results of acomprehensive experimental characterization of the precipitate microstructureat several stages of the casting, austenitization and quality heat treatmentprocesses are briefly reviewed. Then, a theoretical model for simulation ofthe precipitation process is outlined and typical simulation results are presented.The interaction of precipitates with dislocations and subgrains and theconsequences for the strength of the material are discussed. The chapterconcludes with a quantitative analysis of the loss of precipitation strengtheningduring service for the example of a typical 9%Cr steel.

10.2 Microstructure analysis of the COST alloy CB8

The experimental test alloy CB8, which is exemplarily used in the followinganalyses, has been designed in the COST programme 522. The COST variant

HEAT TREATMENT CREEP LOAD

10.1 Schematic evolution of the microstructure during heat treatmentand service. Note that all images relate to the same length scale.

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Precipitation during heat treatment and service 307

CB8 has been selected because the most complete experimental picture ofthe precipitate evolution is available for this melt. These data are finally usedfor verification of the computer simulations. It is a typical 9%Cr steel forcast application, which showed excellent creep properties in short-term tests.For this reason, the variant CB8 has also been extensively investigated atlonger times. However, a significant drop in creep strength has been observedafter approximately 10 000 h of service exposure. The most important resultsof the experimental characterization are summarized subsequently. Thechemical analysis of the steel CB8 is given in Table 10.1.

10.2.1 Precipitate evolution during manufacturing

In extensive work of Sonderegger9 and Plimon,10 the evolution of precipitateshas been investigated for the first time at different positions (Pos. 1-4) in theheat treatment process. The thermal profile of this treatment is shown in Fig.10.2. It is typical for industrial components and already applied for other 9–12%Cr cast steels.

In the ‘as-cast’ condition (Pos. 1), the specimen microstructure consistsmostly of martensite. Only small amounts of retained austenite could bedetected by X-ray diffraction (0.8%). The size of the primary austenite grainsis between 0.5 and 2.0 mm. The martensite lath width as well as the subgrain

Table 10.1 Chemical composition of steel CB8–heat 173 (in wt%)

C Si Mn Cr Ni Mo W V Nb Co Al B(ppm) N0.17 0.27 0.2 10.72 0.16 10.40 – 0.21 0.060 2.92 0.028 112 0.0319

T (

°C)

1400

1200

1000

800

600

400

200

00 50 100 150 200 250 300

t (h)

Pos. 1 Pos. 2 Pos. 3 Pos. 4

Solidification

Austenitizing1080°C/8h

Tempering730°C/10 h 730°C/12 h 730°C/14 h

Stress relieving

Furnace(45K/h)

FA

SA SA

10.2 Heat treatment of COST alloy CB8 with specimen positions forexperimental characterization.

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size is approximately 1 µm. These parameters remain almost constantthroughout the entire heat treatment. Figure 10.3 shows an optical micrographof the as-cast microstructure, with prior austenite grain boundaries clearlyvisible as thin white lines.

On detailed investigation of the prior austenite grain boundaries (PAGB),elongated precipitates have been detected (see Fig. 10.4). The chemistry andcrystal structure of the precipitates indicates that these are large Mo-richprecipitates (probably Mo3B2) and Cr-rich precipitates (probably M7C3 andM23C6) with diameters of approximately 200 nm. Precipitates of a similartype also occur in the strongly segregated interdendritic regions (dark regionsin Fig. 10.3). The interior of the prior austenite grains is otherwise more orless free from precipitates (see Fe-jump ratio TEM image in Fig. 10.5).

After austenitization (Fig. 10.6), a certain degree of homogenization ofthe segregated concentration peaks is observed. The Mo3B2 precipitates atthe PAGB have more or less disappeared. Moreover, NbC precipitates, whichare randomly distributed in the matrix, and small needle-shaped Mo-richprecipitates are identified. The latter disappear again during further heattreatment. After the first quality heat treatment (Pos. 3), a significant increasein the number of precipitates is observed (Fig. 10.7). M23C6 precipitates arefound in great quantities as well as VN particles. Both types of precipitateappear preferentially at the martensite lath and subgrain boundaries.

In the sample corresponding to the ‘as-received’ condition (Pos. 4), the

1 mm

10.3 Optical micrograph of microstructure in the as-cast condition.Segregated regions from solidification are clearly observed as wellas prior austenite grain boundaries.

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10 µm

10.4 Elongated precipitates along the PAGB (scanning electronmicrograph, SEM).

10.5 CB8 in as-cast condition. Fe jump-ratio image (transmissionelectron micrograph, TEM).

0.5 µm

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0.5 µm

10.6 CB8 after austenitizing (TEM bright field).

0.5 µm

10.7 CB8 after first heat treatment (TEM bright field).

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first indications of slight coarsening of M23C6 precipitates are observed (Fig.10.8). The volume fraction of VN and the mean precipitate radius havesignificantly increased. NbC is also still present in this heat treatment condition.According to Dimmler et al.,11 the first sparse Laves phase precipitation canbe found in the as-received condition.

10.2.2 Precipitate evolution during service

In the analysis of the development of the microstructural parameters in thecourse of long-term service of CB8 at 650°C for 2000 and 7000 hours,Sonderegger9 found that the martensite lath width and the subgrain sizeremain almost constant in samples without mechanical loading. In sampleswhich have been exposed to creep loading, the subgrain size slightly increasesfrom 0.7 µm to 1.0 µm. Moreover, clear indications of coarsening of M23C6precipitates are observed in the heat-treated and creep-loaded samples. Incontrast, the mean precipitate radius of VN remains almost constant in bothsamples, while the number density increases visibly.

The phase fraction of Laves phase increases significantly during service.Dimmler11 found that the phase fraction increases from 0.4 to 0.8% forsamples analysed after 50 and 16 000 h, respectively. Simultaneously, themean radius also increases. In the first 1 000 h, an increase of the number

0.5 µm

10.8 CB8 in as-received condition (TEM bright field).

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density of Laves phase precipitates is observed, indicating significant nucleationof new precipitates of this phase. After 1000 h, the number density remainsconstant.

After 16 000 hours of service, the number density of the VN precipitatessuddenly reduces drastically. This tendency is stronger in the creep-loadedsamples, which indicates that the external load enhances the microstructureevolution. The reason for the VN dissolution is found in the appearance ofthe modified Z-phase, which has been detected in the sample with16000 hours’ service time. The Z-phase has a higher thermodynamic stabilitycompared to VN, however, nucleation of this phase is very difficult, whichis the reason that the Z-phase was not observed earlier. The role of the Z-phase in the drop of creep resistance in many 9–12%Cr steels is also discussedin Chapter 9 by Abe.

10.3 Modelling precipitation in complex systems

Since the interaction of the various different precipitates in the differentstages of the thermal treatment of advanced materials is rather complex, atheoretical approach has recently been developed for simulation of precipitationin these multi-component, multi-phase systems, taking into account allthermodynamic and kinetic interactions of the different alloying elements.This approach is briefly reviewed here and parameters relevant for the actualprecipitation simulation in the test alloy CB8 are discussed.

10.3.1 The precipitation kinetics model

Consider a unit volume of a multi-component alloy. Allow an arbitrary numberof spherical precipitates to nucleate and grow on random locations in thisvolume. The corresponding situation is sketched in Fig. 10.9.

The total Gibbs energy of this thermodynamic system can be written as:12

G N ci

n

i i k

mk

k i

n

ki ki k

m

k k = + 4

3 + + 4

=1 0 0 =1

3

=1 =12Σ Σ Σ Σµ

µ

πρλ πρ γ [10.1]

where N0i is the number of moles of component i in the matrix phase, µ0i isthe corresponding chemical potential, λk is the energy contribution due tovolumetric misfit, ρk is the radius of the precipitate with index k, cki is theconcentration of component i, µki the corresponding chemical potential andγk is the interfacial energy.

In thermodynamic equilibrium, the Gibbs free energy is a minimum. Sincereal systems during heat treatments are in a highly non-equilibrated state,driving forces exist for evolution of the precipitate microstructure such thatG is minimized. With each microstructural process that occurs in the system,part of the free energy is dissipated. In the model, three dissipative processes

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have been considered: (i) Migration of interfaces with a mobility Mk, (ii)diffusion of atoms in the precipitates and (iii) diffusion of atoms in thematrix. Detailed expressions for these quantities are described in Svoboda etal.12 and Kozeschnik et al.13

With the total Gibbs free energy and the corresponding dissipation terms,the thermodynamic extremal principle14,15 can be applied and a linear systemof rate equations for the change of radius and chemical composition of eachindividual precipitate is obtained. To find the evolution of the entire precipitatepopulation, the rate equations are integrated numerically under the constraintof mass conservation. The corresponding algorithm has been implementedin the MatCalc software,16 which is used in the following sections to computethe precipitate evolution in the steel CB8.

10.3.2 Treatment of multi-component nucleation

A most important step in modelling precipitation is the accurate treatment ofthe precipitate birth process, that is the nucleation stage. In the presentmodel, nucleation of precipitates is dealt with in the framework of an extensionof classical nucleation theory (CNT). According to this theory, the nucleationrate J, which describes the frequency of creation of new precipitates in unittime and unit volume, is given by:

10.9 Modelling precipitation in complex materials. Sphericalprecipitates in a multi-component matrix.

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J Z N GkT t

= exp – exp – *C

*β τ

[10.2]

where Z is the Zeldovich factor, NC is the number of potential nucleationsites, β* is the atomic attachment rate, G* is the energy barrier to form acritical nucleus, k is the Boltzmann constant, T is the absolute temperatureand τ is the incubation time.

The term ‘extended’ CNT emphasizes that, in multi-component systems,some quantities in Equation [10.2] must be reformulated to apply to multi-component situations. These expressions have been summarized and discussedrecently by Kozeschnik et al.17

10.3.3 Linking microstructure and precipitate nucleation

Most quantities in Equation [10.2] for the nucleation rate are eitherthermodynamic quantities or kinetic quantities related to the diffusivity ofthe atoms. They are global parameters and can be obtained from independentthermodynamic and kinetic databases. In contrast, the number of potentialnucleation sites, NC, strongly depends on the microstructure of the materialand the type of heterogeneous nucleation site which is preferred by theparticular precipitate. This feature allows for a straightforward considerationof the material microstructure, that is, the grain and subgrain size and dislocationdensity, in the nucleation stage of the precipitation simulation.

In order to obtain the correct number of available nucleation sites, asimple and yet realistic representation of the real microstructure is desirable.In the software MatCalc, the grain and subgrain structure of the matrix isapproximated by an arrangement of tetrakaidecahedrons (Fig. 10.10), whichare space-filling objects with 14 surface elements. In the symmetric geometry,

d

H

D

dλ

10.10: Tetrakaidecahedrons representing the matrix microstructureand determining the number of potential nucleation sites in theprecipitation simulation.

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these objects resemble globular grains. If this structure is elongated, a goodrepresentation of martensite laths can be given.

On the basis of the real, experimentally observed microstructure, thenumber of potential nucleation sites, for example, on prior austenite grainboundaries or subgrain boundaries, can be evaluated. The mathematicalexpressions for these quantities are described by Rajek.18 Table 10.2 presentsan overview of the number of potential nucleation sites at differentheterogeneous lattice positions, if each atomic position is considered torepresent a potential nucleation site.

10.4 Computer simulation of the precipitate

evolution in CB8

In this section, the kinetic model, which has been outlined previously, is appliedto the simulation of the precipitate evolution during the entire heat treatmentand service of steel CB8. First, a thermodynamic equilibrium analysis of thissteel is performed, which provides an overview of the type and amount ofdifferent phases that can be expected to occur at a given chemical compositionand temperature. Then, the results of the kinetic simulation are discussed.

10.4.1 Thermodynamic equilibrium analysis

The thermodynamic equilibrium analysis is an important step in comprehensivematerial characterization. By this method, information can be obtained about:

• the type and number of phases which occur in this material underequilibrium conditions,

Table 10.2 Number of nucleation sites (m–3) at 650°C in a stretchedtetrakaidecahedron with a subgrain structure. Dislocation density in austeniteis assumed to be ρ = 1011m–2, the dislocation density in ferrite ρ = 1014m–2,the austenite and ferrite grain sizes are 100 µm and the ferrite subgrain size is0.1 µm with an elongation factor s = 100. s0 is the lattice constant, afcc andabcc are the mean atomic distances

Nucleation sites (m–3) Ferrite (bcc) Austenite (fcc)a0 = 2.87 × 10–10 m a0 = 3.515 × 10–10 mabcc = 2.27 × 10-–0 m afcc = 2.214 × 10–10 m

Bulk (B) 8.27 × 1028 8.37 × 1028

Dislocations (D) 4.36 × 1023 4.37 × 1020

Grain boundary (GB) 7.63 × 1023 7.70 × 1023

Grain boundary edge (E) 5.13 × 1018 5.15 × 1018

Grain boundary corner (C) 1.68 × 1013 1.68 × 1013

Subgrain boundary (SGB) 2.00 × 1026 /

bcc = body centred cubic; fcc = face centred cubic

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• chemical composition of the equilibrium phases,• equilibrium solution temperatures of the individual phases, giving

indications of temperatures for optimum heat treatment conditions,• equilibrium transformation temperatures at which allotropic

transformations (e.g. fcc (face centred cubic)–bcc (body centred cubic)/bct (body centred tetragonal) transition) occur.

Figure 10.11 shows calculated phase diagrams for the steel CB8 as a functionof carbon and nitrogen content. All simulations are based on the thermodynamicdatabase TCFE3 and the diffusion database Mobility_v21 from ThermoCalcAB, Stockholm, Sweden. To account for the stabilizing effect of silicon onthe Laves phase, the corresponding parameters have been modified as suggestedby Dimmler.11 Moreover, although not considered in the phase diagrams ofFig. 10.11, a revised thermodynamic description for the modified Z-phase19

has been added to this database, which is a further development of the initialassessment of Danielsen and Hald.20 These values have been used in thekinetic simulations presented in the following section.

10.4.2 Precipitation simulation in CB8

With the theoretical model for multi-component precipitation kinetics andthe thermodynamic and kinetic data described in the previous section, theentire heat treatment of the alloy CB8 has been studied. Figure 10.12summarizes the results of the simulation.

When looking at the temperature profile of the heat treatment (top image(a) in Fig. 10.12), several individual steps can be distinguished. The simulationstarts at 1400°C, which is just below the solidus temperature of this alloy. Itis assumed that all elements are homogeneously distributed in the matrix atthis time and no precipitates exist. The material then cools linearly to atemperature of 350°C. This temperature corresponds to the observedaustenite to martensite transformation start temperature. In the simulation, itis assumed that this transformation occurs instantaneously and the parentand target phases have identical chemical composition. It is further assumedthat no diffusive processes and, consequently, no precipitation occurs belowthis temperature. At this point, the matrix phase is changed from a facecentred cubic (fcc) austenite to a body centred cubic (bcc) ferrite structure.Since no separate thermodynamic description is available for the bct martensitephase, the bct phase is substituted by the bcc phase in the simulations.

In the next step, the material is reheated for austenitization. At theexperimentally observed transformation temperature of 847°C, the ferritematrix is changed to austenite again. After austenitization, the transformationto martensite/ferrite is performed again at 350°C. After three quality heattreatments, service at 650°C for 100 000 h is simulated.

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T (

°C)

1600

1400

1200

1000

800

600

400

0 0.02 0.04 0.06 0.08 0.1XN(%)

1

2 3

45

67

1098

12 13 14 1915

202524

28

2927

23

26

21

2216

17

11

1. Liquid2. Liquid, δ3. Liquid, δ, γ4. Liquid, δ, γ, M2B tetr.5. δ, γ, M2B tetr.6. γ, M2B tetr.7. γ, M2B tetr., AIN8. γ, M2B tetr., NbC9. γ, M2B tetr., AIN, NbC

10. γ, M2B tetr., AIN, VN11. γ, M2B tetr., AIN, NbC, VN12. γ, M2B tetr., NbC, M23C613. γ, M2B tetr., AIN, NbC, M23C614. γ, M2B tetr., AIN, NbC, VN, M23C615. γ, M2B tetr., AIN, VN, M23C6

16. α, γ , M2B tetr., NbC, M23C617. α, γ , M2B tetr., AIN, NbC, M23C618. α, γ , M2B tetr., AIN, NbC, VN, M23C619. α, γ , M2B tetr., AIN, VN, M23C620. α, γ , M2B tetr., AIN, VN, M23C6, Cr2N21. α, M2B tetr., NbC, M23C622. α, M2B tetr., AIN, NbC, M23C623. α, M2B tetr., AIN, NbC, VN, M23C624. α, M2B tetr., AIN, VN, M23C625. α, M2B tetr., AIN, VN, M23C6, Cr2N26. α, M2B tetr., AIN, NbC, M23C6, Laves27. α, M2B tetr., AIN, NbC, VN, M23C6, Laves28. α, M2B tetr., AIN, VN, M23C6, Laves29. α, M2B tetr., AIN, VN, M23C6, Cr2N, Laves

(a)

10.11 Phase diagrams for steel CB8 obtained from computationalthermodynamics as function of (a) C and (b) N, respectively. Themodified Z-phase, which replaces VN and, eventually, NbCprecipitates at long times, has been neglected; the boride phases areincluded.

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The three other plots in Fig. 10.12 display the evolution of the phasefraction, Fig. 10.12(b), the mean precipitate radius, Fig. 10.12(c), and thenumber density, Fig. 10.12(d), of each precipitate type. The phase fractionsof M23C6 and Laves-phase are multiplied by a factor of 1/10 to give a bettervisual representation of the results.

During cooling from 1400°C, various precipitate phases nucleate at the

10.11 (Continued)

1. Liquid2. Liquid, δ3. Liquid, δ, γ4. Liquid, δ, M2B tetr.5. Liquid, γ6. Liquid, γ, M2B tetr.7. δ, M2B tetr.8. Liquid, δ, γ, M2B tetr.9. δ, γ, M2B tetr.

10. γ, M2B tetr.11. γ, M2B tetr., NbC12. δ, γ, M2B tetr., AIN13. γ, M2B tetr., AIN14. γ, M2B tetr., NbC, AIN

15. γ , M2B tetr., AIN, VN16. α/δ, γ, M2B tetr., AIN, VN17. γ , M2B tetr., NbC, AIN, M23C618. γ , M2B tetr., NbC, AIN, VN19. γ , M2B tetr., NbC, AIN, VN, M23C620. α/δ, γ, M2B tetr., NbC, AIN, VN21. α/δ, M2B tetr., AIN, VN22. α/δ, γ, M2B tetr., NbC, AIN, VN, M23C623. α/δ, M2B tetr., NbC, AIN, VN24. α/δ, M2B tetr., NbC, AIN, VN, M23C625. α/δ, M2B tetr., AIN, VN, Laves26. α/δ, M2B tetr., NbC, AIN, VN, Laves27. α/δ, M2B tetr., NbC, AIN, VN, M23C6, Laves

T (

°C)

1600

1400

1200

1000

800

600

400

0 0.05 0.1 0.15 0.2 0.25 0.3 0.35 0.4XN(%)

1

2 6

109 8

16

14

19

15

20

25

24

27

23

26

18

11

21 22

17

5

3

12 13

7 4

(b)

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10.12 Kinetic simulation of the precipitate evolution in CB8 during(a) heat treatment and service; (b) f, phase fraction (c) R, meanprecipitation radius; and (d) N, number density. The phase fractionsof M23C6 and Laves-phase are multiplied by a factor of 1/10 to givea better visual representation of the results.

1 10 100 1000 10000 100000t (h)

N(m

–3)

1×1024

1×1022

1×1020

1×1018

1×1016

1×1014

(d)

1 10 100 1000 10000 100000t (h)

R (

nm

)

100

10

1

0.1

(c)

1 10 100 1000 10000 100000t (h)

f(%

)

0.4

0.3

0.2

0.1

0

(b)

1 10 100 1000 10000 100000t (h)

1200

1000

800

600

400

T(°

C)

M23C6(×0.1)

M7C3Laves-phase(×0.1)NbCVN

Z-PhaseP

Z-PhaseMM23C6exp. (×0.1)MXexp.VNexp.

(a)

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austenite grain boundaries. At room temperature, NbC, VN, M23C6, M7Cand Laves-phase precipitates are observed. After changing the matrix phasefrom austenite to ferrite and reheating, severe precipitation of all phases setsin again. Most phases, except NbC, dissolve again during further heating toand holding at the austenitization temperature of 1080°C. During subsequentcooling and the first of the three quality heat treatments, nucleation of variousprecipitates continues. Slight coarsening of some precipitates, particularlyM23C6 is already observed.

During service, the simulation predicts significant precipitation of theLaves-phase, which is in accordance with the observations of Dimmler.21

After several thousand hours of service, the phase fraction of the modifiedZ-phase gradually increases and, simultaneously, the thermodynamically lessstable VN precipitates start to dissolve. Again, this is consistent withexperimental observation.9 It should be noted that, owing to the differentpossibilities of Z-phase nucleation (Z-phase can be formed by heterogeneousnucleation in the matrix as well as by direct transformation of VN into Z-phase),22 two Z-phase populations are introduced in the simulation.

Generally, the simulation results are in reasonable agreement with theexperimental data for CB8. In view of the complexity of the problem and thehigh degree of abstraction of the theoretical model, the overall performanceof the simulation is excellent. It is particularly important to emphasize thatthe simulations have been performed on the basis of independentthermodynamic and diffusion databases and no general fit parameters havebeen used. The few necessary modifications of the original thermodynamicdatabase and the correction for the estimated interfacial energy for the Laves-phase are well founded and described in detail by Rajek.18

In the next section, the interaction between precipitates and microstructureis analysed and in the last section, a prediction of the loss of precipitationstrengthening over the lifetime of a component made from CB8 is attempted.

10.5 Microstructure–property relationships

As already pointed out in the Introduction (Section 10.1), precipitates act asmicrostructure-stabilizing components and as efficient obstacles for dislocationmovement. In this section, the interaction forces are analysed on a quantitativebasis.

10.5.1 Precipitate–dislocation interaction

Precipitates and mobile dislocations can interact in one of the followingways (see e.g. McLean):23

1. A dislocation can pass coherent precipitates by cutting (breaking) theprecipitate. A stacking fault is left in the precipitate;

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2. A dislocation can pass precipitates by bending between them and closingthe bent lines to loops. A dislocation loop is left around the by-passedprecipitate (Orowan mechanism);

3. A dislocation can pass the precipitates by climbing;4. A dislocation can drag the precipitates with it. This mechanism is possible

only for very small precipitates. The velocity-determining factor in thiscase is the mobility of the dragged precipitates.

The precipitate–dislocation interaction mechanism in operation dependson a number of factors, among which the availability of glide planes, theheight of the local forces and the hardness of the precipitates are mostimportant. Owing to the physical nature of the four processes, mechanisms1 and 2 are considerably faster than mechanisms 3 and 4. If the latter are theoperating creep mechanisms, the creep rate is significantly lower than thecreep rate based on mechanisms 1 and 2. For details, the interested reader isreferred to References (23)–(27).

If the precipitates are sufficiently hard, such that dislocations cannot bypassthem by cutting, the upper limit of the pinning force is determined by theOrowan stress. According to Ashby,28 the Orowan stress τ0 is given by:

τ λξ

00

= ln C Gbr

[10.3]

where C is a constant (C = 0.159 for screw dislocations and C = 0.227 foredge dislocations), G is the shear modulus, b is the Burgers vector, λ0 is themean particle distance, r0 is the ‘inner cut-off radius’ and ξ the ‘outer cut-offradius’ of the dislocation.28

10.5.2 Precipitate–subgrain boundary interaction

The pinning force of precipitates upon subgrain boundaries can be describedbased on a suggestion of Zener (reference (3) (private communication)) inSmith,29 which originally describes the drag force on grain boundaries duringgrain growth in the presence of precipitates. The basic idea is that a precipitate,which is located on a grain boundary, reduces the effective grain boundaryarea. On leaving the particle behind, this area must be re-established. Thisprocess requires energy and thus acts against boundary migration. A numberof modifications of the original theory have since been developed, whichhave been reviewed by Manohar et al.30

McLean31 has pointed out that Zener’s ideas can likewise be applied tosubgrain pinning. Accordingly, the mobility of subgrain boundaries is stronglyreduced in the presence of precipitates, which has been observed experimentally,for instance, in the TEM investigations of Eggeler.32 The critical subgrainradius Rcrit during subgrain coarsening, above which the coarsening processcomes to a stop, is given by:

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Rn rcrit = 0.846 1

⋅ [10.4]

where n is the number density of the precipitates and r is the mean precipitateradius.

The effective retarding force of subgrains upon dislocation movement canbe estimated according to Gladman33 and McElroy and Szkopiak34 based onthe grain size-dependent contribution to the Hall–Petch relation. Taking intoaccount Equation [10.4], the strength contribution ∆τsubgr from subgrains isgiven as:

∆τ subgrd

s=

kd m [10.5]

where kd represents the subgrain strengthening coefficient,33 ds is the subgrainsize and m is an exponent, which is typically in the order of m = 1/2.According to Gladman,33 kd is usually much smaller than the coefficient forgrains. Kosik et al.35 have considered the effect of subgrains as beingcomparable to cold-working. However, it has been found9 that the subgrainsize changes only slightly during heat treatment and thermal ageing.Consequently, the stress contribution of subgrains remains more or less constantduring service and it is therefore not considered further here.

10.5.3 Precipitate–grain boundary interaction

Precipitates similarly affect the mobility of grain boundaries; however, themechanism of pinning is different from that for subgrain boundaries. Thelatter are small angle boundaries and thus represent an array of dislocations,which accommodate the small lattice misfit between the two subgrains. Agrain boundary is a randomly oriented high-angle boundary. In the case of9–12%Cr steels, the fraction of grain boundaries is very small comparedwith subgrain boundaries because typical grain sizes for cast materials are inthe order of millimetres, whereas subgrain sizes are typically in the order ofmicrometres or less. The influence of grain boundaries is therefore alsoneglected in the further analysis.

10.6 The back-stress concept

If an external force is acting on a microstructure, it is frequently assumedthat the external force acts on each representative volume of the microstructuresimultaneously. If the external force is high enough, plastic deformation ofthe material occurs via movement of dislocations and grain (subgrain)boundaries. At elevated temperatures, and if only a small external load isapplied, which are conditions that are typical for creep deformation, the

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situation is different. In this case, part of the external driving pressure σe iscounteracted by heterogeneous internal microstructural constituents, such asprecipitates and interfaces. Consequently, the entire external load cannot beassumed to represent the driving force for the creep process; only the part ofthe external stress σex, which exceeds the amount of inner stress σi from thecounteracting microstructure, effectively contributes to the creep process.Since the inner stress reduces the effect of the external stress, this approachis commonly denoted as the back-stress concept. The effective creep stressσeff can be expressed as:

σeff = σex – σi [10.6]

In a recent treatment by Dimmler,21 the inner stress σi has been expressed asa superposition of individual contributions from dislocations and precipitates.When also taking into account the contribution from subgrain boundaries,the inner stress is:

σi = Mτi = M(τdisl + τprec + τsgb) [10.7]

where M is the Taylor factor (usually between 2 and 3, see Dimmler)21 andτ is the shear stress. The subscripts in the terms in the parentheses denotecontributions from dislocations, precipitates and subgrain boundaries,respectively. When taking into account the inner stress, the general Nortoncreep law (see e.g. Čadek)36 can be rewritten as:

ε σ σ σ = ( – ) = ex i effA An n⋅ ⋅ [10.8]

where A and n are constants. When examining the individual contributions ofthe different mechanisms to the back-stress based on Equation [10.7], accordingto Taylor (cited by Dimmler),21 the part stemming from dislocations can beexpressed as:

τ α ρdisl m= ⋅ ⋅ ⋅G b [10.9]

where ρm denotes the density of mobile dislocations and the value of α isbetween 0.84 and 1 (see e.g. Weinert).37 The quantities G and b are definedin Equation [10.3].

The contribution of precipitates to the total back stress has already beendiscussed in Section 10.5.1 and is described by the critical Orowan stress τ0

(equation [10.3]). This quantity denotes the maximum back stress caused bya random distribution of precipitates using a mean distance λ between theprecipitates. The latter can be estimated using the assumption that eachprecipitate consumes approximately the same bulk volume. In this case, λ isgiven by:

λ π = 6

3 ⋅ ni[10.10]

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where ni represents the number density of precipitates in units of m–3. Whencombining Equations [10.3] and [10.10], the total back-stress contributionfrom a bulk distribution of precipitates reads:

τ π ξprec

3

0=

6

ln CGbn

ri⋅

[10.11]

The quantity τprec represents the maximum back-stress caused by precipitates.If the external load reduced by the back-stress contribution of the otherstrengthening mechanisms is below this threshold, the dislocations areeffectively pinned and can only pass the precipitates by the climb mechanism.Since dislocation climb is a diffusional process, the effective creep rates areusually low. If the threshold stress is exceeded, the dislocations can bypassthe precipitates by the Orowan mechanism, which is a much faster processcompared to climb. When this change in mechanism occurs, the exponent inthe Norton creep law increases significantly and creep deformation is stronglyenhanced.

The selection of the operative creep mechanism is mainly determined bythe height of the external load. However, a transition from dislocation climbto the Orowan mechanism can also be caused by a decrease in back-stressduring service. If this transition occurs, for instance, owing to coarsening ofprecipitates or thermodynamic instability of a precipitation strengtheningphase, extrapolation of the creep strength from short-term experiments tolong-term service behaviour can result in fatal overestimation of the residuallifetime of components. This aspect has been discussed in detail by Dimmlerand co-workers.21,38

10.7 Loss of precipitation strengthening during

service of CB8

In this section, the theory and methodology described previously is appliedto a prediction of the evolution of the strength contribution of precipitatesduring service of the steel CB8. The same calculation as shown previouslyin Section 10.4.2 is used as a basis and evaluated in terms of the maximumOrowan stress, that is the back-stress contribution from precipitates.

For the evaluation of the back-stress during service at 650°C, the constantsin Equation [10.11] are assumed to be C = 0.19 and G = 62.3 GPa (convertedfrom the data in Guntz et al.39 The outer cut-off radius ξ is assumed to havea value of twice the precipitate radius ξ = 2rprec, the inner cut-off radius istaken as twice the Burger’s vector with r0 ~ 2b ~ 0.5 nm. Figure 10.13 showsthe predicted evolution of the back-stress during service as calculated by themodel described in this chapter. The graphs show the predicted contributionsfor each phase separately as well as the total back stress including the combinedeffect of all precipitates with and without the effect of the Z-phase.

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In the as-received condition, after all heat treatments, the total strengthcontribution from precipitation hardening is estimated to be in the order of90 MPa. This quantity varies strongly with selection of input values andshould therefore not be considered in terms of absolute values. Owing to theinevitable effect of Ostwald ripening (coarsening), the density of precipitatesis reduced, which is reflected in the gradual decrease of the total back-stressup to times of approximately 10 000 h. When the modified Z-phase is included,the total back stress shows a drastic depression between 10 000 and 20 000 h.This effect is due to the enhanced nucleation and subsequent growth of theZ-phase, which causes dissolution of the VN precipitates. In the later stages,the decrease in back stress continues in a steady manner again, however, ata much lower level than before. Comparison of the curves for the integratedback stress in Fig. 10.13 indicates that Z-phase formation causes an additionalback stress reduction at 100 000 h approximately 20 MPa. This effect isassumed to be responsible for the drop in creep strength of various differentferritic/martensitic creep-resistant 9–12%Cr steels during long-term creepexposure (compare with Chapter 9).

10.8 Summary and outlook

Precipitation strengthening is a key mechanism for improving mechanicalproperties of creep resistant materials. To capture the evolution of the precipitatemicrostructure during heat treatment and service in these complex materials,experimental characterization must be performed with state-of-the-arttechniques and methodologies. If these are complemented with suitable

10.13 Predicted degradation of the back stress during service of CB8including the modified Z-phase contribution (solid bold line) andartificially suppressing this phase (dashed bold line). Subscript Pdenotes precipitates of Z-phase nucleated on existing VN particles,the subscript M denotes Z-phase precipitates nucleated in the matrix.

1000 10000 100000t (h)

Heattreatment

Service

100

80

60

40

20

0

Bac

k-st

ress

(N

mm

–2) VN

Totalwith Z-Phase

Totalwithout Z-PhaseLaves-phase

M23C6

NbC

Z-PhaseP

Z-PhaseM

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modelling and simulation approaches, a profound understanding of theinteractions of all microstructural constituents can be achieved and predictionsof the influence of variations of chemical composition and process parameterscan be attempted.

A corresponding methodology is introduced in this chapter. With thesimulation software MatCalc, which is based on independent thermodynamicand kinetic databases and on a novel theoretical approach to model multi-component multi-phase precipitation kinetics, the strengthening contributionfrom precipitates is predicted for the entire life time of a sample for theexample of the 9–12%Cr steel experimental test alloy COST CB8. Theinfluence of Z-phase precipitation on the degradation of the creep strength isquantified and a significant drop of the creep strength is confirmed at thetime when the VN precipitates dissolve.

In further optimization of the mechanical properties of creep-resistantmaterials, improved understanding of the thermodynamic and kineticinteractions between the elements and phases in the microstructure of thesecomplex alloy systems is necessary. Based on these improved input data andprovided that the theoretical models describing the physical processes arerealistic and accurate, kinetic modelling and simulation can aid even betterin optimizing existing alloy concepts and production routes or even help toidentify promising concepts for new materials for the next generation ofcreep-resistant materials.

10.9 References

1 F. Kauffmann, G. Zies, D. Willer, C. Scheu, K. Maile, K.H. Mayer and S. Straub,‘Microstructural investigation of the boron containing TAF steel and the correlationto the creep strength’, 31. MPA-Seminar Werkstoff- und Bauteilverhalten in der Energie-& Anlagentechnik, Stuttgart, 2005, 13–14 October.

2 A. Kostka, K.-G. Tak, R.J. Helling, Y. Estrin and G. Eggeler, ‘On the contribution ofcarbides and micrograin boundaries to the creep strength of tempered martensiteferritic steels’, Acta Mater., 2007, 55, 539–550.

3 J. Eliasson, A. Gustafson and R. Sandström, ‘Kinetic modelling of the influence ofparticles on creep strength’, Key Eng. Mater., 2000, 171–174, 277–284.

4 R. Lagneborg, ‘Effect of grain size and precipitation of carbides on creep propertiesin Fe–20%Cr–35%Ni alloys’, J. Iron and Steel Institute, 1969, 1503–1506.

5 F.R.N. Nabarro, ‘Grain size, stress, and creep in polycrystalline solids’, Phys. SolidState, 2000, 42, 1456–1459.

6 K.E. Amin and J.E. Dorn, ‘Creep of a dispersion strengthened steel’, Acta Metallurgica,1969, 7, 1429–1434.

7 K. Maruyama, K. Sawada and J. Koike, ‘Strengthening mechanisms of creep resistanttempered martensitic steel’, ISIJ Int., 2001, 41, 641–653.

8 K. Sawada, K. Kubo and F. Abe, ‘Creep behaviour and stability of MX precipitatesat high temperature in 9Cr–0.5Mo–1.8W–VNb steel’, Mater. Sci. Eng., 2001, A,319–321, 784–787.

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9 B. Sonderegger, Characterisation of the Substructure of Modern Power Plant Steelsusing the EBSD-Method, PhD Thesis, Graz University of Technology, 2005 (in German).

10 S.W. Plimon, Simulation of an Industrial Heat Treatment and AccompanyingMicrostructural Investigation of a Modern 9–12% Chromium Steel, Master Thesis,Graz University of Technology, 2004 (in German).

11 G. Dimmler, P. Weinert, E. Kozeschnik and H. Cerjak: ‘Quantification of Laves-phase in advanced 9–12% Cr-steels using a standard SEM’, Mater. Character., 2003,51, 341–352.

12 J. Svoboda, F.D. Fischer, P. Fratzl and E. Kozeschnik, ‘Modelling of kinetics inmulti-component multi-phase systems with spherical precipitates I – Theory’, Mater.Sci. Eng. A, 2004, 385 (1–2), 166–174.

13 E. Kozeschnik, J. Svoboda and F.D. Fischer, ‘Modified evolution equations for theprecipitation kinetics of complex phases in multi-component systems’, CALPHAD,2005, 28 (4), 379–382.

14 L. Onsager, ‘Reciprocal relations in irreversible processes’, Phys. Rev., 1931, 37,405–426, 1931, 38, 2265–2279.

15 J. Svoboda, I. Turek and F.D. Fischer, ‘Application of the thermodynamic extremalprinciple to modeling of thermodynamic processes in material sciences’, Phil. Mag.,2005, 85 (31), 3699–3707.

16 E. Kozeschnik, J. Svoboda, P. Fratzl and F.D. Fischer, ‘Modelling of kinetics inmulti-component multi-phase systems with spherical precipitates II – Numericalsolution and application’, Mater. Sci. Eng. A, 2004, 385 (1–2), 157–165.

17 E. Kozeschnik, J. Svoboda and F.D. Fischer, ‘On the role of chemical composition inmulti-component nucleation’, invited paper in Proceedings International ConferenceSolid-Solid Phase Transformations in Inorganic Materials, PTM 2005, Pointe HiltonSquaw Peak Resort, Phoenix, AZ, USA, 2005, 29.5.–3.6, 301–310.

18 J. Rajek, Computer Simulation of Precipitation Kinetics in Solid Metals and Applicationto the Complex Power Plant Steel CB8, PhD Thesis, Graz University of Technology,2005.

19 H.K. Danielsen, private communication.20 H.K. Danielsen and J. Hald, Z-phase in 9–12%Cr Steels, Internal Report 863,

Värmeforsk AB, 2004.21 G. Dimmler, Quantification of Creep Resistance and Creep Fracture Strength of 9–

12%Cr Steel on Microstructural Basis, PhD Thesis, Graz University of Technology,2003 (in German).

22 H.K. Danielsen, J. Hald, F.G. Grumsen and M.A.J. Somers, ‘On the crystal structureof Z-phase Cr(V, Nb)N’, Metall. Mater. Trans., 2006, 37A, 2633–2640.

23 M. McLean, ‘On the threshold stress for dislocation creep in particle strengthenedalloys’, Acta Metallurgica, 1985, 33, 545–556.

24 L.M. Brown and R.K. Ham, in Strengthening Methods in Crystals, A. Kelly and R.B.Nicholson (eds), Elsevier, Amsterdam 1971.

25 R. Lagneborg, ‘Bypassing of dislocations past particles by a climb mechanism’,Scripta Metallurgica, 1973, 7, 605–614.

26 E. Arzt and M.F. Ashby, ‘Threshold stresses in materials containing dispersed particles’,Scripta Metallurgica, 1982, 16, 1285–1290.

27 J.D. Verhoeven, Fundamentals of Physical Metallurgy, John Wiley & Sons, NewYork, 1975.

28 M. Ashby, ‘The theory of the critical shear stress and work hardening of dispersion-

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hardened crystals’, G.S. Ansell, T.D. Cooper and F.V. Lenel (eds), MetallurgicalSociety Conference, Vol. 47, Gordon and Breach, New York, 1968, 143–205.

29 C.S. Smith, ‘Grains, phases and interfaces: an interpretation of microstructure’, Trans.AIME, 1948, 175, 15–51.

30 P.A. Manohar, M. Ferry and T. Chandra, ‘Five decades of the Zener equation’, ISIJInt., 1998, 38, 913–924.

31 D. McLean, ‘Resistance to hot deformation’, Trans. Metall. Soc. AIME 1968, 242,1193–1203.

32 G. Eggeler, ‘The effect of long-term creep on particle coarsening in tempered martensiteferritic steels’, Acta Metallurgica, 1989, 37, 3225–3234.

33 T. Gladman, The Physical Metallurgy of Microalloyed Steels, The Institute of Materials,London, 1997.

34 R.J. McElroy and Z.C. Szkopiak, ‘Dislocation-substructure-strengthening andmechanical–thermal treatment of metals’, Int. Metall. Rev., 1972, 17, 175–202.

35 O. Kosik, D.J. Abson and J.J. Jonas, ‘Strengthening effect of hot work subgrains atroom temperature’, J. Iron and Steel Inst., 1971, 209, (88), 624–629.

36 J. Čadek, Creep in Metallic Materials, Elsevier, Amsterdam 1988.37 P. Weinert, Modelling of the Creep Behaviour of Ferritic/Martensitic 9–12% Cr

Steels on a Microstructural Base, PhD Thesis, Graz University of Technology, 2001(in German).

38 G. Dimmler, P. Weinert and H. Cerjak, ‘Extrapolation of short-term creep rupturedata – The potential risk of over-estimation’, Proceedings Creep Conference, 12–14September 2005, London, 165–176.

39 G. Guntz, M. Julien, G. Kottmann, F. Pellicani, A. Pouilly and J.C. Vaillant, The T 91Book – Ferritic Tubes and Pipe for High Temperature Use in Boilers, Vallourec andMannesmann Tubes, France, 1991.

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329

11Grain boundaries in creep-resistant steels

R . G . F A U L K N E R , Loughborough University, UK

11.1 Introduction

The consequences of the grain boundary topology and the grain boundarystructure in creep-resistant steels have been given limited prominence inproportion to their importance. Much creep deformation takes place alonggrain boundaries and the importance of identifying the ideal grain boundariesfrom the viewpoint of grain boundary engineering is still not fully appreciated.Grain boundary precipitation is also valuable in reducing grain boundarymobility and thus stabilising the grain size during heat treatment and thermalexposure. Grain boundary engineering (Palumbo, 1998) is a phase developedin the last decade to encompass all attempts to engineer the grain boundarystructure to give a preferred boundary type. This type of grain boundary istypically a ∑3 twin-type boundary and it can lead to superior mechanicalproperties in alloys that have been tested. These alloys still remain in thecopper and nickel alloy sector. An example study on IN718 (see Table 11.1for composition) is discussed by Randle (2003), but very little work has beendone on steels.

Bearing in mind that creep properties are much more dominated by boundaryeffects, grain boundary engineering for creep strength in high alloy steels isrequired to consider much more than the crystallographic misorientationproperties of the boundaries. Precipitation which is so prevalent in all highalloy creep resistant steels can have substantial effects on creep strength(Sun et al., 1992). This is not only because of the roughening effect thatprecipitates must have on grain boundary sliding but also because theprecipitates will change the vacancy sink and emissive properties of theboundaries at high temperature.

Furthermore, substantial segregation effects are seen on grain boundaries(Faulkner, 1996). These can be controlled by heat treatment but are rarelytaken into consideration when considering the dynamics of creep deformationat high temperatures. One element that excites attention in this respect isboron. Boron is increasingly being accepted as playing a role in creep by

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being segregated to grain boundaries in both austenitic and ferritic steels(Abe, 2006; Williams et al., 1976) and, as a result, slowing down atomtransport mechanisms along the boundary plane.

Dynamic recrystallisation effects at grain boundaries have been observedin ferritic steels with the result that creep deformation around the grainboundaries is more dispersed and creep life lengthened. The final aspect ofgrain boundaries that can lead to creep property control is the grain sizeitself. Traditionally, fine grain sizes have been thought to encourage creepdeformation but nanocrystalline materials studies have opened up a newarena in which many of the accepted views on grain size effects on creep arebeing questioned. The central issue is that for many metallic nanocrystallinematerials the strength decreases with decreasing grain size. This contradictsthe Hall–Petch relationship but it does offer a weaker grain option in manymaterials that is the requirement for high creep strength materials. The originof the anomaly lies in the very high triple junction density in thesenanocrystalline materials (Suryanarayana et al., 1992).

This chapter will review studies that have attempted to take into considerationthe above-mentioned grain boundary effects in determining creep-strengthsin creep resistant steels. The discussion will be divided between ferritic (9–12%Cr) steels and austenitic (18%Cr–12%Ni) steels.

A last consideration must concern the control of chromium behaviour ongrain boundaries in high Cr (> 12%Cr) steels. The proliferation of chromiumcarbide on grain boundaries in these steels results in substantial Cr depletionwithin the sub-micrometre region close to the boundaries. These Cr depletedzones allow accelerated oxidation to take place from the surface along grainboundary planes. This process is called intergranular cracking and is thebasis of other accelerated corrosion effects seen in these steels, such asintergranular stress corrosion cracking (Faulkner et al., 2005 and irradiation-assisted stress corrosion cracking (Bruemmer, 1999). The subject is extensiveand occurs at high temperatures, but it is not directly connected with creepproperties. For this reason the subject will not be discussed further here butthe interested reader is referred to a typical review (Newman, 2001).

11.2 Ferritic steels

11.2.1 General

A typical ferritic high creep strength steel contains 9–12%Cr and additionsof B, W, Mo, C, V, N, Mn and Si. The alloys are described in more detailelsewhere in this book. The microstructure of such a steel is shownschematically in Fig. 11.1.

It can be seen that there are various grain boundary types. First the prioraustenite grain boundaries define the grain structure that existed before thealloy was quenched into the martensite regime. After the quench, martensite

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Grain boundaries in creep-resistant steels 331

is formed in small, sub-micrometre sized plates contained in packets of 10–20 plates. These plates are often called laths and the bundles of laths arereferred to as lath packets. Whilst the mechanical strengths of these lathboundaries and prior austenite boundaries are quite different, the ability ofall types of boundary to nucleate chromium carbide (M23C6) is reasonablysimilar. Also segregation of elements like B, Hf, P and Mo is observed tooccur to all types of boundary (Morgan et al., 1992) with equal intensity.

11.2.2 Grain boundary precipitation

Grain boundary precipitates observed in P92 alloy (see Table 11.1 forcomposition) are mainly M23C6, where M stands mainly for Cr with someMo, Fe and W (Czyrska-Filemonowicz et al., 2003; Hattestrand and Andren,2001). The grain boundary precipitates provide creep strength to the alloyand further contribute to creep strength by pinning grain boundaries so thatthe grain size is stabilised during high temperature exposure.

The experimental studies of grain boundary precipitates in this alloy havebeen performed by transmission electron microscopy on alloys that haveseen a variety of tempering treatments in the temperature range 750–760°Cfor up to 2 h, followed by ageing at temperatures between 550 and 700°C fortimes up to 10 000 h (about 1 year). These studies show that most of thegrain boundary precipitate is nucleated by the end of the tempering treatmentand that there is a regime of stable precipitate size and spacing out to about10 000 h at 650°C, after which coarsening occurs. It is at this stage that thecreep strength is lost and tertiary creep rapidly sets in. Although VN precipitateswithin the grains, the intergranular M23C6 plays a substantial role in controllingthe creep strength. The thermodynamics and kinetics of grain boundaryprecipitate evolution in ferritic steels has been modelled using a dedicated

11.1 Schematic view of high chromium ferritic steel microstructure.

10 micrometres

Lath boundary

Prior Austenite GB

Carbides,nitrides, Zphase, and

Laves phase

Lath packet boundary

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332

Table 11.1 Alloy compositions (wt%: Fe bal except where specified)

C Mn Si P S Cr Ni Mo V Al Nb W N B Ti Y

E911 0.12 0.35 0.19 0.007 0.003 9.10 0.22 1.00 0.23 0.006 0.069 0.98 0.07P92 0.11 0.35 0.19 0.008 0.003 8.96 0.06 0.47 0.20 0.07 1.84 0.05 0.001IN718 0.04 0.06 0.40 18.2 Bal 3.00 0.50 5.20 0.004MA957 14.0 0.3 1.0 0.25Eurofer 0.11 0.42 0.06 8.9 0.2 1.1 0.31.4914 0.11 0.35 0.45 11.3 0.7 0.5 0.3 0.25 0.029 0.00720/25 0.049 20.0 25.0 0.52PE 16 0.08 0.05 17.1 42.5 3.1 1.3 1.2304LN 0.027 1.46 0.72 0.012 0.012 18.8 10.5 0.06 0.177 0.002316LN 0.021 0.44 0.28 0.012 0.019 17.7 12.1 2.2 0.25 0.003

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Monte Carlo modelling method (Yin and Faulkner 2005). The predictions ofthe model have been shown to give good fits with the grain boundary precipitateexperimental data for P92 (See Fig. 11.2).

Various workers have used Avrami-DICTRA type relationships to forecastthe evolution of M23C6 (Golpayegani et al., 2003), and to particularlydemonstrate that the effect of B is to reduce the rate of coarsening in P91type compositions by slowing down the transport of Cr between the M23C6

particles during coarsening.

11.2.3 Grain boundary segregation

There are two main types of grain boundary segregation (Faulkner, 1996).The first is Gibbs equilibrium segregation, which is caused by misfitting

10–2 10–1 100 101 102 103 104 105

Time, t (h)(a)

10–2 10–1 100 101 102 103 104 105

Time, t (h)(b)

Eq

uiv

alen

t ci

rcle

rad

ius,

L (

nm

) 100

10

1

Eq

uiv

alen

t ci

rcle

rad

ius,

L (

nm

) 100

10

1

11.2 Precipitation kinetics in P92 at 600°C (a) and 650°C (b) (see Table11.1 for composition). Lines are model predictions (dotted,intragranular; solid, intergranular). Symbols are experimental data.

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impurity atoms seeing energy benefits from sitting on a disordered grainboundary plane as opposed to residing in the grain. Atoms diffuse under thisdriving force towards the boundary and the kinetics are controlled by thediffusion of impurity atoms in the ferritic steel matrix (McLean, 1957). Thenet result is that at high temperatures (> 600°C in ferritic steel) the energybenefit seen by the impurity becomes very limited, whereas at lowertemperatures, diffusion limits the process. Therefore there is an optimumtemperature where the most prevalent equilibrium segregation takes place.This is typically around 500°C in ferritic steels.

The second type of grain boundary segregation is non-equilibriumsegregation (Aust et al., 1967). This process occurs when the steel is rapidlyquenched from a temperature like the normalising or solution treatmenttemperature. In these circumstances, the grain boundaries are able to preservetheir equilibrium vacancy concentrations because of the good vacancy sinkefficiency of the grain boundary. On the other hand, in regions away fromthe boundaries, in the grain centres, the vacancy concentration is maintainedat high, non-equilibrium levels. Consequently, in quenched material, thereexists a positive vacancy concentration gradient towards the grain boundaryplane, down which the vacancies diffuse during subsequent heat treatmentafter the quench. It appears that in ferritic steels this gradient is similar forlath, lath packet and prior austenite grain boundaries. Impurity atoms withlarge misfit will be preferentially attracted to these vacancies diffusing towardsthe grain boundaries and will be dragged with them towards the grain boundary.This results in an accumulation of impurity on the boundary. It is a non-equilibrium process and so if the system is allowed to equilibrate by heatingat an intermediate temperature the segregating impurity atoms will diffuseback down their concentration gradients to even out the segregationconcentration profile surrounding the boundary. Another way of looking atthis is to represent the return to equilibrium conditions by very slow quenchrates. At the other end of the timescale, very fast quench rates will not allowsufficient time for the impurity drag mechanism to accumulate large amountsof impurity on the boundary. This leads to the conclusion that there is acritical quench rate at which the maximum non-equilibrium segregation tothe grain boundary is expected. This effect is displayed for Nb and Si in aGerman ferritic steel 1.4914 (see Table 11.1 for composition) in Fig. 11.3,showing the effect of quench rate and starting temperature on the impurityenrichment expected on the grain boundary (Faulkner, 1987).

11.2.4 Dynamic recrystallisation at grain boundaries

Abe (2006) has observed recrystallisation to occur near to prior austeniteboundaries in ferritic steels (see Fig. 11.4). He has ascribed this phenomenonto the onset of tertiary creep in such materials. Consequently, any elemental

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addition or prior working treatment that can be applied to postpone the onsetof this transformation will be valuable in improving the creep strength. B hasbeen recognised as one element that can do this because B additions to P91,

Si in 1.49141.61.41.21.00.80.60.40.2

15.0

14.5

14.0

13.5

13.0

Temperature (K × 100)

1.61.41.21.00.80.60.40.2

F

Cooling ra

te (s–1 × 1000)

12

34

56

78

9

Nb in 1.4914

7654321

15.0

14.5

14.0

13.5

13.0

Temperature (K × 100)

7

654321

Cooling ra

te (s–1 × 1000)

0.10.3

0.50.7

0.91.1

1.3

F ×

10

11.3 Extent of grain boundary segregation, F, as a function of startingtemperature and cooling rate for Si and Nb in 1.4914 (see Table 11.1for composition) ferritic/martensitic steel.

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together with some adjustments to the W/Mo ratios to create P92, haveresulted in improved creep strength. Presumably the effectiveness of theaddition element must be contained in its ability to restrict grain boundarymovement at high temperature. For this reason, elements and thermal treatmentswhich produce excessive precipitate or segregation on grain boundaries couldbe useful in increasing creep life.

11.2.5 Effect of boron

The ability of B to interact with diffusing atoms on grain boundary planessuch that grain boundary diffusion is curtailed is one of the attractive featuresof this element. Abe (2006) has shown that grain boundary diffusion rates inan Fe–9Cr–3Co–3W alloy with small amounts of nitrogen are reduced, withthe effect being manifested in a reduced M23X6 precipitation rate, and acorresponding increase in creep life is obtained. Boron additionally assists inretarding the formation of fine grain size regions in welds (Abe, 2006).These are responsible for type IV cracking and when B is added to the abovesteel, the occurrence of type IV cracking can be eliminated. This effect ispresumably connected with the influence of boron on the dynamicrecrystallisation in the boundary neighbourhood mentioned in the previoussection. To complete the boron story, the intragranular precipitation may be

1 µm

11.4. Dynamic recrystallisation in modified 9Cr–1Mo ferritic steeltested at 100 MPa and 600°C for 34 141 h. After Abe (2006).

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stabilised against coarsening by limitation of diffusion between coarseningparticles (Golpayegani et al., 2003).

11.2.6 Ductile to brittle transition temperature DBTT

One of the important distinctions between ferritic steels and austenitic steelsis the difference in their crystal structure. Ferritic steels possess the bodycentred cubic crystal structure, whereas the austenitic steels are face centredcubic. Body centred cubic systems have fewer slip systems available and soat lower temperatures their capability to deform plastically becomes severelylimited. This is represented on a curve of yield strength, σy, versus temperatureas shown in Fig. 11.5. Reviewing the temperature dependence of the cleavageor grain boundary failure strength indicates a line with a much lower slope,so that the picture for grain boundary fracture stress, óF, as opposed to matrixfracture is summarised as in Fig. 11.5.

If segregation, dynamic recrystallisation, or any change in grain boundarystructure occurs, this will act to reduce the level of the grain boundaryfracture line in Fig. 11.5. The temperature at which the crossover betweenmatrix yield and grain boundary or cleavage fracture is known as the ductileto brittle transition temperature (DBTT). Any change in the grain boundarystate which lowers the grain boundary line or any strengthening that increasesthe matrix strength will increase the DBTT. To explain more fully, an increasein the matrix yield stress, ∆σy, will cause a shift in the DBTT of ∆Ty, assuming

Str

ess

∆σy

Matrix yield stress

Grain boundaryfracture stress

∆σF

∆TF

∆Ty

Temperature

11.5 Stress-temperature diagram, showing the contributions ofcleavage, grain boundary fracture stress and matrix failure stress tothe ductile to brittle transition temperature (DBTT) shift. An increasein the matrix yield stress, ∆σy, will cause a shift in the DBTT of ∆Ty,assuming that the grain boundary fracture stress remains the same.If the grain boundary fracture stress decreases by ∆σF, while thematrix yield stress remains the same, then the DBTT shift will be ∆TF.If both the matrix yield stress and grain boundary fracture stresschange simultaneously, the total DBTT shift will be ∆Ty + ∆TF.

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that the grain boundary fracture stress remains the same. If the grain boundaryfracture stress decreases by ∆σF, while the matrix yield stress remains thesame, then the DBTT shift will be ∆TF. If both the matrix yield stress andgrain boundary fracture stress change simultaneously, the total DBTT shiftwill be ∆Ty + ∆TF. Materials possessing a low DBTT are more favourablebecause the grain boundary failure is commonly associated with fast,uncontrollable fracture. If the DBTT is above room temperature, it is generallyrecognised that the material will be unsafe. Therefore it is especially importantfor creep-resistant ferritic steels to possess a low DBTT as well as a highcreep strength.

The most important factor controlling the grain boundary strength line inFig. 11.5 is the segregation. Elements like S and P present in quantitiesgreater than about 0.01% combined in an ferritic alloy with no strong carbide-forming elements like Ta or Nb are likely to have low grain boundary strengths.The grain boundary structure is also important: high angle random grainboundaries being the worst. The total contributions from all these effectshave recently been calculated and summed in a theory of grain boundaryfracture (Faulkner, 2005). Thus, although DBTT is not a creep-dependentproperty, it is a property that must be taken into consideration for a creep-resistant ferritic steel and there are a set of criteria which will enable the steeldesigner to decide whether a chosen creep-resistant ferritic steel will be fitfor purpose in a creep application.

11.2.7 Hafnium

One of the success stories of modelling of creep-resistant ferritic steels is theforecast that was made about the beneficial properties of hafnium. The additionof this element in quantities of up to 1% has been shown to convey resistanceto intergranular segregation effects concerning phosphorus and chromium(Lu et al., 2006). Furthermore the Hf addition forms HfC preferentially tochromium carbide on the grain boundaries (and in the grains), thus replacinga relatively coarsening-prone phase by a very stable grain boundary precipitatephase. The initial work has been done on E911 (Yin and Faulkner, 2005 (seeTable 11.1 for composition). The net result is to prolong the time at whichthe microstructure will remain stable at a given temperature, thus producinglonger creep life or similar creep life at higher temperatures (see Fig. 11.6).Currently alloy development is underway to test the predictions of themodelling.

11.2.8 Oxide dispersion strengthening

Attempts to improve the strength of ferritic steels at high temperature usuallybecome unsuccessful when the temperature rises above 625°C. This is because

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the strengthening components in the alloys become insignificant above thistemperature. The poor oxidation resistance above this temperature shouldalso be borne in mind. However, if the application is in protected circumstances,like vacuum operation or if coatings are being considered, it should bepossible to produce a ferritic steel that will have good creep strength up to800°C. The means of achieving this objective are through oxide dispersionstrengthening.

The most common ferritic alloy is MA957, a 14%Cr ferritic steel (seeTable 11.1 for composition) with a 0.25 wt% yttria dispersion. More recentlythe international fusion reactor programme has developed a special ferritic9%Cr oxide dispersion strengthened (ODS) steel called Eurofer ODS (seeTable 11.1 for composition) with microstructure and properties similar to the9–12%Cr steels discussed so far in this chapter, but with creep resistance upto 800°C. The microstructure consists of the typical tempered martensitegrain structure with yttria particles situated on intragranular sites. Studiesare being made of the grain boundary structure in this material (Lu andFaulkner, 2007) and the field emission gun scanning electron microscope(FEGSEM) with electron back scatter diffraction (EBSD) is proving aninvaluable technique in determining the real grain size and the distributionof grain boundary misorientations as a function of processing parameters.For example, there is a clear effect of quench rate from the solution treatmenttemperature on the percentage of low angle boundaries seen in Eurofer97ODS (see Fig. 11.7). The prevalence of low angle boundaries provides moreductility in general yielding and this leads to lower DBTTs, which, as discussedearlier, is an obvious design advantage for these brands of ferritic steel.

0 200000 400000 600000 800000Time, t (h)

100

80

60

40

20

0

Str

ain

, ε (

%)

50 years

M2N VN HfC

11.6 Model predictions for creep strain rate using various phases asparticle strengtheners. Microstructural predictions have been linkedwith the continuum damage mechanics models to produce thisfigure. Alloy P91 at 600°C.

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30 µm

(a)

11.7 (a) EBSD picture and grain boundary misorientation plot ofEurofer97 ODS as a function of cooling rate from the solutiontreatment temperature. Note the lack of low angle boundaries inmaterial that has been furnace cooled (FC).

WQACFC

Nu

mb

er f

ract

ion

0.25

0.20

0.15

0.10

0.05

0.0010 20 30 40 50 60

Misorientation angle (degrees)(b)

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11.3 Austenitic steels

11.3.1 General

Referring to Fig. 11.1, the main difference between the microstructures ofaustenitic and ferritic steels is that the remanents of the martensitictransformation are absent. This is because the Ni stabilises the austenitephase at all temperatures. The only grain boundaries are the so-called prioraustenite boundaries of the ferritic microstructure. Other differences betweenthe two types of steel are that austenitic steels possess a face centred cubiccrystal structure and this results in a more compact lattice than that of theferritic steel. The consequences of this are reduced diffusion rates and thermalconductivity, and higher thermal expansion coefficients. The higher solubilityof carbon in austenite means that there is a greater potential for formingcarbides but this is offset because the coarsening rates of such carbides aremuch greater than those contained within a ferritric matrix. Consequentlycarbide strengthening for creep purposes must rely on other phase strengtheningmechanisms than carbide precipitation. Nevertheless the precipitation hardeningcomponent of creep life is still high because of the lower diffusion rates andaustenitic steels can usually operate in creep environments with highertemperatures than those typically allowed for ferritic steels.

11.3.2 Precipitation

The main type of grain boundary precipitate in austenitic steels is M23C6.TiC and NbC can form in higher temperature windows than those typical ofservice life for the austenitics. TiC will predominate on grain boundaries inthe 800–1000°C regime, and at temperatures from 1000–1300°C, NbC willpreferentially form. A typical set of isothermal precipitation curves fo NbCand M23C6 in a typical stainless steel based on the austenitic composition(see Table 11.1) are shown in Fig. 11.8 (after Faulkner, 1979).

These results come from earlier modelling exercises based on analyticalcalculations of grain boundary precipitate behaviour (Carolan and Faulkner,1988). Examples of grain boundary precipitation in austenitic steel are givenin Fig. 11.9.

11.3.3 Segregation

The austenitic steel composition has probably received more attention in relationto intergranular segregation than any other except perhaps Cu–Bi. Studies ofboron segregation have been explored theoretically and autoradiographicallyby Williams et al. (1976) and by atom probe field ion microscopy by Karlssonand Norden (1988). Boron clearly segregates and it is thought that the mainmechanism is non-equilibrium segregation. This is demonstrated by a series

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TheoryExperimental points fromAdamson, 1972

10 100 1000 10 000Time (s)

M23C6(0.1 vol%C)

(0.05 vol% C)

M23C6(0.01 vol%C)

NbC (0.01 vol%C)

NbC (0.1 vol%C)

Precipitate size = 2.5 × 100nm diameter

1450

1350

1250

1150

1050

950

Tem

per

atu

re (

K)

Tem

per

atu

re (

°C)

1300

1200

1100

1000

900

800

700

600

11.8 Isothermal precipitation curves for 20C/25Ni austenitic steel (seeTable 11.1 for composition), showing the ranges of temperature overwhich NbC and M23C6 are expected.

of autoradiographs showing B grain boundary segregation as a function ofstarting temperature and cooling rate (Fig. 11.10). The trends shown stronglysupport the proposed mechanism of non-equilibrium segregation discussed inthe segregation section applied to ferritic steels in that the segregation enrichmentpasses through a maximum at an intermediate cooling rate.

Another interesting grain boundary effect seen in austenitic steels concernsthe segregation behaviour of chromium. If sensitised in the 600–700°Cregime then chromium depletion will occur caused by the grain boundarycarbide precipitation. If air cooled and not subsequently aged, chromiumsegregates to the boundaries (Flewitt and Vorlicek, 1993). The dependenceof the depletion effect on grain boundary structure has been most intensivelystudied by Laws and Goodhew, (1991), who looked at 51 boundaries in thetransmission electron microscope after sensitising Type 316 steel. Generallythose boundaries with low fitting, high ∑ values, lead to wide depletionzones and, generally, lower minimum chromium concentrations ∑ =3,11,13a,13b and 29a have narrow depletion zones and higher minimumboundary chromium concentrations. The depletion is also correlated with theextent of chromium carbide precipitation on the boundary.

11.3.4 Boron, hafnium and zirconium

There is evidence (Hattestrand and Andren, 2001) that creep-resistantaustenitic steels like Type 316 (see Table 11.1 for an alloy of similar

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0.5 µm

(a)

(b)

11.9 Examples of intergranular precipitation in austenitic steels. (a)M23C6 and M6C in Type 304 LN (see Table 11.1 for composition),after 28 840 h at 650°C (after Vodarek et al., 1998); (b) TiC in NimonicPE16 (see Table 11.1 for composition), after 1 min at 825°C (afterFaulkner and Caisley, 1977).

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344

Ti = 1350°C Ti = 1250°C Ti = 1150°C

Ti = 1200°C Ti = 1200°C

(b) (c)(a)

(d) (e)

11.10 Autoradiographs of 316 LN steel (see Table 11.1 for composition) showing the effects of solution treatmenttemperature (a) 1350°C, (b) 1250°C, (c) 1150°C, and cooling rate from the solution treatment temperature of 1200°C at (d)500 s–1 and (e) 5000 s–1 on B intergranular segregation.

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composition) are improved if B is added. The reason for this is that thediffusion of chromium between grain boundary carbides is limited by thepresence of boron in the neighbourhood of the boundaries. This reduces thecoarsening rate with a resultant improvement in long-term creep life. Heavyelements like Zr and Hf are thought to limit grain boundary sliding duringhigh temperature deformation. The most evidence for this is found insuperalloy compositions.

11.4 Grain boundary properties and constitutive

creep design equations

Constitutive equations exist for forecasting the damage accumulation in avariety of forms, and features connected with the grain boundary propertiesare extremely important parameters. The creep rupture life is usually calculatedon the basis of the Hull–Rimmer equation, with adjustments made by Rajand Ashby (1975). These models calculate the vacancy flux to grain boundariesand elucidate how rapidly a given grain size structure will cavitate at inclusionsto the point, or time tr, that cavity linkage will occur and the material will failby grain boundary decohesion. The Raj and Ashby equation is given by:

t kTD

F

F

Af AA

A

rb

3/2V

B3/2= 3

32

( )

( )

d( )

min

maxπδ σ ρ

ααΩ ∞ ∫ [11.1]

where σ∞ is the applied tensile stress, δDb is the grain boundary diffusivity,k is Boltzmann’s constant, T is the absolute temperature, α is the angle thatdefines the shape of the cavity (the angle between the tangent to the cavityand the grain boundary plane), FB is the shape factor determining the area ofvoid at the inclusion–matrix interface, FV is a similar shape factor determiningthe volume of the void adjacent to the inclusion, Ω is the molar volume andρ is the number of voids per unit area of boundary.

The problem is in determining the integral involving the fraction of cavitatedboundaries needed to cause fracture in the Raj expression, equation 11.1. Itis found to be dependent upon whether the supply of vacancies is constant orlimited. Therefore a more understandable form of predictive equation remainsthe Hull–Rimmer formulation:

tkT

Dr

3

b F=

ηδ σΩ

[11.2]

where k is the Boltzmann constant, T is the absolute temperature, η is theinterparticle or cavity nucleation site spacing on the grain boundary, Ω is theatomic volume, δ is the grain boundary width, Db is the grain boundary self-diffusion coefficient and σF is the applied stress. In both of the approaches,

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the grain boundary properties like diffusion, width and grain boundary nucleantspacing are all central to the time to rupture problem solution.

Furthermore, grain boundaries feature strongly in damage-based parametricequations for secondary creep rate, dε/dt, where ε is strain and t is time.Traditionally a Monkman–Grant type analysis is used. This can be used torelate the minimum secondary creep rate to the rupture life. Abe (2006) usesan expression of the following form:

t r

min

= constd ln

d˙ ˙ε ε

ε

[11.3]

where ε min is the minimum creep rate and the differential is the accelerationof creep rate in the acceleration creep regime.

There are several versions of the equation depending upon which regimeof temperature–stress space is occupied. The grain boundary propertiesdominate in the high temperature–low stress domain. When in this regionCoble creep operates and the working equation commonly used is as follows:

ε σmin

4

3 = ADbkTd

[11.4]

where

D DDD

= 1 + dl

b

l

πδ

[11.5]

where Dl is the lattice diffusion coefficient and Db is the grain boundarydiffusion coefficient. As before, δ is the grain boundary width, b is theoperational dislocation Burgers vector, σ is the applied stress, k is the Boltzmannconstant, T is the absolute temperature and d is the grain size. Clearly thegrain size and grain boundary diffusion coefficient must be known in orderto forecast properly the minimum creep rate in operating conditions whereCoble creep is dominant.

11.5 Future trends

As mentioned earlier it is normally accepted that fine grain sizes lead to poorcreep properties because there is more opportunity for grain boundary slidingand cavity nucleation. This is the normally accepted reason for making singlecrystal superalloys for high creep strength applications in gas turbines. Thereis however evidence that the normal grain boundary controlled creep behaviouris altered when very small, nanosized grains are present. Under thesecircumstances the nature of the grain boundary becomes much less planarand the intergrain boundary spacing becomes less than the average dislocationpile-up length. Both of these factors could have a large effect on the ability

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of the materials to deform around the boundary plane. But there is in principlea larger effect concerned with the very high density of triple junctions expectedin nanocrystalline materials. Galina et al. (1987), have shown that when thetopological class of boundaries, according to the Von Neumann–Mullinscriterion, is less than 6, the triple junction mobility reaches a rather lowequilibrium state, implying that there is a strong limit to grain growth in agrain structure characterised by the n<6 criterion when the density of triplejunctions is high, that is, in a nanocrystalline material. This means that manynanocrystalline materials could possess the kind of microstructure that wouldcontain triple junction drag up to high temperatures and thus produce highercreep strength. Ultra-fine grain size steels are currently being explored fromthis viewpoint.

Reference has already been made to the effects of Hf on the potentialcreep properties of ferritic high creep strength alloys. Further work is inprogress exploring the beneficial properties of Hf in austenitic steels andsuperalloys. In all these types of materials, Hf seems to slow down migrationkinetics in the neighbourhood of the boundary and to produce more stablecarbide and nitride phases on the prior austenite and lath boundaries. Both ofthese characteristics lead to higher creep strengths. The removal of chromium-rich phases and substitution of hafnium-rich phases on the grain boundariesfurthermore reduces chromium depletion effects that commonly lead toaccelerated intergranular oxidation at high temperatures. This means that Hfwill reduce the intergranular corrosion susceptibility in many austenitic steels.

Finally, the modelling of grain boundaries and associated precipitation,deformation and segregation effects is set to undergo a considerabletransformation in the near future. Molecular dynamics (MD) modelling,which goes back to first principles associated with the interatomic potentials,is becoming more powerful with the advent of parallel processing computersand the use of tricks to extend the timescale over which MD calculationsoperate. This means that atom movements will be simulated from first principlescalculations to predict the approach of a non-equilibrium material to itsequilibrium state. This, with a superimposed stress, is effectively a descriptionof the creep process. It is anticipated that soon all analytical equation approachesto describing creep will be superseded by the MD simulation method. It isalready being used successfully to predict the effects of neutron irradiationdamage and dislocation–obstacle interaction. High temperature microstructuralevolution and the associated creep property prediction could well be the nextmajor application of the MD method.

11.6 References

Abe F. (2006), ‘Metallurgy of long term stabilisation of ferritic steels for thick sectionboiler components in USC power plant at 650°C’, Proceedings Conference on ‘Materials

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for Advanced Power Engineering, J. Lecompte-Beckers, M. Carton, F. Schubert andP.J. Ennis (eds), Scriften des Forschungszentrum Juhlich, Germany, 965.

Adamson J.P. (1972), Precipitation Behaviour in Austenitic Steels, D.Phil Thesis, OxfordUniversity.

Aust K.T., Armijo S.J., Koch E.F. and Westbrook J.A. (1967), ‘Vacancy-driven grainboundary segregation’, Trans. ASM, 60, 360.

Bruemmer S.M. (1999), ‘Grain boundary composition and its effects on environmentaldegradation’, Mater. Sci. Forum., 294–296, 75.

Carolan R.A. and Faulkner R.G. (1988), ‘Grain boundary precipitation in alloy 800’,Acta. Metall, 36, 257.

Czyrska-Filemonowicz A., Bryla K., Spiradek-Hahn K., Firganek H., Zieliunska-Lipiec A.and Ennis P.J. (2003), ‘Role of boron in 9% cr steels for steam power plant’, ProceedingsParsons 2003: Engineering Issues in Turbine Machinery, Power Plant and Renewables,Strang A., Conroy R.D., Banks W.M., Blackler M., Leggett J., McColvin G.M., SimpsonS., Smith M., Starr F. and Vanstone R.W. (eds), Maney, London, p. 365.

Faulkner R.G. (1979), ‘Inter-granular precipitation in austenitic alloys’, J. Mater. Sci.Lett., 14, 2249.

Faulkner R.G. (1987), ‘Combined grain boundary equilibrium and non-equilibriumsegregation in ferritic-martensitic steels’, Acta. Metall, 35, 2905.

Faulkner R.G. (1996), ‘Segregation to boundaries and interfaces in solids’, Int. Mater.Rev., 41, 198.

Faulkner R.G. (2005), ‘Grain boundary segregation and fracture’, Zeitschrift fur Metallk,96, 1213.

Faulkner R.G. and Caisley J. (1977) ‘Kinetics of grain boundary precipitation in nimonicPE16’, Metal Sci. J, 11, 200.

Faulkner R.G., Yin Y., Cintas J. and Montes J.M. (2005), ‘Modelling and experimentalstudies of inter-granular corrosion in austenitic steels used in light water reactorsystems’, 12th Conference on Environmental Degradation of Materials in NuclearPower Systems, Allen T., King P. and Nelson L. (eds), Salt Lake City, TMS publication,Warrendale, PA, USA p. 135.

Flewitt P.E.J. and Vorlicek V. (1993), ‘Cooling induced segregation of impurity elementsto grain boundaries in Fe–3Ni alloys, 2 1/4Cr 1Mo steel, and submerged arc weldmetal’, Nuclear Electric Report, TD/SEB/REP/2050/93.

Galina A.V., Fradkov V.E. and Shvindlerman L.S. (1987), ‘Grain boundary mobility andits relation to topological class’, Phys. Met. Metalloved., 63, 165.

Golpayegani A., Hattestrand M. and Andren H-O. (2003), ‘Effect of boron on precipitation,growth and coarsening in martensitic Cr steels’, Proceedings Parsons 2003, StrangA., Conroy R.D., Banks W.M., Blackler M., Leggett J., McColvin G.M., Simpson S.,Smith M., Starr F. and Vanstone R.W. (eds), Maney, London, p. 347.

Hattestrand M. and Andren H-O. (2001), ‘Influence of strain on precipitation reactionsduring creep of an advanced 9% Cr steel’, Acta Materialia, 49, 2123.

Karlsson L. and Norden H. (1988), ‘Non-equilibrium segregation of boron in austeniticsteels, Acta. Metall., 36, 13.

Laws M.S. and Goodhew P.J. (1991), ‘Grain boundary structure and Cr segregation in a316 stainless steel’, Acta. Metall., 39, 1525.

Lu Z. and Faulkner R.G. (2008), to be published in J. Nucl. Mater.Lu Z., Faulkner R.G., Sakaguchi N., Kinoshita H., Takahashi H. and Flewitt P.E.J. (2006),

‘Effect of hafnium on radiation-induced segregation in ferritic steel’, J. Nuclear Mater.,351, 155.

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McLean D. (1957), Grain Boundaries in Metals, Oxford University Press, p.131.T.S. Morgan, E.A. Little, R.G. Faulkner and J.M. Titchmarsh (1992), ‘Interfacial segregation

in fast reactor-irradiated 12% Cr martensitic steel’, Effects of Radiation on Materials:15th International Symposium, Nashville, Stoller R.E., Kumar A.S. and Gelles D.S.(eds), (1992) ASTM STP 1125, p. 633.

Newman R.C. (2001), ‘Environmentally assisted corrosion’, Corrosion, 57, 1030–1041.Palumbo G., LeHockey E.M. and Lin P. (1998), ‘Grain boundary engineering’, J. Metals,

50, 40.Raj R. and Ashby M.F. (1975), ‘Inter-granular fracture at elevated temperature’, Acta.

Metall, 23, 653.Randal V. (2003), ‘Grain boundary engineering in IN718’, Scripta Metall., 44, 2789.Sun W.P., Militzer M. and Jonas J.J. (1992), ‘Grain boundary precipitation and high

temperature deformation in steels’, Metall. Trans., 23A, 821.Suryaranayana C., Mukohopadhyay D., Patankar S.N. and Froes F.H. (1992), ‘The Hall-

Petch relationship in nanocrystalline materials’, J. Mater. Res., 7, 25.Vodarek V., Sobotkova M. and Sobotkova J. (1998), ‘Effect of microstructural evolution

on the creep rupture behaviour of CrNi (Mo)N austenitic steels’, MicrostructuralStability of Creep Resistant Alloys for High Temperature Applications, A. Strang (ed.),Institute of Materials, London, p. 69.

Williams T.M., Stoneham A.M. and Harries D.R. (1976), ‘Grain boundary segregation ofboron in 316 steel’, Metal Sci., 10, 14.

Yin Y. and Faulkner R.G. (2005), ‘Creep damage and grain boundary precipitation inpower plant steels’, MST, 21, 1239.

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350

12Fracture mechanism map and fundamental

aspects of creep fracture

K . M A R U YA M A , Tohoku University, Japan

12.1 Introduction

Suppose a structural component made of steel is loaded at room temperature.The component deforms elastically or plastically depending on the stresslevel, but the deformation under constant load (or stress) always stops withina short period of time. At elevated temperature, on the other hand, thecomponent deforms continuously even under constant stress, and finallybreaks after a certain period of time. This is the consequence of creepdeformation and creep fracture. The time-dependent deformation and fractureare characteristics of high temperature deformation. Resistance to thedeformation and fracture is an important property to be considered in thestructural design of elevated temperature plants made of steel. The stress tocause fracture in 105 h is especially important, since it usually gives theallowable stress of steels. The value is found from short-term creep data byformulating creep data based on, for example, the Larson–Miller or Orr–Sherby–Dorn equations. The formulation should be done on the sound physicalbasis of creep fracture. This chapter deals with the physical basis of creepfracture.

The creep fracture mechanism often changes from transgranular fracturein short-term creep to intergranular fracture in long-term creep. The changesin fracture mechanisms are summarized in fracture mechanism maps. Thetransition of fracture modes is often accompanied by a decrease in stressexponent and activation energy, namely breakdown of creep strength. Sincethe change in activation energy causes serious problems in the prediction oflong-term properties, one should be careful about this change. This chapterwill explain why and how creep fracture mechanisms and creep ruptureproperties change with creep testing conditions. Several examples are providedof how the changes bring about overestimation of the long-term rupture life.A multi region analysis of creep rupture data is proposed to preventoverestimation.

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Fracture mechanism map and fundamental aspects 351

12.2 Fracture mechanisms and ductility of materials

A material exhibits several modes of fracture depending on creep testingconditions, namely stress and temperature. The fracture modes are classifiedinto three types as shown in Fig. 12.1:1 (a) brittle intergranular fracture, (b)intergranular or transgranular fracture with some ductility, and (c) ruptureafter 100% reduction in area. The fracture ductility changes substantiallywith the fracture mechanisms. There are two types of brittle intergranularcreep fracture (Fig.12.1 (a)). One is brought about by the coalescence ofcavities formed on grain boundaries aligned perpendicular to the appliedstress (see Section 12.3.1). The other is wedge cracking formed at triplegrain boundary junctions by stress concentration that is created by grainboundary sliding. The intragranular creep cavities are nucleated at inclusions(solid circles) as shown in the lower part of Fig. 12.1 (b) and their coalescenceby deformation brings about the final fracture. Recrystallization occurs inthe necking part during the rupture (Fig. 12.1 (c)).

Ashby and his co-workers have compiled fracture mechanism maps andFig. 12.21 provides an example. The map can foretell the fracture modeoperative under a given creep condition. The ductile fracture appearing atthe high stress is typical of tensile deformation at room temperature. In Fig.12.2, testing temperature is given as a fraction of melting temperature Tm.The rupture takes place at very high temperature only. Structural materials

Brittle Ductile

σ

σ

(a) (b) (c)

12.1 Representative fracture modes at high temperature. (a) Brittleintergranular fracture by wedge cracking and diffusional cavitygrowth, (b) intergranular and transgranular cavity growth due toplastic deformation and (c) rupture after 100% reduction of area. Thesolid circles in (b) represent inclusions.

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are usually used around 50% of their melting temperature, and transgranularcreep fracture and intergranular creep fracture caused by the coalescence ofcavities are important mechanisms under such creep conditions. The formerappears at higher stress (shorter term creep) and the latter at lower stress(longer term creep). Wedge cracking is sometimes observed in the transitionregion between the two fracture mechanisms.

12.3 Stress and temperature dependence of rupture

life

12.3.1 Cavity growth controlled by grain boundarydiffusion

The rupture life tr of a nickel based alloy is plotted in Fig. 12.2 as a functionof creep stress σ. The rupture life is expressed as:

tr = t0 σ–n exp (Q/RT) [12.1]

where t0 is a material constant, n is the stress exponent, Q is the activationenergy for rupture life, R is the universal gas constant and T is the absolutetemperature. As evident in the figure, the stress exponent (n = dlntr/dlnσ)and activation energy for rupture life change with the fracture modes: fromhigher values of the transgranular fracture to low values of the intergranular

100 102 104 106 108 1010

Rupture life, tr (s)

Ductile fracture 0.50 Tm Nimonic 80A

Wedgecrack

Trans–granularfracture

RuptureIntergranular

fracture

0.78

Cavity

0.66

W + C 0.620.59

0.53

No

rmal

ized

str

ess,

σ/ E

10–2

10–3

10–4

12.2 Fracture mechanism map of Nimonic 80A. Creep stress σ isnormalized with Young’s modulus E. Testing temperatures are givenas a fraction of melting temperature Tm.

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fracture controlled by diffusional cavity growth. The different lengths of thehorizontal arrows drawn in the figure confirm the decrease in activationenergy. Section 12.4 will examine how the fracture modes change and whythe change brings about decreases in the stress exponent and activationenergy. This section is focused on the stress and temperature dependence ofrupture life, and provides physical bases for the discussion in Section 12.4.

A typical process of intergranular fracture consists of three stages: (1)nucleation of cavities at grain boundaries, (2) their growth and coalescence,resulting in a facet crack of single grain size, and (3) propagation of thecrack to the final fracture. The rupture life tr of a material can be expressedas:

tr = tn + tg + tp ≅ tg [12.2]

where tn, tg and tp are the time taken for cavity nucleation, cavity growth andcrack propagation, respectively. The second term is usually the longest increep of heat-resistant steels and determines their rupture lives. Supposecreep cavities are nucleated on a grain boundary at an interspace of 2S (seeFig. 12.3), then the duration tg of cavity coalescence is given by:

t r t rr

S

g–1= (d /d ) d

0∫ [12.3]

where r0 is the initial radius, r is the current radius of curvature of cavitiesand (dr/dt) is the growth rate of cavities. The cavity diameter is usually closeto 2r. One can easily calculate rupture life, when (dr/dt) is known. Severalformulae for the growth rate have been proposed to calculate tg and anexample is given below.

σ

σ

2S

Cavity

Cavitysurface

diffusion

Grainboundarydiffusion

12.3 Vacancy flow paths during diffusional growth of cavities ongrain boundaries aligned perpendicular to the applied stressdirection.

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Suppose we have cavities on a grain boundary aligned perpendicular tothe applied stress (see Fig. 12.3). Concentrations of vacancies on the grainboundary and on the cavity surface with radius of curvature of r are given by:

Cgb = Cv exp(Ω σ/k T) [12.4]

Ccavity = Cv exp(2Γs Ω/r k T) [12.5]

where Cv is the concentration of vacancies in matrix, Ω is the volume of avacancy, Γs is the surface energy and k is Boltzmann’s constant. Since σ tobe used in Equation [12.4] is the stress component normal to the grainboundary plane, vacancy concentration is highest on the boundaries alignedperpendicular to the applied stress. Therefore, creep cavities are more frequentlyformed on grain boundaries aligned normal to the applied stress direction.When

σ > 2 Γs/r [12.6]

vacancies flow from grain boundary (high vacancy concentration of Cgb) to cavitysurface (low concentration of Ccavity) by diffusion through the grain boundaryand then the cavity surface, resulting in cavity growth. When grain boundarydiffusion controls the cavity growth rate, the growth rate is expressed as:2

dr/dt ∝ (Dgb δgb/r2)(Ω σ/kT) [12.7]

where Dgb is the grain boundary diffusion coefficient and δgb is the thicknessof the grain boundary zone. Substituting this equation into Equation [12.3],one can obtain the following equation describing stress and temperaturedependence of rupture life:

tr = tg ∝ (S3/Dogb δgb)(k T/Ω σ) exp(Qgb/R T) [12.8]

where Dogb is a material constant and Qgb is the activation energy for grainboundary diffusion. This equation predicts a stress exponent of n = 1 and anactivation energy of Q = Qgb.

12.3.2 Cavity growth controlled by cavity surface diffusionand by creep deformation

Growth rates of grain boundary cavities have been formulated in other cases:cavity growth controlled by cavity surface diffusion and by creep deformation.Rupture lives of these cases can easily be derived by adopting a similarprocedure to the one mentioned in Section 12.3.1. When the cavity surfacediffusion controls cavity growth, the rupture life is given by:2,3

tr ∝ (S Γ s2 /Dos δs)(k T/Ω σ n) exp(Qs/R T) [12.9]

where Dos is a material constant, δs is the thickness of surface zone and Qs

is the activation energy for surface diffusion. The stress exponent n for

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rupture life is 1.5 or 3 depending on the detailed assumptions about thecavity growth process.

At higher stress or higher temperature, creep deformation of the matrixsurrounding a cavity controls the growth rate of the cavity in grain boundaries.In such a case, the rupture life is given by:4

tr ∝ (1/ ε ) ∝ (1/Dol σ n) exp(Ql/RT) [12.10]

where ε is the creep rate of the matrix, Dol is a material constant and Ql isthe activation energy for lattice diffusion. The stress exponent n is greaterthan 3. The value of Ql is about twice those of Qgb and Qs. The growth rateof the wedge crack at triple grain boundary junctions is also controlled bycreep deformation in the interior of grains and their rupture lives take thesame form of Equation [12.10]. The same equation is applicable to transgranularfracture.

12.4 Fracture mechanism maps

As mentioned in Section 12.3, the growth rate of creep cavities is controlledby (1) grain boundary diffusion (2) cavity surface diffusion or (3) creepdeformation in the body surrounding cavities. The stress and temperaturedependences of the rupture life differ among cavity growth mechanisms asrepresented by Equations [12.8]–[12.10]. It can be postulated that the fastestmechanism among them, namely the one giving the shortest rupture lifegoverns creep fracture of a material under a given creep condition. Thisassumption provides the stress and temperature dependences of rupture lifedepicted in Fig. 12.4 (a) and (b), respectively. The cavity growth mechanismwith low stress exponent (cavity growth controlled by grain boundary diffusion,n = 1) appears at low stress, whereas the cavity growth mechanisms controlledby creep deformation within grains (cavity growth within grains (transgranularfracture), wedge cracking and grain boundary cavity growth controlled bycreep deformation, n > 3) are dominant at high stress. The cavity growthcontrolled by cavity surface diffusion takes place in the medium stress range,and may be absent at high temperature.5 Since the activation energies forgrain boundary diffusion Qgb and surface diffusion Qs are about a half thatof lattice diffusion Ql, the cavity growth mechanisms controlled by theseshort-circuit diffusion are operative at low temperatures.

As aforementioned, creep fracture theories predict a decrease in the stressexponent with decreasing stress (Fig. 12.4 (a)). The transgranular fracture(with a large n value) and intergranular fracture (with a low n value) shouldoperate at high stress and low stress, respectively. The fracture mechanismmap shown in Fig. 12.2 clearly confirms these predictions. Under a givenstress, the fracture mechanism changes from transgranular fracture tointergranular fracture with decreasing testing temperature. The fracture

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mechanism change is accompanied by a decrease in the activation energyfrom Ql to Qgb or Qs (Fig. 12.4 (b)). The different length of the arrows in Fig.12.2 confirms the change in activation energy.

12.5 Influence of fracture mechanism change on

creep rupture strength

A set of creep rupture data for type 316 stainless steel is given in Fig. 12.5,6

together with its fracture mechanism map. Three regions with different n and

log

σ

log tr(a)

1

3–1.5

n Grain boundary cavity growth

By cavitysurface

diffusion

By grainboundarydiffusion

GB cavity growth bycreep deformation

Wedge cracking

Transgranularfracture

1/T(b)

Grain boundary cavitygrowth

By cavity surfacediffusion

By grain boundarydiffusion

GB cavity growth bycreep deformation

Wedge cracking

Transgranularfracture

log

tr

Q1

Qs or Qgb

12.4 (a) Stress and (b) temperature dependence of rupture life. Thestress exponent and the activation energy are different depending onthe fracture mechanisms that are operative.

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Q values appear in the figure, and the dash–dot lines are the boundariesbetween the three regions: n = 9.7 and Q = 440 kJ mol–1 in the high stressregion H, n = 5.8 and Q = 390 kJ mol–1 in the medium stress region M andn = 2.1 and Q = 280 kJ mol–1 in the low stress region L. In practical materialsprecipitates are often the major strengthener and the amount of precipitateusually decreases with increasing temperature. Therefore, the apparentactivation energy for creep rupture life is often greater than the one predictedby creep fracture theories. This is the case in type 316 stainless steel. Theactivation energy for lattice self-diffusion in γ iron is Q1 = 285 kJ mol–1. Thevalue of Q in region L (280 kJ mol–1) can be below Q1, if the apparentcontribution of the microstructural change is eliminated. In Fig. 12.5 thestress exponent and the activation energy for rupture life decrease withincreasing rupture life. This fact is consistent with the fracture mechanismmap discussed in Section 12.4. The stress exponent is less than three inregion L in addition to the low activation energy, suggesting intergranularfracture controlled by diffusional growth of grain boundary cavities. Thestress exponents in regions M and H are greater than three, suggestingtransgranular fracture or intergranular fracture caused by deformation controlledgrowth of grain boundary cavities.

Fracture mechanisms of the steel are indicated by the letters T, W, C andσ in Fig. 12.5. Transgranular fracture (T) occurs in the high stress region,

102 103 104 105

Rupture life (h)

Str

ess

(MP

a)

500

300

200

100

70

50

30

20

Type 316 steel600 °C

650 °C

700 °C

750 °C

H

W

TM

C

L

σ

12.5 Fracture mechanism map for type 316 stainless steel togetherwith its creep rupture data. T: transgranular fracture, W: wedgecracking, C: cavity nucleation at the grain boundary carbides,σ: cavity nucleation at the grain boundary σ phase.

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corresponding to region H. Wedge cracking (W, controlled by creepdeformation) and grain boundary cavity formation at carbides (C) correspondto the medium stress range M. Grain boundary cavities are formed by σphase particles and intergranular fracture occurs in the σ region. This regioncorresponds to the low stress region L. The changes in fracture modes coincidevery well with the changes in n and Q values explained above. Grain boundarycavity is usually formed at the intersection of a particle surface and a grainboundary, and σ phase particles are responsible for cavity formation and theconsequent appearance of region L. Figure 12.5 suggests that rupture lifecan be extended up to the extension of the dashed lines, if we can prevent theformation of the σ phase.

12.6 Influence of microstructural degradation on

creep rupture strength

Advanced high Cr ferritic steels for boiler applications are normalized andthen tempered at 750–800°C for a few hours and have a tempered martensiticlath structure. Figure 12.67 is an example of 9Cr–2W–0.4Mo–1Cu–VNbsteel tempered at 770°C for 2 h. Vickers hardness values were measured ina grip portion of crept specimens at room temperature and they are plotted inFig. 12.6 as a function of rupture life (in other words, aging time). Theconstant hardness for short-term exposure demonstrates that the microstructureis partly stabilized and remains as it is up to 4 × 103 h exposure at 650°C.Since the microstructure is not in equilibrium, agglomeration of precipitatesand recovery of the lath structure become evident after long-term exposure.A drop in hardness starts at 4 × 103 h at 650°C and 600 h at 700°C. Thesoftening curve can be expressed as:

ta = to exp(–Hv/Hvo) exp(Qa/R T) [12.11]

where ta is the aging time, to is a constant, Hv is the Vickers hardness, Hvo isa material constant, and Qa is the activation energy. The softening curvesprovides Qa = 295 kJ mol–1, being in agreement with the activation energyfor lattice diffusion in the temperature range of interest (in ferromagneticregion). The dashed curve drawn in Fig. 12.6 (a) was predicted at 600°Cbased on the activation energy.

Creep rupture data for the same steel are given in Fig. 12.6 (b).7 There arethree regions with a different stress exponent n and an activation energy Q:n = 15 and Q = 775 kJ mol–1 in the high stress and short-term region H1, n= 10 and Q is the same in the medium stress and short-term region H2, andn = 5.3 and Q = 565 kJ mol–1 in the low stress and long-term region L. Thedifferent lengths of the horizontal arrows confirms the decrease in activationenergy. The dash–dot line is the boundary between regions H and L. Thedecrease in n and Q namely the breakdown of creep strength in long-term

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creep, is a serious problem in advanced high Cr ferritic steels. Several causeshave been proposed to explain the breakdown: formation of the Z phase atthe expense of MX precipitates, enhanced recovery of the lath structurealong grain boundaries, a change from transgranular fracture to intergranularfracture, and so on. The downward arrows in the figure indicate the onset ofhardness drop. There is a close correlation between the hardness drop andthe breakdown of creep strength. A good correlation has been confirmed inseveral advanced high Cr ferritic steels.7 This fact suggests that themicrostructural degradation can be a major cause of breakdown and theaccompanying decrease in n and Q values. If we can postpone microstructuraldegradation (hardness drop) by some means, we may be able to improvecreep strength in region L up to the dashed lines in Fig. 12.6 (b).

12.7 Change in creep rupture properties at athermal

yield stress

As mentioned in Chapter 8, Section 8.4, instantaneous plastic deformationtakes place upon loading when a material is creep tested above the athermal

12.6 (a) Vickers hardness in the grip portion and (b) creep rupturedata for advanced high Cr ferritic steel (9Cr–2W–0.4Mo–1Cu–VNbsteel). Downward arrows: onset of hardness drop at eachtemperature.

600°C650°C700°C

600 °C

H1

H2L

100 101 102 103 104 105

Rupture life (h)

(a)

(b)

240

220

200

180

200

100

60

40

Har

dn

ess,

Hv

Str

ess

(MP

a)

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yield stress. Since instantaneous plastic deformation introduces dislocationsinto a specimen and alters its dislocation substructure, creep deformationbehavior above the athermal yield stress should be different from that belowit (see Chapter 8, Section 8.6). Creep tests are always performed below theathermal yield stress in the case of high strength steels such as advancedhigh Cr ferritic steels with a tempered martensitic lath structure. On the otherhand, the creep test conditions for austenitic stainless steels are often abovethe athermal yield stress, since the steels are creep-tested after solutiontreatment. A substantial amount of plastic deformation takes palce uponloading above the athermal yield stress. Structural materials are used belowthe athermal yield stress in engineering plants. We should know what thechanges in creep rupture properties are introduced at the athermal yieldstress.

In this section creep rupture data for 2.25Cr–1Mo steel are taken as anexample. The steel was normalized at 930°C and then tempered at 720°C. Itscreep deformation behavior is described in Chapter 8, Section 8.6, and creeprupture data for the steel are plotted in Fig. 12.7.8 The dotted line indicatesthe athermal yield stress level σa. The creep stress is normalized by Young’smodulus E, since σa is truly independent of the testing temperature whennormalized by E.9 Creep tests were conducted both above and below theathermal yield stress. Usually the slope of the log tr versus log(σ/E) curvedecreases with decreasing stress. However, it increases discontinuously belowσa, indicating threshold-like behavior. The same threshold-like behavior ofthe creep rate is recognized (see Fig. 8.10). When the curves determinedabove σa are extrapolated to lower stress (dashed line), the extrapolationunderestimates the creep rupture strength below σa. Therefore, we should

450°C475°C500°C

525°C550°C600°C650°C

σa

101 103 105

Rupture life (h)

No

rmal

ized

str

ess,

σ/1

0–3 E

4

2

1

0.6

0.4

0.2

0.1

12.7 Creep rupture life of normalized and tempered 2.25Cr–1Mosteel. Creep stress σ is normalized by Young’s modulus E. The dottedline represents the athermal yield stress of the steel.

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carry out accelerated creep tests below the athermal yield stress in ordercorrectly to evaluate long-term creep properties under service conditions.Despite the change in the stress exponent, the activation energy for rupturelife does not change in this case: Q = 360 kJ mol–1 both above and below σa.The value is close to the activation energy for lattice diffusion of α iron inthe ferromagnetic temperature range a little below the Curie temperature.

12.8 Multi-region analysis of creep rupture data

The long-term creep strength of a material is evaluated from short-termcreep data by means of a time–temperature parameter (TTP) method, such asthe Orr–Sherby–Dorn (OSD) and Larson–Miller parameters:10

log tr – Q log e/R T = f(σ) [12.12]

(log tr + C ) T = f(σ) [12.13]

where C is the Larson–Miller constant, f(σ) is a function of creep stress σ ande has its usual meaning. The parameters assume a linear relation between logtr and 1/T as shown in Fig. 12.8 (a). The conventional TTP analyses supposethat the temperature dependence of rupture life, in other words, the activationenergy Q for rupture life or the Larson–Miller constant C is unchangd in therupture data set. Therefore, they adopt the linear extrapolation and predict therupture life given by the open circles in Fig. 12.8 (a). However, the activation

T1 T2 T1 T21/T

Experimentaldata

ln t

r

σ3

σ2

σ1

′Q

R

QR

σ3

σ2

σ1

I

II

(a) (b)

12.8 (a) The basic assumption of the TTP methods and (b) thedecrease in activation energy from Q to Q’ sometimes observed inreal creep rupture data. The square symbols and the solid linesrepresent available experimental data, and circles and dotted linesare predictions.

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energy sometimes decreases to a smaller value of Q′ in long-term creep(Fig. 12.8 (b)), as demonstrated experimentally in Figs. 12.2, 12.5 and 12.6 (b).If this is the case, the open circles overestimate the true lives given by thesolid circles in Fig. 12.8 (b). Such overestimation is a serious problem inaustenitic stainless steels11 and advanced high Cr ferritic steels.12,13

Multi-region analysis of creep rupture data14 has been proposed to avoidoverestimation of long-term creep properties. In the analysis, a set of creeprupture data is divided into several data sets, so that the value of Q is uniquein each divided data set. Then conventional TTP analysis based on Equations[12.12] or [12.13] is applied to each divided data set. Creep rupture data forGr.122 steel are given in Fig. 12.9.12 The data set is divided into two regionsH and L at the boundary represented by the thin dash–dot curve. Multi-region analysis (OSD analysis of each divided data set) gives the regressioncurves represented by thick solid lines. The solid symbols in Fig. 12.9 werenot used in the regression analyses. The activation energies are QH = 710 kJmol–1 and QL = 370 kJ mol–1 in regions H and L, respectively.

Conventional TTP analysis assumes a unique Q value for the whole dataset. This conventional single regional analysis provides the regression curvesrepresented by the dashed lines. The slope of the straight lines in Fig. 12.9(b) corresponds to an activation energy of 620 kJ mol–1. The standard errorof estimate in terms of log tr is 0.05 in multi-region analysis and 0.15 insingle region analysis. In the multi-region analysis regression curves providea better fit to the data points than those of conventional single region analysis.In the multi-region analysis regression curves describe very well the trendsof data point in Fig. 12.9, and can predict the solid symbols adequately. Thisfact assures proper evaluation of long-term properties by multi-region analysis.On the other hand, the regression curve at 60 MPa in Fig. 12.9 (b) given byconventional single region analysis (dashed line) overestimates the solidsymbol measured at 650°C. This is the overestimation often made byconventional TTP analyses.

In the example given in Fig. 12.7 the activation energy does not obviouslychange up to 105 h. In this case one can correctly predict long-term creeprupture life from short-term data by conventional linear extrapolation.

12.9 Summary

Creep tests are carried out over a wide range of stress and temperatureconditions and the creep rupture behavior of a material changes with the testconditions. For example stress exponent n and activation energy Q for rupturelife decrease from high values for short-term creep to low values for long-term creep at low stress or low temperature. The change in Q values is oftenaccompanied by a change in the fracture mechanism from transgranularfracture in short-term creep to intergranular fracture in long-term creep.

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Str

ess

(MP

a)

400

200

100

60

40

20

101 102 103 104 105

Rupture life (h)(a)

650°C675°C700°C 750°C

L

H

600°C625°C

0.95 1 1.05 1.1 1.151000/T (K–1)

(b)

σ (MPa)40

6090

140Gr. 122

QH = 710kJ/mol

620kJ/mol

H

QL = 370kJ/mol L

Ru

ptu

re l

ife

(h)

105

104

103

102

101

12.9 (a) Stress versus rupture life plot and (b) temperaturedependence of rupture life of Gr.122 steel. Solid regression curves:multi-region analysis taking account of the change in activationenergy Q between regions H and L. Dashed regression curves: singleregion analysis neglecting the change in Q.

Creep fracture theories can explain rationally the change in fracture mechanismand the consequent decrease in Q and n values. Of course, these changesmay not appear in some materials within an appropriate range of test duration.Microstructural degradation during high temperature exposure is anotherprobable cause of the decrease of Q in high Cr ferritic steels.

The decreases in Q and n values result in a quick drop in creep strengthand are recognized as breakdown of creep strength. Breakdown is a seriousproblem in austenitic stainless steels and advanced high Cr ferritic steels.When the onset of intergranular fracture is a cause of the breakdown, one

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can achieve significant improvement in creep strength by preventing thefracture mode from changing. Stabilization of the microstructure is a properway to postpone breakdown caused by microstructural degradation.

We usually evaluate long-term creep rupture life by linear extrapolation ofa log tr versus 1/T relation assuming a constant Q value. The decrease inactivation energy is the major cause of the overestimation of long-term creepstrength evaluated by conventional time–temperature parameter methods. Multi-regional analysis of creep rupture data is necessary to avoid overestimation.

Creep rupture behavior above the athermal yield stress is different from thatbelow it. We should carry out accelerated creep tests below the athermal yieldstress to evaluate properly long-term creep strength under service conditions.

12.10 References

1 Ashby M F, Gandhi C and Taplin D M R (1979), ‘Fracture mechanism maps andtheir construction for F.C.C metals and alloys’, Acta Metall, 27, 699–729.

2 Nix W D, Yu K S, and Wang J S (1983), ‘The effects of segregation on the kineticsof intergranular cavity growth during creep condition’, Metall Trans, 14A, 563–70.

3 Chuang T J, Kagawa K I, Rice J R and Sills L B (1979), ‘Non-equilibrium modelsfor diffusive cavitation of grain interfaces’, Acta Metall, 27, 265–84.

4 Beere W and Speight M V (1978), ‘Creep cavitation by vacancy diffusion in plasticallydeforming solid’, Metal Sci, 12, 172–6.

5 Cocks A C F and Ashby M F (1982), ‘On creep fracture by void growth’, ProgMater Sci, 27, 189–244.

6 Nakakuki H, Maruyama K, Oikawa H and Yagi K (1995), ‘Collective evaluation oftemperature and stress dependence of creep rupture life in austenitic stainless steels’,Tetsu-to-Hagane, 81, 220–224.

7 Ghassemi Armaki H, Maruyama K, Yoshizawa M and Igarashi M (2007), ‘Predictionof breakdown transition of creep strength in advanced high Cr ferritic steels byhardness measurement of aged structures at high temperature’, Key Eng Mater, 345–346, 553–556.

8 Maruyama K, Kushima H and Watanabe T (1990), ‘Prediction of long term creepcurve and rupture life of 2.25Cr–1Mo steel’, ISIJ Int, 30, 817–22.

9 Frost H J, and Ashby M F (1982), Deformation Mechanism Maps, Pergamon Press,Oxford.

10 Viswanathan R (1989), Damage Mechanisms and Life Assessment of High TemperatureComponents, ASM International Metals Park.

11 Maruyama K, Ghassemi Armaki H and Yoshimi K (2007), ‘Multiregion analysis ofcreep rupture data of 316 stainless steel’, Int J Press Vess Piping, 84, 171–176.

12 Maruyama K and Yoshimi K (2007), ‘Influence of data analysis method and allowablestress criterion on allowable stress of Gr.122 heat resistant steel’, Trans ASME, JPress Vess Technol, 129, (3), 449–453.

13 Maruyama K and Yoshimi K (2007), ‘Methodology of creep data analysis for advancedhigh Cr ferritic steel’, Proc of 8th int conf on Creep and fatigue at elevated temperatures,San Antonio, USA, 2007, ASME, New York, Paper No. 26510, 1–6.

14 Maruyama K, Baba E, Yokokawa K, Kushima H and Yagi K (1994), ‘Errors of creeprupture life extrapolated by time-temperature parameter methods’, Tetsu-to-Hagane,80, 336–341.

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365

13Mechanisms of creep deformation in steel

W. B L U M, University of Erlangen-Nuernberg, Germany

13.1 Introduction

The mechanisms of creep can be grouped into diffusive flow (matter transportby atomic diffusion), sliding on high-angle grain boundaries andcrystallographic slip. As the latter is dominant in producing strain undermost conditions of creep,1 we will focus attention on this mechanism.

Crystallographic slip is conveniently described by the motion of dislocations,representing the borderlines of slipped regions. Owing to their stress fields,dislocations interact with each other, with solute atoms and with particles offoreign phases. These interactions are the clue to understanding creepmechanisms. To a good approximation it is sufficient to describe themicrostructure by frequency distributions and the mean values of itscharacteristic spacings, in particular the spacings between grain and subgrainboundaries, free dislocations inside the subgrains, dislocations in low-angleboundaries, solute atoms and particles along dislocation lines. Even such asimplified view of the microstructure leads to considerable complexity. Itshould be taken before working on the lower scale of discrete dislocationswith much higher computational effort. While complementing observationson various length scales from the atomic to the macroscopic level are necessaryto obtain a consistent picture, the microstructural level with spatial averagesof characteristic spacings is indispensable to achieving a basic understandingof the observed creep behavior and providing a basis for a phenomenologicaldescription of creep properties.

The present treatment begins with the initial state of steels after productionunder typical conditions (Section 13.2). The fundamental difference betweenaustenitic and ferritic steels with regard to the solid state phase transformationfrom austenite to ferrite leads to significant differences between their initialstructures. These are reflected in the variation of creep resistance with strainin monotonic creep at constant stress as displayed in the creep rate–straincurves (Section 13.3). Section 13.4 summarizes information on creepmechanisms which can be obtained from the transient material response to

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sudden changes in stress. This information is also useful in understandingthe acceleration and deceleration of creep resulting from cyclic variation ofload (Section 13.5). In Section 13.6, observations on creep behavior presentedin the preceding sections are given a semi-quantitative microstructuralinterpretation. Section 13.7 reports on the quantitative description of creepby dislocation models coupling laws of deformation kinetics and structureevolution. Section 13.8 compares the model to in situ observations of dislocationprocesses during straining in a transmission electron microscope (TEM) atelevated temperature. The prospects for dislocation models in providing aquantitative fit of the observed macroscopic creep properties on amicrostructural basis and in guiding alloy development are briefly discussedin Section 13.9.

13.2 Initial microstructure


Austenitic steels do not undergo a phase transformation during their production.Their initial structure is typical of a recrystallized material with a low contentof dislocations within the grains. The high solubility of interstitial elementsin the matrix reduces the tendency for precipitation of carbonitrides. However,depending on the concentrations of the interstitial elements C and N, there isheterogeneous precipitation of carbonitrides, preferably at grain boundariesand dislocations.2,3 High contents of alloying elements like Cr and Ni in thematrix lead to solid solution strengthening.

13.2.2 Ferritic steels

During cooling from the melt, the matrix of ferritic steels transforms fromaustenite to ferrite. The transformation is accompanied by a strong reductionin solubility of foreign atoms. The martensitic transformation is a particularlyimportant case generating a class of tempered martensite steels. It comprisescooperative shearing of the crystal lattice (see e.g. Wayman4 and Haasen5)and thus causes intensive local deformation of the matrix, resulting in strongwork hardening due to a cellular dislocation structure with high dislocationdensity. Tempering of the martensitic structure leads to precipitation of soluteatoms and to recovery of the dislocation cell structure6–8 resulting in a subgrainstructure, characterized by the frequency distributions of misorientationsand of boundary spacings to be determined along test lines.

The subgrains are bounded by the boundaries of the prior austenite grains,of blocks of martensite laths of similar orientation, and of martensite lathsand of subgrains within the laths.9,10 The first two kinds of boundaries are ofthe high-angle type with misorientations across the boundaries generally

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Mechanisms of creep deformation in steel 367

lying above 10–15°;10–12 the latter ones are of the low-angle type with lowermisorientations. In contrast to low-angle boundaries constituting planardislocation networks, high-angle boundaries interrupt the coherency of thelattices of the neighboring crystallites, cannot in general be penetrated bydislocations, allow grains to slide relative to each other, and provide aparticularly effective short circuit path for diffusion of atoms.

The subgrain size w (mean linear intercept) yields the overall boundaryarea per volume, 2/w.13 If the subgrains are not equiaxed, one may differentiatebetween the mean subgrain intercepts w|| parallel and w⊥ perpendicular to thedirection of subgrain elongation. It was found that 2/w ≈ 1/w|| + 1/w⊥.14

Other methods of quantifying the subgrain size, for instance, from the meansubgrain area, lead to results which differ to a limited extent from the meansubgrain intercept length w, and vary in proportion to it.15 It is known that wapproaches a steady state value in the course of deformation which scales ininverse proportion to stress:

w∞ = kwb G/σ [13.1]

Here b is the length of the Burgers vector, G is the elastic shear modulusand kw is a numerical factor which for steels was reported to lie at about10.16–18 A typical value of the initial subgrain size w0 in tempered martensiteis 0.4 µm.6,17,19 According to Equation [13.1] this value corresponds to astress level 380 MPa, representing an estimate of the maximal local stresseswhich have acted in the steels during martensitic transformation.

Owing to the qualitative differences in the properties of high- and low-angle boundaries mentioned above, the areal fraction fhb = w/d of high-angleboundaries (d is the spacing of high-angle boundaries) in the subgrain structureof tempered martensites may be of importance. In the initial state fhb is toosmall to cause a significant influence of high-angle boundaries on the creepresistance. The situation may change, however, when the subgrain structurecoarsens during creep at low σ so that fhb increases significantly along with w.

The relatively low solubility of interstitial atoms in the ferritic grainsdrives precipitation of carbonitrides during and after martensitic transformation.Nucleation of most carbonitrides is heterogeneous and preferentially takesplace at boundaries of relatively high misorientation.6,20 These precipitatesgrow to large sizes dp corresponding to the low number of competing nuclei.Some carbonitrides, in particular those of type MX (where M are metal andX interstitial atoms), also precipitate more homogeneously inside the subgrains,probably owing to lattice coherence facilitating nucleation. They remaindistinctly smaller than the heterogeneous ones and therefore have a highstrengthening potential in spite of their low volume fraction fp. This followsfrom the fact the strengthening potential scales in inverse proportion to thespacing of precipitates along dislocation lines which in turn is proportional

to d fp p/ for an isotropic distribution of particles (see e.g. Reppich21).

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13.3 Creep at constant stress

Creep tests require a specimen to be heated to test temperature T and thestress to be raised from zero to the intended final value. During loading therate ε of mechanical (elastic plus inelastic) strain is relatively high (typically≈ 10–3 s–1). To sustain the applied stress the specimens work harden. Theinelastic strain accumulated during loading is similar to that measured intests at constant ε = 10–3 s–1 when the flow stress has reached the level ofcreep stress.

After loading the specimen creeps at constant stress. As the elastic strainremains constant, the creep strain is purely inelastic in nature. Continuingwork hardening at constant stress causes ε to decrease in the so-calledprimary stage of creep down to a minimal value εmin in the secondary creepstage. As metals in general do not fracture in uniaxial compression, a tertiarystage with ε increasing by tensile fracture, coupled with external and internalnecking, does not occur here. Investigation of compressive creep is thus ameans of identifying loss of creep resistance owing to global changes(coarsening) of the microstructure (see Section 13.3.2).


Being relatively soft in their coarse-grained, recrystallized initial state, austeniticsteels may display large inelastic loading strains in creep tests at elevatedtemperature. These may be suppressed by prior work hardening (increase indislocation density) in predeformation.22,23

Figure 13.1 displays a set of creep rate ε–strain ε curves for the alloy800H at 1073 K. The creep stress of 156 MPa exceeds the yield stress of≈ 100 MPa at ε ≈ 10–3 s–1 and induces a loading strain ≈ 0.03 which hardensthe material at an average rate ∆σ/∆ε ≈ 0.04 G, where G ≈ 50 GPa24 is theelastic shear modulus of alloy 800H at the test temperature. This rate issimilar to the work hardening rate M2 × G/300 ≈ 0.03G expected forpolycrystals (M = 3: the Taylor factor) from the work hardening rate ofsingle crystals at Stage II, resulting from storage of dislocations in the absenceof pronounced dynamic recovery.25,26 Subsequently the creep rate decreasessteeply at an average rate | d ln ε /dε| > 950. Using the stress exponent n =d log ε /d log σ ≈ ∆log ε /∆log σ ≈ 11 of the creep rate derived from Fig.13.4, the work hardening rate at the beginning of creep at constant σ isestimated as dσ/dε = σd ln σ/dε = (σ/n) d ln ε /d ε > 0.21 G, exceeding theformer one by a factor > 5. Within a subsequent small strain interval ≈ 0.01,the work hardening rate abruptly drops to low values near zero before workhardening at a low rate < 0.004 G recommences.

The sequence of an extraordinary relative maximum in work hardeningrate immediately followed by a relative minimum at the beginning of primary

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creep is even more pronounced for the austenitic Cr–Ni steel shown in Fig.13.2; here the minimal work hardening rate even becomes negative so that atransient period of work softening develops before work hardening isresumed. These features can hardly be explained in terms of storage andannihilation of dislocations alone, which in pure materials lead to acomparatively slow, monotonic transition to the steady state.27 They will bediscussed in Section 13.6.3 in terms of interaction between dislocations andforeign atoms.

Beyond the extrema of work hardening rate, primary creep proceeds at agradually decreasing rate and enters the secondary stage of creep at a minimumcreep rate εmin which is close to the rate of steady state creep, where workhardening and recovery of dislocations compensate and the dislocation structureis in a state of dynamic equilibrium. The tensile creep curves of Fig. 13.1 endwith an extended tertiary stage where creep is strongly accelerating at increasingrate d log ε /dε. As this stage is absent in compression, it must be related totensile fracture.

13.3.2 Ferritic steels

From their thermal-mechanical production history described above, ferriticsteels have a relatively high yield stress owing to hardening by precipitatesand dislocations and therefore do not generally need significant work hardeningto carry the creep stress (Fig. 13.2). This is particularly pronounced in temperedmartensites where hardening by fine dislocation structures is particularlystrong. Here the primary stage of creep relatively quickly leads into the

ε (s

)–1

10–3

10–4

10–5

10–6

10–7

10–8

0.0 0.1 0.2 0.3 0.4ε

σ/MPa = 58 800H1073 K

156

138125

103

76

13.1 Creep rate versus strain curves for alloy 800H (Germandesignation X10NiCrAlTi32 20, composition in mass %: ≈ 32 Cr,≈ 20 Ni, 0.1 C) in tension. From Portella.24

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secondary stage with a relative minimum εmin of creep rate which isimmediately followed by a tertiary stage with increasing ε .

In contrast to the typical behavior of austenitic steels, a tertiary stageappears not only in tension, but also in compression of tempered martensites,and displays a rate d log ε /dε of increase of log ε that decreases withcompressive strain ε (Fig. 13.2). As mentioned above, the reason for tertiarycreep in compression is microstructural coarsening. When the coarseningrate has sufficiently declined in relation to the rate at which strain is progressing,the compressive creep rate ε attains a constant steady state value at relativelylarge strains (Fig. 13.2), which by far exceeds the minimum creep rate εmin .The latter must therefore be kept quite distinct from the steady state creeprate.28 It simply characterizes the maximum deformation resistance at thepoint where the effects of initial work hardening and softening bymicrostructural coarsening happen to compensate. In tension of temperedmartensites the steady state of creep is generally not attained owing to interferingfracture processes.16

13.4 Transient response to stress changes

13.4.1 Small transient strains

Sudden changes in stress during creep cause materials to respond by elastic(time-independent, reversible), anelastic (time-dependent, reversible) and

ε (s

)–1

10–3

10–4

10–5

0.0 0.1 0.2 0.3 0.4 0.5ε

X3 CrNi 18 9923 K, 250 MPa

X20 CrMoV 12 1915 K, 235 MPa

13.2 Creep rate versus strain curve for tempered martensite 12Cr–1Mo–V steel in uniaxial compression. Thin line: austenitic 18Cr–9Nisteel for comparison. From work by Hofmann,92 and his collegueStraub.

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plastic (time-dependent, irreversible) deformation. Figure 13.3 shows thesereactions in a test on a tempered martensite steel where the uniaxial compressivestress σ was reduced in steps. As creep decelerates with decreasing σ, alogarithmic time scale was chosen to display the progress of ε.

Each change ∆σ of the initial stress σ0 to a new level σ = σ0 + ∆σ leadsto an instantaneous change in elastic strain ∆εel. If ∆εel is measured free frommachine contributions, the ratio E′ = ∆σ/∆εel equals the static elastic (Young’s)

εmech,0 σ0 = 294 MPa

ε 0 = 1.3 · 10–4/s

ε mec

hσ

(MP

a)

0.1915

0.1910

0.1905

0.1900

0.1895

0.1890

300

250

200

150

100

50102 103 104 105

(t – 5200 s) (s)

13.3 Mechanical strain εmech as function of time t for an experimentaltempered martensite 12Cr–2W–5Co steel at 923 K in response tostepwise stress reduction from the initial state after creep for εmech,0

at stress σ0 to creep rate ε0 . Measurement in uniaxial compressionon a specimen of 6 mm height and 5 mm width. Horizontal dashesmark end of elastic reaction to unloading. From Backes.31

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modulus Estat of the steel specimen. In the test for Fig. 13.3, E′ is lower thanthe dynamic modulus E of the steel by a factor of 2, which is probablymainly due to reversible machine contributions to the measured length changes.Tensile tests on long specimens are better suited for the correct determinationof Estat than tests on short compression specimens (see e.g. An et al.).29

The plastic response to a stress reduction | ∆σ | < σ0 is forward strainingat a rate depending on acting stress σ and the evolving structure. For small| ∆σ |, the creep rate immediately after stress reduction is positive as plasticstrain accumulation dominates. This allows one to measure the σ-dependenceof the plastic deformation rate at nearly constant dislocation structure. As thelatter is directly proportional to the average velocity of gliding dislocations,these tests give access to the dislocation glide velocity at the given structure.

The anelastic response to sudden stress changes is forward or back strainingfor stress increases and stress reductions, respectively, mainly owing toreversible motion of dislocations towards or away from the obstacles exertinga back stress on them. At the lowest σ-levels shown in Fig. 13.3 the dominanceof anelastic back straining shortly after the stress reduction is clearly seen.One notices that the anelastic reaction is smaller than the elastic reaction(this holds also when the strain signal is corrected for elastic machinecontributions).

At a certain intermediate value of stress reduction, plastic forward strainand anelastic back strain exactly compensate for some time and the averageinelastic (anelastic plus plastic) strain rate is zero immediately after thestress reduction. This stress reduction is an estimate of the effective stress σ*for glide needed for modeling (see Section 13.7.1). However, the test in Fig.13.3 is not optimal for this task, as the constancy of the microstructure in thesequence of stress reductions is not guaranteed. It was designed to measurethe (positive) creep rate ε at approximately constant total strain ε as functionof σ; this function allows one to calculate the stress relaxation behavior.30,31

Stress reductions for estimating σ*32–34 must be performed when a definedmicrostructural state with a certain deformation resistance, given by the rateε 0 of plastic deformation at the stress σ0, is established. Therefore suchstress reductions have to start from a point where σ = σ0 and ˙ ˙ε ε = 0 .

13.4.2 Large transient strains

When creep is monitored for large strain intervals after a σ-change, structuralchanges following the change in σ become visible in the changes of ε withε as the dislocation structure evolves towards the new steady state correspondingto the new stress. Figures 13.4 and 13.5 show that ε varies in a non-monotonicmanner with ε after large stress reductions and, for 800H, also after largestress increases. The transients after stress reduction, with work hardeningfollowed by work softening, resemble the initial primary creep reaction of

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austenites described above. In addition, one sees that the steels creep fasterat the higher stress after an intermediate period of slow creep over a relativelylong time and, analogously, more slowly at the lower stress, when creep timeis saved by an intermediate period of relatively fast creep at the higher stress.This indicates that there are also time-dependent softening processes occurringin both steels during creep.

13.5 Transient ε ε– response of tempered martensite 12 C–1Mo steelto changes in compressive stress. Thin lines: reference curves atconstant stress. From Straub.17

X20 CrMoV 12 1873 K

σ/MPa = 196

150

0.00 0.05 0.10 0.15 0.20 0.25 0.30ε

ε (

s)

–1

10–4

10–5

10–6

10–7

10–8

13.4 Transient ε ε– response of alloy 800H to sudden reductions intensile stress at 1073 K. Thin lines: reference curves at constantstress. After Portella.24

ε (

s)

–1

10–3

10–4

10–5

10–6

10–7

10–8

0.0 0.1 0.2 0.3ε

80

123

σ/MPa =

800H1073 K

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13.5 Cyclic creep

The progress of inelastic strain under cyclic variation of stress is calledcyclic creep. In a simple case a time period ∆tc of cyclic creep consists oftwo phases where a maximum stress σ and a reduced stress Rσ (0 ≤ R ≤ 1)are acting in time intervals ∆tc,σ and ∆tc – ∆tc,σ > 0, respectively. Owing tothe strong stress dependence of creep rate, progress of creep at significantlyreduced stress is often negligible, so that the net creep rate per cycle may beestimated to be reduced by the ratio ∆tc,σ/∆tc. However, this is not generallyfound. Rather, one observes cyclic creep which is accelerated or deceleratedcompared to the naive expectation.

13.5.1 Cyclic creep deceleration

An example of repeated changes of creep mode of alloy 800H from monotoniccreep at constant stress σ to cyclic creep and back is presented in Fig. 13.6.

For the last range of cyclic creep with the largest period ∆tc = 250 s thecreep strain at σ in the time interval ∆tc,σ = 0.5 ∆tc is expected to be 1.25 ×10–3 from a rate of 10–5 s–1 for monotonic creep just before the mode change.This is larger by a factor of about 1.5 than the change in elastic strain relatedto the stress change (taking E ≈ 150 GPa). The net cyclic creep rate shownin the figure drops by a factor of 2.1 compared to the level of monotoniccreep rate. This is close to the naively expected time ratio ∆tc/∆tc,σ = 2.

However, for the first range of cyclic creep with the smallest cycle period

ε (

s)

–1

10–3

10–4

10–5

10–6

10–7

10–8

0.0 0.1 0.2 0.3ε

800H

R = 1

1073 K124 MPa

1 1 1

0.0450 s

0.04250 s

R = 0.04∆tc = 10 s

∆tcσ = ∆tc/2

13.6 Transient ε ε– response of alloy 800H to sudden change frommonotonic creep (R = 1) to cyclic creep (R = 0.04, period ∆tc) at net

rate ε . From Portella.24

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∆tc = 10 s, where the plastic creep strain is expected to be only 5% of theelastic strain, the net cyclic creep rate drops by a factor of ≈ 100 comparedto the preceding rate of monotonic creep instead of the expected factor of 2.This means that the plastic creep rate at σ under conditions of cyclicallyinterrupted loading is much smaller than at constant stress σ. In addition,there is a pronounced transient response after the mode change which closelyresembles the transient response following the first stress reduction shown inFig. 13.4. This indicates microstructural changes after a sudden change frommonotonic to cyclic creep, which are similar to those following stressreductions. We conclude that a change from monotonic to cyclic creep withintermittent unloading at sufficiently high frequency may act like a stressreduction. This phenomenon is known as cyclic creep deceleration.

13.5.2 Cyclic creep acceleration

When the frequencies of cycling are relatively low so that the creep strain atmaximum stress σ is larger than the elastic (and the anelastic) strain, creeptends to be accelerated by the cyclic unloading.35 One reason for accelerationis the recovery of the dislocation structure in the reduced stress phase whichmay cause significant acceleration of creep at full stress in materials withnormal transient creep behavior,35,27 in particular in pure materials wherenormal creep transients are most pronounced and anelastic deformation isrelatively fast, as solute friction is missing. Another obvious possible reasonfor cyclic creep acceleration lies in time-dependent softening processes.36–38

As described in Section 13.4, these processes tend to accelerate creep at themaximum stress σ. However, they also tend to decelerate it at the reducedstress, so that the acceleration is partially compensated.

The creep cycles reflecting loading conditions of real components, forinstance of power stations, are quite complicated. In a phenomenologicaldescription, the total creep acceleration factor results from a weighted sumof acceleration factors for each phase of cyclic creep.37 The experimentaleffort needed in determining the acceleration and weighting factors cycle bycycle is high; an accompanying modeling of cyclic creep on a microstructuralbasis would be most useful in saving costs and time.

13.6 Microstructural interpretation of creep rate

13.6.1 Basic equations

The dominant process of strain generation in metals is crystallographic slipby glide of free dislocations. The average rate ε M , at which the resolvedshear strain increases in polycrystals, is directly proportional to the productof density (length per volume) ρf and average glide velocity v of free

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dislocations, equalling the slipped area per volume and time (Orowanequation):

εM = b ρf v [13.2]

(Note: The term ‘free dislocations’ covers all dislocations lying in the interiorof subgrains, regardless of their mobility. The free dislocation structure isalso addressed as the three-dimensional Frank network.) The obstacles todislocation glide are found in the dislocation structure itself, in solute atomsand in hard particles of foreign phase. It is common to assume linearsuperposition of the strengthening components σp from particles and σ fromdislocations:

σ = σp + σρ [13.3]

based on the idea that particles lead to a general reduction in the stressavailable to move dislocations in the matrix between the particles.

Following classical approaches (see Seeger39 and Nes40) the obstaclesposed by the dislocation structure itself, for example forest dislocations, areassumed to consist of two parts, an athermal one, which can only be overcomeby mechanical force, and a thermal peak which is susceptible to thermalactivation. The athermal part requires a stress component σG,ρ, depending ontemperature only through the elastic shear modulus G(T). The thermal partrequires an additive thermal stress component σ*, also called the effectivestress, to be overcome at a certain speed. Assuming linear superposition oneobtains:

σρ = σG,ρ + σ* [13.4]

σG,ρ may be written in the simplest form as the sum of contributions fromfree dislocations (subscript f) and dislocations forming low-angle boundaries(subscript b):

σG,ρ = σG,f + σG,b [13.5]

σG,f varies in inverse proportion to the average spacing ρf–0.5 of free dislocations:

σG,f = Cf α M b Gρf–0.5 [13.6]

α ≈ 0.3 is the dislocation interaction constant. The factor Cf increases from0 to 1 as the dislocations move past their obstacles when the stress is applied.Similarly σG,b must be built up by motion of free dislocations against thesubgrain boundaries.

13.6.2 Solid solution hardening

The velocity v at which dislocations glide over the thermal obstacles dependson σ* and the nature of the obstacles. It is a monotonic function of σ* with

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v = 0 at σ* = 0. Within a sufficiently limited interval of temperature andstress the v(σ*) relation can be well approximated by a power law:

v = B σ*m B BQRT

= exp –0g

[13.7]

with effective stress exponent m, apparent activation energy Qg for glide anda constant B0. It is important to note that the function v(σ*) differs dependingon whether or not the dislocations are surrounded by a cloud of solute atomsduring glide. Such clouds with enhanced solute concentration tend to formaround the (edge) dislocations, as solutes are attracted by the dislocationstress field. When the dislocations start gliding, the clouds of solutes exert adrag on the dislocations. The solute drag causes glide of dislocations to beviscous with m = 1.41 The expression for the parameter B in Equation [13.7]derived by Cottrell and Jaswon42 for interaction between solute atoms anddislocations on the basis of the relative atomic size misfit εa of solutesreads.41,43

Bk T

M G b r rDc a

= 9

ln ( / ) B

2 42 1

sol

02⋅

ε[13.8]

where kB is Boltzmann’s constant and r1 and r2 are the internal and externalcut-off radii of dislocation stress field. c0 and Dsol are the atomic concentrationand the diffusion coefficient of the cloud of solutes respectively (see Chapter7 for details). The temperature dependence of Dsol makes a determiningcontribution to Qg.

If a cloud of solutes does not form or if dislocations break away from theirclouds, glide occurs in a jerky manner where fixed obstacles, formed bydislocations in combination with solutes, are overcome after a certain waitingtime with support by thermal activation. The exponent m is >> 1 in this case.

The hardening effect of solutes is strong, if the factor B in Equation [13.7]is low so that the dislocations move slowly. As seen from Equation [13.8],solute atoms with both high concentration in the matrix and strong interactionwith dislocations, for instance by large atomic misfit, are strong hardeners.In the case of viscous glide, a low diffusivity Dsol is an additional conditionfor strong hardening. Mo atoms turn out to be particularly effective solidsolution hardeners of steel. However, Equations [13.7] and [13.2] show thatthe effect of solid solution hardening becomes small under conditions ofslow long-term creep, because the effective stress diminishes in directproportion to the dimishing dislocation velocity v. Therefore the long-termcreep resistance must be guaranteed by particle hardening rather than bysolute hardening. The schematical Fig. 13.7 illustrates the fact that jerkyglide is faster than viscous glide at high σ* and correspondingly low T. Thebreak away of the dislocations from their clouds at a certain maximum of

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effective stress for viscous glide is connected to the increase in v. The formationof clouds at a certain minimum value of effective stress for jerky glide leadsto a decrease in v. Processes related to the formation of and break away fromclouds are known by the term dynamic strain aging. Far from being a peculiarity,this addresses a ubiquitous phenomenon in the deformation of alloys. Themanifold of different interstitial and substitutional alloying elements and thefact that the activation energies of creep and diffusion of solutes often aresimilar produces dynamic strain aging effects in a wide range of temperaturesand strain rates. Creep tests are particularly well suited to revealing theeffects of dynamic strain aging, because the creep rate for obvious reasons ismuch more sensitive to changes in the thermal stress component σ* than theflow stress.

With this qualitative knowledge of v(σ*) and using Equations [13.2] to[13.7] it is possible to gain a qualitative understanding of the processesduring creep by considering the variations in the microstructural parametersσf, ρb and σρ with strain.

log

ν

Jerky

Break-away

Thinning ofclouds

1

2

3

4

Cloud formation

Viscous

log σ∗

13.7 Dislocation velocity v as function of effective stress σ* forjerky and viscous glide at given T, with jerky ↔ viscous transition(schematic); upward arrow: viscously gliding dislocations breakaway from their solute clouds; downward arrow: jerkily movingdislocations slow down owing to formation of a solute cloud.

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13.6.3 Loading strain and initial primary creep

Austenitic steels with low initial dislocation density

Primary creep transients of the type described in Section 13.3.1 have alsobeen found in highly alloyed Ni–Cr base alloys29 and in Ti alloys.44,45 Theirinterpretation in terms of dynamic strain aging applies to steels as well.

A simple interpretation neglecting the frequency distribution of thedislocation velocity and cooperative effects in formation and break awayfrom clouds is as follows: During work hardening in the loading phasewhere stress increases, few dislocations move at high velocity v in the jerkymode without cloud (point 1 in Fig. 13.7). Increase of dislocation densityleads to reduction of v at approximately constant ε (see Equation [13.2])without change of mode (point 2). When the increase of flow stress stops andcreep begins, continuing work hardening makes σ* and v fall. This enablessolute atoms to attach themselves to the dislocations, form a cloud andsignificantly reduce v to the level of viscous glide (point 3). The abruptdecrease in v reduces the rate of generation of dislocations so strongly thatthe rate of recovery of dislocations, which as a diffusion-controlled processis rather insensitive to solute atoms, now exceeds the generation rate. Thisleads to coarsening of the dislocation structure, decrease of σG,ρ and increaseof σ* and v, until the steady state density of free dislocations is eventuallyreached at point 4. This consideration explains the pronounced relativemaximum of apparent work hardening rate (see Section 13.2.1) as a directconsequence of the jerky–viscous transition (from point 2 to point 3). Thesubsequent minimum in work hardening rate is due to the recovery-inducedtransition from point 3 to 4.

The present discussion of dynamic strain aging was restricted to interactionbetween dislocations and solutes by solute drag. It may be noted that shortrange ordering, which has been reported, for example for Cr in Fe (seeSchönfeld),46 may cause similar effects; the destruction of short-range orderby fast glide corresponds to the break away phenomenon; the rebuilding ofshort range order at low creep rate corresponds to cloud formation.

If the interpretation of the beginning of primary creep in terms of dynamicstrain is correct, the effective stress for viscous glide must be smaller thanthe available effective stress. This condition can be checked. Taking literaturedata for Ni,29,47 combining Equations [13.2], [13.7] and [13.8], andapproximating ρf by (σ/(Gb))2 yields σ*/σ = 0.6 for creep at 76 MPa. Thisis a reasonable value in view of the numerous uncertainties and approximationsinvolved. The effective stresses σ v , max

* for break away from clouds werereported by Oikawa and Langdon41 to range between 10–4 and 10–3 G. Thismeans that the highest creep stress ≈ 150 MPa = 3 × 10–3 G is close to theviscous–jerky transition. Thus the proposal of dynamic strain aging in creepof alloy 800H is viable.

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The preceding consideration does not cover the slow work hardening inthe extended primary creep stage preceding the steady state of creep (Fig.13.1). This work hardening is easily explained in the usual manner by gradualformation of a subgrain structure which is superimposed on the evolution offree dislocations and has been observed, for instance, in a Ni–Cr alloy,27,29,48

which is similar in behavior to alloy 800.In addition to solid solution strengthening, the strengthening by precipitates

and dispersed particles reduces the primary creep rate.24 Coarsening of theprecipitates with time increases the spacing of precipitates along dislocationlines and causes gradual loss of particle strengthening. This process maysuperimpose on the ones described above.

Ferritic steels with high initial dislocation density

The processes described in the preceding section tend to be suppressed inferritic steels, which in their initial state generally are more strongly hardenedby precipitates and dislocations inherited from phase transformation and inwhich the contents of alloying elements in solid solution tend to be smaller.Often the dislocation density need not increase during loading, as the initialdensity already exceeds the steady state density under creep conditions,6,17,49

and the loading strain is negligible.A primary stage of creep with decrease of ε is nevertheless observed

(Figs. 13.2 and 13.5). The reason is that the athermal dislocation hardeningterm σG,ρ is low at the beginning of creep. This can be explained in thefollowing way: The deformation paths differ in phase transformation andcreep. The high initial dislocation content inherited from martensitictransformation comprises a significant fraction of dislocations which areeasily glissile under the combined influences of inherited internal stressesand applied creep stress. These dislocations glide past obstacles. As they doso, the forces exerted by the gliding dislocations on the obstacles rise. Inconsequence, the factor Cf in Equation [13.6] rises up to its maximum valueof 1. Assuming a total length per volume, ρ, of dislocations to move by adistance in the order of an average dislocation spacing ρ–0.5, before beingstopped at obstacles in the dislocation structure, means a strain of (b/M) ρρ–0.5 (see Equation [13.2]), amounting to 0.002 if the total dislocation densityis estimated from Equation [13.6] by setting σG,f equal to the estimate of thelocal stress 380 MPa acting during martensitic transformation (see Section13.2.2). This rough estimate lies in the order of magnitude of the observedprimary creep strains. It is concluded that the fast creep in the primary stagecan be explained in terms of build-up of dislocation interactions opposingthe applied stress.

The description of primary creep given above resembles the idea that siteswith easy strain generation are gradually exhausted (exhaustion creep).50

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The theory of exhaustion creep predicts a decrease of creep rate with time taccording to ε ∝ t–p. p = 2/3 corresponds to the empirical Andrade creeplaw.51,52 This type of time-dependent decay of primary creep rate is in factobserved in steels.52–55

13.6.4 Transient creep response to changes in creepconditions

Stress changes

The interpretation of primary creep transients in terms of dynamic strainaging given in Section 13.6.3 can be transferred to the transients of Figs.13.4 and 13.5 following abrupt changes in stress. In this interpretation theinitial work hardening (decrease of ε with ε) after stress reduction is connectedto the densification of the clouds of solutes leading to significant reductionin dislocation velocity. In the subsequent range of work softening (increaseof ε ) the dislocation structure coarsens until a new balance between generationof dislocations in the course of glide and annihilation of dislocations in thecourse of recovery is re-established. Analogously, the reaction after stressincrease is explained by thinning of the clouds and refinement of the dislocationstructure.

In addition to dynamic strain aging, recovery of dislocations will initiallyadd to the observed decrease of ε.56,57 This is seen from the followingconsideration. The effective stress was determined by stress reduction teststo be about 25 MPa in creep of 800H at 123 MPa.24 This means that forwardglide associated with generation of dislocations ceases after the stress reductionin Fig. 13.4, because σ* is no longer positive. In this situation creep strain isproduced as a by-product of recovery, for instance by migrating subgrainboundaries.58 However, the strain rate associated with recovery falls quicklyas recovery slows down and glide of free dislocations becomes dominantagain.

For tempered martensite steels the situation is similar to that discussed foraustenites with regard to solute effects and recovery-associated strain. Anestimate similar to the one in Section 13.6.3 based on thermodynamic data47,59,60

yielded relative effective stresses σ*/σ = 0.37 for viscous drag of a cloud ofMo at 200 MPa and 873 K. The stress of 200 MPa corresponds to 3 × 10–3

G, which is again close to the break away of dislocations from clouds.A difference between Figs. 13.4 and 13.5 lies in the fact that the stress

changes in the tempered martensite are performed in the tertiary stage ofcreep, in contrast to those in the austenitic alloy. In both cases the transientsjoin the general trend of increase and decrease of ε with ε. The fact that thetransients keep a distance from the curves measured without stress changeshas already been related to time-dependent softening (see Section 13.4).

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Loss of particle hardening offers itself as the most probable mechanism (seeSection 13.6.6).

Changes from monotonic to cyclic creep

The intermediate unloading in each cycle of cyclic creep allows the dislocationsto come to rest, apart from anelastic backward glide motion. In this situationmore solutes have the chance to attach themselves to the dislocations. Thus,a change from monotonic to cyclic creep acts in a similar way as a limitedstress reduction. The net creep rate therefore changes with strain in the modechange test of Fig. 13.6 in close similarity to the stress change test (Fig.13.4); for instance, repeated unloading for 5 s followed by creep at full loadfor 5 s causes a transient creep reaction similar to lowering the stress from124 MPa to 80 MPa.

13.6.5 Tertiary creep caused by subgrain coarsening

In Section 13.6.3 formation and refinement of subgrains have been invokedto explain primary creep of steels with low initial dislocation density.Correspondingly, coarsening of a pre-existing subgrain structure will lead tosoftening, that is, to increase in creep rate with time, which phenomenologicallyis considered as tertiary creep, ending the secondary stage of creep. Subgraincoarsening occurs in creep when the initial subgrain size is smaller than thesteady state subgrain size according to Equation [13.1]. This is the case inmost applications of tempered martensite steels.

Reduction of creep strength by subgrain coarsening was confirmed intests where subgrains were made to coarsen within a relatively short time,where changes in the precipitate structure are negligible.61–63 The coarseningwas achieved by strain-controlled cyclic straining at elevated temperature.When the maximum stress acting in cyclic deformation is sufficiently low,the subgrains grow fast with accumulating inelastic strain towards theinstantaneous stress-dependent value of the steady state subgrain size (Equation[13.1]). The subgrain coarsening to 1.5 µm caused the minimum creep rateto increase by about an order of magnitude.61,62

Softening caused by subgrain coarsening, shows up in the tertiary stage ofcreep of tempered martensite steels in tension as well as in compression.17 Incontrast to austenitic steels, the softening begins at low strains, as soon asthe initial primary creep related to build-up of forces on the obstacles (seeSection 13.6.3) has ceased. Subgrain strengthening is extremely importantfor the relatively high creep resistance of the tempered martensites not onlyby reducing the minimum creep rate compared to the steady state creep rate,but also by shifting the minimum to low strains. These two advantagescounteract the disadvantage of ferritic steels compared to conventional

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austenitic ones regarding higher diffusivity, causing higher rates of structuralcoarsening and creep.

It must be noted that subgrain strengthening is not a pure dislocationeffect. The energy stored in fine structures provides a high driving force forstructural coarsening. It needs the pinning effect of precipitates to stabilizethe subgrain structure of tempered martensites. This was convincinglydemonstrated by an experiment by Kostka et al.64 These authors produceda subgrain structure in a ferritic Fe–Cr alloy by severe plastic predeformationin equal channel angular pressing followed by annealing. The initial subgrainstructure was quite similar to that of tempered martensites. But the lack ofprecipitates at the subgrain boundaries caused fast coarsening of the subgrainstructure by dynamic recrystallization during creep. In effect, the unstabilizedsubgrain structure induced by severe plastic predeformation led to a creepcurve at 100 MPa and 873 K where the relative minimum in ε had nearlydisappeared so that the actual value of the minimum creep rate had increasedby a factor of 104 compared to the tempered martensite. This result clearlyshows the essential effect of structural stabilization by particles. The particlesreduce the velocity of migration of the subgrain boundaries and thus inhibitthe formation of stable nuclei for recrystallization. The subgrain structurein tempered martensite steels is therefore particle stabilized.6

13.6.6 Increase of creep rate caused by degradation ofparticle hardening

Degradation of particle hardening caused by an increase in the mean spacingof particles along dislocations is unavoidable at elevated temperatures. Itgenerally occurs by Ostwald ripening of the existing precipitates. Duringripening the average precipitate volume increases linearly with time t spentat elevated temperature T:

d d k tp3

p,03

p = + [13.9]

so that the spacing of precipitates along dislocations increases and particlehardening is diminished.16,17 This time-dependent softening counteracts thework hardening in primary creep and contributes to the increase of creep ratein the tertiary stage of creep.16,17 The growth constant kp is proportional tothe specific energy of the phase boundary of precipitates and depends on theatoms forming the precipitates. Atoms with a high concentration in theprecipitates, a low concentration in the matrix and low diffusivity inthe matrix are useful in reducing kp.

65–67 The time effects on creep rateobserved in stress change tests (Section 13.4) are consistent with Ostwaldripening.

An example of relatively fast Ostwald ripening is given by the phaseM2X in tempered martensite steels. This phase precipitates in a fine

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distribution within the volume, hardens the subgrain interior and providesa high short-term creep resistance. However, it also coarsens rapidly, preferablyat subgrain boundaries, while the fine precipitates inside the subgrainsdissolve. The accompanying loss of precipitation hardening of subgrainsleads to relatively fast degradation of creep resistance.63,68,69

The ripening of precipitates becomes more complicated when a newlynucleating phase like the Z-phase disturbs the diffusion currents set up betweenthe precipitates. A stability advantage of the new phase may cause existingphases to dissolve in favor of the new phase. The result may be similar tofast ripening of existing phases and lead to fast degradation of an initiallyhigh creep resistance. It is therefore of importance to discourage the nucleationof more stable phases. This is difficult to achieve, as there are no simplerecipes for inhibiting formation of more stable phases, because heterogeneousnucleation is a local event depending on the properties of the individualnucleation sites.

The effect of coarsening of the precipitate structure may degrade thecreep resistance of an initially creep-resistant steel to an extent that its long-term creep resistance is lower than that of a steel with relatively low initialcreep resistance, but more stable particles.68,69

13.6.7 Deceleration and acceleration of creep by cyclicvariation of stress

The large reduction in net creep rate caused by cyclic unloading at highfrequency (Fig. 13.6) cannot be explained by dislocation–solute interactionalone. It needs an explanation in terms of anelastic deformation.36,70,71 Werecall that cyclic creep is decelerated when the expected plastic strain in thephase of maximum stress is small compared to the elastic strain. In this casethe anelastic strain may be small, too, and may have been interrupted beforeit reached its final value corresponding to the applied stress. Anelastic glidemotion of dislocations is necessary to build up the forces on the dislocationobstacles which are necessary for their irreversible overcoming producingplastic deformation. This is described by the factor Cf in the exampleof Equation [13.6] and holds for other obstacles like particles as well.Interruption of anelastic deformation means that the forces on the obstaclesare lower than in monotonic deformation. This explains why the plasticdeformation rate may be strongly reduced by cyclic variation of stress athigh frequency.

Acceleration of creep by cyclic stressing has nothing to do with anelasticity,but is explained by structural coarsening in the unloading phases as alreadymentioned in Section 13.5.2.

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13.7 Dislocation models of creep

13.7.1 Kinetics of dislocation glide

An important step in development of a microstructural model of creep sketchedin Section 13.6.1 is the formulation of the dislocation glide velocity v asfunction of the effective stress σ* at given T. This requires an educatedguess, as neither v nor σ* can be fixed exactly. By definition of effectivestress, v(σ*) must vanish at σ* = 0. Using this condition, σ* can in principlebe determined in tests where σ is suddenly reduced to a level where thestrain increment ∆ε within a small time increment ∆t is zero. For technicalmaterials like steels this is comparatively easy. The stresses and strains arerelatively high so that their changes become relatively large. This eases theproblems34 encountered with pure materials where the strain changes may betoo small to be resolved. However, the problem of systematic errors, caused,for instance, by strains associated with recovery, remains. Nevertheless thestress reduction technique allows the order of magnitude of σ*/σ to be fixedand thus facilitates reasonable model assumptions.

As an example, Fig. 13.8 shows the results of a series of stress changetests for a tempered martensite steel. δε is the difference between the maximumback strain and the elastic back strain. For stress reductions ∆σ < 0.05 σ0 thedifference δε was zero, that is, no net inelastic back flow was observed,indicating that the effective stress for glide remained ≥ 0 after the stress

0.5 0.6 0.7 0.8 0.9 1.0(σ0 – ∆σ )/σ 0

873 Kσ0 = 320 MPa

X20 CrMoV 12

δε (

10–4

)

0

–1

–2

–3

13.8 Difference δε between maximum back strain and elastic backstrain as a function of relative reduced stress for the temperedmartensite steel from Fig. 13.2. Stress reduction tests were performedstarting from σ0 at strains 0.035 < ε0 < 0.44 and initial creep rates of 4× 10–5 s < ε0 < 4 × 10–4 s in uniaxial compression. The estimatedexperimental error of individual δε values is ±10–4. After Goblirsch.67

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reduction. For stress reductions ∆σ > 0.24 σ0 the inelastic back flow |δε| hasgrown to 2 × 10–4 indicating net back flow under σ* < 0. A straightforwardconclusion is that the effective stress at σ = σ0 should lie between 0.05 and0.24 σ0 in the investigated case. However, Fig. 13.8 shows that this conclusionis reliable only, if the uncertainty in δε due to uncertainty in strain measurementand systematic errors is distinctly less than 10–4.

If creep rate ε , dislocation density ρf and effective stress σ* are known,it is possible to derive the dislocation velocity v as function of effectivestress σ* (Equation [13.7]) from Equation [13.2]. A complete set of data wasprovided by Milička for alloy 800 at 975 K.22 It yields the v(σ*) relationshown in Fig. 13.9. The effective stress exponent m of v lies near 3. Thisvalue corresponds to the range of the curve in Fig. 13.7 where the solutecloud is thinning with increasing v before break away occurs. Wolf and co-workers29,53 performed a similar analysis for a Ni–Cr alloy. Their result, alsoshown in Fig. 13.9, is similar to that for alloy 800, if the difference in T isneglected. Polcik49 proposed the curve of Fig. 13.9 for a tempered martensitesteel at 923 K. The lower and the upper branch of this curve were interpretedin terms of viscous and jerky glide, respectively.

ν (m

s–1)

10–6

10–7

10–8

10–9

10–10

10–11

10–12

10 100σ* (MPa)

9-12% Cr steels, 8

73 KAlloy 800, 9

75 KNiCr22Co12M0, 1073 K

13.9 Dislocation velocity v as a function of effective stress σ* for alloy800 derived from data of Milička,22 for typical 9–12%Cr steels49 andfor a Ni-base alloy.53,29

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The overall consistency of the data displayed in Fig. 13.9 is encouraging.Even though no definite conclusion is possible at present regarding the exactform of the v(σ*) relation, investigations of the kind described are a usefulguideline in fitting the v(σ*) relation, needed for modeling, to experimentaldata.

13.7.2 Evolution of dislocation structure

It is desirable to model the evolution of dislocation structure on the basis ofexpressions for the rates of generation and loss of dislocations. Significantprogress has been made in developing these expressions.40,43,72,73 However,there have been only few efforts in the field of steel (see e.g. Weinert)74 anda widespread direct application of the structure evolution approach to steelis still lacking.

Fortunately, the exact system of differential equations for dislocationstructure evolution can be approximated by a system which is based onexperimental knowledge of the steady state dislocation structure.17,18 Knowingthe initial values X0 and the steady state values X∞ of the characteristicspacings X, the evolution of the dislocation structure with strain can beapproximated by integrating rate equations of the type

dd

= – – X X XkXε

∞ [13.10]

with adequately chosen values of kX. At constant stress and temperature,Equation [13.10] predicts X to approach X∞ in an exponential fashion. Thisphenomenological law of structure evolution has successfully been appliedto tempered martensite steels.16–18,49 It was shown to be applicable not onlyfor creep, but also for subgrain size evolution in strain controlled cyclicdeformation at elevated temperature.62,69 The modeling is simplified by thefact that the steady state values of the spacings X = w and ρf

–0.5 can beapproximated as unique functions of normalized stress σ/G and that themean spacing s of dislocations in subgrain boundaries of low-angle type isapproximately constant.17 The latter approximation is probably not valid forsteels without pre-existing subgrain structures and should be modified accordingto proposals for strain controlled decrease of s (increase of average boundarymisorientation b/s).75

13.7.3 Particle hardening

As mentioned before, particle hardening and stabilization of the dislocationstructure by particles may have a decisive influence on the creep resistance.In order to model creep of particle-hardened steels one needs to know themechanism of interaction between precipitates and dislocations and theevolution of the particle structure during creep.

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The simplest approach to interaction stress σp between free dislocationsand particles is to assume that the dislocations by-pass the particles by glide.In this case σp takes its maximum value given by the Orowan stress. As seenin Section 13.8, the assumption that σp is near to its maximum value duringdeformation at elevated temperature is supported by electron microscopicstudies. The Orowan stress can be calculated for particles which are relativelyhomogeneously distributed in the volume, like MX-precipitates. However, itdoes not consider the effect of climb over particles during slow creep. Incase of climb over particles, the interaction stress may well be given by theattractive interaction between particles and free dislocations.76 In this caseσp would be smaller than the Orowan stress and might also have a thermalcomponent.

The stabilization of low-angle boundaries (dislocation networks) againstmigration under stress must be described in a different manner. Polcik49 usedan approach based on Zener drag resulting from attractive interaction betweenboundaries and particles, which is another form of attractive dislocation–particle interaction. At present there is no safe procedure for modelingprecipitation hardening. Therefore it is necessary to rely on fitting modelingapproaches containing particle structure parameters to experimental results.An example will be given below (see Appendix).

In any case the coarsening of the particle structure and the related softeningis a fact which needs description. The parameters entering the descriptionare the locations of the particles (at boundaries or distributed over the volume)and the sizes dp and volume fractions fp of particles. As different phasesbehave differently, a phase specific description of evolution is necessary.The simplest approach is to assume that particles evolve via diffusive processeswith time. At the present state of knowledge this appears to be justified.68 Inspite of repeated claims (e.g. Eggeler77 and Cerri et al.78) a clear influenceof concurrent creep straining on ripening of particles has not been unequivocallyproven.68,79 However, there are indications that particle ripening is enhancedat low-angle boundaries where short-circuit diffusion along dislocations maybe important.6,68,69,79

Despite the complexity of the situation with different phases at differentlocations, which may be enhanced by the appearance of new phases,characterization of the particle structure appears possible on the basis ofexperimental observations combined with the knowledge provided bythermodynamic data banks.67,80 The next step leading from the particle structureto expressions for particle hardening needs educated guesses and fitting.

13.7.4 Composite model

As an example of a statistical dislocation model of creep of steels, we considerthe (iso-strain) composite model of plastic deformation proposed by Mughrabi,81

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expanded for inclusion of strain rates by Blum and co-workers82–84 andapplied to steel by Straub and co-workers85 and Polcik49 and to Ni-phasealloys with dominant solute hardening by Meier and Blum.48 In this modelthe material is viewed as a composite of soft subgrains surrounded by hardboundaries. The equations used in fitting the model to a tempered martensitesteel are given in the Appendix. The composite model comprises:

• dislocation structure evolution,• particle structure evolution for individual precipitate phases in the subgrain

interior and at the subgrain boundaries,• hardening by free dislocations and subgrain boundaries,• particle hardening of subgrain interiors and boundaries,• solute hardening influencing the dislocation glide velocity.

In the case of inhomogeneous formation of subgrains in solid solutionhardened fcc metals, the deformation resistances of the volume fractionswith and without subgrain structure were averaged using an iso-stress compositeapproach.48

The composite model applied to tempered martensite steels provides ε ε–curves which reproduce experimentally observed features in that the creeprate reaches a minimum with strain and then increases owing to strain-dependent coarsening of the dislocation structure and time-dependentcoarsening of the particle structure.49,85,86 There is a sufficient number ofparameters in the model to fit the quantitative result to experimentalobservations. Anelastic deformation caused by interaction between softsubgrain interiors and hard boundaries is implicit in the model. The softeningwith both strain and time is reproduced by the model. The stress dependenceof the minimum creep rate can also be reproduced to a reasonable extent.Little effort has been spent so far to improve the fit between model andexperiment leaving opportunity for improvement.

13.8 In situ transmission electron microscope

observations of dislocation activity

As mentioned in the introduction, creep is all about dislocation activity. Onewould therefore like to see dislocations moving in a steel under stress atelevated temperature. This is in principle possible by deforming specimensof steels inside a transmission electron microscope (TEM) and recording themotion of dislocations. Such in situ TEM observations have recently beenmade by Messerschmidt et al.87 on an experimental tempered martensitesteel. The subgrain structure in this steel had been coarsened to w ≈ 1.5 µmby prior cyclic deformation at 873 K.62 This facilitated the observations.

Figure 13.10 shows the interior and some of the boundaries of a subgrain.Most of the boundaries appear to be pinned by precipitates. A free dislocation

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is held up at two precipitates. It is relatively strongly bowed. This means thatthe stress needed for the dislocation to overcome the precipitates in thesubgrain interior is not far from the maximum given by the Orowan stress.

Figure 13.11 illustrates why subgrain boundaries in tempered martensitesteels are regarded as relatively hard regions with relatively high densities ofdislocations and particles. The subgrain boundary in the lower half of thefigure identifies itself by the linear black–white contrast (fringes of equalthickness) and the sharply delineated intersection with the foil surface. It ispinned by large precipitates; note the slight bowing of the boundary betweenprecipitates. Below the boundary dislocations are out of contrast. Above theboundary bowed dislocation are seen which apparently come out of theboundary. This means that dislocations are able to cross boundaries by movinginto it from one side and, some time later, leaving it towards the other side.This is in line with the notion that the mean free path of free dislocations in

13.10 In situ observations in TEM with 1 MV acceleration voltage of athin area of an experimental tempered martensite 12Cr–2W–5Co steelduring slow elongation at 923 K by Messerschmidt et al.87 reportedby Chilukuru:69 subgrain with boundaries at top, bottom and right(see dislocation networks) and free dislocations bowed atprecipitates (arrows). Still picture.

2 µm

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subgrain structures is distinctly larger than the subgrain size, see Blum27 andNes.40 It is unlikely that the boundary is completely dissolving, because thesubgrain structure is already relatively coarse and essentially maintains its averageboundary spacing during deformation at the elevated temperature.16,49,62,88

Figure 13.12 shows video frames of a subgrain taken at different times.Dislocation loops are generated at the bottom subgrain boundary. They expandinto the subgrain. Some dislocations overtake others. Many dislocations areresting. These dislocations are relatively straight, probably because they feelonly a low resolved shear stress acting on them. When bowed dislocationsare released from strong obstacles, probably precipitates, they jump forwardat high speed.

In situ TEM observations can be used to test the validity of themicrostructural model of creep. The lower part of the subgrain is crossed byabout six dislocations within a time interval ∆t = 50 s. This means a shearstrain of ∆γ = 6b/tfoil and a strain rate ε = M–1 ∆γ/∆t ≈ 10–5 s–1 (tfoil ≈ 0.5 µmfor 1 MV-TEM). This rate corresponds to a stress of about 174 MPa (datafrom Dubey et al.62 for steel 2A extrapolated with a stress exponent of 8).From the data given in the Appendix one expects an average spacing of free

2 µm

13.11 As Fig. 13.10: free dislocations leave the subgrain boundary.Still picture.

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dislocations of 0.35 µm at this stress. This is compatible with the spacingsseen in Fig. 13.12. Assuming an effective stress fraction σ*/σ = 0.2 givesσ* = 35 MPa.

The curve for the tempered martensite steel in Fig. 13.9(b) shows that theaverage velocity vg of viscously gliding dislocations should be 2 × 10–4 µms–1 at 873 K and, with the activation energy given in the Appendix, 0.6 µms–1 at 823 K. The dislocation seen jumping between the positions in Fig.13.12(a) and (b) moves particularly fast, at a velocity above 2 µm s–1. Onaverage, the velocities seen in situ are lower by about a factor of 10. This isconsistent with the average velocity estimated from Fig. 13.9. The radius r≈ MbG/σloc of curvature of resting dislocations gives information about thelocally acting stress σloc.

81 In the subgrain interior typical values of r lie near

13.12 In situ observation of configuration of gliding dislocations insubgrains of the steel in Fig. 13.10 (frames from video ofMesserschmidt et al.)87: (a) dislocations have been emitted from asource in the subgrain boundary at the bottom of this TEM picture,(b) within less than 0.1 s from (a), dislocation 4 has jumped a distanceof 0.2 µm, much faster than average, (c) 3 s later dislocation 5 is heldup at an obstacle; bowing has reached a maximum, (d) within ≈ 0.1 sfrom (c) dislocation 5 has moved for 0.1 µm after release from theobstacle into a new position where it slows down again.

(b)(a)

(d)(c)

0.5 µm

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0.7 µm corresponding to a local stress σloc = 33 MPa which is distinctly lessthan the estimated stress of 174 MPa acting in the investigated area. At thesubgrain boundaries, r is found to be as small as 0.08 µm. This correspondsto a local stress of 290 µm in the immediate vicinity of the subgrain boundarieswhich significantly exceeds the acting stress. These observations are consistentwith the concentration of stress at the subgrain boundaries postulated in thecomposite model. The strong curvature of the dislocation in Fig. 13.11(c)shortly before release from the obstacle (probably a precipitate) also indicateslocally enhanced stress which corresponds to the relatively high velocity ofthe dislocation shortly after release. We conclude that the in situ observationssupport the validity of the microstructural model of creep.

13.9 Discussion and outlook

As shown in the preceding, rather simple expressions for structure evolutionand dislocation velocity are available which explain a variety of creep featuresand even yield a quantitative formulation for creep after fitting to experimentalresults. Applying them to specific steels gives a feeling of the relativeimportance of hardening contributions.

It is common to ask for the rate-determining mechanism of creep. Creepcontrol by either climb or glide of dislocations are popular alternatives.However, the model described above teaches that creep is a complex plasticityphenomenon where a number of subprocesses interact as illustrated by Blumet al.43 Being a special case of crystal plasticity, creep must be described byplasticity models which include the description of dislocation motion as wellas the evolution of the dislocation structure, which is connected to dislocationmotion, and the evolution of the hardening phases.89 Glide and climb ofdislocations enter the description of the subprocesses of creep, but do notcontrol creep in an autonomous manner.

As another example of coupling of subprocesses, we consider the motionsof free dislocations and of low-angle boundaries. Subgrain boundary migrationmakes a certain constant relative contribution to strain lying in the order of10% for pure metals.58,90 In this sense, subgrain boundary migration maycontrol the rate of creep by glide of free dislocations. However, according tothe composite model it is not an autonomous process, as the glide of freedislocations leads to stress concentration at the subgrain boundaries.

One of the weakest points of microstructural creep models is thequantification of the particle influence on the motion of free dislocations andlow-angle boundaries under stress. Particle hardening is essential in creep ofsteels. It enhances the creep resistance not only directly by adding a stresscomponent (Equation [13.3]), but also indirectly by refining the subgrainstructure caused by the enhanced stress level (Equation [13.1]) and stabilizingit against strain-controlled coarsening and recrystallization. It is notable that

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the apparent activation energies of creep are much higher than those ofatomic diffusion in tempered martensite steels. This indicates a deficit in thepresent model which may have to do with the present description of particlestrengthening. Approximation of the particle hardening in the subgrain interiorby the Orowan stress may not be sufficient. Thermal activation by climb ofdislocations over large particles and thermally activated detachment fromsmall particles may need to be taken into account. This will lead to a temperaturedependence of the particle hardening term in addition to that following fromthe temperature dependence of the solubility of alloying elements. There isalso the possibility that the description of particle hardening of the subgrainboundaries by the Zener stress is insufficient. The Zener stress addressespinning of the boundaries by particles. However, the motion of free dislocationsthrough low-angle subgrain boundaries is the primary process to be consideredin the composite model. This motion may be strongly impeded by densearrays of large particles at the boundaries which probably have to be overcomeby climb.

In the model proposed in Section 13.6.1, particle hardening reduces theglide velocity simply by reducing the effective stress for glide between theparticles. A better way to arrive at the average glide velocity of free dislocationsmay be explicitly to treat the stop-and-go mode of dislocation glide, withwaiting at particles and relatively fast glide between the particles. This waywas followed in deriving the creep rate of a mechanically alloyed Fe-basealloy.91

The formulation of the velocity of glide of free dislocations in the hardregion (Equation [A.4] of the Appendix) poses a basic problem. A concentratedstress σh is needed by the free dislocations to cross the dense dislocationsarray of the low-angle subgrain boundaries. Once this is achieved, thedislocations bow out between the precipitates into the subgrain interior wherethe local stress is relatively small. This is seen from in situ pictures like thosein Fig. 13.10 which show only limited bowing under relatively low stressalthough the dislocations are still linked to precipitates at the subgrainboundaries. This observation indicates that the model should be modified inthe sense that the precipitates at the boundaries contribute to the particlehardening experienced by the free dislocations in the soft subgrain interior.The pinning effect exerted by the boundary particles on the subgrain boundarieswill still be maintained. It leads to suppression of subgrain boundary migrationwhich explains the relatively low value of the growth constant kw.

Despite the existing deficiencies, modeling of creep on a microstructuralbasis is quite promising to be useful in guiding alloy development and givingqualitative or even semi-quantitative explanations for material properties.The formulation of the model in terms of differential equations allows complexloading paths similar to those occurring in practical applications to be simulated,provided that the relevant subprocesses of creep and the relevant microstructural

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elements enter the model in a reasonable manner. Coupling the deformationmodel with a thermodynamic data base has been attempted successfully80

and is a promising further step in model development.Industrial development of steels which resist creep will have to concentrate

on restricting dislocation motion by enriching the steel matrix with stronglyinteracting, slowly diffusing solutes and a stable structure of fine, hard,slowly coarsening particles stabilizing the dislocation structure by opposingdislocation motion. As the creep resistance increases, the danger of creepfracture increases, too. Therefore the resistance of the material againstnucleation of pores and microcracks requires additional attention.

13.10 Acknowledgments

Thanks are due to Dr.-Ing. S. Straub and Dr.-Ing. P. Polcik, who explored themicrostructure of the steels and applied the composite model to them, to Dr.-Ing. H. Chilukuru and R. Agamennone for further scientific progress reachedin their works, to H. Chilukuru and W. Wranik for support in preparing thefigures, to Prof. U. Messerschmidt, Dr. M. Bartsch and their co-workersfrom MPI für Mikrostrukturphysik Halle for performing in situ TEMobservations, to Dr. J. Granacher, Dr. A. Scholz and Prof. C. Berger from TUDarstadt for providing long-term crept specimens and to the DeutscheForschungsgemeinschaft, the Bundesministerium für Wirtschaft and theassociated group of industrial companies for continuous financial support.

13.11 References

1 B. Wilshire, ‘On the evidence for diffusional creep processes’, In B. Wilshire and R.W. Evans (eds), Creep and Fracture of Engineering Materials and Structures, TheInstitute of Metals, London, 1990, 1–9.

2 F. R. Beckitt and B. R. Clark, ‘The shape and mechanism of formation of M23C6

carbide in austenite’, Acta Metall., 1967, 15, 113–129.3 B. Sasmal. ‘Mechanism of the formation of M23C6 plates around undissolved NbC

particles in a stabilised austenitic stainless steel’, J. Mater. Sci. 1997, 32, 5439–5444.

4 C. M. Wayman, Introduction to the Crystallography of Martensitic Transformations,Macmillan Series in Material Science, The Macmillan Company, New York, 1964.

5 P. Haasen, Physical Metallurgy, Cambridge University Press, Cambridge, 2nd edition,1986.

6 G. Eggeler, N. Nilsvang and B. Ilschner, ‘Microstructural changes in a 12% chromiumsteel during creep’, Steel Res. 1987, 58, 97–103.

7 F. Abe, H. Araki and T. Noda, ‘Microstructural evolution in bainite martensite, andδ ferrite of low activation Cr-2W ferritic steels’, Material Science and Technology,1990, 6, 714–723.

8 P. J. Ennis, A. Zielińska-Lipiec and A. Czyrska-Filemonowicz, ‘Influence of heattreatments on microstructural parameters and mechanical properties of P92 steel’,Materials Science and Technology, 2000, 16(10), 1226–1232.

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9 F. Yoshida, D. Terada, H. Nakashima, H. Abe, H. Hayakawa and S. Zaefferer,‘Microstructure change during creep deformation of modified 9Cr-1Mo steel’, InAdvances in Materials Technology for Fossil Power Plants, Proceedings of the 3rdConference, R. Viswanathan, W.T. Bakker and J. D. Parker (eds), held at Universityof Wales Swansea, 5–6 April The Institute of Materials, London, 2001, 143–151.

10 A. Dronhofer, J. Pešic∨ka, A. Dlouhý and G. Eggeler, ‘On the nature of internalinterfaces in tempered martensite ferritic steels’, Z. Metallkd., 2003, 94(5), 511–520.

11 A. H. Cottrell, An Introduction to Metallurgy, Edward Arnold London, 1968.12 M. Winning, G. Gottstein and L. S. Shvindlerman, ‘Migration of grain boundaries under

the influence of an external shear stress. Mater. Sci. Eng., 2001, A317(1–2), 17–20.13 E. E. Underwood, Quantitative Stereology, Addison-Wesley, Reading, MA, 1970.14 H. Chilukuru, W. Blum, M. Schwienheer and A. Scholz, ‘Creep-fatigue interaction

in martensitic tempered steels by dynamic subgrain growth’, In Langzeitverhaltenwarmfester Stähle und Hochtemperaturwerkstoffe, Beiträge zur 26.Vortragsveranstaltung der Arbeitsgemeinschaft für warmfeste Stähle und fürHochtemperaturwerkstoffe am 28. November in Düsseldorf, Stahlinstitut VDEh,2003, 65–74.

15 W. Blum and G. Götz. Evolution of dislocation structure in martensitic CrMoV-steels: the subgrain size as a sensor for creep strain and residual creep life. Steel Res,1999, 70, 274–278.

16 S. Straub, M. Meier, J. Ostermann and W. Blum, ‘Development of microstructureand strengthening in the ferritic steel X20 CrMoV12 1 at 823 K during long-termcreep tests and during annealing’, VGB Kraftwerkstechnik, 1993, 73, 646–653.

17 S. Straub, Verformungsverhalten und Mikrostruktur warmfester martensitischer 12%-Chromstähle, Fortschr.-Ber. VDI Reihe 5 Nr. 405. VDI Verlag, Düsseldorf, 1995.

18 S. Straub and W. Blum, ‘Microstructural development and deformation kinetics iniron-and nickel-base alloys’, In Proceedings of the International Symposium on HotWorkability of Steels and Light Alloys-Composites, H. J. McQueen, E. V. Konoplevaand N. D. Ryan, editors, Montréal, 1996, 189–203.

19 P. Polcik, Mikrostruktur und Verformungsverhalten des Stahles X22 CrMoV12 1nach Zeitstandbeanspruchung im 105 h Bereich, PhD Thesis, Universität Erlangen-Nürnberg, 1993.

20 F. Abe, ‘Effect of quenching, tempering, and cold rolling on creep deformationbehavior of a tempered martensitic 9Cr-1W steel’, Metall. Mater. Trans. A, 2003,34A, 913–925.

21 B. Reppich, ‘Particle strengthening’, In Plastic Deformation and Fracture of Materials,H. Mughrabi (ed), Volume 6 of Materials Science and Technology, (ed. by Cahn, R.W. and Haasen, P. and Kramer, E. J.), VCH Verlagsge-sellschaft, Weinheim, 1993,311–357.

22 K. Mili c∨ka, ‘Internal stress and structure in creep of cold prestrained Fe-21Cr-32Nialloy at 975 K. Metal Sci., 1982, 16, 419–424.

23 B. Wilshire and R. M. Willis, ‘Creep and creep fracture of prestrained type 316Hstain-less steel’, In Proceedings of the 10th Joint International Conference on Creep& Fracture of Engineering Materials and Structures, Creep Resistant MetallicMaterials, M. Filip, V. Foldyna, R. Gladiš, A. Jakobová, Z. Kuboň, J. Purmenskýand J. Sobotka, editors, 8–11 April 2001, Vitkovice-Research and Development andTERIS 2002, Prague, Czech Republic, 2001, 6–15.

24 P. D. Portella, Monotones und zyklisches Kriechverhalten der Legierung 800H bei800°C, PhD Thesis, University of Erlangen-Nürnberg, 1984.

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25 H. Mecking and U. F. Kocks, ‘Kinetics of flow and strain hardening,’ Acta Metall.,1986, 29, 1865–1875.

26 U.F. Kocks and H. Mecking, ‘Physics and phenomenology of strain hardening: theFCC case. Progr. Mater. Sci., 2003 48(3), 171–273.

27 W. Blum, ‘High-temperature deformation and creep of crystalline solids’, In H.Mughrabi, editor, Plastic Deformation and Fracture of Materials, Volume 6 ofMaterials Science and Technology, ed. by Cahn, R. W. and Haasen, P. and Kramer,E. J., VCH Verlagsgesellschaft, Weinheim, 1993, 359–405.

28 W. Blum, S. Straub and S. Vogler, ‘Creep of pure materials and alloys. In D. G.Brandon, R. Chaim and A. Rosen, editors, Proceedings 9th International Conferenceon the Strength of Metals and Alloys (ICSMA 9), Vol. I, Freund Publishing House, 12London, 1991, 111–126.

29 S. U. An, H. Wolf, S. Vogler and W. Blum, ‘Verification of the effective stress modelfor creep of Inconel 617 at 800°C. In B. Wilshire and R.W. Evans, editors, Creep andFracture of Engineering Materials and Structures, The Institute of Metals, London,1990, 81–95.

30 W. Blum and F. Breutinger, ‘New method of determining stress relaxation behaviorin creep machines by controlled unloading’, Z. Metallkd., 2003, 93(7), 649–653.

31 B. Backes, Ermittlung des Verformungswiderstandes eines neuen, hochwarmfesten12% Cr-Stahls durch Spannungsrelaxation, Master’s Thesis, University of Erlangen-Nürnberg, 2003.

32 G. B. Gibbs, ‘Creep and stress relaxation studies with polycrystalline magnesium’,Philos. Mag. A, 1966, 317–329.

33 C. N. Ahlquist and W. D. Nix, ‘A technique for measuring mean internal stressduring high temperature creep’, Scripta Metall., 1969, 679–682.

34 W. Blum and A. Finkel, ‘New technique for evaluating long range internal backstresses’, Acta Metall., 1982, 30, 1705–1715.

35 W. Blum, P. D. Portella and R. Feilhauer, ‘Zyklisches Kriechverhalte’, In B. Ilschner,editor, Festigkeit und Verformung bei hoher Temperatur, Deutsche Gesellschaft fürMetallkunde, Oberursel, 1983, 41–59.

36 H. Wolf, M. Schießl and W. Blum, ‘Acceleration of creep due to cyclic loading of aCrMo-steel’, In H. J. McQueen, J. P. Bailon, J. I. Dickson, J. J. Jonas and M. G.Akben, editors, Proceedings 7th International Conference on the Strength of Metalsand Alloys (ICSMA 7), Pergamon Press, Oxford, Montreal, Canada, 1985, 607–612.

37 W. Blum and J. Granacher, ‘Cyclic creep of heat resistant steels’, In D. G. Brandon,R. Chaim and A. Rosen, editors, Proceedings 9th International Conference on theStrength of Metals and Alloys (ICSMA 9), Vol. I, Freund Publishing House, London,1991, 429–436.

38 S. Straub, P. Polcik, D. Henes and W. Blum, ‘Simulation of the long-term cycliccreep behaviour of a low alloyed ferritic chromium steel’, Mater. Sci. Eng., 1997,A234–236, 1037–1040.

39 A. Seeger, Dislocations and Mechanical Properties of Crystals, Wiley, New York,1957.

40 E. Nes, ‘Modelling work hardening and stress saturation in FCC metals,’ Progr.Mater. Sci., 1998, 41(3), 129–193.

41 H. Oikawa and T. G. Langdon, ‘The creep characteristics of pure metals and metallicsolid solution alloys’, In B. Wilshire and R. W. Evans, editors, Creep Behaviour ofCrystalline Solids, Swansea, Pineridge Press, 1985, 33–82.

42 A. Cottrell and M. Jaswon, ‘Distribution of solute atoms round a slow dislocation’,

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Proc. Roy. Soc. London, Series A, Mathematical and Physical Sci. (1934–1990),1949, 199(1056), 104–114.

43 W. Blum, P. Eisenlohr and F. Breutinger, ‘Understanding creep – a review’, Metall.Trans., 2002, 33A, 291–303.

44 F. Breutinger and W. Blum, ‘Effect of dynamic strain ageing on creep of commerciallypure titanium. In J.D. Parker, editor, Proceedings 9th International Conference onCreep and Fracture of Engineering Materials and Structures, The Institute of Metals,London, 2001, 39–48.

45 W. Blum, Y. J. Li and F. Breutinger, ‘Deformation kinetics of coarse-grained andultrafine-grained commercially pure Ti’. Mater. Sci. Eng. A, 2007, 462, 275–278.

46 B. Schönfeld, ‘Local atomic arrangements in binary alloys’, Prog. Mater. Sci., 1999,44, 435–543.

47 H. Stöcker, Taschenbuch der Physik, Verlag Harri Deutsch, Thun und Frankfurt amMain, 1994.

48 M. Meier and W. Blum, ‘Accounting for subgrain hardening in NiCr22Co12Mo9with the composite model. In B. Wilshire and R. W. Evans, editors, Creep andFracture of Engineering Materials and Structures, The Institute of Materials, London,1993, 167–178.

49 P. Polcik, Modellierung des Verformungsverhaltens der warmfesten 9–12%Chromstähleim Temperaturbereich von 550–650°C. D29, Dissertation, Universität Erlangen-Nürnberg, Shaker Verlag, Aachen, 1999.

50 B. Garofalo, Fundamentals of Creep and Creep Rupture, Macmillan, New York,1965.

51 E. N. da C. Andrade, ‘On the viscous flow in metals and allied phenomena’, Proc.Roy. Soc. Lond., A, 1910, 84, 1–12.

52 B. Ilschner, Hochtemperaturplastizität, Springer, Berlin, 1973.53 H. Wolf, Kriechen der Legierungen NiCr22Co12Mo und 10CrMo9 10 bei konstanter

und zyklischer Beanspruchung, PhD Thesis, University of Erlangen-Nürnberg, 1990.54 F. Abe, ‘Creep rates and strengthening mechanisms in tungsten-strengthened 9cr

steels’, Mater. Sci. Eng., 2001, A319–321, 770–773.55 F. Abe, ‘Strengthening mechanisms in steel for creep and creep rupture’, In Creep

Resistant Steels, F. Abe, U. Kern and R. Viswanathan (eds), Chapter 9. WoodheadPublishing, Cambridge, 2007 (this book).

56 W. Blum, J. Hausselt and G. König, ‘Transient creep and recovery after stressreduction during steady state creep of AlZn. Acta Metall., 1976, 24(4), 293–297.

57 W. Blum, ‘On the evolution of the dislocation structure during work hardening andCreep’, Scripta Metall., 1984, 18(12), 1383–1388.

58 W. Müller, M. Biberger and W. Blum, ‘Subgrain boundary migration during creep ofLiF, III. Stress reduction experiments. Phil. Mag. A, 1992, 66, 717–728.

59 M. Hättestrand, M. Schwind and H. O. Andrén, ‘Mycroanalysis of two creep resistant9–12% chromium steels’, Mater. Sci. Eng. A, 1998, 250, 27–36.

60 J. Hald. Long-term stability of 9–12% Cr Steels – current understanding and futureperspectives. In Werkstoffe und Qualitätssicherung 2004, VGB, Dortmund, March 2004.

61 W. Blum, ‘Creep Simulation,’ In L.-Q. Chen, F. Barlat, F. Roters and D. Raabe,editors, Continuum Scale Simulation of Engineering Materials Fundamentals,Microstructures, Process Applications. Wiley-VCH, Weinheim, 2004.

62 J. S. Dubey, H. Chilukuru, J. K. Chakravartty and W. Blum. On effects of cyclicdeformation and hold periods on subgrain structure and creep behaviour in temperedmartensitic 9–12%CrMoV-steels. Mater. Sci. Eng. A, 2005, 406, 152–159.

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63 W. Blum and H. Chilukuru, ‘Tertiary creep of tempered martensite 9–12 microstructuralorigin of creep life limitation’, In Pressure Vessels and Piping: Materials and Properties,Proceedings International Conference on Pressure Vessels and Piping (OPE 2006–Chennai, Feb 7–9, 2006), B. Raj, B. K. Choudhary and A. Kumar, (eds), NarosaPublishing House, Chennai, ASM International, Alpha Science, 2007.

64 A. Kostka, K.-G. Tak, R. J. Hellmig, Y. Estrin and G. Eggeler, ‘On the contributionof carbides and micrograin boundaries to the creep strength of tempered martensiteferritic steels, Acta Mater., 2007, 55, 539–550.

65 A. Umantsev and G. B. Olson, ‘Ostwald ripening in multicomponent alloys, ScriptaMetall. Mater., 1993, 29(8), 1135–1140.

66 J. Ågren, M. T. Clavagura-Mora, J. Golcheski, G. Inden, H. Kumar, and C. Sigli.Application of computational thermodynamic to phase transformation nucleationand coarsening. Calphad, 2000, 24(1), 41–54.

67 D. Goblirsch, Verformung von X20 CrMoV 12 1 bei konstanter Druckspannungunter besonderer Berücksichtigung der Übergangsverformung nach spannungswechsel,Master’s Thesis, University of Erlangen-Nürnberg, Erlangen, Germany, 1990.

68 R. Agamennone, W. Blum, C. Gupta and J. K. Chakravartty, ‘Evolution ofmicrostructure and deformation resistance in creep of tempered martensitic 9–12%Cr–2%W–5%Co steels. Acta Mater., 2006, 54, 3003–3014.

69 H. Chilukuru. On the Microstructural Basis of Creep Strength and Creep–fatigueInteraction in 9–12% Cr Steels for Application in Power Plants, PhD Thesis, Universityof Erlangen-Nürnberg, Erlangen, Germany, 2007.

70 P. D. Portella and W. Blum, ‘Cyclic creep of Incoloy 800 H at 1073 K’, In J. B.Bilde-Sørensen, N. Hansen, A. Horsewell, T. Leffers and H. Lilholt, editors,Deformation of Multi-phase and Particle Containing Materials, Risø NationalLaboratory, Roskilde, Denmark, 1983, 493–498.

71 H. Wolf and W. Blum, ‘Acceleration and deceleration of creep of a 21/4 Cr–1Mosteel by cyclic stressing at 550°C’, In B. Wilshire and R. W. Evans, editors, Proceedings3rd International Conference on Creep and Fracture of Engineering Materials andStructures, The Institute of Metals, London, 1987, 649–662.

72 P. Eisenlohr, On the Role of Dislocation Dipoles in Unidirectional Deformation ofCrystals. Dr.-Ing. Thesis, Universität Erlangen-Nürnberg, 2004.

73 P. Eisenlohr and W. Blum, ‘Bridging steady-state deformation behavior at low andhigh temperature by considering dislocation dipole annihilation’, Mater. Sci. Eng. A,2005, 400–401, 175–181.

74 P. Weinert, Modellierung des Kriechens von Ferritisch/Martensitischen 9–12%Cr–Stählen auf Mikrostruktureller Basis. PhD Thesis, Technische Universität Graz, Graz,Austria, 2001.

75 D. A. Hughes, N. Hansen and D. J. Bammann. Geometrically necessary boundaries,incidental dislocation boundaries and geometrically necessary dislocations. ScriptaMater., 2003, 48, 147–153.

76 J. Rösler and E. Arzt, ‘A new model-based creep equation for dispersion strengthenedmaterials’, Acta Metall. Mater., 1990, 38, 671–683.

77 G. Eggeler, ‘The effect of long-term creep on particle coarsening in tempered martensiteferritic steels’, Acta Metall., 1989, 37, 3225–3234.

78 E. Cerri, E. Evangelista, S. Spigarelli and P. Bianchi, ‘Evolution of the microstructurein a modified 9Cr–1Mo steel during short term creep’, Mater. Sci. Eng. A, 1998,245(2), 285–292.

79 H. Chilukuru, K. Durst, M. Göken and W. Blum, ‘On the roles of M2X and Z-phase

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in tempered martensitic 9–12% Cr steels’, In J. Lecomte Beckers, M. Carton, F.Schubert and P. J. Ennis, editors, Materials for Advanced Power Engineering,Proceedings of the 8th Liege Conference, volume Part III, 2006, 1181–1190.

80 T. Barkar and J. Agren, ‘Creep simulation of 9–12% cr steels using the compositemodel with thermodynamically calculated input’, Mater. Sci. Eng. A, 2005, 395,110–115.

81 H. Mughrabi, ‘Dislocation walls and cell structures and long-range internal stressesin deformed metal crystals’, Acta Metall., 1983, 31(9), 1367–1379.

82 W. Blum, A. Rosen, A. Cegielska and J. L. Martin, ‘Two mechanisms of dislocationmotion during creep’, Acta Metall., 1989, 37, 2439–2453.

83 S. Vogler and W. Blum, ‘Micromechanical modelling of creep in terms of the compositemodel’, In B. Wilshire and R. W. Evans, editors, Creep and Fracture of EngineeringMaterials and Structures, The Institute of Metals, London, 1990, 65–79.

84 R. Sedlác∨ek and W. Blum, ‘Microstructure-based constitutive law of plasticdeformation’, Comp. Mater. Sci., 2002, 25(1–2), 200–206.

85 D. Henes, H. Möhlig, S. Straub, J. Granacher, W. Blum and C. Berger, ‘Microstructure-based modelling of the long-term monotonic and cyclic creep of the martensitic steelX20(22) CrMoV12 1’, In Microstructure and Mechanical Properties of MetallicHigh-Temperature Materials, H. Mughrabi, G. Gottstein, H. Mecking, H. Riedel andJ. Tobolski, (eds) Wiley-VCH, Weinheim, 1999, 179–191.

86 P. Polcik, S. Straub and W. Blum, ‘Simulation of the deformation behaviour ofmartensitic 9–12% chromium steels on a microstructural basis’, In The 4th EuropeanConference on Advanced Materials and Processes, Associatione Italiana diMetallurgia, Mailand, 1995, 313–318.

87 U. Messerschmidt, M. Bartsch, C. Dietzsch, R. Agamennone, C. Gupta and W.Blum, unpublished results.

88 Y. Qin, G. Götz and W. Blum, ‘Subgrain structure during annealing and creep of thecast martensitic Cr-steel G-X12CrMoWVNbN10-1-1, Mater. Sci. Eng. A, 2003, 341,211–215.

89 B. Holmedal, K. Marthinsen and E. Nes, ‘A unified microstructural metal plasticitymodel applied in testing, processing, and forming of aluminium alloys. Z. Metallkd.,2005, 96(6), 532–545.

90 S. F. Exell and D. H. Warrington, ‘Sub-grain boundary migration in aluminum’,Phil. Mag. A, 1972, 26, 1121–1136.

91 R. Herzog, H. Schuster, M. Weiße and W, ‘Blum. Mikrostruktur und mechanischeEigen-schaften der ODS-Eisenbasislegierung PM 2000 bei quasistationärer undinstationärer Verformung’, In Werkstoffwoche ’96, Symposium 7,Materialwissenschaftliche Grundlagen, F. Aldinger and H. Mughrabi (eds), DGM-Informationsgesellschaft, Frankfurt, 1997, 243–248.

92 U. Hofmann, Kriechverhalten des austenitischen stahles X3 CrNi 18 9 bei 923 K,Master’s Thesis, Universität Erlangen-Nürnberg, Erlangen, Germany, 1991.

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13.12 Appendix: Microstructural model Mikora

Soft (s) subgrains surrounded by hard (h) boundaries (thickness a) deforminelastically (inel) at equal mechanical (mech) strain so that the local stressesdiffer.84

ε ε σ ε σmech inel,s

sinel,h

h = + = + E E

[A.1]

σ = fsσs + fhσh, fs = 1 – fh, f awh = 2 [A.2]

where f are the volume fractions. The local rates of deformation are describedby Equation [13.2]. For steels from the group of 9–12%Cr steels the localdislocation velocities were formulated as:

v AQRTs s

ss*

s* = exp – exp(( ) )

⋅σ βσ ξ [A.3]

v AQRT

b sMk Th h

h2

h p,h

B = exp – sinh

( – )

σ σ[A.4]

σ σ α ρ σs*

s f0.5

p,s p,s = – – MbG C [A.5]

Particle hardening was expressed by Polcik49 by adding the Orowan stressesin the soft region and the Zener stresses in the hard region:

σ p,sp,s,k0.5

k = 3.32 Gb

f

dkΣ [A.6]

σ p,h1 p,h,k

k =

kb

fdk

Σ k1 = 7.7 × 10–9 N. [A.7]

where fp,s,k and fp,h,k are the local volume fractions of the precipitate phasesk existing in the soft and hard regions, respectively. The dislocation interactionparameter α evolves with strain from α = 0 at εinel = 0 as:

dd

= 0.25 – ( / )inel,s f

0.5α

εα

ρb M[A.8]

The empirical factor Cp,s ≤ 1 was introduced by Polcik49 to take intoaccount that the particle strengthening term falls below the Orowan stressσs,Or owing to climb over particles and to ensure that σ s

* > 0 even for smallσ.

The characteristic dislocation spacings w, f–0.5ρ and a evolve with strain

according to Equation [13.10] from their starting values w0 f,0–0.5, ρ = 100

nm17 and a0 = 0.025 w0 towards their steady state values given by Equation[13.1], δ∞ = 0.39 w∞ and a∞ = 0.025 w∞ with kw = 0.12, kδ = 0.0005 and ka

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= 0.003. s was set equal to 7.6 nm. The precipitate sizes evolve with time taccording to Equation [13.9]; see Agamennone et al.68 and Chilukuru69 forvalues of growth constants. The volume fractions fp are assumed to remainconstant except for the Laves phase which precipitates during creep. Thestarting values w0, fp,0 and dp,0 are the microstructural input parameters.

Material parameters: b = 0.248 nm, M = 3, E = 165 GPa, G = 62 GPa.Constants: Qs = Qh = 562 kJ mol–1, As = 3.6 × 1015 m s–1 Pa–1, Ah = 4.4

× 1020 m s–1, β (MPa) and ξ are (i) 0.043 and 1.55 and (ii) 0.48 and 0.45 forσ s

* below and above 62 MPa, respectively.

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403

14Constitutive equations for creep curves and

predicting service life

S . R . H O L D S W O R T H , EMPA – MaterialsScience & Technology, Switzerland

14.1 Introduction

Creep strain curves are determined from the results of continuous-measurementor interrupted tests involving the application of a constant load (or stress) toa uniaxial testpiece held at constant temperature (Fig. 14.1). In continuous-measurement tests, the creep strain, εf, is monitored without interruption bymeans of an extensometer attached to the gauge length of the testpiece(EN 10291, 2000). In interrupted tests, the total plastic strain, εp, is determinedfrom optical measurements of εper at room temperature during plannedinterruptions (where εper = εp – εk i.e. εper = εi + εf – εk, Fig. 14.2, and εp =εper when εk ≈ 0). A list of the symbols and terms used in this chapter is givenin the Nomenclature Section 14.7.

The creep curve data collected in this way may then be modelled by aconstitutive equation. Depending on the nature of the creep model application,the analysis will be of several εf(t) or εp(t) curves determined for a single castor several casts of the specified material. The creep strain curves may havebeen determined from a matrix of t(T, σ) tests for which T and σ are: (1)relatively homogeneously distributed or (2) inhomogeneously distributed interms of T, σ and cast of material. Case (1) is the ideal situation and generallyarises within R&D projects or well co-ordinated data generation activities.Case (2) is more typical of large multi-national datasets, comprising informationfrom many casts, gathered to produce alloy representative creep strengthvalues for Standards (e.g. Holdsworth et al., 2005).

A constitutive equation is a relation between two or more physical quantitieswhich may be simply phenomenological or be directly derived from firstprinciples with a physical basis. The requirement for a representative descriptionof a material’s εf(t, T, σ) creep strain behaviour is no longer just for scientificinterest and metallurgical understanding. The creep deformation behaviourof engineering components is now routinely evaluated using PC-based finite-element analysis tools. Design engineers require the parameters for modelequations to describe the long-time creep behaviour of a specified alloy type

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(not simply the characteristics of a single cast), typically in the primary andsecondary deformation regimes (i.e. the P and S regimes in Fig. 14.1). Incontrast, remaining life assessment engineers are more likely to require thebest model description for a single cast of material, in the secondary andtertiary creep regimes.

The following sections contain a review of creep constitutive equations(Section 14.2) and an approach for assessing model-fitting effectiveness(Section 14.3). In Section 14.4, the application of constitutive equations for

Str

ain

Time

T = constantσ0 = constant

P

S

T

Creep regimes

Au

t u

Str

ain

σ0

εe εi

εt

εp

εf

t = 0t = t1

ET

T = constant

Strain

εper εk εe

14.1 Schematic representation of creep curve showing primary (P),secondary (S) and tertiary (T) deformation regimes.

14.2 Schematic representation of strains generated during a creeptest.

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Constitutive equations for creep curves 405

service life prediction is examined. Finally, there is a consideration of futuretrends (Section 14.5) and some concluding remarks (Section 14.6).

14.2 Constitutive equations

A wide range of creep model equations are in use today to represent the hightemperature-time dependent deformation behaviour of engineering materials(e.g. Table 14.1). Many of these comprise components originating from asmall number of classical representations of primary, secondary and/or tertiarycreep deformation (e.g. Table 14.2, with reference to Fig. 14.1). Typically,the effect of temperature is acknowledged by incorporating an Arhenius, Aexp (–Q/RT), function into the equation (examples are given in Table 14.1).

As a generality, logarithmic creep only occurs at lower temperatures (i.e.below 0.3Tm). At higher temperatures, primary creep is more typicallyrepresented by power, exponential or sinh functions of time (Table 14.2).Similarly, secondary creep rate can be represented by power, exponentialand sinh functions of stress. For secondary creep rate, a sinh formulationreduces to a power law at low stresses and an exponential law at higherstresses. Tertiary creep or creep rate is typically modelled by power orexponential representations, with or without a damage accumulation function,see for example Kachanov (1958). More recently there has been a tendency,in particular for precipitation strengthened alloys, to replace σ in certain ofthe models listed in Table 14.1 by (σ – σi) to acknowledge the existence ofa friction stress, see for example McLean (1980).

The list of constitutive equations contained in Table 14.1 is not exhaustive,but is representative of the range of creep deformation models currentlyemployed within the power generation sector in Europe, as reviewed recentlyby the European Creep Collaborative Committee (ECCC, 2005b).

14.3 Constitutive equation selection

No single constitutive equation effectively represents the creep deformationcharacteristics of all materials over their entire temperature application range.The effectiveness of a constitutive equation to model primary, secondaryand/or tertiary creep deformation for specific applications can vary withmaterial characteristics and source data distribution. In particular, not allmodel equations and fitting procedures are suitable for the prediction ofalloy-mean long-time creep strength behaviour (Holdsworth et al., 2005).

14.3.1 Model fitting effectiveness

The ability of a constitutive equation effectively to characterise the creepdeformation behaviour of a material depends not only on model characteristics

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Table 14.1 Review of creep equations used in ECCC assessmentintercomparisons

Model reference Creep equation

Norton, (Norton, 1929) ε σf,min 1 = exp (– / )d Q RT n

Modified Norton ε f,min = b1 exp(–QB/RT)σn + c1 exp(QC/RT )σn

Norton–Bailey εf = d1σntp

Bartsch εf = e1 exp (–Q1/RT)σ exp(–b1σ)t p

(Bartsch, 1995) + e2 exp(–Q2/RT)σ exp(b2σ)tGarofalo, (Garofalo, 1965) εf = εt[1 – exp(–b1t)] +

ε f,min t

Modified Garofalo εf = εf1[1 – exp(–g1(t/t12)u)

(Granacher, et al., 2001) + ε f,min t + c23(t/t23)f ]

BJF εf = n1[1 – exp(–t)]β + n2t(Jones and Bagley, 1996) where t = (σ/A1)

n exp (–Q/RT)

Li–Akulov model

εε ε ε

εεf

f,min i f,min

f,mins= ln 1 +

– (1 – exp (– )) +

˙ ˙ ˙

˙˙

kkt t

(Li, 1963; Akulov, 1964) + εT(exp (t/tt) – 1)

Theta εf = θ1[1 – exp(–θ2t)] + θ3[exp(θ4t) – 1]

(Evans and Wilshire, 1985) where log(θi) = ai + biT + ciσ + diσTModified Theta εf = θ1[1 – exp(–θ2t)] + θmt + θ3[exp(θ4t) – 1]

where θm = Aσnexp(–Q/RT)

Graham–Walles εf = at1/3 + dt + ft3

(Graham and Walles 1955)Modified Graham–Walles

ε

σ εω

εf( / )= 10

(1 + )1 +

1 1

1

1e Q T A

n

m

+ 10

(1 + )1 +

( / )2 2

2

e Q T A

nσ εω

where ω = e(–QD/T)10AD(σ(1 + ε))nDεmD

Rabotnov–Kachanov ˙ ˙ε σ

ωω σ

ω

ν

ζ =

(1 – ) =

(1 – )1 1h kn

(Kachanov, 1986)Dyson and McClean,

ε ε

σσ ωf 0 d

0 p= (1 + )exp(– / )sinh

(1 – )(1 – )(1 – )

′

D Q RT

HD

(Dyson and McClean, 1998)

Baker–Cane model εf = Atm + εp + φεs + εs(λ – φ)

lt t

– / – 1 –

u

1–

–φφ

φλ φ

(Baker and O’Donnell, 2003) where l = εu/εs, εs = ε m ut and φ = tp/tu

Mech. E (CSWP, 1983) Ru/t/T = (a1 + b1/ε – c1ε2)Rε/t/T + d1 + e1/ε + f1/ε2 – g1ε2

Characteristic strain εf(σ) = ε(Ru/t/T/Rε/t/T – 1)/(Ru/t/T/σ – 1)model (Bolton, 2005a)MHG model, (Grounes, 1969) tε = exp(TF(ε, σ) + C) where the F(ε,σ) function is

(Holmström and freely selected from multilinear combinations of σAuerkari, 2004) and ε with an optimised value of COmega, (Prager, 1995)

˙ ˙ ˙ε ε ε Ωf f,min f,min= /(1 – )t

Modified Omega ε f

tr

= 1 – 12Ω C

(– ln(tu – t) + ln(tu))

(Merckling, 2002) + Ctr(1 – exp(mtrt))

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but also on model fitting approach. In an ECCC creep strain assessmentintercomparison activity, it became clear that model fitting effectiveness canbe strongly influenced by the rigour of the analyst and the model fittingprocedure applied. Three generic approaches were adopted in the ECCCεf(T, σ, t) model-fitting activity:

Approach-1

• model-fitting individual experimental creep curves with the selectedconstitutive equation (simultaneously or consecutively for individualdeformation regimes) to establish the model parameters for specificconditions of T and σ,

• determination of the temperature and stress dependence of the selectedmodel parameters to define the material mean master equation for allεf(T, σ, t).

Approach-2

• determination of specific εf(T, σ, t) coordinates from individualexperimental creep curves either directly (unconstrained by a formalmodel description) or as a result of model fitting (with a model differentto that used for final fitting),

• parametric model-fitting of the specific coordinates to establishparameters to define the mean master equation for the material in theform of either εf(T, σ, t) or εf(T, σ).

Table 14.2 Classical representations of primary, secondary and tertiary creep

Model equation Source reference

Primary creepLogarithmic: εf = a log(1 + bt) Phillips (1905)

Power: εf = at b Graham and Walles (1955)

Exponential: εf = a(1 – exp(–bt)) McVetty (1933)

Sinh: εf = a sinh (btc) Conway and Mullikin (1962)Secondary creepPower:

ε σf,min = d n Norton (1929)Exponential:

ε σf,min = exp ( )d eSinh:

ε σf,min = sinh ( )d e Nadai (1938)Tertiary creepPower: εf = ft g Graham and Walles (1955)

Exponential: εf = f (exp(–gt)–1) McHenry (1943)

ε

σω

f = (1 – )

a n

q where

ω

σω

= (1 – )

c k

r Kachanov (1958)

Rabotnov (1969)Omega:

˙ ˙ε ε εf 0 = exp ( )Ω Prager (1995)

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Approach-3

• derivation of mean σε(T, t) relationships from specific observed tε(T, σ)coordinates from individual experimental creep curves,

• model-fitting with derived σε(T, t) relationships to establish parametersfor the selected model to define the mean master equation for thematerial.

Using these approaches, ECCC investigated ways of quantifying theeffectiveness of a number of constitutive equations to represent the creepdeformation characteristics of large datasets of three alloys, namely 21/4CrMolow alloy ferritic steel (P22), 9CrMoVNb martensitic stainless steel (P91)and 18Cr8Ni austenitic stainless steel (TP316) (Holdsworth et al., 2005).The initial focus was on a large single-laboratory/single-cast multi-temperaturet(T, σ) creep curve dataset for 21/4CrMo steel (Holdsworth and Merckling,2003). The T, σ test conditions responsible for the creep curves in thisdataset were relatively homogeneously distributed to give six rupture timesin the range ~100 to ~3000 h at five temperatures between 510 and 600°C.Moreover, each creep curve comprised approximately the same number ofε(t) observations (~400). The study involved the application of 16 constitutiveequations by nine analysts and resulted in the development of the Z-parameterto provide a measure of model-fitting effectiveness in specific creep strainregimes. It was concluded that as a generality, specific model equations arebetter suited to representing creep strain accumulation characteristics for agiven material in either the primary/secondary regimes or the secondary/tertiary regimes (e.g. Fig. 14.3), although some models can be suitable forboth (e.g. Fig. 14.4). The Z-parameter provides a means of quantifying model-fitting effectiveness (Equation [14.1], Figs 14.3 and 14.4).

log( ) = log ( ) 2.5s = log( ) log( )p / /*

p / / A–RLT p / /t t t ZT T Tε σ ε σ ε σ± ±[14.1]

For a normal distribution, almost 99% of the observed times to specificstrain values would be expected to lie within the boundary lines defined byEquation 14.1. A perfect prediction of tpε/σ/T by the master equation isrepresented by Z equal to zero. Ideally Z should be ≤2 (Holdsworth andMerckling, 2003).

The multi-source/multi-cast multi-temperature datasets for the P91 andTP316 steels were much larger, each containing almost 100 creep curveswith rupture durations ranging from ~100 to >50 000 h (Fig. 14.5). Thetemperatures and stresses for which the creep curves were generated werenot uniformly distributed and the number of ε(t) observations per curveranged between 2 and 250, the creep curves having been generated by bothinterrupted and continuous measurement testing. These datasets were moretypical of those used to form the basis of mean alloy (rather than single

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cast) long time strength values for European Product Standards (Holdsworthet al., 2005). In the case of the P91 and TP316 datasets, 11 constitutiveequations were applied by 11 analysts. For such large inhomogeneouslydistributed datasets, Approaches 2 or 3 appeared to be the most appropriate

14.3 Example comparisons of predicted and observed times to (a)0.2% and (b) 1.0% plastic strain for a constitutive equation providinga poor fit to the experimental data in the low primary/secondarystrain range regime (Z >> 2), but a good fit to the experimental dataat high strains in the secondary/tertiary deformation regime (Z = 2).

510C540C565C580C600C

510C540C565C580C600C

Theta model

z = 2

Eqn. 14.1

Log

pre

dic

ted

tim

e (h

)

3

2

1

0

–1–1 0 1 2 3

Log observed time (h)(a)

Theta model

Eqn. 14.1

Log

pre

dic

ted

tim

e (h

)

3

2

1

0

–1–1 0 1 2 3

Log observed time (h)(b)

z = 2

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for model fitting, both effectively averaging the multi-cast εf(T, σ, t) datain a balanced way prior to final model fitting. Even so, a criterion of Z ≤7 appeared to be more reasonable as a pragmatic model-fit criterion ofacceptability. However, subsequently it became apparent that Z ≤ 2 is still

510C540C565C580C600C

510C540C565C580C600C

Garofalo modifiedmodel

z = 2Eqn. 14.1

Log

pre

dic

ted

tim

e (h

)

3

2

1

0

–1–1 0 1 2 3

Log observed time (h)(a)

Eqn. 14.1

Log

pre

dic

ted

tim

e (h

)

3

2

1

0

–1–1 0 1 2 3

Log observed time (h)(b)

z = 2

Garofalo modifiedmodel

14.4 Example comparisons of predicted and observed times to (a)0.2% and (b) 1.0% plastic strain for a constitutive equation providinggood fits to the experimental data in the primary, secondary andtertiary creep regimes (Z~2).

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attainable for strains in the secondary to tertiary creep regimes, even forsuch complex alloy datasets when the assessment involves a quantitativeconsideration of cast-to-cast variability, for example, by strength compensation(Bolton, 2005a).

123b123e123m123o

H1 (CT)H2 (CT)H3 (CT)H4 (CT)H5 (IT)H6 (IT)

1 101 102 103 104 105

Time (h)(a)

Str

ain

(%

)

102

101

1

10–1

10–2

170

180

165 150164

140

140

120

115100

90

80

80

72

5056

10–1 1 101 102 103 104 105

Time (h)(b)

Str

ain

(%

)

102

101

1

10–1

10–2

190123125

75

165

100 8090 78

62

49

40

98

90 6298

6312550

14.5 Examples of the distribution (at one temperature only) of largemulti-source/multi-cast multi-temperature datasets typical of thoseused for determining long time creep strength values for EuropeanProduct Standards, (a) P91 at 600°C, and (b) TP316 at 700°C. (CTmeans continuous measurement test, IT means interruptedmeasurement test.)

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14.3.2 Model selection

The selection of constitutive equation and model-fitting approach can dependon a number of factors including material characteristics, data distributionand practical application.

Material characteristics

The effectiveness of a creep equation in representing a material’s characteristicε(t) curve shape can depend on features such as the relative proportions ofprimary, secondary and tertiary creep as fractions of the strain and time atrupture (e.g. Fig. 14.6(a)) and the way in which they vary over the T, σregime of interest (Fig. 14.6(b)). The Z-parameter (Equation [14.1]) providesa useful indication of model effectiveness in these circumstances.

Data distribution

Model selection and the choice of model-fitting approach can be influencedby the distribution of the data to be assessed. A creep dataset typicallyconsists of a number of ε(t, T, σ) curves (creep test records), eachcomprising a number of ε(t) observations, for example Fig. 14.5. Both theε(t, T, σ) curve and ε(t) distribution characteristics are influential (see Section14.3.1).

Practical application

Model selection can also depend on the purpose for which the material’screep strain description is required. Typically, the priority of the scientist isfor a creep constitutive equation to have a sound physical basis in the primary,secondary and tertiary regimes. As a generality, it is more important fordesign and assessment engineers for the constitutive equation to be simple toimplement and effective in its description of creep deformation at long times.For design engineers, effective modelling is more important in the relativelylow strain primary/secondary creep regimes whereas for remaining lifeassessment engineers, the priority is more likely to be an accurate knowledgeof secondary/tertiary behaviour.

14.4 Predicting service life

There is no universally preferred constitutive equation for predicting servicelife. In practice, the selection often depends on which model best representsthe high temperature deformation characteristics of the material and thepreference of the analyst and/or the requirements of the available application

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tools. An approach for testing model-fitting effectiveness is described in theprevious section.

The life assessment of most engineering structures involves considerationof the multiaxial stress state at a critical location and it is conventionally

450C, 520MPa450C, 430MPa650C, 120MPa650C, 80MPa

0.0 0.2 0.4 0.6 0.8 1.0Normalised time

(a)

Steel-91600°C

21/4CrMo540°C

Type-316650°C

No

rmal

ised

pla

stic

str

ain

1.0

0.8

0.6

0.4

0.2

0.0

0.0 0.2 0.4 0.6 0.8 1.0Normalised time

(b)

No

rmal

ised

pla

stic

str

ain

1.0

0.8

0.6

0.4

0.2

0.0

14.6 Examples of creep curve shape variations for (a) 21/4CrMo,Steel-91 and Type-316 at typical application temperatures, and (b)Steel-91 over a wide T, σ range.

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assumed that multi-axial creep deformation is characterised by the von-Mises effective stress, σVM . For a given σVM , multiaxial loading does notexert a strong influence on deformation characteristics during primary andsecondary creep (Fig. 14.7), but has a strong effect on tertiary creep deformationbehaviour and time to rupture (Dyson and Osgerby, 1993) and in particularthe magnitude of the rupture ductility (Rice and Tracey, 1969; Cocks andAshby, 1980; Shammas and Marchant, 1986).

σ σ σ σ σ σ σVM 1 22

2 32

3 12 = 1

2[( – ) + ( – ) + ( – ) ] [14.2]

When the end-of-life criterion is rupture and the critical location in thecomponent is subject to multi-axial loading, the simple use of σVM as therepresentative stress is not always appropriate (see Section 14.4.2).

14.4.1 End-of-life criteria

Constitutive creep equations may be used to predict service life directly in adefect-free design assessment of a component subject to steady loading athigh temperature. In such circumstances, an important consideration is theend-of-life criterion. For most engineering applications, this is unlikely to berupture (or the attainment of rupture strain as determined in a uniaxial test).More usually, the end-of-life is a strain limit, for example to avoid loss ofclearance or interference during service duty (Bolton, 2005b). Alternatively,it is a stress-state dependent rupture strain (Rice and Tracey, 1969; Cocks

14.7 Schematic representation of influence of stress state on creepdeformation behaviour.

Cre

ep s

trai

n

Time

T, σVM

Multi-axial

Uniaxialtension

Torsion

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and Ashby, 1980; Spindler, 2003), with an applied safety factor, or aconservative estimate of uniaxial rupture strain such as that provided by theMonkman–Grant relationship ε f,min ut (Monkman and Grant, 1956).

14.4.2 Multi-axial stress rupture

When rupture (or a safe fraction of the rupture time) is the end-of-life criterion,one approach for the determination of service life in components subject tomultiaxial loading is to use an appropriate creep constitutive equation first todetermine the steady-state stress distribution. With a knowledge of themultiaxial stress rupture criterion obeyed by the material, an appropriaterepresentative stress can then be computed for critical locations in thecomponent (see below). Ultimately, the service life may be based on uniaxialtu(T, σ) creep-rupture relationships (e.g. ECCC, 2005a) for which the stressinput is the calculated representative stress state.

The following examples do not provide an exhaustive list of representativestress formulations (see also Webster et al., 2004). Equations [14.3] to [14.6]are respectively due to Sdobyrev (1958), Othman and Hayhurst (1993), Cane(1979) and Huddleston (1985):

σrep = ασ1 + (1 – α) σVM (with 0 ≤ α ≤ 1) [14.3]

σrep = ασ1 + 3βσm + (1 – α – β) σVM (with 0 ≤ α + β ≤ 1) [14.4]

σ σ σ σλrep 1 VM

/VM = ( / ) n (with 0 ≤ γ ≤ n) [14.5]

σ σ σrep VM m s = exp [ (3 / – 1)]C S (with = + + )s 12

22

32s σ σ σ[14.6]

where α, β and γ are parameters reflecting the multi-axial rupture criterionobeyed by the material and are determined from the results of uniaxial andmulti-axial rupture tests, see for example Webster et al., 2004). In Equation[14.6], C is a material constant (e.g. 0.24 for austenitic stainless steels). Suchformulations acknowledge that creep rupture may be controlled by themaximum principal stress, σ1, the von-Mises effective stress, σVM , and/orthe hydrostatic stress, σm.

Creep life assessment can also be determined by analysing the stress–strain state of the structure in its entirety using detailed finite element solutionswith demanding material property input data requirements or by using moresimplified approaches involving some form of reference stress, e.g. skeletalpoint stress (Calladine, 1969), equivalent yield stress (Goodall et al., 1979),structural stress (Bolton, 2005b). The application of such reference stressapproaches is usually limited to creep ductile materials for which bothdeformation and failure are controlled by σVM .

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14.4.3 Defect assessment

Service lives determined on the basis of a high temperature defect assessmentprocedure also require the implementation of appropriate creep constitutiveequations, e.g. (R5, 2003). For assessments involving the application of hightemperature failure assessment diagrams (HTFADs), or creep crack initiationor creep crack growth rate model equations, it is first necessary to determinethe steady-state stress distribution and creep deformation response at criticallocations in the component.

14.5 Future trends

The continuously increasing availability of higher performance, lower costcomputer technology will lead to the wider practical application of the mosteffective constitutive equations for predicting service life. In principle, therewill be no limit to the complexity of the constitutive equations which can beimplemented. In practice, the most commonly adopted equations will bethose requiring readily available and/or determinable stress and temperature-dependent material property parameters.

There will continue to be research effort aimed at developing constitutiveequations that require results from short duration creep tests to predict longtime service lives. While this is a laudable objective, it will require morevision and advanced computing power than is currently being applied toaddress the difficulty of reliably extrapolating beyond the mechanism regime(s)for which there exist experimentally determined material property data.


The chapter has reviewed the formulation of a wide range of creep constitutiveequations currently adopted for scientific research and predicting servicelife. No single constitutive equation effectively represents the creep deformationof all materials over their entire temperature application range. An approachfor quantifying model-fitting effectiveness has been reviewed. Importantconsiderations for the prediction of service life at high temperatures are theselection of representative stress to characterise deformation and rupture dueto multiaxial loading in the component material, and the end-of-life criteria.A number of options have been examined.

14.7 Nomenclature

ECCC European Creep Collaborative CommitteeMSRC multiaxial stress rupture criterionn stress exponent

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Q activation energy for creepR universal gas constantRpε/t/T, Ru/t/T creep strength and rupture strength for a given time and

temperaturesA-RLT standard deviation of all residual log timest timetu, tu,max observed time to rupture, maximum observed time to rupturetpε/σ/T, t*pε/σ/T observed and predicted times to given plastic straintεf(T,σ) time to a specific creep strain as a function of temperature

and stressT, Tm temperature, melting temperature of materialZ Parameter quantifying effectiveness of master creep

equation to predict times to specific strains(see Equation [14.1])

ε, εe, εi strain, elastic strain, instantaneous plastic strainεf, εp, εk, εper creep strain, plastic strain, anelastic strain, permanent strain˙ ˙ ˙ε ε ε, ,f,min ave strain rate, minimum creep strain rate, average strain rateεu creep rupture ductilityσ stressσi friction stressσrep representative rupture stressσ1, σ2, σ3 principal stresses (where σ1>σ2>σ3)σm hydrostatic stress, i.e. σm = (σ1 + σ2 + σ3)/3σVM Von-Mises effective stressσε(t, T) stress to give a specific strain as a function of time and

temperatureω ω, ˙ damage, rate of damage accumulationECCC terms and terminology recommendations are given in Volume 2 (ECCC,2005b)

14.8 References

Akulov N S (1964), ‘On dislocation kinetics’, Acta Metall., 12, 1195.Baker A J and O’Donnell M P (2003), ‘R5 high temperature structural integrity assessment

of a cracked dissimilar metal weld vessel test’, in Proceedings 2nd InternationalConference on Integrity of High Temperature Welds, 10–12 November 2003, IOM andI. Mech. E, London.

Bartsch H (1995), ‘A new creep equation for ferritic and martensitic steels’, Steel Res., 66(9), 384–388.

Bolton J L (2005a), ‘A ‘characteristic-strain’ model for creep’, in Proceedings ECCCCreep Conference on Creep & Fracture in High Temperature Components – Design& Life Assessment Issues, Shibli I A, Holdsworth S R and Merckling G, (eds) I. Mech.E., 12–14 September 2005, London, 465–477.

Bolton J L (2005b), ‘Analysis of structures based on a characteristic-strain model of

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creep’, in Proceedings ECCC Creep Conference on Creep & Fracture in HighTemperature Components – Design & Life Assessment Issues, Shibli I A, HoldsworthS R and Merckling G, (eds), I. Mech. E., 12–14/9/05, London, 1032–1045.

Calladine C R (1969), ‘Time scales for redistribution of stress in creep of structures’,Proc. Roy. Soc. London., A309, 363–375.

Cane B J (1979), ‘Creep cavitation and rupture in 21/4CrMo steel under uniaxial andmultiaxial stresses’ in Proceedings 2nd International Conference on MechanicalBehaviour of Materials, Miller K J and Smith R F (eds) Pergamon Press, Oxford, 2,173–182.

Cocks A C F and Ashby M F (1980), ‘Intergranular fracture during power-law creepunder multi-axial stress’, Met Sci., 14, 395–402.

Conway J B and Mullikin M J (1962), ‘An evaluation of various first stage creep equations’,in Proceedings AIME Conference, Detroit, Michigan.

Creep of Steels Working Party (CSWP) (1983), High Temperature Design Data for FerriticPressure Vessel Steels, Institute of Mechanical Engineers, London.

Dyson B F and McClean M (1998), ‘Microstructural evolution and its effects on the creepperformance of high temperature alloys’, in Strang A and McClean M (eds),Microstructural Stability of Creep Resistant Alloys for High Temperature Applications,IOM, 371–393.

Dyson B F and Osgerby S (1993), Modelling and Analysis of Creep Deformation andFracture in a 1Cr1/2Mo Ferritic Steel’, NPL Report DMM(A)116.

ECCC (2005a), Data Sheets for Rupture Strength, Creep Strength and Relaxation StrengthValues for Carbon–Manganese, Low Alloy Ferritic, High Alloy Ferritic and AusteniticSteels, Nickel Base Alloys and High Temperature Bolting Steels/Alloys, Robertson DG and Holdsworth S R (eds), ECCC(ETD) publishers.

ECCC (2005b), Recommendations for Creep Data Validation and Assessment Procedures,Holdsworth S R, Brown T B, Buchmayr B, Bullough C K, Calvano F, et al. (eds) Vol.1: Overview, Vol. 2: Terms and terminology, Vol. 3: Data acceptability criteria, Datageneration, Vol. 4: Data exchange and collation, Vol. 5: Data assessment, Vol. 6:Characterisation of microstructure and physical damage for remaining life assessment,Vol. 7: Data assessment – creep crack initiation, Vol. 8: Data assessment – multi-axial,Vol. 9: Component assessment, ECCC(ETD) publishers.

EN 10291 (2000), Metallic Materials, Uniaxial Creep Testing in Tension, Method of Test,European Norm.

Evans R W and Wilshire B (1985), Creep of Metals and Alloys, Institute of Metals,London.

Garofalo F (1965), Fundamentals of Creep and Creep Rupture in Metals, New York,Macmillan.

Graham A and Walles K F A (1955), ‘Relations between long and short time propertiesof commercial alloys’, JISI, 179, 105–120.

Granacher J, Möhlig H, Schwienheer M and Berger C (2001), ‘Creep equation for hightemperature material’, in Proceedings 7th International Conference on Creep & Fatigueat Elevated Temperatures (Creep 7), 3–8/6/01, Tsukuba, NRIM, 609–616.

Grounes M (1969), ‘A reaction rate treatment of the extrapolation methods in creeptesting’, J. Basic Eng., Series D, Trans ASME.

Goodall J W, Leckie F A, Ponter A R S and Townley C H A (1979), ‘The development ofhigh temperature design methods based on reference stresses and design methods’, J.Eng. Mat. Tech., 101, 349–355.

Holdsworth S R and Merckling G (2003), ‘ECCC developments in the assessment of

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creep-rupture data’, in Proceedings 6th International Charles Parsons Conference onEngineering Issues in Turbine Machinery, Power Plant & Renewables, Trinity College,Dublin, 16–18 September 2003.

Holdsworth S R, Askins M, Baker A, Gariboldi E, Holmström S, Klenk A, RingelM, Merckling G, Sandström R, Schwienheer M. and Spigarelli S (2005), ‘Factorsinfluencing creep model equation selection’, in Proceedings. ECCC Conference onCreep & Fracture in High Temperature Components – Design & Life AssessmentIssues, 12–14 September 2005, eds. Shibli I A, Holdsworth S R and Merckling G,(eds), I. Mech. E., London.

Holmström S and Auerkari P (2004), ‘Prediction of creep strain and creep strength offerritic steels for power plant applications, in Proceedings Baltica Conference on LifeManagement and Maintenance for Power Plants, VTT Symposium 234, 8–10 June2004, Espoo.

Huddleston R L (1985), ‘An improved multi-axial creep-rupture strength criterion’, TransASME, J. Press. Vessel Technol., 107, 412–429.

Jones D I G and Bagley D L (1996), ‘A renewal theory of high temperature creep andinelasticity’, in Proceedings Conference on Creep and Fracture: Design and LifeAssessment at High Temperature, London, 15–17/4/96, MEP, 81–90, London.

Kachanov L M (1958), ‘Time to failure under creep conditions’, Izv. Akad. Navk. SSR.Otd Teck. Nauk., 8, 26–31.

Kachanov L M (1986), Introduction to Continuum Damage Mechanics, Kluwer Academic,Dordrecht.

Li J C M (1963), ‘A dislocation mechanism for transient creep’, Acta Metall., 11, 1269.McHenry D (1943), ‘A new aspect of creep in concrete and its application to design’,

Proc. ASTM, 43, 1069.McLean M (1980), ‘Friction stress and recovery during high-temperature creep: interpretation

of creep transients following a stress reduction’, Proc. Roy. Soc. London, Series A,Mathematical Phys. Sci., 371 (1745), 279–294.

McVetty P G (1933), ‘Factors affecting the choice of working stresses for high temperatureservice’, Trans ASME, 55, 99.

Merckling G (2002), ‘Metodi di calcolo a confronto per la previsione dellulteriore esercibilitàin regime di scorrimento viscoso’, in Proceedings Conference on Fitness for Service,Giornata di Studio CESI-CONCERT, Milan, 28 November 2002.

Monkman F C and Grant N J (1956), ‘An empirical relationship between rupture life andminimum creep rate in creep-rupture tests’, Proc. ASTM, 56, 593–620.

Nadai A (1938), ‘The influence of time upon creep, The hyperbolic sine creep law’, inStephen Timoshenko Anniversary Volume, Macmillan, New York.

Norton F N (1929), The Creep of Steel at High Temperature, McGraw-Hill.Othman A M and Hayhurst D R (1993), ‘Determination of large strain multi-axial creep

rupture criterion using notched bar data’, Int. J. Damage Mech., 2, 16–52.Phillips F (1905),‘The slow stretch in india rubber, glass and metal wire when subjected

to a constant pull’, Phil. Mag., 9, 513.Prager M (1995), ‘Development of the MPC Omega method for life assessment in the

creep range’, ASME J. Pressure Vessel Technol., 117, May, 95–103.Rabotnov Yu N (1969), Creep Problems in Structural Members, North-Holland, Amsterdam.Rice J R and Tracey D M (1969), ‘On the ductility enlargement of voids in triaxial stress

fields’, J. Mech. Phys. Solids, 17, 201–217.R5 (2003), Assessment Procedure for the High Temperature Response of Structures,

Procedure R5 Issue 3, British Energy, Gloucester, UK.

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Sdobyrev V P (1958), ‘Long term strength of alloy EI-437B under complex stresses’, Izv.Akad. Nauk. SSR. Otd. Teck. Nauk., 4, 92–97.

Shammas M S and Marchant K D (1986), ‘Torsion testing in an inert atmosphere’, inTechniques for Multiaxial Creep Testing, Gooch D J and How I M. (eds), ElsevierApplied Science, London.

Spindler, M W (2003), ‘The multi-axial creep ductility of austenitic stainless steel’,Fatigue Fract. Eng. Mater. Struct., 27, 273–281.

Webster G A, Holdsworth S R, Loveday M S, Nikbin K., Perrin I J, Purper H, Skelton RP and Spindler M W (2004), ‘A code of practice for conducting notched bar creep testsand for interpreting the data, Issue 3’, Fatigue and Fract. Eng. Mater. Struct., 27(4),319–342.

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421

15Creep strain analysis for steel

B . W I L S H I R E and H . B U R T University of WalesSwansea, UK

15.1 Introduction

When a tensile stress (σ) is applied to a metal or alloy at a temperature (T),the total strain (εtot) which accumulates over a time (t) can be expressed as:

εtot = ε0 + ε [15.1]

where ε0 is the initial strain on loading and ε is the subsequent creep strain.The loading strain depends on stress and temperature as:

ε0 = ƒ1(σ, T) [15.2]

allowing the magnitude of ε0 to be discussed by reference to the values ofYoung’s modulus (E), the yield or 0.2% proof stress (σY) and the ultimatetensile stress (σTS) determined from high-strain-rate tests (~10–3s–1) at the creeptemperature. Hence, ε0 is predominantly elastic when σ < σY (with ε0 = σ/E),whereas ε0 has elastic and significant plastic components when σ > σY. Incontrast to ε0 (Equation [15.2]), ε varies with stress, temperature and time as:

ε = ƒ2 (σ, T, t) [15.3]

with the creep strain/time characteristics being dependent on the T/Tm ratio,where Tm is the absolute melting point. When T < 0.4Tm, the total creepstrains are low and failure rarely occurs. Conversely, when diffusion cantake place at around 0.4Tm and above, the creep strains can be large, leadingto eventual failure. Thus, while creep at low temperatures is normally oflittle practical significance, creep and creep fracture are often the life-limitingphenomena during component service in power plant and other high-temperature applications.

For design of large-scale components in power plant, a knowledge isusually required of stresses which the relevant steels can sustain at the operatingtemperatures without failure occurring in 100 000 h. To provide thisinformation, major experimental programmes involving tests lasting up to30 000 h or more are currently completed but, for many steel grades, results

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have already been obtained for stress–temperature conditions giving failuretimes exceeding 100 000 h.

To reduce the scale and cost of these programmes, reliable techniquesmust be evolved to allow accurate estimation of long-term properties byextrapolation of short-term measurements. However, confidence in anypredictive methodology is improved when the analysis procedures have asound theoretical basis. For this reason, a variety of theoretical and practicalapproaches to high-temperature creep and creep fracture are now summarizedin relation to straightforward descriptions of the changes in creep strain/timebehaviour as the stress and temperature conditions are altered (Equation[15.3]). In this way, a coherent foundation is laid for consideration of thefactors affecting creep strain accumulation, as well as the times and strainsto failure. The industrial implications of these concepts are then addressedby defining physically meaningful techniques for rationalization of thebehaviour patterns observed for power plant steels, before introducing easilyapplicable extrapolation procedures for long-term design data prediction.

To ensure that the approaches adopted and the results achieved are opento scrutiny, the present analyses use only the internationally respected datasets1,2 produced by the National Institute for Materials Science (NIMS),Japan. In particular, attention is focused on two 9% Cr steels, namely, Grade91 (9 Cr–1Mo–V–Nb) and Grade 92 (9 Cr–0.5Mo–1.8W–V–Nb).

15.2 Creep-induced strain

Under a sustained tensile stress at temperatures above 0.4Tm, most metalsand alloys exhibit normal creep strain/time curves. Thus, following the initialloading strain, the creep strain rate ( ε = dε/dt) decreases continuously withtime during the primary stage, reaching a minimum or secondary rate ( )mεbefore accelerating during the tertiary stage which leads to fracture after atime (tf). The product, ε m ft , is often but not always a constant (M), signifyingthat creep failure is strain controlled because tf increases as ε m falls withdecreasing stress and temperature.

Numerous relationships have been proposed to quantify the variations increep strain with time. Several of these equations seek to describe only theearly stages of the creep curves, while others attempt to define the shape ofthe entire ε/t trajectories, but no agreement has been reached on the relationshipswhich should be used. Despite the distinctive shape of normal curves, it hastherefore become common practice to ignore the primary and tertiary stages,assuming that the secondary rate remains constant with increasing time andstrain. Equation [15.3] then reduces to:

ε m = ƒ3 (σ, T) [15.4]

with a further simplification giving:

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Creep strain analysis for steel 423

ε m = ƒ4 (σ)ƒ5(T) [15.5]

so the variables are treated as separate and independent.In Equation [15.5], at a fixed temperature, the stress dependence of ε m

can be defined as:

ε m ∝ ƒ4 (σ) ∝ σn [15.6]

where n is the stress exponent. Alternatively:

ε m ∝ ƒ4 (σ) ∝ exp σ [15.7]

giving an exponential dependence of ε m on stress. Conversely, at a fixedstress, the temperature dependence of ε m in Equation [15.5] is generallyrepresented by an Arrhenius equation of the form:

ε m ∝ ƒ5 (T) ∝ exp (–Qc/RT) [15.8]

where Qc is the activation energy for creep in units of J mol–1 when the gasconstant, R = 8.314 J mol–1 K–1. Combining Equations [15.5], [15.7] and[15.8] then gives:

M/tf = ε m = B exp [–(Qc – V σ)/RT] [15.9]

where B and V are treated as constants. However, in most theoretical andpractical studies carried out over the last half century, Equations [15.5],[15.6] and [15.8] have been combined to obtain the standard power lawrelationship:

M/tf = ε m = A σn exp (–Qc/RT) [15.10]

but the values of the parameter, A, as well as n and Qc, vary in differentstress/temperature regimes.

15.2.1 Parametric approaches to data analysis

Although NIMS Creep Data Sheet No. 43 (1996) details only the stressrupture properties,1 results available from other sources3,4 allow the creeplives to be considered in relation to the creep rate characteristics of tubesamples of Gr. 91 steel. Thus, using Equation [15.10], the log ε m /log σ plotsin Fig. 15.1 can be represented3 by a set of straight lines showing a decreasefrom n ≅ 16 at 848 K to n ≅ 9 at 923 K. Similarly, the stress/creep liferelationships determined over extended stress ranges at 773–973 K for multiplebatches of Gr. 91 tube1 reveal gradient changes corresponding to decreasesfrom n ≅ 17 to n ≅ 4.5 with increasing temperature (Fig. 15.2). With Qc

ranging from 600 to 700 kJ mol–1, these anomalously large values of n andQc are typical of the behaviour patterns reported for power plant steels andother particle-hardened alloys.

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Because of the complex stress and temperature dependences of ε m and tf

(Figs. 15.1 and 15.2), estimation of long-term properties by extrapolation ofshort-term measurements involves the continued use of various parametricmethods introduced in the 1950s.5–7 These empirical approaches define‘correlation parameters’, incorporating both creep life and temperature, whichcan be plotted as functions of stress to superimpose multi-batch results ontoa single ‘master curve’ for a given steel. Unfortunately, no one parametric

15.1 Stress dependence of the minimum creep rates recorded for Gr.91 tube steel3 at 848–923 K.

848K : n ≅ 16873K : n ≅ 11898K : n ≅ 10923K : n ≅ 9

Min

imu

m c

reep

rat

e (s

–1)

10–5

10–6

10–7

10–8

10–9

10–10

75 100 150 200 300Stress (MPa)

773K823K873K923K973K

Str

ess

(MP

a)

500

300

200

100

50

20106 107 108

Time to fracture (s)

15.2 Stress dependence of the creep lives recorded for multiplebatches of Gr. 91 tube steel1 at 773–973 K.

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method has proved capable of fitting the experimental data reported for themajority of power plant steels and, even when the best fitting procedure isselected, the accuracies achieved are not always satisfactory.8

One limitation inherent in parametric methods is linked to the ‘variableconstants’ encountered with standard power law relationships (Equation[15.10]). For instance, Equation [15.10] can be re-written to give the parameter(POSD) proposed by Orr, et al.,7 as:

POSD = log tf – (P1/T ) [15.11]

where P1 includes Qc. Hence, variations in Qc ensure that the superimposedparametric plots are non-linear. The curvatures of plots based on Equation[15.11] are not then removed by replacing Equation [15.10] with Equation[15.9]. In fact, Equation [15.9] can be rearranged to give the parameter(PLM) proposed by Larson and Miller5 as

PLM = T(log tf + P2) [15.12]

where P2 now contains Qc.Combined with the broad scatter bands generally associated with multi-

batch data sets, the unknown curvatures of traditional parametric plots limitextrapolation to only about three times the longest reliable test measurementsavailable. For this reason, tests lasting up to 30 000 h and more must usuallybe completed to estimate 100 000 h rupture strengths.

15.2.2 Alternative procedures for data rationalization

For pure metals, rearranging Equation [15.10] and using the activation energiesexpected for diffusion, the stress/creep rate relationships recorded at differenttemperatures are superimposed9 simply by plotting the dependences on (σ/E) of the temperature-compensated creep rate, ε m exp(Qc/RT). Similarly,the corresponding stress rupture properties are rationalized by plotting thetemperature-compensated creep lives, tf exp (–Qc/RT), against (σ/E). However,this approach is not applicable to power plant steels, because minor variationsin the thermomechanical processing conditions selected for componentmanufacture can alter the resulting microstructures.

Elastic moduli are temperature dependent but do not vary markedly withchanges in microstructure, whereas the creep and fracture properties of particle-hardened alloys are both temperature and microstructure sensitive, as are σY

and σTS. Hence, early rationalization procedures based on normalizing σthrough E9 are up-dated by normalizing σ through σY or σTS.10–13 In thisway, using the σTS values measured for each batch of Gr. 91 steel investigated,1

the multi-batch stress rupture data in Fig. 15.2 are superimposed in Fig. 15.3using a modified power law expression:11–13

M/tf = ε m = A*(σ/σTS)n exp (– / )c*Q RT [15.13]

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where A* ≠ A and Qc* is obtained from the temperature dependence of ε m at

constant (σ/σTS) rather than at constant σ as in the determination of Qc in

Equation [15.10]. From Fig. 15.3, Qc* = 300 kJ mol–1, a value close to that

for lattice diffusion in the alloy steel matrix. With Gr. 91 steel, as with othermetals and alloys,10–14 creep property sets are also rationalized effectivelyby normalizing σ through σY, so that (σ/σTS) can be replaced by (σ/σY) inEquation [15.13]. Irrespective of whether σY or σTS is chosen, the resulting‘master curve’ in Fig. 15.3 is at least as impressive as that obtained usingparametric methods. Moreover, with Equation [15.13], the empirical termsin parametric relationships (Equations [15.11] and [15.12] are replaced byphysically meaningful properties, namely, a sensible activation energy andthe measured σY and σTS values.

15.2.3 Interpretation of power law behaviour

Equation [15.13] avoids the large and variable Qc values observed when datasets for Gr. 91 steel are described using Equation [15.10] (Fig. 15.2), butdoes not eliminate a decrease from n ≅ 20 to n ≅ 4 (Fig. 15.3), a trendgenerally expected to continue towards n ≅ 1 or less as the test duration andtemperature increase.

One early attempt to explain the anomalously large n values suggested15

that creep occurs not under the full applied stress (σ) but under a reducedstress (σ – σo), such that

773K823K873K923K973K

σ/σ T

S

0.8

0.6

0.4

0.2

0.110–15 10–14 10–13 10–12 10–11 10–10 10–9 10–8

tf exp (–Qc*/RT) (s)

15.3 Dependence of the temperature-compensated creep life on(σ/σTS), using the σTS values for each batch of Gr. 91 tube steelinvestigated,1 with Q c

* = 300 kJ mol–1.

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ε m ∝ (σ – σo)m [15.14]

where m = 4, with σo now called a ‘threshold stress’. Comparing Equations[15.10] and [15.14], n ≅ m ≅ 4 when σo ≅ 0 or when σo ∝ σ, whereas n > mwhen σo is large. This approach has been widely applied to creep of particle-hardened alloys,16 but little progress has been made because σo cannot bemeasured or predicted reliably.

More commonly, n value variations have been interpreted on the basisthat different creep mechanisms become dominant in different stress–temperature regimes. Thus, while the superimposed results in Fig. 15.3suggest that the gradient changes continuously as (σ/σTS) decreases, suchcurves could be approximated by a series of straight line segmentscorresponding to n >> 4, n ≅ 4 and eventually n ≅ 1, with each gradientchange linked to a mechanism transition. However, no agreement has beenreached on the detailed processes involved and, from a practical viewpoint,current theories do not allow prediction of the creep and creep fractureproperties of engineering steels. Moreover, accepting multi-mechanismconcepts, analysis of results recorded in one mechanism regime would notallow prediction of properties in another regime, necessitating the completionof long-term test programmes. In this context, it therefore seems reasonableto consider whether standard power law relationships offer a valid basis forrepresentation and interpretation of creep properties.

With Equation [15.10], the fact that n and Qc are themselves functions ofstress and temperature means that, in the simplification of Equation [15.4] toobtain Equation [15.5], the variables are not separate and independent. Inaddition, the assumption that the variations in creep strain with time, stressand temperature (Equation [15.3]) can be quantified adequately through thestress and temperature dependences of a secondary or ‘steady state’ creeprate (Equation [15.4]) is highly questionable. For these reasons, withoutinvoking mechanism transitions, an alternative approach17 contends that thebehaviour patterns displayed by power plant steels (Figs. 15.1, 15.2 and15.3) are easily understood through Equation [15.3] by considering thesystematic changes in creep curve shape which occur with increasing testduration and temperature.

15.3 Patterns of creep strain accumulation

Although conventional ε/t plots give the impression that normal creep curvesexhibit clearly defined primary, secondary and tertiary stages prior to failure,this view is contradicted by the available experimental evidence. Instead, formost metals and alloys,4,10–14,17 plotting the variations of the tensile creepstrain rate ( ε ) with time or strain demonstrates that a minimum rate ( ε m )rather than a secondary or ‘steady state’ value is reached when the decaying

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primary rate is offset by the tertiary acceleration, as illustrated for Gr. 91tube steel4 in Fig. 15.4. Hence, the creep rate at any instant ( ε ) should beexpressed as:

˙ ˙ ˙ε ε ε = + p t [15.15]

where the primary rate ( )pε decays and the tertiary creep rate ( )tε accelerates

with time. The ε m and tf data derived from sets of normal creep curvesshould then be interpreted in terms of the deformation mechanisms controllingstrain accumulation and the damage processes causing the creep rate toaccelerate during the tertiary stage.

In line with Equation [15.15], inspection of the ε /t trajectories in Fig.15.4 emphasizes that much information is lost by ignoring the primary andtertiary characteristics when creep mechanism studies have traditionally focusedon ‘steady state’ behaviour. Thus, most equations introduced to describecreep curve shapes feature a ‘steady-state’ term. Yet, by explicitly quantifyingthe decaying primary and accelerating tertiary components (Equation [15.15]),the θ Projection Concept17 provides accurate descriptions of the curve shapedependence on stress and temperature, allowing extrapolation of short-termdata to provide long-term property estimates.10,12,17 However, thecomprehensive sets of high-precision constant-stress creep curves requiredto carry out valid θ analyses have been produced for relatively few materials.Even so, the dominant deformation and damage processes can be clarifiedsimply by monitoring some basic quantities defining the shapes of normalcreep curves.

Cre

ep r

ate

(s–1

)

10–5

10–6

10–7

10–8

10–9

10–10

10–11

103 104 105 106 107 108

Time (s)

100MPa110MPa

120MPa

200MPa 160MPa 140MPa

15.4 Variations of the creep strain rate ( ε ) with time (t) for tubesamples of Gr. 91 steel4 over a stress range from 100 to 200 MPa at873 K.

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15.3.1 Variations in creep curve shape

With power plant steels, creep and stress rupture tests are generally performedunder tensile stresses less than σY so, from Equation [15.1], the essentiallyelastic loading strain (εo) is small in relation to the total creep strain tofailure, often termed the creep ductility (εf). In turn, εf can be regarded as thesum of the primary creep strain (εp) accumulated over the period from t = 0to t = tm (where tm is the time to the minimum rate), plus the tertiary strain(εt), where εt = (εf – εp). Using these easily visualized quantities, several keytrends become evident from the ε /t trajectories in Fig. 15.4 as the stress isreduced from 200 to 100 MPa at 873 K for Gr. 91 steel.4

The εf value falls from over 0.3 towards 0.2 as the creep life increasesabove about 10 000 h at stresses below 140 MPa, accompanied by a gradualdecrease from εp ≅ 0.03 towards εp < 0.01 with increasing test duration (Fig.15.5). The decaying primary rate is then offset by the tertiary acceleration ata progressively earlier fraction of the creep life, so the tm/tf ratio decreasesas tf increases. The creep curves therefore become increasingly tertiarydominated as the primary stage becomes less significant with decreasingapplied stress.

With Gr. 91 steel at 873 K (Fig. 15.5), the modest εp values fall graduallywith decreasing stress, a trend which would be expected when creep occursby movement of dislocations in the alloy steel matrix. This view is consistent

with the present observation that Qc* = 300 kJ mol–1, (Fig. 15.3). In fact, for

Gr. 91 and other power plant steels, detailed microstructural studies haveindicated that dislocation processes are dominant under long-term testconditions4 and even under the low stress levels experienced during power-plant service.18

εfεm · tf

εp

.

ε f

0.5

0.4

0.3

0.2

0.1

0.080 100 120 140 160 180 200 220

Stress (MPa)

0.06

0.04

0.02

0.00

εm

ft &

εp

15.5 Variations of the primary strain (εp), the creep ductility (εf) andthe product, εm ft with stress at 873 K for Gr. 91 tube steel (from Fig.15.4).

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While no major change appears to occur in the dislocation mechanismscontrolling strain accumulation over extended stress–temperature ranges,several different processes can start the tertiary acceleration. These includecavity and crack development, neck formation and evolution of themicrostructures of particle-hardened alloys. In addition, once the tertiarystage begins, more than one process can influence the subsequent rates ofstrain accumulation, while the damage mechanisms initiating tertiary creepand causing failure may also differ. The relative importance of differentdamage processes can then vary as the test conditions alter, causing the creepcurve shape to change in the manner illustrated in Fig. 15.4.

15.3.2 Creep damage tolerance values

An informative method of quantifying the tertiary characteristics is introducedbecause the product, ε m ft , is linked to the creep ductility, εf, as:

λ ε ε = /f m f˙ t [15.16]

where λ is the creep damage tolerance parameter.19 Strictly, because damagedevelopment influences the tertiary not the primary stage, εf should be replacedby εt in Equation [15.16] when the primary strains are substantial.20,21

In practice, the contribution of εp to εf can usually be neglected becauseεf >> εp (Fig. 15.5), allowing λ to be calculated from the values of ε m , tf andεf reported for power plant steels. The measured λ values then relate to thedamage processes initiating tertiary creep.

In many instances, the tertiary acceleration has been modelled by defininga damage parameter (ω), which is zero for undamaged material.22 The valueof ω increases as the damage levels increase, so:

˙ ˙ε ε ωt o = (1 + ) [15.17]

with the rate of damage accumulation ( )ω increasing as ε t accelerates froman initial value ( )oε when t = 0. Adopting this approach, modelling exercises23

have predicted that λ ≅ 1.5 to 2.5 when tertiary creep and fracture are attributableto cavitation, with higher values expected when the tertiary stage beginsas a consequence of necking (with λ > 2.5) or precipitate coarsening (withλ > 5).

Unfortunately, experimental studies covering the tertiary behaviour ofvarious metals and alloys20,21 have shown that the λ value alone does notallow unambiguous identification of the dominant damage mode. When λprogressively exceeds about 1.5 with increasing test duration and temperature,the tertiary acceleration is due to microstructure evolution, for example thestrength loss associated with precipitate coarsening. In contrast, whenλ ≅ 1.5, tertiary creep can begin by intergranular or transgranular cracking

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and/or necking, which must be distinguished through use of additionalapproaches such as microstructural studies, profiling of fractured testpieces,and so on.

As evident from Fig. 15.5, although εf decreases, the accompanying fallin ε m ft means that an increase from λ ≅ 5 towards λ > 15 occurs as the testduration increases from about 40 to almost 40 000 h at 873 K. This result, aswell as the patterns of strain accumulation for Gr. 91 steel (Fig. 15.4), supportsthe outcomes of microstructural investigations4 showing that the tertiaryacceleration starts through evolution of the initial tempered martensite structure.Although dislocation processes appear to govern creep strain accumulationat all stress levels, the creep curve shape changes because microstructureevolution causes the minimum rate to occur at a progressively earlier fractionof the creep life with increasing test duration and temperature (Fig. 15.4).The gradual decrease in n value (Figs. 15.1, 15.2 and 15.3) is then attributable,not to creep mechanism transitions, but to the complex dependence of ε m

and tf on the systematic variations in the creep curve shape as (σ/σTS) decreases.However, the curve shape variations may also be accompanied by changes inthe creep ductility (εf), as indicated in Fig. 15.5.

Under tensile creep conditions, the creep life can be considered convenientlyas the time taken for the creep strain to reach the limiting creep ductility, thatis, t → tf as ε → εf. An abrupt fall in εf during long-term exposure could thenresult in a reduction in the failure times expected by projection of short-termdata. Hence, in seeking valid extrapolation procedures, it is necessary toclarify the manner in which εf varies over stress–temperature ranges leadingto fracture in times up to 100 000 h or more. Yet, compared with the emphasisplaced on measurement and interpretation of ε m and tf data, comparativelylittle attention has been devoted to the factors affecting creep ductility (εf)and the related reduction of cross-sectional area at fracture (RoA).

15.3.3 Creep ductility

With power plant steels at a fixed creep temperature, εf is often observed todecrease with increasing test duration, as illustrated in Fig. 15.5, but detailedductility trends may be masked by the broad scatter bands encountered withmulti-batch εf measurements. Hence, the εf and RoA values recorded for Gr.91 tube samples1 are plotted, not against stress (Fig. 15.5), but against thetemperature-compensated creep life, tf exp (– / )c

*Q RT , with Qc* = 300 kJ

mol–1 (Fig. 15.6). Although a fall in εf is not immediately apparent becauseof scatter, there is an obvious decrease in the corresponding RoA values withincreasing test duration and temperature (Fig. 15.6). In fact, this RoA decreaseoccurs with failure times of less than 100 000 h at 883 K, the upper servicetemperature for Gr. 91 steel.

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The decrease in εf as tf increases (Fig. 15.5) has been linked to the loss ofcreep strength caused by progressive evolution of the martensitic microstructure,leading to the suggestion that extrapolation of short-term tf measurementsoverestimates long-term performance.4 Noting a related change in the stressdependence of the creep lives, it has been recommended that results obtainedunder the higher stress ranges covered at each test temperature should beexcluded from data projection exercises,8,24 making long-term test programmesmandatory. However, this conclusion appears to be negated by several featuresof the ductility characteristics displayed by Gr. 91 steel:

• Under all test conditions investigated,1 the results presented in Fig. 15.6reveal that the measured reduction of area at fracture (RoA) is significantlygreater than the creep ductility (εf). Fracture is therefore preceded byneck formation, indicating that failure occurs in a relatively ductile manner,even during long-term creep exposure.

• As evident from Fig. 15.4, even in the test lasting almost 40 000 h at 873K, the creep rate is accelerating rapidly late in the creep life. With mostof the creep strain accumulating just prior to failure, only a major decreasein εf would cause a marked reduction in tf.

• The magnitude of λ (Equation [15.16]) is important in practical situationswhen high strains develop in regions, say, where a change in componentcross-section leads to stress concentrations. With λ values of 5 or moreconsidered to be adequate, the present estimates suggest that the localizedstress concentrations typically encountered during service should notlead to premature cracking of Gr. 91 pipework, even though problemsare being experienced with type IV failure of weldments.25

773K823K873K923K973K

Open symbols: εfClosed symbols: RoA

ε f a

nd

Ro

A

1.0

0.8

0.6

0.4

0.2

0.010–17 10–16 10–15 10–14 10–13 10–12 10–11 10–10 10–9 10–8


15.6 Changes in creep ductility (εf) and reduction of area at fracture(RoA) as a function of the temperature-compensated creep life, tf

exp (– / )c*Q RT , with Q c

* = 300 kJ mol–1, using data obtained1 formultiple batches of Gr. 91 tube steel at 723–923 K.

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These observations suggest that, despite the property trends shown in Fig.15.6, the failure mode is essentially unchanged over extended stress–temperature ranges. Hence, with no definitive evidence for transitions in thedominant creep mechanism, there seems to be no a priori reason why analysisof short-term measurements should not allow prediction of long-term properties.Indeed, the procedures shown to rationalize multi-batch tf values (Figs. 15.2and 15.3) introduce straightforward extrapolation methods for estimation of100 000 h design data.12,13

15.4 Practical implications of creep strain analysis

Even without conceiving new extrapolation procedures, replacing Equation[15.10] by Equation [15.13] allows standard multi-batch tf measurements(Fig. 15.2) to be superimposed onto master curves (Fig. 15.3), suggestingthat the total number of tests required to supply 100000 h creep ruptureestimates can be reduced substantially.10 Thus, tf measurements could bemade for one batch of steel at the planned service temperature for stressescausing failure in times up to 100 000 h. By determining σTS at differenttemperatures for several batches, only a limited number of additional testswould be required to confirm that an activation energy of 300 kJ mol–1

allows effective data rationalization. Standard stress rupture plots (Fig. 15.2)can then be computed from the resulting ‘master curves’ (Fig. 15.3). However,not only the number of tests but also the maximum test durations can bereduced through alternative relationships12,13 introduced to describe thedependence of the temperature-compensated creep lives on (σ/σTS).

15.4.1 Creep life extrapolation for Gr. 91 steel

With σTS representing the maximum stress which can be applied at the creeptemperature, the failure times approach zero as σ → σTS, whereas infinitecreep lives must be recorded when σ → 0. These criteria are met by replacingEquation [15.13] with a modified version of Equation [15.9],12,13 giving:

( / ) = exp – [ exp (– / )] TS 1 f c*σ σ k t Q RT u [15.18]

where k1 and u are evaluated by plotting ln [ exp (– / )]f c*t Q RT against

ln[–ln(σ/σTS)], with Qc* = 300 kJ mol–1, as illustrated in Fig. 15.7. In this

case, k1 and u were determined for results acquired1 only under test conditionsgiving creep lives of less than 5000 h, although it is evident from Fig. 15.7that virtually identical values would be derived by including all measurementspresented in Fig. 15.3. Using Equation [15.18], the ‘master curve’ in Fig.15.8 then demonstrates that tf → 0 as (σ/σTS) → 1 and tf → ∞ as (σ/σTS) →0, with a point of inflection occurring at about 0.5 σY, providing an impressivedescription of the stress rupture properties reported1 for Gr. 91 steel.

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15.7 Dependence of ln [tf exp (– / )c*Q RT ] on ln [−ln (σ/σTS)], using the

σTS values recorded for each batch of Gr. 91 tube steel investigated,1

with Q c* = 300 kJ mol–1. Test conditions where tf < 5000 h are shown

as closed symbols, while results for tf > 5000 h are included as opensymbols.

773K823K873K923K973K

ln [

t f ex

p (

– Qc*

/RT )

] (s

)

–20

–22

–24

–26

–28

–30

–32

–34

k1 = 26.37u = 0.134

–1.5 –1.0 –0.5 0.0 0.5 1.0ln [–ln(σ/σTS)]

773K823K873K873K923K973K

σ/σ T

S

0.8

0.6

0.4

0.2

0.010–16 10–15 10–14 10–13 10–12 10–11 10–10 10–9 10–8 10–7


15.8 Dependence of log [tf exp (– / )c*Q RT ] on (σ/σTS), using the σTS

values recorded for each batch of Gr. 91 tube steel investigated,1

with Q c* = 300 kJ mol–1. Results for tests carried out at stresses

greater than 0.5 σY are presented as closed symbols and for tests atstresses less than 0.5 σY as open symbols, with the behaviourpattern predicted from Equation [15.18] shown as a solid line.

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Knowing k1 and u (Fig. 15.7), plus the σTS values for the different batchesof Gr. 91 tube,1 Equation [15.18] also allows tf values to be computed overextended stress ranges at various creep temperatures. These predicted curvesare included in Fig. 15.9, together with recently reported multi-batch tf values,produced through NIMS for test conditions giving creep lives up to 100 000h for Gr. 91 plates, pipes, forgings and tubes.26

Clearly, the current projections derived using Equation [15.18] to describetf values for tests lasting less than 5000 h fit well with the general propertytrends in Fig. 15.9, despite the inevitable scatter in the long-term measurements.However, to assess the predictive accuracy of Equation [15.18], noting thatthe present estimates were obtained for Gr. 91 tube samples only, the predictionsfrom Fig. 15.9 are compared in Table 15.1 with recent 100 000 h rupturestrength estimates from other sources.26–28 Several features of the results inTable 15.1 then merit comment:

• In two cases,26,27 the estimates in Table 15.1 were obtained using theLarson–Miller relationship (Equation [15.12]). Both analyses highlightthe fact that very different strength estimates are produced, dependingon the values chosen for the empirical constant.

• In Table 15.1, the 100000 h rupture strengths were estimated by Kimura26

using the latest long-term NIMS data for Gr. 91 steel (Fig. 15.9). Yet,contrary to recent suggestions,4,8,24 the property projections now derivedby applying Equation [15.18] to the short-term NIMS measurements

773K823K873K923K973K

103 104 105 106 107 108 109


Str

ess

(MP

a)

500

300

200

100

50

20

Time to fracture (h)100 101 102 103 104 105

15.9 Stress/creep life behaviour predicted from Equation [15.18] forGr. 91 tube steel (solid lines) compared with long-term multi-batchstress rupture data recorded26 for plate, pipe, forging and tubesamples at 773–973 K.

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(i.e. tf < 5000 h) slightly underestimate rather than overestimate theresults achieved by excluding the high-stress tf values at each testtemperature from the extrapolation exercise based on the long-termproperties.26

• Given the spread in the 100000 h strength values obtained using parametricmethods (Table 15.1), the present predictions appear reasonable, agreeingwell with the outcomes of extensive long-term test programmes28

undertaken through the European Collaborative Creep Committee (ECCC).

Having considered the creep and creep fracture properties of Gr. 91 steel,as a further check on the general applicability of Equation [15.18], the samedata analysis methods are now adopted to discuss the stress rupture behaviourof Gr. 92 steel. Again, for this material, microstructure studies combinedwith parametric extrapolation exercises have suggested that the inclusion ofresults from test lasting less than about 3000 h cause overestimation of long-term performance.29,30

15.4.2 Creep life extrapolation for Grade 92 steel

The data sets reported for Gr. 92 tube2 essentially replicate the behaviourpatterns shown for Gr. 91 tube in Figs. 15.1 and 15.2, with Equation [15.13]again superimposing the stress/creep life relationships at different testtemperatures, in line with Fig. 15.3. Hence, to evaluate the predictivecapabilities of Equation [15.18], the stress–rupture properties are plotted inFig. 15.10 to show the variation of ln[ exp (– / )]1 c

*t Q RT as a function of

Table 15.1 100000 h creep rupture estimates (MPa) for Gr. 91 steel at 823, 873and 923 K

Temp (K) Present Ref 26a Ref 27b Ref 28c

estimatesA B A B

823 154 150 132 159 153 150873 87 98 83 100 86 85923 43 42 48 56 46 44

aEstimates obtained using the Larson–Miller relationship to analyse data produced atstresses less than 180, 130 and 90 MPa at 823, 873 and 923 K, respectively whereas,by describing the full data sets from tests lasting up to almost 100000 h, significantlylarger estimates were determined for the 100000 h strengths.bEstimates obtained using the Larson–Miller relationship with constants of 20 (Set A)and 36 (Set B).cEstimates obtained for tubes and pipes, with the longest test having a creep lifeexceeding 110 000 h under a stress of 150 MPa at 823 K (Set A). Set B quotes the 2005estimates produced through the European Collaborative Creep Committee (ECCC).

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ln[–ln(σ/σTS)]. In contrast to the single-line fit achieved for Gr. 91 (Fig.15.7), it is evident from Fig. 15.10 that the values of k1 and u in Equation[15.18] change as the temperature-compensated creep life increases with theGr. 92 samples. Specifically, linear extrapolation of the results obtainedunder high stresses at low temperatures overestimates the creep rupture strengthsunder low stresses at high temperatures (Fig. 15.10).

These changes in k1 and u appear to support the proposal26,29,30 that short-term tf values should not be included when seeking to predict long-termstress-rupture properties. However, it is apparent from Fig. 15.10 that thelong-term behaviour of Gr. 92 steel can be determined through Equation[15.18] using only the data recorded under stress–temperature conditionsgiving maximum creep lives of around 5000 h or so. In fact, the k1 and uvalues in Fig. 15.10 were calculated on this basis. Essentially, the tf valuesdetermined at 823 and 873 K represent the failure properties when (σ/σTS) ishigh, while the results at 973 and 1023 K describe the stress rupturecharacteristics when (σ/σTS) is low. In agreement with the results presentedfor Gr. 91 steel (Fig. 15.8), Fig. 15.11 then shows that incorporating thederived k1 and u values into Equation [15.18] provides a sensible descriptionof the tf data for Gr. 92 steel, that is, tf → 0 as (σ/σTS) → 1, again with a pointof inflection at σ ≅ 0.5 σY ensuring that tf → ∞ as (σ/σTS) → 0.

Knowing the k1 and u values (Fig. 15.10), Equation [15.18] also allowsprediction of the trends in the log tf/log σ relationships obtained2 from testslasting up to 30 000 h at different temperatures for multiple batches of Gr. 92tube and pipe (Fig. 15.12). In addition, as found for the Gr. 91 material

15.10 Dependence of ln[tf exp (– / )c*Q RT ] on ln[–ln(σ/σTS)], with

Q c* = 300 kJ mol–1, using data reported2 for Gr. 92 tube steel. Results

are presented for creep lives up to about 5000 h or so as closedsymbols and for longer term data as open symbols.

823K873K923K973K1023K

ln [

t f ex

p (

– Qc*

/RT)

] (s

)

–16

–18

–20

–22

–24

–26

–28

–30

–32–1.0 –0.8 –0.6 –0.4 –0.2 0.0 0.2 0.4 0.6 0.8

ln [–ln(σ/σTS)]

k1 = 8.35u = 0.097

k1 = 37.96u = 0.16

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823K873K923K973K1023K

σ/σ T

S

0.8

0.6

0.4

0.2

0.010–15 10–14 10–13 10–12 10–11 10–10 10–9 10–8 10–7


Gr.91

15.11 Dependence of log[tf exp (– / )c*Q RT ] on (σ/σTS)] using the σTS

values recorded for each batch of Gr. 92 tube steel investigated,2

with Q c* = 300 kJ mol–1. Results for tests carried out at stresses

greater than 0.5 σY are presented as closed symbols and for tests atstresses less than 0.5 σY as open symbols, with the behaviourpattern predicted from Equation [15.18] shown as a solid line. Forcomparison purposes, the behaviour predicted for Gr. 91 tube steel(Fig. 15.8) is included as a broken line.

823K873K923K973K1023K

104 105 106 107 108 109


Str

ess

(MP

a)

500

200

100

50

20

10

Time to fracture (h)101 102 103 104 105

15.12 Stress/creep life behaviour predicted from Equation [15.18],using the k1 and u values derived in Fig. 15.10 (solid lines),compared with the multi-batch stress rupture data recorded2 at823–1023 K for tube and pipe samples of Gr. 92 steel.

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(Table 15.1), the present estimates of the 100 000 h creep rupture strengthsat 823, 873 and 923 K are certainly reasonable in relation to the ranges ofvalues derived27,29 when the detailed analysis procedures differ using theLarson–Miller relationship (Table 15.2). However, the current predictionsare of particular interest when compared with the strength estimates reportedfor Gr. 92 steel over the last decade.

In 1995, a 100 000 h strength value of 132 MPa at 873 K was determinedthrough Larson–Miller assessments31,32 of results from 284 tests at 823to 1023 K, with the longest test duration being 43 513 h at 873 K.Using these and additional results, four different analysis methods were thenapplied33–36 to 704 creep life measurements, giving revised 873 K strengthvalues of 116 to 129 MPa, with a figure of 123 MPa presented in the 1999ECCC data sheet. In 2005, again adopting several different procedures toanalyse over 800 data points recorded at 823 to 973 K (with a maximumcreep life exceeding 110 000 h), the 873 K strength estimate was furtherreduced30 to within the range 107 to 118 MPa (Table 15.2).

This progressive reduction in the 100 000 h strength values seems toconfirm the importance of long-term test measurements, simultaneouslyjustifying the omission of short-term results from the data analysis exercises.Yet, the present predictions obtained by applying Equation [15.18] only to tfvalues up to around 5000 h or so are in excellent agreement with the strengthsrecently quoted30 for creep lives of 100, 1000, 10 000 and 100000 h at 823,873 and 923 K in Table 15.3. On this basis, rather than focussing attentionon extending the test durations and selective editing of short-termmeasurements, it seems that the parametric procedures still widely adoptedfor data analysis should be seriously reevaluated.

Table 15.2 100000 h creep rupture estimates (MPa) for Gr. 92 steel at 823, 873and 923 K

Temp (K) Present Ref 29a Ref 27b Ref 30c

estimatesA B A B

823 182 183 192 141 186 187 to 190873 104 112 128 91 120 107 to 118923 53 51 70 52 70 53 to 62

aUsing the Larson–Miller relationship to describe the stress rupture data recorded intests giving creep lives up to 70 000 h, the estimated 100000 h creep rupture strengthsderived by excluding tf measurements less than 3000 h (Set A) are substantially lowerthan the predictions obtained by analysing all results available (Set B).bEstimates obtained using the Larson–Miller relationship with constants of 20 (Set A)and 36 (Set B).cEstimates obtained using several different procedures to analyse extensive datasets, with a maximum creep life of 110 450 h at 873 K. See Table 15.3.

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15.4.3 Appraisal of proposed data analysis procedures

With both Gr. 91 and Gr. 92 steels, microstructural studies4,26,28,29 haveestablished that the patterns of creep strain accumulation and failure arestrongly influenced as the tempered martensite structures of the as-receivedbatches evolve with increasing time and temperature. Yet, while the overallproperty trends seem similar (Figs. 15.8 and 15.11), the effects of the strengthloss caused by microstructure evolution appear to be more severe with Gr.92 than with Gr. 91 steel, as reflected by the fact that the k and u valueschange in Fig. 15.10 but not in Fig. 15.7. These data comparisons thenindicate that, using Equation [15.18], detailed differences in long-termbehaviour can be diagnosed by analysis of short-term property sets. Giventhe scale of attention currently being directed to identification of reliablemethodologies for prediction of long-term stress rupture properties,33–41 it istherefore useful to consider several key features of the approaches nowproposed for data rationalization and extrapolation.

• By plotting the temperature-compensated creep lives as a function ofeither (σ/σTS) or (σ/σY), stress rupture measurements obtained over broadstress–temperature ranges are superimposed onto well-fitted ‘mastercurves’, using the activation energy for lattice diffusion in the alloy steelmatrices (i.e. Qc

* = 300 kJ mol–1). In this way, data rationalization can

Table 15.3 Comparisons of the present predictions obtained using Equation[15.18] with estimates of the stresses (MPa) causing creep failure of Gr. 92steel in 100, 1000, 10 000 and 100 000 h at 823–923 K derived by applyingseveral different analysis procedures to long-term data30

Temperature Analysis Creep rupture strength (MPa) for times of(K) method 100 h 1000 h 10 000 h 100 000 h

823 Graphical 276 251 221 190ISO 6303 284 254 223 188Larson–Miller 294 257 221 187Present 272 245 215 182estimate

873 Graphical 209 183 149 107ISO 6303 213 184 150 111Larson–Miller 223 186 151 118NIMS 150 109Present estimate 202 177 148 104

923 Graphical 147 122 90 53ISO 6303 153 121 86 56Larson–Miller 157 122 89 62NIMS 86 55Present estimate 144 121 86 53

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be achieved through Equation [15.13] (Fig. 15.3) or Equation [15.18](Fig. 15.8). In either case, without assuming distinctive mechanismtransitions, the results can be explained in terms of the dislocation processescontrolling creep strain accumulation and the damage processes causingthe tertiary acceleration and eventual failure.

• Avoiding the problems associated with the choice of the empirical constantsin parametric relationships (Tables 15.1 and 15.2), the coefficients (k1

and u) in Equation [15.18] are easily determined (Figs. 15.7 and 15.10),allowing reasonable prediction of long-term stress rupture properties byanalysis of tf measurements from tests lasting only about 5000 h (Figs.15.9 and 15.12). Of course, with any forecasting method, the accuracyof the predictions must inevitably improve as the extent of the extrapolationdecreases. Because forecasts are generally based on continuity of trends,the shorter the extent of the extrapolation, the lower is the risk of unforeseenevents influencing the predictions. Even so, the use of short-term dataoffers advantages with steels prone to severe oxidation, that is thepredictions are not significantly affected by the stress intensificationassociated with the loss of cross-sectional area during long-term creepexposure.10

• Inspection of the data sets superimposed using Equation [15.18] revealsthat a point of inflection in the sigmoidal curves occurs at around 0.5 σYfor both the Gr. 91 and Gr. 92 steels (Figs. 15.8 and 15.11). This observationcoincides with the proposal by Kimura8 that only data determined atstresses less than 0.5 σY should be incorporated in extrapolation exercisesbased on traditional parametric relationships. However, the generalapplicability of the 0.5 σY inflection in Figs. 15.8 and 15.11, as well asits interpretation, would require analysis of data sets for a range of othersteels.

In seeking to validate the proposed rationalization and extrapolationprocedures, the means of assessment are already available through the NIMSCreep Data Sheets. These comprehensive reports detail not only the stressrupture properties but also a wide range of other relevant information, includingthe measured σY and σTS values for each batch of steel investigated.Additionally, these documents should even allow evaluation of other dataanalysis concepts now suggested for future study.

15.5 Future data analysis options

Earlier investigations12,13 have demonstrated that, as with the rationalizationand extrapolation of stress rupture properties through Equation [15.18], ε m

data recorded over extended stress-temperature ranges can also be analysedas:

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( / ) = exp – [ exp( / )] TS 2 m c*σ σ εk Q RT v˙ [15.19]

with the coefficients (k2 and v) again determined by plotting the temperaturecompensated creep rate as a function of either (σ/σTS) or (σ/σY). The sigmoidalcurves produced through Equation [15.19] then satisfy the essential criterionthat ε m → ∞ as tf → 0 when (σ/σTS) → 1, with a point of inflection againensuring that ε m 0→ as tf → ∞ when (σ/σTS) → 0.

In a similar manner, the times to reach certain pre-defined creep strains(tε) can be described as:

(σ/σTS) = exp –k3 [tεexp (– / )c*Q RT ]w [15.20]

with a set of k3 and w values quantifying the behaviour at different strains.Provided that full creep curves are recorded, a series of tε exp (– / )c

*Q RTagainst (σ/σTS) plots at various creep strains must map onto the appropriatetf exp (– / )c

*Q RT against (σ/σTS) relationship when ε = εf so tε = tf (Figs. 15.8and 15.11).

Adopting these procedures, detailed computer-efficient descriptions ofthe changes in creep strain (or strain rate) with time, stress and temperature(Equation [15.3]) could be made available for incorporation into the modernfinite element codes developed for high-temperature engineering design.Moreover, this form of property representation would link directly with plotsof εf and RoA measurements as functions of either the temperature-compensatedcreep lives (Fig. 15.6) or the (σ/σTS) and (σ/σY) values, clarifying long-termcreep ductility trends. Overall, these new approaches should then provide aneffective development avenue for future analysis, interpretation andextrapolation of creep and creep fracture data for power plant steels andother creep-resistant alloys.

15.6 References

1 NIMS Creep Data Sheet No. 43, Data sheets on the elevated-temperature propertiesof 9Cr-1Mo-V-Nb steel tubes for boilers and heat exchangers (ASME SA-213/SA-213M Grade T91) and 9Cr–1Mo–V–Nb steel plates for boilers and pressure vessels(ASME SA-387/SA-387M Grade 91), 1996.

2 NIMS Creep Data Sheet No. 48, Data sheets on the elevated-temperature propertiesof 9Cr-0.5Mo-1.8W-V-Nb steel tube for power boilers (ASME SA-213/SA-213MGrade T92) and 9Cr–0.5Mo–1.8W–V–Nb steel pipe for high temperature service(ASME SA-335/SA-335M Grade P92), 2002.

3 Spigarelli S, Cerri E, Bianchi P and Evangelista E, ‘Interpretation of creep behaviourof a 9Cr–Mo–Nb–V–N (T91) steel using threshold stress concept’, Mater Sci Tech,1999, 15, 1433–1440.

4 Kimura K, Kushima H and Abe F, ‘Heterogeneous changes in microstructure anddegradation behaviour of 9Cr–1Mo–V–Nb steel during long term creep’, Key EngMater, 2000, 171–174, 483–490.

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5 Larson F R and Miller J, ‘A time-temperature relationship for rupture and creepstresses’, Trans ASME, 1952, 74, 765–775.

6 Manson S S and Haferd A M, ‘A linear time-temperature relation for extrapolationof creep and stress rupture data’, NASA TN 2890, 1953.

7 Orr R L, Sherby O D and Dorn J E, ‘Correlations of rupture data for metals atelevated temperature’, Trans ASM, 1954, 46, 113–128.

8 Kimura K, ‘Present status and future prospect of NIMS creep data sheet’, CreepDeformation and Fracture, Design and Life Extension, R S Mishra, J C Earthman,S V Raj and R Viswanathan (eds), MS&T, Pittsburgh, 2005, 97–106.

9 Barrett C R, Ardell A J and Sherby O D, ‘Influence of modulus on the temperaturedependence of the activation energy for creep at high temperatures’, Trans AIME,1964, 230, 200–204.

10 Wilshire B and Burt H, ‘Yield stress rationalization of creep and creep fractureproperties’, Scripta Mater, 2005, 53, 909–914.

11 Wilshire B and Burt H, ‘A unified theoretical and practical approach to creep andcreep fracture’, Creep Deformation and Fracture, Design and Life Extension, R SMishra, J C Earthman, S V Raj and R Viswanathan (eds), MS&T, Pittsburgh, 2005,3–12.

12 Burt H and Wilshire B, ‘Theoretical and practical implications of creep curve shapeanalyses for 7010 and 7075’, Metall Mater Trans A, 2006, 37A, 1005–1015.

13 Wilshire B and Burt H, ‘Creep data rationalization for power plant steels’, MaterialScience Forum, 2007, 539–543, 254–261.

14 Wilshire B and Battenbough A J, ‘Creep and creep fracture of polycrystalline copper’,Mater Sci Eng A, 2007, 443, 156–166.

15 Williams K R and Wilshire B, ‘On the stress and temperature dependence of creepof Nimonic 80A’, Metal Sci J, 1973, 7, 176–179.

16 Arzt E, ‘Creep of dispersion strengthened materials: A critical assessment’, ResMech, 1991, 31, 399–453.

17 Evans R W and Wilshire B, Creep of Metals and Alloys, Institute of Metals, London,1985.

18 Williams K R and Wilshire B, ‘Effects of microstructural instability on the creep andfracture behaviour of ferritic steels’, Mater Sci Eng, 1977, 28, 289–296.

19 Leckie F A and Hayhurst D R, ‘Constitutive equation for creep rupture’, Acta Metall,1977, 25, 1059–1070.

20 Wilshire B and Burt H, ‘Tertiary creep of metals and alloys’, Z Metallkd, 2005, 96,552–557.

21 Wilshire B and Burt H, ‘Damage evolution during creep of steels’, Creep and Fracturein High Temperature Components – Design and Life Assessment Issues, I A Shibli,S R Holdsworth and G Merckling (eds), DESTech Publ, London, 2005, 191–200.

22 Kachanov L M, ‘On the time to failure under creep conditions’, Izv Acad NaukUSSR, Otd TeKd Nauk, 1957, 8, 26–38.

23 Ashby M F and Dyson B F, ‘Creep damage mechanics and micromechanisms’,Advances in Fracture Research, S R Valluri (ed.), Pergamon Press, Oxford, 1993,Volume 1, 3–30.

24 Foldyna V, Kuboň Z, Jakobová A and Vodarek V, ‘Development of advanced highchromium ferritic steels; microstructural development and stability’, High ChromiumFerritic Power Plant Steels, A Strang and D J Gooch (eds), Institute of Materials,London, 1997, 73–92.

25 Brett S J, Oates D L and Johnston C, ‘In-service type IV cracking in a modified 9Cr

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(Grade 91) header’, Creep and Fracture in High Temperature Components – Designand Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds),DEStech London, 2005, 563–572.

26 Kimura K, ‘Review of allowable stress and new guideline of long-term creep strengthassessment for high Cr ferritic creep resistant steels’, Creep and Fracture in HighTemperature Components – Design and Life Assessment Issues, I A Shibli, S RHoldsworth and G Merckling (eds), DEStech London, 2005, 1009–1022.

27 Masuyama F, ‘Creep rupture life and design factors for high strength ferritic steels’,Creep and Fracture in High Temperature Components – Design and Life AssessmentIssues, I A Shibli, S R Holdsworth and G Merckling (eds), DEStech London, 2005,983–996.

28 Di Gianfrancesco A, Cipolla L and Cirilli F, ‘Microstructural stability and creep dataassessment of Tenaris Grades 91 and 911’, 1st International Conference, Super-HighStrength Steels, Rome, Italy, (AIM), 2005 CD-Rom.

29 Ennis P, ‘The significance of microstructural changes and steam oxidation for theservice life of chromium steel components’, Creep and Fracture in High TemperatureComponents – Design and Life Assessment Issues, I A Shibli, S R Holdsworth andG Merckling (eds), DEStech, London, 2005, Volume 2, 79–87.

30 Bendick W and Gabrel J, ‘Assessment of creep rupture strength for the new martensitic9% chromium steels E911 and T/P92’, Creep and Fracture in High TemperatureComponents – Design and Life Assessment Issues, I A Shibli, S R Holdsworth andG Merckling (eds), DEStech London, 2005, 406–418.

31 Naoi H, Mimura H, Ohgami M, Morimoto H, Tanaka T, Yazoki Y and Fujita T,‘NF616 pipe production and properties and welding consumable development’, EPRI/National Power Conference, London, 1995, 8–29.

32 Masuyama F, ‘ASME code approval for NF616 and HCM 12A’, EPRI/NationalPower Conference, London, 1995, 98–113.

33 Granacher J and Monsees M, ECCC DESA CRDA Procedure Document, ECCCRecommendations, 2001, Vol 5, Appendix D2.

34 Holdsworth S R, Bullough C K and Orr J, BS PD6605 Creep Rupture AssessmentProcedure, ECCC Recommendations, 2001, Vol. 5, Appendix D3.

35 Orr J, ECCC ISO CRDA Procedure Document, ECCC Recommendations, 2001, Vol5, Appendix D1a.

36 Granacher J and Schwienheer M, ECCC Procedure Document for Graphical Multi-Heat Averaging and Cross Plotting Method, ECCC Recommendations, 2001, Vol 5,Appendix D4.

37 Merckling G, ‘Long term creep rupture strength assessment: the development of theEuropean Collaborative Creep Committee post assessment tests’, Creep and Fracturein High Temperature Components – Design and Life Assessment Issues, I A Shibli,S R Holdsworth and G Merckling (eds), DESTech, London, 2005, 3–19.

38 Yagi K, ‘Acquisitions of long term creep data and knowledge for new applications’,Creep and Fracture in High Temperature Components – Design and Life AssessmentIssues, I A Shibli, S R Holdsworth and G Merckling (eds), DESTech, London, 2005,31–45.

39 Swindeman R W and Swindeman M J, ‘A comparison of creep models for nickelbase alloys for advanced energy systems’, Creep and Fracture in High TemperatureComponents – Design and Life Assessment Issues, I A Shibli, S R Holdsworth andG Merckling (eds), DESTech, London, 2005, 361–371.

40 Holdsworth S R, Askins M, Baker A, Gariboldi E, Holmstrom S, Klenk A, Ringel M,

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Merckling G, Sandstrom R, Schwienheer M and Spigarelli S, ‘Factors influencingcreep model equation selection’, Creep and Fracture in High Temperature Components– Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling(eds), DESTech, London, 2005, 380–393.

41 Abe F, ‘Stress to produce minimum creep rate of 10–5 %/h and stress to cause ruptureat 105h for ferritic and austenitic steels and superalloys’, Creep and Fracture in HighTemperature Components – Design and Life Assessment Issues, I A Shibli, S RHoldsworth and G Merckling (eds), DESTech, London, 2005, 997–1008.

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446

16Creep fatigue behaviour and

crack growth of steels

C . B E R G E R, A . S C H O L Z, F . M U E L L E Rand M . S C H W I E N H E E R, Darmstadt University of

Technology, Germany

16.1 Introduction

The lifetime of components of power plants depends in most cases on variableloading conditions. This concerns fatigue and creep–fatigue as well as crackinitiation and crack propagation. The normal variable service conditionswhich encompass phases of start-up, full load, partial load and shutdowncause variable stress–strain distributions and temperature transients and leadto a large variety of combined static (primary load) and variable loading(secondary load) situations. A wide range of loading parameters have beenintroduced in order to simulate complex loading of components. Conventionaltesting of temperature-induced loading deals with fatigue experimentsrepresenting the cyclic loading at the heated surface of components, whilecreep experiments represent the quasi-static loading phases of components.

A key problem deals with the identification and interpretation of differentphysical damage mechanisms. Therefore, phenomenological solutions weretraditionally introduced and put into practical use. Fatigue cracks occur atthe surface of a component (Fig. 16.1), but creep damage is initiated bycreep cavities and micro cracks preferably at grain boundaries. Their interactionwas intended for conventional heat-resistant steels, but their consideration inlife estimation methods has not been realised satisfactorily. A clear influenceon endurances caused by the superposition of fatigue and creep has beenobserved.1 Increasing tensile hold time decreases the fatigue endurance limit.Thus, the failure mode changes from a fatigue dominant to a creep dominantcondition. In the high strain regime, damage is characterized as fatiguedominated. With decreasing strain, creep damage predominates. It is knownfrom the literature2 that hold times at the tension and compression stage canhave an individual material influence on endurance. Additionally, oxidationeffects can contribute significantly to the reduction of endurance limits.2

Summarizing, either fatigue damage or creep damage prevails, or they mayinteract under variations of strain range, tension and/or compression holdtime, frequency, temperature and ductility of the material.

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Creep fatigue behaviour and crack growth of steels 447

The initial stage of fatigue failure is characterized by dislocation processeswhich lead to surface defects. This is followed by growth processes by bulkdeformation and is completed by tearing up a small remaining ligament.Creep damage owing to nucleation is much more difficult to define. It isassumed that microstructure damage is nucleated during the creep life andcan be interpreted as the onset of tertiary creep.

16.2 Creep–fatigue experiments

Generally, conventional mechanical experiments are needed to measure thematerial properties in order to provide a consistent basis for quality controlpurposes and for design data. In addition, complex experiments may representservice conditions and contribute to verification in life prediction methods.

The creep–fatigue behaviour at the heated surface of heavy components,such as turbine rotors, normally is investigated by conventional low-cyclefatigue (LCF) experiments with standard (LCF) cycles (Fig. 16.2(a)) andcreep–fatigue cycles with dwell periods at maximum and minimum strain(Fig. 16.2(b)). In contrast, a single-stage service-type strain cycle (Fig 16.2(c)to (e)) was developed,1,3 which is characterized by a compressive strain holdphase 1 simulating the start-up condition, a zero strain hold phase 2, withapproximate temperature equilibration at constant loading, a tensile strainhold phase 3, simulating shut-down conditions and an additional zero strainhold phase 4, which characterizes a zero loading condition.

Anisothermal experiments (type ‘an’) approximate the service conditionsmore closely than isothermal experiments. They were carried out up to failuretimes of 8000 h and gave only insignificantly smaller numbers of cycles to

(a)

(b)

(c)To

tal

stra

in r

ang

e

Creep-dominatedfailure (c)

Creep–fatigueinteraction(b)

Increasin

g

tensile dwell

Creep–fatigueinteraction(b)

Pure fatigue and compressive dwells (a)

Number of cycles to failure Nf

16.1 Failure modes at fatigue, creep fatigue and creep-dominatedmaterial behaviour of heat-resistant steels, schematically.2

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(a) T= const.

(b) T= const.

ε

0t

ε r

tht

3

t

ε r

thc

1

0

(c)(d)

(e)

ε

0

T

0

σ

0

ε r ε r

Holdphase j

1 2

ε r = 6%/min

ε1 = –2∆ε/3 ε3 = ∆ε/3 σ

ε3 3 4

ε1Total strain

range ∆ε

Cycle period tp

Tmax

Temperaturerange ∆T

Tmin

t

t

t1 = 0.075 tp t2 = 0.700 tpt3 = 0.150 tp t4 = 0.075 tp

σm

σ

σmax

ε ∆σeff

∆εeff

+

σxσeff

σi σm

εStressrange

∆σ+

16.2 Different strain cycles simulating the conditions at the heated surface of heavy components. Standard cycle without(a) and with (b) hold times as well as service-type cycle (c), the stress–strain path (d) and (e) show mean stress σm,stress σmax, internal stress σi, effective stress σeff and corresponding effective values.

ε

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failure than comparable isothermal tests (type ‘iso’). Considering the designlife of power plants of up to 200 000 h or more, long-term strain cycling isof interest. This could be realised by an isothermal package-type testingprocedure (‘pa’) (Fig. 16.3). It is composed of packages of strain cycles withshort hold times which are periodically inserted into much longer creeppackages. Maximum test durations of 70 000 h have been achieved.

16.3 Stress–strain behaviour

In order to develop a creep–fatigue life estimation procedure, the stress–strain path of the service-type experiments were analysed. Deformation analysesled to the determination of an effective stress concept, σeff = σ – σi (Fig.16.4). The internal stress σi for any time of the measured hysteresis loop, forexample point A, stress σx, is defined as the centre of the hypotheticalelastic–plastic flank curve loop which is inserted in the flank curve loopenveloping the whole measured loop. The flank curves are derived from acyclic or quasi-static yield curve by multiplying the latter by a factor of two.The cyclic yield curve can be experimentally determined by a strain cyclewithout hold times which is inserted into the service-type strain cycling.Owing to the different relaxation phases during the hold times, the meanstress σm varies from zero (Fig. 16.4) and has to be considered for thelifetime estimation. For this purpose a mean stress factor νσ which includesthe Smith–Watson–Topper parameter4 can be used.

16.4 Creep–fatigue interaction, life estimation

For life estimation under creep fatigue loading the life fraction rule5,6 iswidely used. Failure is determined by the summation of fatigue damage Lf asa cycle fraction and creep damage Lc as a time fraction up to a critical creepfatigue value L, whereby L depends on the material. In this paper L is usedin the sense of a damage fraction.

The damage mechanisms during creep–fatigue conditions are mainlyinfluenced by microstructural coarsening7 and cavity growth. Typicalmechanisms of cavity growth have been identified in 1CrMoNiV steels during

∆ε

σε

∆ε

ε 23

4

1

tp

Service type cycleStrain cycle package + Creep package

n – times (n > 10)

24

t1

t t

16.3 Simulation of long-term isothermal service-type testing byisothermal package type tests.

3

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450

16.4 Presentation of a hysteresis loop of an isothermal service-type strain cycle according to Fig. 16.2(c) and definitionof internal stress σi, mean stress σm and the mean stress factor νσ.

A A′σ

σi

∆εelf

σmε

σmax

Measured loop

Flank curve loop(twice the yield curve)

Inserted loop defining theinternal stress σi of point A

Course of internal stress σi

PSWT

Nf

P E

NN

SWT max

f max eff

f eff eff

= /2

= ( , )

( /2, )

σ ∆ε

ν σ ∆ε∆σ ∆εσ

⋅ ⋅

σeff

∆σeff

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stress relaxation.8 Grain boundaries may be cavitated in tensile hold time,but cavitation is usually not found in a balanced cycle containing hold timeof equal duration. In service-type strain cycling, (Fig. 16.2(c)) cavitationdominates.

As indicated by Scholz and Berger,9 a high deformation rate can be associatedwith transgranular damage, while a low deformation rate leads to intergranulardamage (Fig. 16.5). Cavities were found at the transition of transgranulardamage to mixed-mode damage on a 1CrMoNiV steel at 525°C.

This assumption was applied to an interaction concept developed for lifeanalysis using the damage accumulation rule, at first for 1CrMoNiV steel(see e.g. Granacher et al.)10 (Fig. 16.6). The identification of a transitiontime ttr is addressed in the separation of dominant fatigue damage at region1 and dominant creep damage at region 2.

A popular rule for creep–fatigue life analysis in the long-term region isthe generalized damage accumulation rule:

Σ Σ Σ∆k j

j ujk

k kt t N N L ( / ) + ( / ) = fo [16.1]

which combines the Miner rule for fatigue damage and the life fraction rulefor creep damage.1,3,8,9 The damage summation over all cycles k includingdamage at hold times j = 1 to 4 leads to creep–fatigue damage L. To calculatecreep damage, the ratio of time increments tj and rupture time tuj is considered.

The structure of this relationship is relatively simple and therefore inpractical use for life monitoring systems.11 Some features of this rule weremodified empirically in order to cover physical aspects of the material, forexample cyclic stress–strain–time behaviour and mechanisms given in

Stress Region 1transgranular

Region 2intergranular

Deformationcavity growth

Constraineddiffusion cavity

growth

ttr Time

16.5 Various mechanisms of cavity growth during tensile dwell in astrain controlled cycle, 1CrMoNiV steel,2 and determination of atransition time tr for application to damage accumulation.

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(Figs 16.4–16.7). The reference value of the number of cycles to failure Nfo:

Nfo = Nf (∆ε, 2tHsta) · νσ [16.2]

is taken from standard strain-cycling tests, with tension and compressionhold times tHsta (Fig. 16.2(b)). In the case of a viscoelastic stress–strain path,at low strain amplitudes where plastic deformation disappears, creep damagedominates owing to intergranular damage (region 2, Figure 16.5) and fatiguedamage was calculated by a fatigue life curve (tp = 0 h). In the full elastic–plastic regime at high values of total strain range, creep–fatigue damagedominates owing to transgranular damage (region 1, Fig. 16.5). Here fatiguedamage is calculated on the basis of a failure life curve based on symmetricalcreep–fatigue experiments with symmetric hold times (Figure 16.2b). Thehold time tH1 = tHsta at compression strain is fully considered as creep–fatigue damage while creep damage is derived from hold phase 2 and 4 andthe remaining time of hold phase 3 (tH3 – tHsta).

Str

ess

Time(a)

TransgranularMixed mode

Intergranular

Str

ess

(b)

Transgranular

Intergranular

Timettr

ttr ttrt

tpσ

ε

∆ε

16.6 Association between failure mode at creep rupture testing (a)and failure mechanisms at stress relaxation (b), steel of type1CrMoNiV.

t

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The reference value of rupture time tu (eq. (1)) is taken from a creeprupture curve for time dependent stresses σeff(t) and is also affected by thehold times tHsta (Fig. 16.2(b)); Fig. 16.7). Thus, cyclic softening effects aretaken into account. In addition, preloading effects can have an influence onthe reference values of Nfo and tuo.1,7 Furthermore, Nfo depends on a meanstress factor νσ (Fig. 16.4):

νσ = Nf (∆σmax, ∆εeff)/Nf (∆σeff/2, ∆εeff) [16.3]

This factor νσ is derived from the Smith–Watson–Topper parameter (PSWT)4

and considers a mean stress σm ≠ 0 in the cycle. Herein the values σmax, ∆εeffand ∆σeff are taken from the flank curve loop (Fig. 16.2(e) and Fig. 16.4).

The main result of a creep-fatigue life analysis as described above is acreep–fatigue damage mean value L . Mean values of L = 0.54 at 500°C andL = 0.52 at 525°C were identified for the conventional European ferritic1CrMoNiV rotor steel.1,3 Higher mean values were obtained for the martensitic12CrMoV steel, where L = 0.75 at 550°C and L = 0.93 at 600°C are reportedin1,3. Finally, for the modern 600°C steel of type 10CrMoWVNbN and its castversion, mean values L = 0.71 and L = 0.65 were found at 600°C.3 As a

tp = 2 · tHsta = 2 · ttr

ε ε ε

σσσ

Fatigue Creep fatigue3/3 min

Creep fatiguetH1/tH1

ε

0

0

σ

∆ε

tH2tH3

ttH1

tH4

Pure elastic Transition regime Elastic plastic

tCreep damage ttr Creep fatigue

damage∆ε<0.36%

0.0

0.36<∆ε<0.44% ∆ε>0.44%

0.0 0/0

3/3 3/3tH1/tH1

Nf

16.7 Association of elastic–plastic deformation and failuremechanisms at stress relaxation for the calculation of creep damageand fatigue damage.

∆ε

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result, all features introduced within this creep–fatigue interaction conceptyielded the smallest scatter band in the creep fatigue life diagram (Fig. 16.8).

Three-stage service-type strain cycling (Fig. 16.9) demonstrates typicalservice loading conditions as cold start, warm start and hot start. Such loading

16.8 Results of creep–fatigue life assessment of service-type straincycling experiments, (a) for 10CrMoWVNbN forged and cast steelsand 9CrMoVNb pipe steel and (b) measured number of cycles tofailure Nf of isothermal service-type strain cycling tests versuspredicted number of cycles to failure N f

**; single heat and materialgrade of 10%CrMo(W)VNbN.3,12

10CrMoWVNbN - forged steel10CrMoWVNbN – cast steel9CrMoVNb - pipe steel

1

0.1

0.01

∑∆ t

/tu

o

Mean value

L

0.01 0.1 1∑N/Nfo

(a)

Nf c

ycle

100000

10000

1000

100

T = 600°C

Extrapolated

εr = 0.06%/min

Single-stage cyclingsingle heat No 1Agrade

Three-stage cyclingsingle heat No 1Agrade

Material:heat No 1A: X12CrMoWVNbN10-1-1grade: 10%CrMo(W)VNbN

100 1000 10000 100000 N f

**

(b)

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Creep fatigue behaviour and crack grow

th of steels455

T

Tmin

C H W C Tm

∆TH = 50°C ∆TW = 100°C∆TC = 300°C

Frequency: 1 Coldstart, 3 Warmstarts, 16 Hotstarts

Collective (1 + 3 + 16) · tc

ε

C1C2

tC

C example for elementary cycles

CH HH HH HH HW WH …H W Rainflow – counting

Rangemean – counting

H2.4 W2 W3 ∆εC

…C

HC

t

C4

C3

∆εW∆εH∆εH W1

H1

H3

W3H3H3

H4 H4

H2

H1

H2

W4

W2

16.9 Three-stage service-type strain cycling and principle of cycle counting.1,10

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sequences are of specific design interest. The frequency is typical for amedium loaded power plant. Three-stage cycling tests have reached longesttimes to failure up to 1.5 × 104 h for 1CrMoNiV steel and 104 h for10CrMoWVNbN steel.

The creep–fatigue damage assessment of the three-stage service-typeexperiments leads to values L within the scatter band of the single-stageservice-type tests (Fig. 16.8). This important result confirms the ability ofthe damage accumulation rule (Equation [16.1]) for multi-stage service-typeloading. In order to transfer knowledge of deformation behaviour and damageassessment obtained in several research programmes, a software tool SARAwas developed for lifetime studies and industrial application as a life-cyclecounter in power plants.3,10 The numerical simulation by SARA envelopesthe input of cycle and material data, the synthesis of stress–strain hysteresisloops according to cycle counting methods, the individual assessment offatigue damage, creep–fatigue damage and finally an output of the predictednumber of cycles-to-failure N f

**, failure time t f** as well as cycle specific

results. Life estimation on the basis of single heat data and material gradedata leads to an acceptable result (Fig. 16.8).

16.5 Multiaxial behaviour

Life estimation concepts either of a conventional type or of a advanced typerequire suitable multiaxial experiments for verification purposes. In additionto tension/torsion or internal pressure experiments developed in the past,experiments with cruciform test pieces (Fig. 16.10) are of great interest.

Investigations by Ohnami and co-workers13 have demonstrated the largeinfluence of loading ratio on number of cycles to crack initiation. Purefatigue tests on a 1%Cr-steel show a factor of 10 in life between the biaxialstrain ratio Φε = –1 and Φε = +1 (Fig. 16.11(a), solid lines).

Long-term service-type creep-fatigue experiments on a 1%CrMoNiV rotorsteel with four hold times (cycle period tp) are performed in the cruciformtesting system. Experiments run under strain-controlled mode with longhold times. The strain ratio is given as Φε = 0.5 and Φε = 1. A total of fivebiaxial service-type creep–fatigue experiments were performed on the1%CrMoNiV rotor steel with relevant test durations of up to about 2000 h.As a first result at strain ratio Φε = 1, a clear influence of superimposed creepat hold times of a factor of two can be observed. Secondly, a strain ratio ofΦε = 0.5 leads to an increase of the number of cycles to crack initiation Ni ofa factor 1.5 compared to Φε = 1. This result confirms the pure fatigue biaxialexperiment results.13 Further, biaxial strain-controlled experiments lead to areduction of number of cycles to crack initiation Ni up to a factor 3 comparedto uniaxial service-type creep fatigue experiments (Fig. 16.11(b)). On theone hand, an increase of total strain range leads to a larger difference in Ni

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th of steels457

(a)(b)

Extensometry

εA, εB

Test zone

y

x

TE

B

A

TE

PEEQ(Ave. Crit.: 75%)

+2.00e–03+1.50e–03+1.00e–03+5.00e–04+0.00e+00

16.10 Scheme of the cruciform specimen (a), and elastic–plastic finite element calculation showing equivalent plasticstrain εp distribution (b); maximum deformation in the test zone (TE thermocouple).

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Tota

l st

rain

ran

ge

∆εx(

%)

1

0.8

0.6

0.4

0.2

0.1102 103 104 105

Number of cycles to crack initiation Ni (cycles)(a)

Tota

l st

rain

ran

ge

∆εx(

%)

1

0.8

0.6

0.4

0.2

0.1102 103 104 105

Number of cycles to crack initiation Ni (cycles)(b)

Φε–1.0–0.50.00.51.0

OhnamiΦε = 1.0, tp = 1.33 h, t1 ≈ 530 h, ∆εx = 0.60%Φε = 1.0, tp = 1.23 h, t1 ≈ 2055 h, ∆εx = 0.42%Φε = 0.5, tp = 1.33 h, t1 ≈ 830 h, ∆εx = 0.60%Φε = 1.0, tp = 3.53 h, t1 ≈ 1016 h, ∆εx = 0.60%Φε = 1.0, tp = 3.43 h, t1 ≈ 2230 h, ∆εx = 0.42%

Uniaxial service-type, tp = 1.0 hUniaxial service-type, tp = 3.2 hΦε = 1.0, tp = 1.33 h, t1 ≈ 530 h, ∆εx = 0.60%Φε = 1.0, tp = 1.23 h, t1 ≈ 2055 h, ∆εx = 0.42%Φε = 1.0, tp = 3.53 h, t1 ≈ 1016 h, ∆εx = 0.60%Φε = 1.0, tp = 3.43 h, t1 ≈ 2230 h, ∆εx = 0.42%

16.11 Influence of biaxial strain ratio Φε on the number of cycles tocrack initiation Ni in fatigue testing, 1%CrMoV, T = 550°C13 and firstresults of long-term service-type creep–fatigue experiments (Fig.16.2(c)) (Φε = 1.0 and Φε = 0.5), dε/dt = 0.06% min–1 (a), comparisonwith uniaxial experiments (b), tp periodic time, ti time to crackinitiation corresponding to a crack depth of about 0.2 mm;1%CrMoNiV, T = 525°C.

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values, but on the other hand, increasing hold time has a more significantinfluence on number of cycles to crack initiation Ni.

Results of experiments with cruciform specimens can be employed incombination with finite element (FE) simulations in order to verify variousmaterial models. These simulations were performed with ABAQUS wherebya constitutive material model is implemented in a user-defined subroutineUMAT. A constitutive material model describes elastic–viscoplastic behaviourfor small deformations and was introduced by Tsakmakis and Reckwerth.14

A key feature is the combination of effective stress with a generalized energyequivalence principle. An undamaged fictitious material is described by meansof effective variables which are the basis of the constitutive material model.The structure of this model can be attributed to Lemaitre and Chaboche.15 Adamage variable D is defined by an approach proposed by Lemaitre additionallyfor another set of variables for the damaged real material. The known behaviourof the undamaged fictitious material is then mapped to the unknown behaviourof the real material with damage. This step is done by substitution of theeffective variables using relations which implicate the damage variable D.

Within current research work on creep–fatigue, the material parameters ofthe constitutive model were determined by a two-step approach using acombination of the neural networks method and the optimization method byNelder–Mead.29 The neural networks method is already established in similarnon-linear problems and can deliver a ‘global’ solution. The method byNelder–Mead is a direct search method without the need for numerical oranalytical gradients and leads only to a ‘local’ solution. This method iscommonly referred to as unconstrained non-linear optimization. In the firststep, the neural network identifies a parameter vector close to the globalsolution within a parameter interval. This result is used subsequently in thesecond step as an initial parameter vector in the Nelder–Mead method forfurther improvement of the solution. In order to identify the parameters ofthe model under examination for 1%CrMoNiV steel by means of neuralnetworks, one-dimensional (1D) calculations depending on parameter variationswere performed. The example in Fig. 16.12 demonstrates the applicability ofthe material model introduced.

16.6 Creep and creep–fatigue crack behaviour

In the case of unavoidable notches, creep–fatigue cracks may be initiated orpropagated by static and/or cyclic high temperature loading. A quantitativedescription is required in order to establish crack initiation and propagationmethodologies.

For a description of crack behaviour based on experiments on standardcompact–tension specimens16 and/or non-standard test specimens17 undercreep–fatigue conditions, the parameter C* and the stress intensity factor KI

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can be applied to heat-resistant steels. The creep fracture mechanics parameterC* is valid for stationary creep in the crack tip environment.16 The stressintensity factor KI is valid only for linear elastic behaviour,18 but it can beused as an approximation if the plastic zone, near the crack tip, is limited.19

A model for creep crack growth was proposed which assumed that crackadvance occurs when the creep ductility is exhausted at the crack tip.20 Thecreep crack growth rate under steady state conditions may be written as:

da/dtNSW = [(n + 1)/ ( )u*A ] · (A/rc)

[1/(n+1)](C*/In)[n/(n+1)] [16.4]

where rc is the size of the creep process zone (usually related to the grain sizeof the material), In represents a non-dimensional function and Au

* is theappropriate (multiaxial) crack tip creep ductility. This model is known as theNikbin–Smith–Webster model (NSW model). For ductile steels it has beenfound that most experimental data approach the Nikbin-Smith-Webster planestress prediction.21 The creep crack growth rate is most sensitive to themultiaxial creep ductility, Au

* . Therefore, the steady state creep crack growthrate, da/dtNSW-A, may be approximated as:22

da/dtNSW-A = 3·(C*)0.85/ Au* [16.5]

where da/dt and C* have the units of mm h–1 and MPa m h–1, respectivelyand Au

* is taken as the uniaxial failure strain, Au, for plane stress conditionsand Au/30 for plane strain.23 This model has been validated for a range ofmaterials.21

If ∆a is the minimum crack extension that can be measured reliably, thenthe initiation time, ti, may be estimated by:

ExperimentSimulation

Φε = 1.0, tp = 1.33 h∆εx = ∆εx = 0.60%

0.00 0.25 0.50 0.75 1.00 1.25 1.50Time (h)

Str

ain

(%

)

0.6

0.3

0.0

–0.3

–0.6

16.12 Finite element simulation of biaxial service-type creep–fatigueexperiments (see Fig. 16.11(a)), comparison of strain versus timecurves from experiment to those of FE model calculated by theconstitutive material model, and the material parameters determined,Φε = 1.0, 1%CrMoNiV steel, T = 525°C, dε/dt = 0.06% min–1.

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ti = ∆a/(da/dt) [16.6]

If the approximate Nikbin–Smith–Webster Equation [16.5] is used, thenEquation [16.6] becomes:

t a A Ci u* 0.85 = /[3 ( *) ]∆ ⋅ ⋅ [16.7]

Predictions from Equation [16.7] will vary under conditions of planestress and plane strain. In the following, the crack initiation is defined by aconstant technical crack initiation length ∆ai = 0.5 mm which is independentof specimen geometry and size. Other definitions are possible, for examplea crack initiation length, which is dependent on geometry and size.

The parameter C* is shown in Fig. 16.13 against the creep crack initiationtime ti for the 10CrMoWVNbN cast steel at 550°C and 600°C. In a short-term regime there is a difference between small scale and large scale specimens.In a long-term regime the influence of specimen geometry and size on crackinitiation time becomes smaller. The plane strain and plane stress initiationtime predictions, given by Equation [16.7] are included in Fig. 16.13. Theplane strain Nikbin–Smith–Webster A model provides a very conservativeestimate of the time to creep crack initiation. Applying the formula for planestress, a better agreement between creep crack initiation data and theapproximate Nikbin–Smith–Webster model can be observed.

The creep crack initiation behaviour of the 10CrMoWVNbN cast steel interms of KI is shown in Fig. 16.14. In this figure the Larson–Miller parameterPLM,24 known from the description of creep data, is used as a time–temperature

101

10–1

10–2

10–3

10–4

10–5

C*

(N m

m–1

h–1

)

Specimen/a0/W/T/environmentCs25/0.55/550°C/airCs25/0.55/660°C/airCs25/0.55/600°C/shielded gasDs60/0.20/550°C/airDs60/0.20/600°C/air

ti = (∆a· Au* )/(3·(C*/1000)0.85)

Au* = Au (NSW-A-stress)

Au* = Au/30 (NSW-A-strain

10CrMoWVNbN-cast steelT = 550 and 600°C

∆ai = 0.5 mm

101 102 103 104 105 106

ti (h)

16.13 Creep crack initiation time ti versus parameter C*;10CrMoWVNbN cast steel, 550°C and 600°C.

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parameter in order to get a temperature-independent representation. There isa significant difference between small scale and large scale specimens. Thecreep crack initiation on large scale specimens occurs later than on smallscale specimens. It is recommended that side-grooved C(T) specimens (B =25 mm) be used for creep crack tests.16 Results from these specimens lead toa conservative estimation of creep crack initiation with regard to large specimensor components.

In the two-criteria-diagram19 the nominal stress σnpl is related to the stresssituation in the ligament, that is in the far field of the creep crack and thefictitious elastic parameter KI at time zero characterizes the crack tip situation.These loading parameters are normalized in a two-criteria-diagram by therespective time- and temperature-dependent values, which indicate materialresistance against crack initiation. The normalized parameters are the stressratio Rσ = σnpl/Ru/t/T for the far field and the stress intensity ratio RK = KI/KIi

for the crack tip, with creep rupture strength Ru/t/T and parameter KIi whichcharacterize the creep crack initiation of the material. This parameter has tobe determined from specimens with high ratio KI/σnpl, preferably side groovedC(T)25-specimens. The two-criteria-diagram (Fig. 16.15) distinguishes threefields of damage mode separated by lines with a constant ratio Rσ/RK. Abovethe line Rσ/RK = 2.0 ligament damage, and below the line Rσ/RK = 0.5 cracktip damage is expected. Between these lines a mixed damage mode is observed.Crack initiation is only expected above a boundary line. In order to transfer

Specimen/a0/W/T/environment

Cs25/0.55/550°C/air Cs25/0.55/600°C/air Cs25/0.55/600°C/shielded gas

Specimen/a0/W/T/environment

Ds60/0.20/550°C/air Ds60/0.20/600°C/air

KIi

(MP

a m

1/2 )

75

50

25

10

7.5

5

10CrMoWVNbN-cast steelT = 550 and 600 °C

∆ai = 0.5 mm

11 12 13 14 15 16 17PLM = T · [13.0 + log(t)]

16.14 Stress intensity factor KIi versus Larson–Miller–Parameter PLM;10CrMoWVNbN cast steel, 550°C and 600°C.

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th of steels463

1

0.5

0

Rσ

= σ n

pl/R

u/t

/T

Bu

rst

Lig

amen

t d

amag

eR σ

/Rk

= 2

Rσ/Rk =

0.5

No crack

Crack tip damageLeakage

0 0.5 1 1.5 2RK = Kl/Kli

Cr

1%

1%

1%

1%

1%

*0.1 - 0.5 mm crack length

Techn. crackinitiation *)With Without

Castings with manufacturingdefectsSmall manufacturing defectsin smooth specimens

CT100, a/W = 0.55CT50, a/W = 0.4 – 0.55DENT, B = 60 mm, a/W = 0.4DENT, B = 60 mm, a/W = 0.2DENT, B = 60 mm, a/W = 0.1

Small specimens, CT25, a/W = 0.55DENT, B = 9-15 mm, a/W = 0.2 - 0.4

Small specimens, CT25, a/W = 0.55DENT, B = 9-15 mm, a/W = 0.2 - 0.4

16.15 Two-criteria-diagram for creep crack initiation; 1 Cr and 12 Cr steels, 550°C. DENT represents the double edgenotched tension specimen. CT represents the compact tension specimen.

Mix

ed d

amag

e

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creep crack initiation data from specimens with different sizes to largecomponents with a similar far field and crack tip situation, the two-criteria-diagram has been developed. The applicability of the two-criteria-diagramwas validated by results of more than 100 small and large scale specimenswith artificial and natural defects.

Creep–fatigue experiments, with increased hold time tH up to 3.0 h, showa significant decrease of crack initiation time (Fig. 16.16). With increasinghold times the data points approach pure creep crack behaviour. Short holdtimes lead to shorter crack initiation times, which decrease due to the influenceof fatigue.

To predict creep-fatigue crack initiation, a modified two-criteria-diagramhas been introduced and validated on 1CrMoNiV and 10CrMoWVNbNsteels.25–27 For practical applications, a reduction of the crack initiation timefrom tic to ticf = 0.6 tic is proposed.25 This characterizes the materials resistancein the modified two-criteria-diagram by a new time-dependent parameterKIicf(ticf). Besides this proposed rule, the structure of the two-criteria-diagramfor creep crack initiation remains unchanged. The validity of the boundaryline is confirmed by the results of creep–fatigue crack tests.

The creep crack growth rate da/dt is plotted against parameter C* in Fig.16.17 for 1CrMoNiV steel and in Fig. 16.18 for 10CrMoWVNbN steel incast and forged conditions. An almost linear correlation between da/dt andparameter C* on the log–log scale can be observed. Furthermore, the dataare compared with the Nikbin–Smith–Webster A model. All data fall more orless close to the plane stress line. The Nikbin–Smith–Webster A plane

Specimen/a0/W/tH

Cs25/0.55/0.3 hCs25/0.55/3.0 hCs25/0.55/∞Cs50/0.55/0.3 hCs50/0.55/3.0 hDs60/0.20/0.3 hDs60/0.20/3.0 h

101 102 103 104 105

ti (h)

KIi

(MP

a m

1/2 )

75

50

25

10

7.5

5

10CrMoWVNbN cast steelT = 600 °C

∆ai = 0.5 mmR = 0.1

Pure creep

1/t H

16.16 Stress intensity factor KIi versus creep-fatigue crack initiationtime ti; 10CrMoWVNbN cast steel, T = 600°C.

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Specimen/a0/WCs25/0.55Cs50/0.40–0.55CT100/0.50

Specimen/a0/WD15/0.4–0.6D30/0.2–0.4D60/0.1–0.4

1CrMoNiV steelT = 550°C

0.5 mm < ∆a < 3.0 mm

da/dt = (3·(C*/1000)0.85)/ Au*


Au* = Au/30 (NSW-A-strain)

10–5 10–4 10–3 10–2 10–1 101

C* (N mm–1 h–1)

da/

dt

(mm

h–1

)100

10–2

10–3

10–4

10–5

10–6

10CrMoWVNbN steelT = 500 and 600°C

0.5 mm < ∆a < 3.0 mm

da/dt = (3·(C*/1000)0.85)/ Au*


Au* = Au/30 (NSW-A-strain)

10–5 10–4 10–3 10–2 10–1 101

C* (N mm–1 h–1)

da/

dt

(mm

h–1

)

100

10–2

10–3

10–4

10–5

10–6

Specimen/a0/W/steel

Cs25/0.55/forgedCs25/0.55/cast

Ds60/0.20/forgedDs60/0.20/cast

16.17 Creep crack growth rate da/dt versus parameter C*; 1CrMoNiVsteel, T = 550°C.

16.18 Creep crack growth rate da/dt versus parameter C*;10CrMoWVNbN steels, T = 550°C and 600°C.

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strain line provides a conservative prediction for the data and thereby forcomponents.

As an example, the fatigue crack growth results for pipe steel 9CrMoVNbare presented in Fig. 16.19. The results of this constant amplitude tests showa common scatter band.

The crack growth under creep–fatigue conditions can be described bydepicting the crack propagation per cycle da/dN against the range of effectivestress intensity ∆KIeff, which relates to the stress ratio R (Fig. 16.20). Withincreasing hold time there is an increase of the creep crack propagation percycle observed. Results for creep–fatigue crack tests with tH ≤ 0.3 h arecomparable to results of tests under pure fatigue conditions.

As mentioned above, fatigue processes dominate the crack growth atshort hold times and creep processes at long hold times. For intermediateloading conditions an accumulative crack growth is assumed, which can bedetermined by the summation of increments of creep crack growth and fatiguecrack growth.21 The creep fatigue crack propagation per cycle is then givenby the accumulation rule:

da/dNcf = da/dNf + tHda/dtc [16.8]

Beginning from the initial crack length a0, creep crack growth incrementswere considered when the accumulated time increments exceeded the creepcrack initiation time tic. Results of such a calculation are represented inFig. 16.21. The predicted creep–fatigue crack length ∆ ′acf is composed ofthe fatigue portion ∆ ′a tf ( ) and the creep portion ∆ ′a tf ( ) . The predicted

Specimen/a0/W

Cs25/0.55D15/0.20D60/0.20

da/

dt

(mm

h–1

)

100

10–2

10–3

10–4

10–5

10–6

5 7.5 10 25 75

∆Kl(MPa m1/2)

9CrMoVNb steelT = 600°C

R = 0.1

16.19 Crack propagation per cycle da/dN versus range of stressintensity ∆KI; 9CrMoVNb steel, T = 600°C, range of effective stressintensity KIeff.

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Specimen/tH (h)/RCs25/0.3/0.1Cs25/3.0/0.1Cs50/0.3/0.1Cs50/3.0/0.1DS60/0.3/0.1Ds60/3.0/0.1Cs25/0.3/0.6

da/

dN

(m

m c

ycle

–1)

100

10–2

10–3

10–4

10–5

10–6

5 7.5 10 25 75∆Kleff(MPa m1/2)

10CrMoWVNbN cast steelT = 600°C

tH

16.20 Creep–fatigue crack propagation da/dN per cycle versus rangeof effective stress intensity ∆KIeff; stress ratio R, 10CrMoWVNbN caststeel, T = 600°C.

Measured ∆acf(t)Calculated ∆acf(t)′

∆acf

(mm

)

4

3

2

1

0

1CrMoNiV steelT = 550°C

Cs25-specimenσn0 = 173 MPa

tH = 1.0 h/R = 0.1

∆ ′a tc ( )

∆ ′a tf ( )

0 2000 4000 6000 8000 10000t(h)

16.21 Experimental value ∆acf and predicted value ∆ ′acf of creep–fatigue crack length as a function of time for an individualexperiment; 1CrMoNiV steel, T = 550°C.

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values meet the measured values in an acceptable scatter. Predicted andmeasured values of the creep-fatigue length of different specimen types andsizes are compared in Fig. 16.22. The crack length is slightly overestimatedby the prediction. The basic crack properties of materials are responsible forthe scatter.28


The multi-stage creep–fatigue behaviour of conventional and modern heat-resistant steels was investigated by service-type experiments and a numericalsimulation. Knowledge of cyclic deformation and creep–fatigue damageassessment has been obtained. The industrial benefit can be summarized asfollows:

• The development of a creep–fatigue interaction concept covers physicalinterpretations of deformation and failure mechanisms.

• Stress–strain path and creep–fatigue life can be predicted by user programSARA on the basis of rules for deformation, relaxation and cyclic stress–strain behaviour including internal stress and mean stress.

• A creep–fatigue life estimation procedure was developed for multi-stageloading, which covers cycle counting methods. The procedure was

16.22 Comparison between the predicted value ∆ ′acf andexperimental value ∆acf of creep–fatigue crack length; 1CrMoNiVsteel, T = 550°C.

Cs25 Cs50 D60tH(h)0.31.03.0

10.0

1 CrMoNiV steelT = 550°C

R = 0.1

0.01 0.1 1 100∆acf (mm)

∆′

acf

(m

m)

100

1

0.1

0.01

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established in power plant applications. An extension to automotive andaero engine applications is of future interest. For verification purposeslong-term service-type creep–fatigue data up to 70 000 h were generated.

• In order to characterize creep and creep fatigue crack behaviour, differentmethods exist, like the two-criteria-diagram for creep crack initiation.The methods consider the linear elastic-parameter K1 as well as thecreep fracture mechanics parameter C*. These methods are validated inlong term regime on numerous materials, mainly on type 1CrMoNiVsteels and in recent years on type 10CrMoWVNbN steels.

• Long-term experiments, both uniaxial as well as multiaxial, are necessaryin order to verify advanced life prediction methods for components.

• Modelling creep–fatigue behaviour in multiaxial and multi-stage loadingwith constitutive models is a current challenge.

• Advanced lifing methods and knowledge of materials as well asmethodologies enables a reduction design efforts, an increase in componentloading and an increase in design quality linked with a reduction intechnical risk and an increase in efficiency and economic benefits.

16.8 Acknowledgements

Thanks are due to the Forschungsvereinigung der Arbeitsgemeinschaften derEisen und Metall verarbeitenden Industrie e.V. (AVIF), Project No. A166,the FKM Forschungskuratorium Maschinenbau e.V., Project No. 052510,the Arbeitsgemeinschaft industrieller Forschungsvereinigungen (AiF), theVDEh-Gesellschaft zur Förderung der Eisenforschung mbH, Project No.11200 N and to the Deutsche Forschungsgemeinschaft (DFG), Projects No.BE1890,13-1 and BE1890, 16-1 for financial support, and the working groupsof power plant industry for their interest and technical support.

16.9 References

1 J. Granacher, A. Scholz and C. Berger, ‘Creep fatigue behaviour of heat resistantturbine rotor steels under service-type strain cycling’, Proceedings of the FourthInternational Charles Parsons Turbine Conference, Newcastle upon Tyne, A. Strang,W. M. Banks, R. D. Conroy and M. J. Goulette (eds), The Institute of Materials,London, 1997, 592–602.

2 D. A. Miller and R. H. Priest, Materials Response to Thermal–Mechanical StrainCycling, High Temperature Fatigue: Properties and Prediction, R. P. Skelton (ed.),Elsevier Applied Science, London and New York, 1987, 113–176.

3 A. Scholz, H. Haase and C. Berger, Simulation of Multi-Stage Creep–Fatigue Behaviour,Fatigue 2002, Proceedings of the Eighth International Fatigue Congress, A. F.Blom (ed.), Volume 5/5, Stockholm, June 2002, 3133–3140.

4 K. N. Smith, P. Watson and T. H. Topper, ‘A stress–strain function for the fatigue ofmetals’, J. Mater., 1970, 4, 767–778.

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5 E. L. Robinson, ‘Effect of temperature variation on the creep strength of steels’,Trans ASME, 1938, 60, 253–259.

6 S. Taira, ‘Lifetime of structures subjected to varying load and temperature’, in Creepin Structures, Hoff, N. J. (ed.), Academic Press, New York, 1962, 96–124.

7 J. S. Dubey, H. Chilukuru, J. K. Chakravartty, M. Schwienheer, A. Scholz and W.Blum, ‘Effects of cyclic deformation on subgrain evolution and creep in 9–12% CrSteels, Mater. Sci. Engng. A, 406 (2005), S. 152–159.

8 K. Yagi, O. Kanemaru, K. Kubo and C. Tanaka, Life prediction of 316 stainless steelunder creep-fatigue loading, Fatigue Fracture Engng. Mater. & Struct., 1987, 9 (6),395–408.

9 A. Scholz and C. Berger, ‘Deformation and life assessment of high temperaturematerials under creep fatigue loading’, First Symposium on Structural Durability inDarmstadt, June 9–10, 2005, Orangerie Darmstadt, Proceedings, C. M. Sonsino(ed.), Fraunhofer IRB Verlag, Stuttgart, 2005, S. 311–328.

10 J. Granacher, A. Scholz, H. Möhlig and C. Berger, ‘Heat resistant power plant steelsunder variable long-term conditions’, 5th International Charles Parsons TurbineConference, Cambridge, Strang, A., Banks, W.M., Conroy, R. D., McColvin, G. M.,Neal, J. C. and Simpson, S. (eds), IOM Communications, London, 2000.

11 B. J. Cane, Life Management of Ageing Steam Turbine Assets, Proceedings, Instituteof Materials, London, 1997, 554–574.

12 R. Znajda, Betriebsähnliches Langzeitdehnwechselverhalten wichtiger Stahlsorten,Doctoral Thesis, D17, Darmstadt University of Technology, Shaker Verlag, Aachen,2007.

13 M. M. Itoh, M. Sakane and M. Ohnami, High temperature multiaxial low cyclefatigue of cruciform specimen, Trans. ASME, JEMT, 116, (1), 1994, 90–98.

14 C. Tsakmakis and D. Reckwerth, ‘The principle of generalized energy equivalencein continuum damage mechanics’, Deformation and Failure of Metallic Continua,Hutter K., Kirchner N. and Baaser H. (eds), Springer Series – Lecture Notes inMechanics, 2003, 381–406.

15 J. Lemaitre and J.-L. Chaboche, Mechanics of Solid Materials, Cambridge UniversityPress, Cambridge, 1990.

16 ASTME 1457-00, Standard Test Method for Measurement of Creep Crack GrowthRates in Metals, Annual Book of ASTM Standards, 2001, 3 (1), 936–950.

17 C. M. Davis, F. Mueller, K. M. Nikbin, K. M. O’Dowd and G. A. Webster, Analysisof creep crack initiation and growth in different geometries for 316H and carbonmanganese steels, J. ASTM Inte., 3, (2), 2006. Paper JAI 13220 available on-line atwww.astm.org.

18 G. R. Irwin, ‘Analysis of stresses and strains near the end of a crack traversing aplate’, Trans. ASME, J. Appl. Mech., 1957, 24, 361–364.

19 J. Ewald and K. H. Keienburg, A two-criteria-diagram for creep crack initiation,International Conference on Creep, JSME, IMechE, ASME, ASTM, Tokyo, 1986,173–78.

20 K. M. Nikbin, D. J. Smith and G. A. Webster, ‘Influence of creep ductility andstate of stress on creep crack growth’, Advances in Life Prediction Methods atElevated Temperatures, Woodford, D. A. and Whitehead, J. R. (eds), ASME, 1983,249–258.

21 G. A. Webster and R. A. Ainsworth, High Temperature Component Life Assessment,1st edition, Chapman and Hall, London, 1994.

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22 K. M. Nikbin, D. J. Smith and G. A. Webster, An Engineering Approach to thePrediction of Creep-Crack Growth, Journal of Engineering Materials and Technology,1986, 186–191.

23 M. Tan, N. J. C. Célard, K. M. Nikbin and G. A. Webster, Comparison of creep crackinitiation and growth in four steels tested in HIDA, Int. J. Pressure Vessels andPiping, 2001, 78(12), 737–747.

24 F. R. Larson and J. A. Miller, Time–temperature relationship for rupture and creepstresses’, Trans. ASME, 1952, 74, 765–775.

25 J. Ewald, S. Sheng, A. Klenk and G. Schellenberg, Engineering guide to assessmentof creep crack initiation on components by Two-Criteria-Diagram, Int. J. PressureVessels and Piping, 2001, 78, (11–12), 937–949.

26 J. Granacher, A. Klenk, M. Tramer, G. Schellenberg, F. Mueller, and J. Ewald,Creep-fatigue crack behaviour of two power plant steels, Int. J. Pressure Vessels andPiping, 2001, 78, (11–12), 909–920.

27. F. Mueller, A. Scholz, C. Berger, A. Klenk, K. Maile and E. Roos, ‘Crack behaviourof 10Cr-steels under creep and creep–fatigue conditions’, ECCC Creep Conference,12–14 September 2005, London.

28 B. Dogan, U. Ceyhan, J. Korous, F. Mueller and R. A. Ainsworth, ‘Sources of scatterin creep/fatigue crack growth testing and their impact of plant assessment’, Proceedingsof FITNET 2006, International Conference on Fitness-for-Service, 17–19 May 2006,Amsterdam.

29 J. A. Nelder and R. Mead, ‘A simplex method for function minimization’, Comput.J., 1965, 7, 308–313.

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472

17Creep strength of welded joints of

ferritic steels

H . C E R J A K and P. M AY R, GrazUniversity of Technology, Austria

17.1 Introduction

Ferritic steel grades are highly valued for fabrication of components in thethermal power generation industry. As well as low chromium ferritic/bainiticsteels for boiler components like waterwalls or for sub-critical steam pipingand headers, steels with a chromium content in the range of 9–12% for ultra-super critical (USC) power plants are of great interest. Welding in all itsvarieties is still the major joining and repair technology for power plantcomponents. Either by repair welds in casting defects, fabrication welds,joints of similar and dissimilar steel grades, connections with small cross-sections for example tube to tube welds, or with large cross-sections forexample pipe to pipe or pipe to casting welds, the microstructure of the joinedmaterials is strongly influenced by the welding process and thereby theirmechanical properties are altered. As for all components exposed to hightemperatures during service, the 100 000 h creep rupture strength of basematerial (BM), weld metal (WM) and cross-welds is still the major designcriterion. Table 17.1 gives an overview of the chemical composition of commonlyused ferritic creep-resistant steels and their 100 000 h creep rupture strength.

Long-term experience of creep exposed welded structures has shown thatthe heat affected zone (HAZ), a narrow zone of base material adjacent to theweld fusion line altered by the weld thermal cycle, in respect to the creepstrength, is often regarded as the weakest link in welded constructions.Considering just plain steam pressure (p) loading in a pipe, the HAZ can bestressed in series and in parallel.1 In the absence of additional mechanical orthermal loading, the hoop stress (σH) is about twice as the longitudinal stress(σL) and series loading of strong and weak zones (WM-HAZ-BM) is moresignificant than parallel loading. Longitudinal seam welds show the highestrisk of failure (Fig. 17.1). If additional loading arises, that is insufficientdead weight support or sagging, the total axial stress (σL) can be greater thanthe hoop stress (σH) and the HAZ in girth welds exhibits at least similar riskof damage as in seam welds.

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Creep strength of w

elded joints of ferritic steels473

Table 17.1 Chemical composition (wt%) and creep rupture strength of widely used creep-resistant ferritic steels for application inthermal power plants

Low Cr steels C Si Mn Cr Mo Cu W Ni V Nb B N Ti TW1 ( C)°

R

m T×10 /2

5W

(MP )a3

13CrMo4-4 min 0.08 0.40 0.70 0.40 550 49(T/P12) max 0.18 0.35 1.00 1.15 0.60 0.30 0.01210CrMo9-10 min 0.08 0.40 2.00 0.90 550 68(T/P22) max 0.14 0.50 0.80 2.50 1.10 0.30 0.0127CrWVMoNb9-6 min 0.04 0.10 1.90 0.05 1.45 0.20 0.02 0.0005 550 130T/P23 max 0.10 0.50 0.60 2.60 0.30 1.75 0.30 0.08 0.0060 0.0307CrMoVTiB10- min 0.05 0.15 0.30 2.20 0.90 0.20 0.0015 0.05 550°C 14710 (T/P24) max 0.10 0.45 0.70 2.60 1.10 0.30 0.0070 0.010 0.10

9% Cr steelsX11CrMo9-1 min 0.08 0.25 0.30 8.00 0.90 550 92(T9) max 0.15 1.00 0.60 10.00 1.10 0.30X10CrMoVNb9-1 min 0.08 0.30 8.00 0.85 0.18 0.06 0.030 600 94(T/P91) max 0.12 0.50 0.60 9.50 1.05 0.30 0.30 0.25 0.10 0.070X11CrMoWV Nb9-1-1 min 0.09 0.10 0.30 8.50 0.90 0.90 0.10 0.18 0.06 0.0005 0.050 600 98(E911) max 0.13 0.50 0.60 9.50 1.10 1.10 0.40 0.25 0.10 0.0050 0.090X10CrWMoVNb9-2 min 0.07 0.30 8.50 0.30 1.50 0.15 0.04 0.0010 0.030 600 113(T/P92) max 0.13 0.50 0.60 9.50 0.60 2.00 0.40 0.25 0.09 0.0060 0.070

12% Cr steelsX20CrMoV12-1 min 0.17 10.00 0.80 0.30 0.25 600 49

max 0.23 0.50 1.00 12.50 1.20 0.80 0.35T/P122 min 0.07 10.00 0.25 0.30 1.50 0.15 0.04 0.0005 0.040 600 103HCM12A max 0.14 0.50 0.70 12.50 0.60 1.70 2.50 0.50 0.30 0.10 0.0050 0.100

1service temperature;2100 000 h creep rupture strength at service temperature;3Software Stahlschlüssel 2004, Verlag Stahlschlüssel Wegst GmbH. Version 4.00.0005; www.stahlschluessel.de;4The T23/T24 Book – New Grades for Waterwalls and Superheaters, Vallourec & Mannesmann Tubes, 2000;5ECCC Data Sheets, 2005, compiled and published by D G Robertson and S R Holdsworth, ETD Ltd;6METI (Ministry of Economy, Trade and Industry) Thermal Power Standard Code, Japan (single phase with no delta ferrite).

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Awareness of the importance of knowledge of creep behaviour in weldedstructures has increased continuously over the last few decades. In the 1990sseveral failures of welded steam piping systems, some of them in a catastrophicmanner (see Fig. 17.2), have backed the need for investigation into the creepbehaviour of welded structures.2 At that time, these failures were not seen asanything else than an anomaly resulting from improper fabrication or installationor inappropriate service conditions. With increasing hours in service, variousother cracks, leaks and even pipe ruptures occurred and the industry realisedthe symptomatic manner of some problems caused by the utilisation of weldedferritic steel components. These failures acted as a driving force for increasedresearch activities on failure characterisation, non-destructive testing methods,remaining life prediction methods, repair technology as well as the developmentof new improved steel grades and welding procedures.

The reason that many power stations are operating beyond their designlife has again boosted the research efforts of manufacturers, operators aswell as academics. Nowadays, an almost similar effort is put into thedetermination of the creep properties of weld metals and weldments as ofbase materials. In the following, the state of the art of science and technologyof creep-exposed ferritic weldments covering topics like influence of weldingprocedures on the microstructure, HAZ simulation, weld metal development,creep behaviour of welded joints, selected damage mechanism in creep-exposed welded structures, implications for the industry using ferritic creep-resistant steels and finally future trends in this field will be discussed.

17.2 Influence of weld thermal cycles on the

microstructure of ferritic heat-resistant steels

Fusion welding, as the most important joining process in power plantconstruction works, strongly affects the joint properties.3 Not only is a new

σH =

2 p D

t××

σL =

p D4 t

××

σH : σL = 2 : 1

σL

σH

17.1 Schematic of different loading conditions acting on the HAZ ofweldments in a pressurised pipe with mean diameter D and wallthickness t.1

p

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Creep strength of welded joints of ferritic steels 475

type of material, the weld metal, deposited between the connected parts butalso the base material is altered by local and very inhomogeneous heattreatment as a result of the weld thermal cycle. Good overviews on weldingmetallurgy are given by Easterling,4 Granjon5 and Schulze.6 This chapterfocuses on phase transformations taking place in the base material of ferriticheat-resistant steels within the HAZ, strongly influencing the resultantmechanical properties of the weldment.

The optimised base materials microstructure and properties, set throughaccurate melting techniques, an exact production process control and properheat treatment by the base metal manufacturer, are changed completely within

Mohave

17.2 Catastrophic failure of seam welded steam pipe in the USA in1985 causing six deaths and 10 injuries. (Picture courtesy of C. D.Lundin, University of Tennessee).2

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the HAZ by the applied weld thermal cycle. In addition, internal stressesgenerated during rapid cooling, which are in the order of the yield strengthof the material, sum up with system stresses later on in service.

In Fig. 17.3, the basic influences of the welding process on the metallurgyin the HAZ are shown and compared to the calculated equilibrium phasediagram of X10CrMoVNb9-1 (P91) steel. Depending on the selected weldingprocess, the base material microstructure within the HAZ will be changed.The resulting microstructure is governed by the heating rate of the weldthermal cycle, the experienced peak temperature (TP), the dwell time, thecooling rate, the effects of multilayer welding and finally by the adjustedpostweld heat treatment (PWHT) parameters.

The heating rate in arc welding processes can be as high as 200–300 Ks–1. As a result, transformation temperatures are shifted to significantly highertemperatures than predicted in the equilibrium phase diagram (T0). For example,ferrite (α) to austenite (γ) phase transformation can occur about 100 K higherat a heating rate of 100 K s–1 than the calculated T0, resulting in considerablesuperheating of the ferrite before transformation. Other parameters affectedby the heating rate are: recrystallisation temperature, coarsening rate of carbidesand nitrides, solution temperature of carbides and nitrides and the mainproportion of grain growth.

Precipitation strengthening is one of the most effective mechanisms activein ferritic creep-resistant steels. Therefore, precipitate stability is a key factorboth in the base material and also in the HAZ of weldments. At higherheating rates, not only phase transformations but also solution temperaturesof precipitates are shifted to higher temperatures. In most cases equilibriumis not reached at high heating rates and short dwell times and superheatingoccurs. It is of great importance to estimate how this may affect themicrostructural evolution, especially the grain growth behaviour, which isstrongly influenced by grain boundary pinning. When welding precipitationstrengthened creep-resistant steels, three scenarios should be considered:

• Peak temperature is too low to have any noticeable effect on theprecipitates;

• Particles only partially dissolve during weld thermal cycle, but coarseningof favoured particles occurs;

• Particles dissolve completely during the thermal cycle and omission ofgrain boundary pinning causes excessive grain growth.

Besides precipitation strengthening, grain size is of major importance,being a key factor for good mechanical properties, such as tensile strength,toughness, creep rupture strength or with respect to determining thesusceptibility of the alloy to several damage mechanisms like cold cracking,reheat cracking or type IV cracking.

In regions where most of the precipitates have been dissolved by the weld

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Creep strength of w

elded joints of ferritic steels477

Heat affected zone

Pea

k te

mp

erat

ure

Tp

Solidified weld

Solid–liquid transition zone

Coarse prior austenite grains +fine prior delta ferrite grains

Grain growth zone

Grain refined zone

Intercritical zoneOver-tempered region

Unaffected BM

FGHAZ

CGHAZ

1400

1200

1000

800

600

Tem

per

atu

re(°

C)

δL + δ

L + γ + δL

L + γδ + γ

γ

α

α + γ

0 0.2 0.4 0.6 0.8 1Carbon (wt %)

17.3 Schematic of the sub-zones of the HAZ corresponding to the calculated equilibrium phase diagram ofX10CrMoVNb9-1 (P91-type) steel.1

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thermal cycle, excessive grain growth may take place. After completion ofα/γ transformation, the stability of newly formed grains is far from equilibrium.With increasing peak temperature, crystallographic favoured grains begin togrow at the expense of smaller grains. The resulting coarse-grainedmicrostructure generally shows low toughness characteristics and in the caseof some low alloyed ferritic/bainitic Cr-steels significant susceptibility toreheat cracking.

In 9–12% Cr steels, according to the equilibrium phase diagram, austenitestarts to transform to delta ferrite (δ) at the highest peak temperatures. Thenucleation of delta ferrite grains at austenite grain boundaries again causesthe overall grain size to decrease. While the lower solubility of carbon andother austenite stabilising elements in ferrite results in the diffusion of theseelements out of the delta ferrite to the remaining austenite regions, ferriteformers, like chromium, are enriched in the ferritic regions. Therefore,segregated regions differ locally in chemical composition and austenitictransformation on cooling may be incomplete, resulting in retained deltaferrite.

Second, high cooling rates, predominant for low heat input weldingprocesses, like electron beam- (EB), laser- and gas tungsten arc welding(GTAW), and thick walled components can also result in an incompletereverse-transformation of delta ferrite back to austenite. Small amounts ofdelta ferrite may be present in the microstructure even at room temperature.Residual amounts of delta ferrite in a martensitic matrix have been shown tohave a negative influence on impact values as well as on creep rupturestrength and therefore are undesirable in these types of steels.7 Multilayerwelding techniques offer more time for diffusion to compensate segregationalprocesses and result in a complete and homogeneous retransformation ofdelta ferrite to austenite.8

In the following paragraphs the HAZ of creep-resistant chromium steelsand their sub-regions are described in more detail.

17.2.1 Heat affected zone (HAZ)

As pointed out above, the welding process strongly influences themicrostructure and properties of the base material. As a result of the severethermal cycle caused by the welding process, the original microstructure isaltered and a so-called heat affected zone (HAZ) is formed (see also Fig.17.3). The HAZ can be divided into a number of sub-zones. No distinctborderline between the different regions is recognisable; it is more a continuousgradient from the fusion line between the deposited weld metal to the unaffectedbase material. Each sub-zone is represented by its characteristic microstructureand properties.

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Grain growth zone (Tp >> Ac3)

This zone adjacent to the fusion line experiences temperatures well abovethe Ac3 transformation temperature. Any precipitates that obstruct growth ofaustenite grains at lower temperatures dissolve, resulting in coarse grains ofaustenite. In 9–12%Cr steels delta ferrite grains may nucleate at the highestpeak temperatures (TP > 1250°C) causing overall grain size to decrease. Oncooling, lower chromium steels form a bainitic/martensitic microstructureand 9–12%Cr steels form a martensitic microstructure. The coarse-grainedzone (CGHAZ) features the highest hardness of the HAZ and generally lowtoughness values are expected. The coarse-grained zone may be vulnerableto reheat cracking during creep loaded service.

Grain refined zone (Tp > Ac3)

Lower peak temperatures of about 1100°C, just above Ac3, result in an improperdevelopment of austenite, following the α/γ transformation during heating,producing small austenitic grains (FGHAZ). In addition, peak temperaturemay not be high enough to dissolve precipitates completely, limiting thegrain growth by pinning the austenite grain boundaries. On cooling, either afine grained bainitic microstructure for lower chromium steels or a martensiticmicrostructure for higher chromium steels is formed. The fine grained regionof the HAZ is regarded as the weakest link in weldments during creep loadedservice. At longer service times and lower stress levels most of weldments ofcreep-resistant ferritic steels fail within this region by the so-called type IVmechanism.

Partially transformed zone – intercritical HAZ (Ac1 < Tp < Ac3)

Peak temperatures lying between Ac1 and Ac3 transformation temperaturesresult in a partial transformation of α into γ on heating. While new austenitegrains nucleate at favoured positions, like prior austenite grain boundaries ormartensite lath boundaries, the untransformed bainitic or tempered martensiticmicrostructure is simply tempered for a second time by this weld thermalcycle. Partial dissolution of precipitates may be experienced in this part ofthe HAZ and coarsening of undissolved precipitates can occur especiallyduring subsequent PWHT. After cooling, a twofold microstructure consistingof newly formed bainite, for low chromium steels, and virgin martensite, forhigher chromium steels, and the tempered original microstructure coexist.The intercritical HAZ shows a small grain size and exhibits the lowest hardnessvalues in weldments. This sub-zone of the HAZ shows similar susceptibilityto type IV cracking as the grain refined zone.

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Over-tempered region

With peak temperatures experienced below Ac1 the microstructure does notundergo any phase transformations but the original microstructure is locallytempered at higher temperatures compared to that of the annealed base material.As a result, coarsening of precipitates may be enhanced by a higher coefficientof diffusion at this temperature. Some alloys show the lowest hardness valuesin this region.

Zone of unchanged base material

The zone of unchanged base materials concerns temperatures up to ca. 700°C,in which changes in morphology of constituents do not appear to occur.Nevertheless, in this region over-tempering effects can occur which mayweaken the creep strength of welded low alloyed quenched and tempered(Q+T) steels, that is 1%CrMoV steels.9

17.2.2 Multilayer welding

Special attention has to be paid to the advantage of multilayer welding techniques.Multilayer welds are formed by subsequent deposition of weld beads on solidifiedformerly deposited runs. Significant changes in the microstructure of multilayerwelds compared to single layer welds are the result of additional heat input intothe weld metal but also the HAZ of the base material. Reed and Bhadeshia10

as well as Cerjak et al.11 characterised the HAZ of multilayer welds bymathematical modelling. This ‘multilayer’ effect is not only representative ofthe HAZ of base material and regions with different microstructure can beformed within the weld metal. In Fig. 17.4, intercritically heated areas withina cross-weld of E911 base material welded with matching filler have beenmarked in black, while the fusion lines are marked by grey lines.12 While acontinuous region of intercritically heated material forms within the HAZ ofthe base material, the HAZ within the weld metal is discontinuously related tothe loading direction. If subsequent weld beads are symmetrically depositedto the weld centreline, a continuous layer of intercritically heated material isalso formed within the weld deposit (see Fig. 17.4 right).13 These heat affectedregions within the weld metal can exhibit the same susceptibility to certaincracking mechanisms as is prominent in the HAZ of the base material. Therefore,in multilayer welds special attention has to be paid to which welding parametersare utilised, especially heat input, joint geometry and weld layer structure.

17.2.3 HAZ simulation

The problem of conducting basic investigations of HAZ structures in actualwelds is the presence of extremely small and inhomogeneous sub-zones (see

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(b)

(a)

17.4 (a) Multilayer weld of E911 base material with matching fillermaterial. Areas exposed to peak temperatures between A1 and A3transformation temperatures (intercritical HAZ) are marked in black.The loading direction is indicated by arrows.12 (b) creep exposed21/4Cr–1Mo SAW weld with symmetrically deposited weld beadsshowing cracking in the refined weld metal at the weld centreline(Picture courtesy of C. D. Lundin, University of Tennessee).13

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Fig. 17.3). For the HAZ simulation the time–temperature profile acting asinput data can either be measured from real welds or be derived by analyticalsolutions of the heat conduction equation or by using more sophisticatednumerical thermal heat source models.14–16 HAZ simulation allows thegeneration of larger volumes of material with uniform microstructure andproperties which represent one specific point within the HAZ. This generatedhomogeneous microstructure can be used for all kinds of metallographicinvestigations and tested by applying common standardised mechanical testingprocedures such as tensile tests, creep and fatigue tests and impact tests.17

Different techniques are currently applied for the HAZ simulation, namely:

• heating in a hot salt bath until the peak temperature is reached, immediatelyfollowed by cooling in a moderated tempered salt bath;

• induction heating with subsequent quenching in an oil bath;• heating in a furnace to peak temperature, followed by cooling in air or

an oil bath;• controlled resistance heating in the specimen using a weld simulator

(GLEEBLETM, SmitweldTM)

Each physical simulation process has its own characteristic advantages anddisadvantages. A good overview is given by Buchmayr.18

17.3 Weld metal development for creep-resistant

steels

Within the last few decades great efforts were put worldwide into thedevelopment of filler materials for newly developed creep-resistant steelgrades. It is of great importance that matching filler metals are developedsimultaneously to base materials.19 The term ‘matching’ is not directly relatedto the chemical composition but more specific to the design-based tensileproperties, in this case to the creep rupture properties of the weld metal,compared to that of the base metal although it is known that similar chemicalcomposition results in similar creep rupture properties.20 This is correct fora great number of filler metals, but special deoxidation practices, additionsof micro alloying elements, improvement of weldability and handling remainthe detailed knowledge of the individual filler material manufacturer.

Baune et al. summarised the objectives for the development of new weldmetals for 2.25%Cr steels (T/P23, T/P24) as follows:21

• adjust alloying additions to the electrodes so that mechanical propertiesof welds are properly balanced, i.e. impact toughness, tensile properties,hardness and creep rupture strength;

• minimise susceptibility to temper embrittlement. Therefore, the amountof residual elements such as P, Sb, Sn, and As should be minimisedthrough careful control of raw materials;

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• maximise toughness in the as-welded condition, whereas impact toughnessvalues higher than 27 J after PWHT are often required by fabricators;

• set tensile requirements of the weld metal at room temperature accordingto the specifications given for the base material;

• ensure hardness values of the weld metal, obtained in as-welded conditionare below 350HV10 and below 248HV10 after PWHT;

• evaluate, finally, the influence of different PWHT conditions on theproperties of weldments.

Matching weld metal for creep-resistant chromium steels solidifies primaryferritic. During solidification, delta ferrite is enriched by chromium andmolybdenum and the residual melt is enriched by nickel. The chemicalcomposition of segregated regions is strongly influenced by diffusion andchanges continuously during cooling. The rapid cooling process inhibits acomplete δ/γ-phase transformation and residual δ-ferrite is often presentafter cooling to room temperature. The detrimental effects of δ-ferrite onimpact toughness and creep rupture properties of ferritic steels are wellknown.7 Therefore, weld metal development especially for 9–12% chromiumsteels, aims to eliminate retained δ-ferrite by modification of the weld metalchemistry. Additions of Ni, as an austenite-stabilising element, showedbeneficial effects on the impact toughness properties of NF616 (Grade 92)weld metal by limiting the retention of δ-ferrite.22 While nickel additionlowers the A1 transformation temperature significantly, the addition of cobalthas almost no influence on the transformation temperatures but can alsoreduce retained δ-ferrite content effectively.12 In the new generation of ferriticchromium steels, tungsten has been added to improve creep rupture strength.However, tungsten is a strong ferrite stabilising element and promotes theretention of delta ferrite in the weld metal. Additions of 1%Ni or 2%Co incombination with 1%W have been shown to be sufficient in eliminatingretained δ-ferrite and improving impact toughness in a modified 9Cr–1Mosteel.23

17.4 Creep behaviour of welded joints

Weldments of ferritic steels exposed to high temperature creep loading showvery similar behaviour in tendency, irrespective of the chemical compositionand other parameters like welding procedure, groove preparation and so on.Numerous creep tests of cross-weld samples showed that at lower temperaturesthere is no big difference between the base metal and the cross-weld creepstrength.24 This difference becomes more prominent as the temperatureincreases and as the applied stress level is lowered. The time of deviation ofcross-weld creep strength from base metal mean creep strength varies withcreep testing conditions, material grades and welding prerequisites, for examplestress level, temperature, welding procedure, PWHT parameters, and so on.

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Figure 17.5 compares the creep rupture strength of cross-welds with themean line of base metal creep rupture strength of 9% Cr steel grade E911. Athigher stress levels and lower testing temperatures the location of failure israther randomly distributed between BM, WM and HAZ. At lower stresslevels and higher testing temperatures the HAZ of creep-resistant steelsappears to be the weakest link, diminishing the creep strength of weldmentsby up to 50%. Generally, as derived from numerous long running creep tests,the creep rupture time is longer in the order of weld metal, base metal, cross-welds and fine grained simulated HAZ material.25

17.5 Selected damage mechanism in creep-exposed

welded joints

In 1974, Schüller et al.26 categorised the types of cracking observed inweldments of heat-resistant steels by a simple scheme (see Fig. 17.6). Crackswere classified depending on their location and orientation within theweldments. Cracks in the deposited weld metal correspond to type I and typeII. They form in the weld metal and develop either in the longitudinal ortransversal direction but rarely may also be completely unoriented. Whiletype I cracks are arrested in the weld metal, type II cracks may propagateinto the HAZ or even into the base material. The other types of cracksdevelop within the HAZ of the weldments. Type III cracks form in thecoarse-grained part of the HAZ close to the fusion line and can prolong in

E911 base materialCrosswelds 600°CCrosswelds 625°CCrosswelds 650°C

Str

ess

(MP

a)

500

300

10080

60

40

10

101 102 103 104 105

Time to rupture (h)

600°C

625°C

650°C

17.5 Comparison of cross-weld creep rupture strength with meanbase material creep rupture strength of E911 steel at differenttemperatures.12

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this zone as well as into the base material, whereas cracks of type IV arerestricted to the fine or intercritical zone adjacent to the unaffected basematerial. Brett27 added a term for type IIIa cracking to this scheme for afailure mechanism taking place close to the fusion line (type III) but in afully refined HAZ structure with higher fracture ductility.

17.5.1 Type I, type II cracking

Defects in the weld metal are commonly associated with the welding processitself or stress relief during PWHT; more rarely weld metal cracks can berelated to creep damage during high temperature service.28,29

Cracks in the weld metal are mostly in the transversal direction althoughlongitudinal failures are also reported. Previous reported cracks were almostintercrystalline and appeared as hot cracks formed during solidification ofthe weld metal. Since weld metal development has improved and the cleanlinessof the weld deposit has increased during the last few decades, the significanceof solidification cracks in ferritic steels has diminished. They are still ofgreat concern in austenitic and nickel base weld metals.

As described above, the weld metal deposit produced by a multilayerwelding technique followed by sub-critical tempering does not exhibit auniform microstructure. In multilayer welding, the solidified weld beadsunderneath experience a significant thermal influence from subsequent weldthermal cycles. Therefore, similar zones, compared to those of the HAZ ofbase metal, develop in a multilayer weld metal. These zones within the weldmetal can be definitely susceptible to different forms of creep damage. Whilecoarse-grained areas in the weld deposit might show a susceptibility to reheatcracking, fine grained regions might fail by type IV cracking with the samemechanisms active as in the analogous part of the HAZ of the base metal.Both failure mechanisms will be described in more detail in the followingparagraphs.

HAZ

IIIaIII

I

II

IV

HAZ

III

IIIaIV

II

I

17.6 Modified schematic of cracking modes in weldments of heat-resistant steels (original by Schüller et al.).26

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17.5.2 Type III cracking – reheat cracking

Reheat or stress relief cracking is defined as intergranular cracking in theHAZ or the weld metal that occurs during the exposure of welded structuresto elevated temperatures produced by PWHT or high temperature service upto about 20 000 h. This type of cracking in welded structures of creep-resistant, precipitation-strengthened alloys has received considerable attentionsince the mid-1950s.30 Initial work was concerned primarily with crackingproblems in austenitic steels (18Cr–12Ni–1Nb) used for power generatingequipment, especially high temperature steam piping. Reheat cracking wasalso reported in some Ni base superalloys.31

In the early 1960s similar cracking problems during stress relief heattreatments and high temperature service were again observed in powergenerating constructions where the ferritic creep resisting 2CrMo and CrMoVweldments in steam pipework and valve assemblies were found to exhibitoccasional cracking. While high alloyed 9–12% chromium steels show almostno susceptibility to reheat cracking, newly developed low alloyed 21/4%chromium steels can exhibit significant susceptibility depending on theiralloying concept.

The mechanism of reheat cracking is generally understood although detailsof the controlling parameters and their mechanistic interpretation still remaina subject for discussion. In general terms, cracking results when the relaxationstrains, occurring with stress relief of residual stresses during PWHT or hightemperature service, exceed the local ductility of the material. Factors affectingthe susceptibility of welded structures to reheat cracking include the chemicalcomposition, the microstructure resulting from the applied welding processand the stress state. The significance of segregation of alloying elements likeAl, B, Mn and impurity elements like P, S, As, Sb, Sn, and so on is underdiscussion in literature. Some authors report a strong influence of elementssegregated at the prior austenite grain boundaries (PAGB) in lowering thecohesive strength of the boundaries whilst others find no interaction betweensegregation and stress relief cracking failures.30–32 Generally, it is evidentthat all mechanisms weakening or embrittling grain boundaries enhance thesusceptibility to reheat cracking.

In the CGHAZ almost all precipitates are dissolved during the appliedweld thermal cycle. During PWHT and high temperature service reprecipitationtakes place at the grain boundaries as well as inside the grains. The graininterior gets strengthened by precipitation of finely dispersed, mainly coherentcarbonitrides while PAGB are energetically favourable for precipitation ofincoherent carbides like Fe3C, M23C6, M6C. Highly diffusive grain boundariesenhance the coarsening process of the latter and result in the depletion ofsolid solution strengthening elements and the dissolution of MX particles inthe vicinity of PAGB.

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As a result of these microstructural changes, precipitation-strengthenedgrain interiors are surrounded by weak grain boundary areas. Hence, reductionof welding residual stresses mainly proceeds by deformation concentrated atthe grain boundaries. Therefore, stress relief cracking appears to be macroscopicas intergranular cracks along PAGB with low rupture ductility.

Examination has shown that damage caused by stress relief cracking formsat the boundaries of large prior austenite grains by a mechanism of creepcavitation.33 Cavities nucleate primarily on PAGB as incoherent precipitates,acting as stress concentrators. Susceptibility to reheat cracking is a functionof precipitation strengthening in the grain interior, strengthening and weakeningthe behaviour of grain boundaries and the relaxation of residual stresses.

In recent research work, the susceptibility of up-to-date 21/4% chromiumsteels for implementation as membrane waterwalls for USC boilers or super-heater and reheater components of conventional boilers and heat recoverysteam generators was investigated.33–36 The Japanese ferritic/bainitic steelHCM2S (7CrWVMoNb9-6, ASTM A213 T/P23) and the German7CrMoVTiB10-10 (ASTM A335 T/P24) showed different susceptibilities toreheat cracking. Isothermal slow strain rate tensile tests on weld simulatedspecimens were used for qualification. After fracture, the reduction of areais measured as an indicator of ductility. The reduction in area is directlyrelated to the cracking susceptibility and classified as shown in Table 17.2.

The results, summarised in Fig. 17.7, show that care must be taken whenwelding T23 and P23. The pipe material P23 is highly susceptible in a broadPWHT temperature range from 600°C to 750°C while tube material showsonly slight susceptibility to reheat cracking at PWHT temperatures higherthan 675°C. In contrast neither T24 nor P24 material shows any susceptibilityto reheat cracking.

Figure 17.8 shows characteristic fracture surfaces of P23 and P24 pipematerial after isothermal slow strain rate testing in an inert atmosphere.While the fracture surface of CGHAZ simulated P23 specimen can becategorised as almost 100% intergranular brittle fracture, P24 pipe materialshows a tendency for intergranular fracture but with a predominant ductilefracture surface. Multi-pass welding, producing a fine grained microstructure

Table 17.2 Classification of reheat cracking susceptibility by reduction in areameasurements on a CGHAZ simulated specimen after isothermal slow strainrate tensile testing30

Susceptibility to reheat cracking % reduction in area

Extremely susceptible < 5Highly susceptible 5–10Slightly susceptible 10–20Not susceptible > 20

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P22T23

P23T24

P24R

edu

ctio

n o

f ar

ea Z

(%

)100

90

80

70

60

50

40

30

20

10

0525 550 575 600 625 650 675 700 725 750 775

Testing temperature (°C)

17.7 Reduction in area as a function of testing temperature forisothermal slow strain rate tensile tests for different CGHAZsimulated 21/4Cr steels. P23 pipe material shows the highestsusceptibility to reheat cracking according the classification in Table17.2.

by the normalising effect on prior deposited layers, resulted in an increase inrupture ductility and eliminated reheat cracking susceptibility.

17.5.3 Cracking in dissimilar welds

For the sake of completeness, cracking in dissimilar welds is mentioned inthis chapter. There are a multitude of different types of dissimilar weldsdepending on the steel grades and welding consumables used. In thermalpower stations, dissimilar welded joints between ferritic and austenitic steelsand high alloyed martensitic and low alloyed ferritic/bainitic steels arecommonly in use. Weldments, joining different alloys, are characterised bya sometimes very sharp transition in microstructure, physical properties,chemical potential and as a result, in mechanical properties. Owing to thediversity of different types of joints, this chapter numerates only basicmechanisms relevant for creep damage in heat-resistant weldments. Thesemechanisms are:

• mismatch of physical properties• decarburisation/carburisation (gradient in chemical composition and

potential).37

First, additional stresses may arise from a mismatch of physical properties.Different coefficients of thermal expansion and heat conductivity introduce

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thermal stresses adding to system stresses. This results in an additional localisedloading caused by the combination of creep and fatigue close to the fusionline leading to premature failures. The more potent mechanisms influencingthe time to rupture and failure location in creep-exposed dissimilar welds are

17.8 Fracture surfaces of weld simulated coarse grained materialafter slow strain rate isothermal tensile testing (a) P23, (b) P24.34

(a)

(b)

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variations within the microstructure owing to gradients in chemicalcomposition.

Formation of localised zones with inferior creep properties can be theresult of diffusion of alloying elements across the interface either initiallyduring PWHT or during service at elevated temperature.38 Investigating weldsbetween a 1%CrMoV cast and a 12%CrMoV forged pipe (X20CrMoV12-1)utilising a 12%CrMoV filler material, Witwer40 showed the formation of acarbide seam in the fusion line area of the high alloyed weld material. In thelow alloyed material, as a result, a carbon depleted soft region is formed. Thewidth of carbide seam and carbon depleted region is strongly influenced byPWHT and service parameters.39 Regarding creep rupture properties, thisnarrow zone weak in creep is surrounded by zones of higher creep strengthand leads to localised premature damage in the decarburised HAZ region ofthe low alloyed steel (Fig. 17.9).40

200 µm

17.9 Cross-weld creep sample of 1%CrMoV cast welded with12%CrMoV filler fractured after 806 h at 180 MPa at 550°C in thedecarburised CGHAZ of 1%CrMoV steel.40 The fracture surface on theleft shows decarburised (bright appearance) areas.

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17.5.4 Type IV cracking

Type IV cracking is defined as the formation and propagation of failures inthe fine grained HAZ and the intercritically heated region of the HAZ. Astrict differentiation between ICHAZ and FGHAZ is generally difficult becauseof very similar microstructural features in both regions. At this time, type IVcracking is considered as the major ‘end of life’ failure mechanism for ferriticcreep-resistant steel weldments in the power generating industry. Therefore,this failure mechanism is of great interest and many researchers haveinvestigated this life limiting phenomenon in welded components. Type IVcracking has been reported in low alloy ferritic/bainitic steels (1/2Cr1/2

Mo1/2V, 1CrMo, 1CrMoV, 11/4 Cr1/2Mo, 2CrMo, T/P22, T/P23, T/P24), aswell as in ferritic/martensitic 9–12%Cr steels (P91, X20CrMoV121, P92,P122, E911).41–52

Time to rupture for failures in the type IV region is sensitive to the stressstate and the applied loading direction. Cracking was observed in seamwelds as well as in girth welds. While girth welds, loaded by system axialstresses, tend to fail by a ‘leak before break’ mechanism, failures in seamwelds, loaded by the hoop stress, can be catastrophic.53 Review papers ontype IV cracking have been written by Middleton and Metcalfe (1990),46

Ellis and Viswanathan (1998)53 and Francis et al. (2006).54

In contrast to creep failures in ferritic weldments at high stress levelswhich take place randomly either in BM, WM or the HAZ, the fracturelocation of weldments exposed to lower stress levels is shifted into the verynarrow FGHAZ or intercritical HAZ region. Generally, type IV cracking canbe seen as the result of a microstructural zone of material weak in creepstrength surrounded by materials that are stronger in creep. This mismatch ofcreep properties leads to highly complex material behaviour.41

Most fractures have a macroscopic appearance typical of a low ductilityfailure. Figure 17.10 shows a cross-weld creep sample of E911 base materialwelded with matching filler and creep tested for 18 000 h at 600°C. Calculationof the failure strain using the overall gauge length results in very low strainvalues, typically below 10%. This leads to the assumption that the fracturemode is relatively brittle. However, significant creep deformation can bemeasured in a narrow part of the HAZ close to the unaffected base material,as investigated by Parker and Stratford.48 This fact indicates a strong strainlocalisation. Parker estimated the longitudinal strain of uniaxial specimensafter creep exposure from measurements of the grains dimensions present inthe fractured region. Changes in grain shape suggest a localised strain ofabout 20–30% and therefore the failures are in fact locally of a highly ductilenature.

The failure mechanism of type IV cracking is governed by creep cavitation.Creep voids generally initiate at the sub-surface and grow by a diffusive

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mechanism. Figure 17.11 shows the formation of a very narrow band ofvoids in the outer region of the HAZ in a cross-weld sample prepared froman E911 pipe welded with matching filler and creep tested for 14 000 h at600°C. The combination of voids forms micro-cracks which in turn coalesceto form macro-cracks finally leading to premature failure.

Preferred nucleation sites for voids are particle/matrix interfaces associatedwith inclusions or second phase particles.55 In the IC/FGHAZ region ofweldments, carbides are only partially dissolved by the applied weld thermalcycle. Precipitation on retained large particles, such as M23C6, is favouredinstead of fine reprecipitation on grain boundaries, in order to decrease theinterfacial energy of the microstructure.56 Therefore, retained carbides coarsenmore rapidly than those in the base material or weld metal during PWHT andare preferred nucleation sites for creep voids. Letofsky57 compared the evolutionof precipitates at different stages of creep testing in the ICHAZ ofG-X12CrMoWVNbN10.1.1 steel with that of the unaffected base materialand matching weld metal using energy filtering transmission electronmicroscopy (EFTEM).57 His work is in agreement with other researchersshowing that kinetics of microstructural changes, for example coarsening ofprecipitates and formation of the Laves- and Z-phase, are significantly fasterin the ICHAZ than in all other regions of weldments (Fig. 17.12).25–62

General agreement also prevails concerning other microstructural featuresof the IC-/FGHAZ. Both regions susceptible to type IV cracking consist offine equiaxed subgrains. A lath-type structure that is generally observed inall other parts of weldments is missing here.56 During creep exposure thesubgrain microstructure tends to coarsen significantly.63,64 TEM investigationsalso revealed a significantly lower dislocation density in the vicinity of the

4 mm

17.10 Cross-weld sample prepared from E911 pipe welded withmatching filler and creep tested at 600°C for 18 000 h. Themacroscopic low deformation fracture is located in the FGHAZ closeto the unaffected base material.1

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FGHAZ compared to the other areas in welded joints. Further recovery ofexcess dislocations takes place during high temperature service.61 All thesemicrostructural changes take place at higher velocity in the IC/FGHAZ thanin other regions of weldments and contribute to the continuous weakening ofthis specific region, finally leading to the breakdown of creep rupture strength.The formation of a ‘soft zone’ with low hardness is not always directlyrelated to low creep rupture stress. While in some investigations the locationof final fracture corresponds with that of the soft zone,12,53,60,65 other researchersclearly distinguish between softened regions and regions with low creeprupture strength.25,56,61

The role of constraint

The role of constraint in the creep weak IC/FGHAZ region from the adjacentstronger CGHAZ and BM is a key factor in clarification of type IV crackingmechanism. Simplifying the load condition to plain tensile loading of thedissimilar regions of the HAZ in series, the weakest part tries to deformtransversely under high strains. If this region is sufficiently thin, it is constrainedfrom doing so by the adjacent stronger material. As a result, a triaxial stressstate predominates and prevents the weaker region from yielding. Under the

G-X12 CrMoWVNbN 10.1.1, HAZ (soft zone)

As received 3.919 h 8.065 h 12.271 h 26.094 h150 MPa 130 MPa 120 MPa 90 MPa

17.12 Evolution of precipitates in the ICHAZ of creep-exposed cross-weld samples of G-X12 steel.62

17.11 Cross-weld sample prepared from E911 pipe welded withmatching filler and creep tested at 600°C for 14 000 h. Localisedformation of voids and their coalescence to macro-cracks at the outerregion of the HAZ, observed by SEM, led to final fracture.1

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creep regime, the situation is more diverse and is discussed intensively in theliterature.66

Li et al. simulated creep behaviour of P122 cross-welds using a FEMmodel, based on Norton’s law representing four regions with different creepproperties. They concluded that the FGHAZ has the highest equivalent strain,high first principal stress and hydrostatic pressure. The constraining effectsof BM and CGHAZ prevent creep deformation in the FGHAZ. Therefore, anarrow FGHAZ is calculated to decrease the equivalent strain in the FGHAZand is considered to be beneficial in decreasing the occurrence of creepvoids.67 Experimental validation is provided by Japanese researchers reducingthe HAZ width of P122 steel by electron beam welding (EBW). The creeprupture strength of EB weldments was improved by a factor of two comparedto weldments produced by a GTAW process, although EBW specimens stillfailed by a type IV mechanism.68 Crack initiation was at lower creep strainin the EB joint because of larger local multi-axiality compared to the GTAWjoint.

Abe et al.25 simulated type IV creep crack growth behaviour in P122weldments using a three-dimensional FE model taking the diffusive growthof creep voids into account. As a result, creep cracks grow faster in thecentre of the specimen thickness than in specimens’ surface areas. Highermultiaxiality in the centre of the specimen thickness leads to a higher vacancyconcentration which is consistent with experimental observation of initialcreep voids in the centre region. This is supported by the acceleration ofvacancy diffusion, formation and the growth of creep voids under multiaxialconditions for a welded joint specimen and as a result, the enhancement ofdamage accumulation by multiaxial stress states.69 Moreover, no mechanicalconstraint throughout the simulated HAZ creep specimen resulted in no voidformation.25

An alternative approach to modelling type IV damage was introduced byKimmins and Smith who suggest that constraint is relaxed by grain boundarysliding.70 Their experimental results suggested that additional grain boundarysliding results in a greater number of cavities. Once sliding is accommodated,the failure time for both, cross-weld and simulated type IV samples, wassimilar. Therefore, they concluded that material weak in creep deformsindependently of adjacent stronger material and rather than using conventionalcontinuum damage models in FE analysis, alternative models involving themechanism of grain boundary sliding require development.41

Detectability

As mentioned earlier, type IV damage initiates as creep cavitation subsurfaceat about half the time of the expected life of the weldments.63 However,cracks may form relatively late in life and crack growth once initiated may

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be very rapid.71 The remaining life for propagation throughout the wall canbe less than 10 000 h.46 Surface bearing cracks do not appear until late inlife. Therefore, surface examination of creep-exposed weldments only byreplication techniques, penetration testing or eddy current testing can bemisleading and severe damage in the subsurface regions can be overlooked.72

A sound residual life investigation on weldments can only be performed byadvanced ultrasonic (UT) inspection or by highly sophisticated methods liketime of flight diffraction (TOFD) which can detect creep voids and micro-cracks even at a life ratio (t/tr) of 0.5. In any case, life assessment methodsfor weldments vulnerable to type IV cracking have to include a qualitativedamage classification scheme and a cavity density-based model.53,73,74

17.6 Implications for industries using welded creep-

resistant steels

Above it was shown that creep rupture strength of base metals and weldmentscan differ significantly. Therefore, it is of great importance to consider aweld strength factor (WSF) or a weld strength reduction factor (SRF) duringthe design stage of new components and the residual life evaluation of existingstructures.75,76 In design codes for nuclear power plants, weld strength reductionfactors have already been implemented. The opposite value of SRF, theWSF, is defined as the ratio of creep rupture stress at a certain time andtemperature between weldments and base material.

The European Creep Collaborative Committee (ECCC) defined WSF andSRF in their ECCC Recommendations as follows:77

WSF t TR

Ru w t T

u t T( , ) = ( )/ /

/ /[17.1]

SRF t TR R

Ru t T u w t T

u t T( , ) =

– / / ( )/ /

/ /[17.2]

where Ru/t/T is the creep rupture strength of base material samples at a timet and temperature T and Ru(w)/t/T is the creep rupture strength of cross-weldsamples at a time t and temperature T. Up until now, a time and temperatureindependent weld strength reduction factor has been defined. As an example,in the German AD2000-Merkblatt B0, a constant WSF of 0.8 is given forcomponents designed using creep rupture strength values. Investigations bymany researchers have proved that this factor cannot be assumed to beconstant over different temperature and stress levels, but rather depends onseveral factors like material type, stress level, temperature and time and canbe either higher but also significantly lower.78,79 Lack of long-term creeprupture data makes the determination of accurate WSF very difficult.Extrapolation of WSF from short-term creep rupture tests holds the risk of

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overestimation of long-term creep properties and is not recommended. Thisemphasises the necessity for long-term creep rupture data of base materialand welded joints.

In Japan a ‘Strength of High Chromium Steel’ (SHC) committee since2004 has systematically collected and analysed creep rupture data for basemetal and welded joints. Takahashi and Tabuchi80 developed creep reductionfactors for welded joints of HCM12A (P122) steel. The serious influence oftemperature can be seen by a strong decrease in WSF from a value of 1 at550°C down to 0.51 at 650°C.80 Tabuchi and Takahashi made similarinvestigations for 9Cr–1Mo (P91) steel.81 At 550°C a weld strength factor of1 predicts no difference in creep rupture strength of welded components andplain base material, while at 650°C a WSF of 0.7 accounts a 30% decreasein creep rupture strength for weldments. In Fig. 17.13, WSF for several steelgrades at different temperatures are shown.82 A reduction of up to 50% increep rupture strength of weldments compared to base material creep rupturestrength emphasises the importance of further research in this field and thecorrect consideration of WSF during the design stage of new structures.

17.7 Future trends

Up to the present, inferior long-term creep properties of weldments of heat-resistant steels can only be taken into account during the design stage ofthermal power plant components. The awareness of designers, engineers andoperators of the risk of extrapolating results from short-term creep tests tolonger times has already contributed to an increase in safety. General acceptanceof the necessity of long-term creep testing data for cross-welds, weld metaland base metal for a reliable material selection is inevitable.83

Wel

d s

tren

gth

fac

tor

1.0

0.9

0.8

0.7

0.6

0.5

0.4400 450 500 550 600 650 700

Temperature (°C)

Carbon steels

1-2%CrMo

G20Mo5

G17CrMoV5–10

P91

P12212CrMoV

17.13 Weld strength factors for the 100000 h creep rupture strengthof different steel grades.82

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Within the last few decades suppliers of welding consumables have improvedtheir filler materials in terms of cleanliness, weldability and mechanicalproperties. As consequence, weld metal cracking has almost disappeared. Inmost cases the HAZ was identified as the weakest link in welded structuresof creep-resistant steels. Shifting of welds to lower creep-stressed regionsmay improve the short time performance of welded components, but failurescannot be avoided over long times. The use of the most advanced non-destructive testing methods is necessary to detect damage as soon as possiblein order to be able to react with appropriate countermeasures.

The HAZ, as part of the base metal, already has to be taken into considerationin the design of new steel alloys. Weldability studies using HAZ simulationtechniques can provide useful information on possible bottlenecks. Recently,in Japan at the National Institute for Materials Science (NIMS) a 9Cr–3W–3Co steel with a reduced nitrogen level and controlled addition of boron wasdeveloped. In contrast to all creep-resistant steels used up until now, thissteel does not show the formation of a fine grained region within the HAZ.Figure 17.14 shows the results of an electron backscatter diffraction pattern(EBSP) analysis of the grain size, as a function of distance from the fusionline, for a conventional P92 steel versus the new 9Cr–3W–3Co material.84

17.14 Electron backscatter diffraction pattern (EBSP) analysis resultsfor 9Cr–3W–3CoNB steel and P92 steel HAZ microstructures.84

EBSP analysis (inverse pole figure map[001])weld microstructure (GTAW width of HAS; 2.5 mm)

90 p

pm

BP

92

Base metal 1.5mm from fusion line 0.5mm from fusion line

100µm 100µm 100µm

100µm 100µm 100µm

111

101101

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By the elimination of fine grains in the HAZ, the formation of creep damageby a type IV mechanism, which is strictly limited to fine grained regions,should be avoided. Up to a duration of 10 000 h, creep tests at 650°C are verypromising and no difference in creep strength between cross-weld specimenand base material is shown. Although the mechanisms active in this steel arenot fully understood, this might be a possible approach for the prevention oftype IV cracking in ferritic creep-resistant steels.

17.8 References

1 Mayr P, Weldability of Modern 9%Cr Steels for Application in USC Power Plants,Doctoral Thesis, Graz University of Technology, 2007.

2 Henry J F, Ellis F V and Lundin C D (1990), The Influence of Flux Composition onthe Elevated Temperature Properties of Cr–Mo Submerged Arc Weldments – WRCBulletin 354, Welding Research Council, New York.

3 Cerjak H H (1992), Welding of Steam Turbine Components, European Communities– Commission, Luxembourg.

4 Easterling K (1983), Introduction to the Physical Metallurgy of Welding, London,Butterworths, London.

5 Granjon H (1991), Fundamentals of Welding Metallurgy, Abington Publishing,Cambridge.

6 Schulze G (2004), Die Metallurgie des Schweißens, Springer, Berlin.7 Kimura K, Sawada K, Kushima H and Toda Y, ‘Degradation behaviour and long-

term creep strength of 12Cr ferritic creep resistant steels’, in 8th InternationalConference Materials for Advanced Power Engineering 2006, ForschungszentrumJülich GmbH, Liege, 2006.

8 Vekeman J, Dhooge A, Huysmans S, Vandenberghe B and Jochum C, ‘Weldabilityand high temperature behavior of 12% Cr-steel for tubes and pipes in power plantswith steam temperatures up to 650°C’, IIW Report XI-863-06, International Instituteof Welding, 2006.

9 Buchmayr B, Cerjak H and Fauland H P, ‘The effect of the precipitation behaviouron the HAZ-properties of 1%Cr-Mo-V steel’, in 2nd International Conference Trendsin Welding Research, Gattinburg, TN, 14–18 May 1989, ASM International, MaterialsPark, OH, 1990.

10 Reed R and Bhadeshia H K D H, ‘A simple model for multipass steel welds’, ActaMetal Mater, 1994, 42 (11), 3663–3678.

11 Cerjak H, Nagel G and Prader R, ‘Quantifizierung der Zähigkeitsverteilung in einerdickwandigen RDB-UP-Viellagenschweißung’, in MPA-Seminar Sicherheit undVerfügbarkeit in der Anlagentechnik mit dem Schwerpunkt Komponenten nuklearerund konventioneller Kraftwerke, Universität Stuttgart, 10–11 October StaatlicheMaterial prüfungsantalt, 1996.

12 Letofsky E, Das Verhalten von Schweißverbindungen moderner Kraftwerkswerktoffe,Doctoral Thesis, Graz University of Technology, 2001.

13 Lundin C D, ‘Weld performance in high energy piping’, in 5th International ConferenceTrends in Welding Research, ASM, Pine Mountain, 1999.

14 Rosenthal D, ‘Mathematical theory of heat distribution during welding and cutting’,Welding Journal, 1941 20 (5), 220–234.

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15 Rykalin N N (1957), Berechnung der Wärmevorgänge beim Schweißen, VEB VerlagTechnik Berlin, Berlin.

16 Goldak J, Chakravarti A and Bibby M, ‘A new finite element model for welding heatsource’, Metall. Trans. B, 1984 15B 299–305.

17 Murdoch A J, Allen D J and Brown S G R, ‘An investigation into weld type Type IVcracking in P91 steel by Gleeble simulation’, in 5th International Conference Conditionand Life Management for Power Plants – Baltica, Heitanen S. (ed), Porvoo, Finland,6–8 June 2001, VTT Manufacturing Technology, 2001.

18 Buchmayr B, ‘Characterisation of the creep behaviour of weldments by HAZ-simulation’, ECCC Recommendations, 2005 3(II) Appendix 2.

19 Baune E, Cerjak H, Caminada S, Jochum C, Mayr P and Pasternak J, ‘Weldabilityand properties of new creep resistant materials for use in ultra supercritical coal firedpower plants’, in 8th International Conference Materials for Advanced PowerEngineering 2006, Forschungszentrum Jülich GmbH, Liege, 2006.

20 Bhadeshia H K D H, ‘Design of creep-resistant ferritic steel welding alloys’, in 2ndInternational Conference Integrity of High Temperature Welds, IOM, London, 2003.

21 Baune E, Leduey B, Bonnet C and Bertoni A, ‘Development of welding consumablesdedicated to the welding of new generation 21/4 Cr-1Mo pipe materials P23 and P24for power generation and hydrogen service’, in 2nd International Conference Integrityof High Temperature Welds, IOM, London, 2003.

22 Naoi H, Mimua H, Ohgami M, Morimoto H, Tanaka T and Yazaki Y, ‘NF616 pipeproduction and properties and welding consumable development’, in ConferenceNew Steels for Advanced Plant up to 620°C, EPRI, London, 1995.

23 Barnes A and Abson D, ‘The effect of composition on microstructural developmentand toughness of weld metals for advanced high temperature 9–13%Cr steels’, in 2ndInternational Conference Integrity of High Temperature Welds, IOM, London, 2003.

24 Masuyama F, Matsui M and Komai N, ‘Creep rupture behaviour of advanced 9–12%Cr steel weldment’, Key Eng Mater, 2000 171–174 99–108.

25 Abe F, Tabuchi M, Kondo M and Tsukamoto S, ‘Improvement of creep strength ofadvanced ferritic steel welded joints’, IIW Report XI-794-04, International Instituteof Welding, 2004.

26 Schüller H J, Hagn L and Woitscheck A, ‘Risse im Schweißnahtbereich vonFormstücken aus Heißdampfleitungen – Werkstoffuntersuchungen’, DerMaschinenschaden, 1974, 47 (1), 1–13.

27 Brett S J, ‘Type IIIa cracking in 1/2CrMoV steam pipework systems’, Sci TechnolWelding Joining, 2004, 9 (1), 41–45.

28 Viswanathan R and Foulds J, ‘Failure experience with seam-welded hot reheat pipesin the USA’, in Conference VGB Conference – Materials and Welding Technology inPower Plants 1994, VGB, Essen, 1994.

29 Lundin C D, Khan K K, Yang D, Hilton S and Zielke W (1990), Failure Analysis ofa Service-Exposed Hot Reheat Steam Line in a Utility Steam Plant – WRC Bulletin354, Welding Research Council, New York.

30 Dhooge A and Vinckier A, ‘Reheat cracking – a review of recent studies’, Weldingin the World, 1986, 24 (5/6), 2–17.

31 Dix A W and Savage W F, ‘Factors influencing strain-age cracking in INCONEL X-750’, Welding Journal, 1971 50 (6), 247–252.

32 Lundin C D and Khan K K, ‘Fundamental studies of the metallurgical causes andmitigation of reheat cracking in 11/4-Cr-1/2Mo and 21/4Cr-1Mo steels’, Welding ResCouncil Bull, 409, 1996.

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33 Nevasmaa P, Salonen J, Holmström S and Caminada S, ‘Heat affected zone toughnessbehaviour and reheat cracking susceptibility of thermally simulated microstructuresin new P23 (7CrWVMoNb9-6) steel’, IIW Doc. IX-A-08-06, International Instituteof Welding, 2006.

34 Dhooge A and Vekeman J, ‘New generation 21/4 Cr steels T/P23 and T/P24 – Weldabilityand high temperature properties’, IIW Doc. XI-810-04, International Institute ofWelding, 2004.

35 Nawrocki J G, DuPont J N, Robino C V and Marder A R, ‘The stress-relief crackingsusceptibility of a new ferritic steel – Part 1: Single-pass heat affected zone simulations’,Welding J, 2000 355–362.

36 Nawrocki J G, DuPont J N, Robino C V and Marder A R, ‘The stress-relief crackingsusceptibility of a new ferritic steel – Part 2: Multiple-pass heat affected zonesimulations’, Welding J, 2001 80 (1), 18–24.

37 Buchmayr B, Cerjak H, Witwer M, Maile K, Theofel H and Eckert W, ‘Experimentaland numerical investigations of the creep behaviour of the dissimilar weldment GS-17CrMoV5-11 and X20CrMoV12-1’, Steel Res, 1990 61 (6), 268–273.

38 Helander T Z, Andersson H C M and Oskarsson M, ‘Structural changes in 12-2.25%Cr weldments – an experimental and theoretical approach’, Materials at HighTemperature, 2000 17 (3), 389–96.

39 Buchmayr B, Cerjak H, Kirkaldy J S and Witwer M, ‘Carbon diffusion andmicrostructure in dissimilar Cr-Mo-V-welds and their influence on the mechanicalproperties’, in 2nd International Conference Trends in Welding Research, ASMInternational, Gatlinburg, 1989.

40 Witwer M, Untersuchungen an Mischschweissverbindungen warmfester CrMoV-Stähle, Doctoral Thesis, Graz University of Technology, 1989.

41 Brett S J, ‘Cracking experience in steam pipework welds in National Power’, inConf VGB Conference – Materials and Welding Technology in Power Plants 1994,VGB, Essen, 1994.

42 Gooch D J and Kimmins S T, ‘Type IV cracking in 1/2Cr1/2Mo1/4V/2 1/4Cr1Moweldments’, in 3rd International Conference Creep and Fracture of EngineeringMaterials and Structures, Maney Publishers, Swansea, 1987.

43 Smith D J, Walker N S and Kimmins S T, ‘Type IV creep cavity accumulation andfailure in steel welds’, International J Pressure Vessels & Piping, 2003, 80, 617–627.

44 Brear J M, Fairman A, Middleton C J and Polding L, ‘Predicting the creep life andfailure location of weldments’, Key Eng Mater, 2000, 171–174, 35–42.

45 Fujibayashi S and Endo T, ‘Creep behaviour of a low alloy ferritic steel weldment’,in 9th International Conference Creep and Fracture of Engineering Materials andStructures, Cambridge, Maney Publishing, 2001.

46 Middleton C J and Metcalfe E, ‘A review of laboratory Type IV cracking data inhigh chromium ferritic steels’, in International Conference Steam Plants for the1990’s, IMechE, London, 1990.

47 Brühl F, Cerjak H, Schwaab P and Weber H, ‘Metallurgical investigation on the basematerial and weldments of the 9% chromium X10CrMoVNb91’, Steel Res, 1991 62(2), 75–82.

48 Parker J D and Stratford G C, ‘Strain localisation in creep testing of samples withheterogeneous microstructures’, International J Pressure Vessels & Piping, 1996 68,135–143.

49 Tezuka H and Sakurai T, ‘A trigger of Type IV damage and a new heat treatment

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procedure to suppress it. Microstructural investigations of long-term ex-service Cr-Mo steel pipe elbows’, International J Pressure Vessels & Piping, 2005, 82, 165–174.

50 Shibli I A and Le Mat Hamata N, ‘Creep and fatigue crack growth in P91 weldments’,in 9th International Conference Creep and Fracture of Engineering Materials andStructures, Woodhead Publishing, Cambridge, 2001.

51 Ennis P J, ‘The mechanical properties and microstructures of 9% chromium steelP92 weldments’, OMMI, 2002, 1 (2), 1–23.

52 Takemasa F, Nonaka I, Ito T, Saitou K, Miyachi Y and Kagiya Y, ‘Type IV creepdamage analysis for full size component test on welded P91 boiler hot reheat piping’,in International Conference Elevated Temperature Design and Analysis, NonlinearAnalysis and Plastic Components, ASME, San Diego, 2004.

53 Ellis F V and Viswanathan R, ‘Review of Type IV cracking in piping welds’, in 1stInternational Conference Integrity of High Temperature Welds, IOM, London, 1998.

54 Francis J A, Mazur W and Bhadeshia H K D H, ‘Type IV cracking in ferritic powerplant steels’, Mater Sci Techn, 2006, 22 (12), 1387–1395.

55 Shinozaki K, Li D, Kuroki H, Harada H and Ohishi K, ‘Analysis of degradation ofcreep strength in heat affected zone of weldment of high Cr heat-resisting steelsbased on void observation’, ISIJ International, 2002, 42 (12), 1578–1584.

56 Hasegawa Y, Muraki T and Ohgami M, ‘Metallurgical investigation of a Type IVdamage at the heat affected zone of weld for tungsten containing martensitic heatresistant steels’, in International Conference Experience with Creep-Strength EnhancedFerritic Steels and New and Emerging Computational Methods, ASME, San Diego,2004.

57 Letofsky E, ‘Microstructure aspects of creep resistant welded joints’, IIW Doc. IX-2055-03, International Institute of Welding, 2003.

58 Albert S K, Matsui M, Watanabe T, Hongo H, Kubo K and Tabuchi M, ‘Microstructuralinvestigations on Type IV cracking in a high Cr steel’, ISIJ International, 2002, 42(12), 1497–1504.

59 Kondo M, Tabuchi M, Tsukamoto S, Yin F and Abe F, ‘Suppression of Type IVfailure in high-B low-N 9Cr-3W-3Co-NbV steel welded joint’, in 4th InternationalConference Advances in Materials Technology for Fossil Power Plants, ASMInternational, Ohio, 2004.

60 Watanabe T, Tabuchi M, Yamazaki M, Hongo H and Tanabe T, ‘Creep damageevaluation of 9Cr-1Mo-V-Nb steel welded joints showing Type IV fracture’,International J Pressure Vessels & Piping, 2006, 83, 63–71.

61 Tabuchi M, Watanabe T, Kubo K, Matsui M, Kinugawa J and Abe F, ‘Creep crackgrowth behaviour in the HAZ of weldments of W containing high Cr steel’, InternationalJ Pressure Vessels & Piping, 2001, 78, 779–784.

62 Letofsky E and Cerjak H, ‘New quantitative microstructural characterisation oncreep tested welded joints’, in 6th International Conference Trends in Welding Research,ASM, Pine Mountain, 2002.

63 Abe F and Tabucshi M, ‘Microstructure and creep strength of welds in advancedferritic power plant steels, Science Technol. Welding Joining, 2004, 9 (1), 22–30.

64 Eggeler G, Ramteke A, Coleman M, Chew B, Peter G, Burblies A, Hald J, JeffereyC, Rantala J, Dewitte M and Mohrmann R, ‘Analysis of creep in a welded P91pressure vessel’, International J Pressure Vessels & Piping, 1994, 60, 237–257.

65 Masuyama F, Komai N and Sasada A, ‘Creep failure experience in welds of advancedsteel boiler components’, IIW Doc. XI-795-04, International Institute of Welding,2004.

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66 Kimmins S T, Coleman M C and Smith D J, ‘An overview of creep failure associatedwith heat affected zones of ferritic weldments’, in 5th International ConferenceCreep and Fracture of Engineering Materials and Structures, IOM, London, 1993.

67 Li D, Shinozaki K and Kuroki H, ‘Stress-strain analysis of creep deterioration inheat affected weld zone in high Cr ferritic heat resistant steel’, Mater Sci Technol,2003, 19, 1253–1260.

68 Tabuchi M, Albert S K, Kondo M, Watanabe T and Abe F, ‘Study on the creepfracture of advanced high Cr steel weldment’, Ultra-Steel Conference: Requirementsfrom New Design of Constructions, Proceedings of 7th Workshop on Ultra Steel, 24–25 June 2003, Tsukuba, Japan, NIMS, Tsukuba, 2003.

69 Perrin I J and Hayhurst D R, ‘Continuum damage mechanics analyses of type IVcreep failure in ferritic steel crossweld specimens’, International J Pressure Vessels& Piping, 1999, 76 (9), 599–617.

70 Kimmins S T and Smith D J, ‘On the relaxation of interface stresses during creep offerritic steel weldments’, J. Strain Anal, 1998, 33 (3), 195–206.

71 Ito T, Nonaka I and Takemasa F, ‘Full size internal pressure creep test and lifeevaluation for welded Gr.91 hot reheat piping’, IIW Doc. XI-796-04, InternationalInstitute of Welding, 2004.

72 Lundin C D and Prager M, ‘A new approach to investigation into Type IV crackingsusceptibility’, in ASME/JSME Pressure Vessels and Piping Conference, 26–30 July,San Diego, CA, ASME, 1998.

73 Parker J and Bisbee L, ‘Girth weld cracking in high temperature headers’, inInternational Conference Creep & Fracture in High Temperature Components –Design & Life Assessment Issues, DEStech Publications, London, 2005.

74 Brett S J, ‘The management of weld cracking in high-temperature CrMoV pipeworksystems’, in International Conference Assuring it’s safe: Integrating Structural Integrity,Inspection and Monitoring into Safety and Risk Assessment, IMEchE, 1998.

75 Allen D J, Harvey B and Brett S J, ‘Fourcrack – An investigation of the creepperformance of advanced high alloy steel welds’, in International Conference Creep& Fracture in High Temperature Components – Design & Life Assessment Issues,DEStech Publications, London, 2005.

76 Tu S T, Segle P and Gong J M, ‘Strength design and life assessment of weldedstructures subjected to high temperature creep’, International J Pressure Vessels &Piping, 1996, 66, 171–186.

77 Morris P F, ‘Terms and Terminology for weld creep testing’, ECCC Recommendations,2001, 2(IIb).

78 Masuyama F, ‘Creep rupture life and design factors for high strength ferritic steels’,in International Conference Creep & Fracture in High Temperature Components –Design & Life Assessment Issues, DEStech Publications, London, 2005.

79 Sandstrom R and Tu S T, ‘The effect of multiaxiality on the evaluation of weldmentstrength reduction factors in high-temperature creep’, Trans ASME, 1994, 116, 76–80.

80 Takahashi Y and Tabuchi M, ‘Evaluation of creep strength reduction factors forwelded joints of HCM12A (P122)’, in 2006 ASME Pressure Vessels and PipingDivision Conference, ASME, Vancouver, 2006.

81 Tabuchi M and Takahashi Y, ‘Evaluation of creep strength reduction factors forwelded joints of modified 9Cr-1Mo steel (P91)’, 2006 ASME Pressure Vessels andPiping Division Conference, ASME, Vancouver, 2006.

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creep rupture strength of welded joints’, in International Conference Creep andFracture in High Temperature Components – Design & Life Assessment Issues,DEStech Publications, London, 2005.

83 Klenk A, ‘Creep testing of weldments: practices and investigations on the effects ofsampling and size on creep test results for weldments’, ECCC Recommendations,2005, 3(II) Appendix 1.

84 Kondo M, Tabuchi M, Tsukamoto S, Yin F and Abe F, ‘Suppressing type IV failurevia modification of heat affected zone microstructures using high boron content in9Cr heat resistant steel welded joints’, Sci Technol Welding and Joining, 2006 11 (2),216–223.

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504

18Fracture mechanics: understanding in

microdimensions

M . T A B U C H I, National Institute for MaterialsScience (NIMS), Japan

18.1 Introduction

For economic reasons, many operators plan to use power and chemical plantsbeyond their originally predicted operating life. At the same time, operatingtemperatures and loading conditions for high temperature components aretending to become more severe in attempts to improve efficiency, save energyand reduce carbon dioxide emissions. Under these circumstances, accuratelypredicting the lifespan and remaining life of high temperature componentsbecomes much more important. Creep voids can form in thick componentsoperating under stress and temperature gradients and these voids grow intomicro and macro cracks, eventually causing fracture. It is therefore importantto be able to predict crack initiation, growth and time to fracture under hightemperature creep conditions.

The application of fracture mechanics to high temperature creep crack growthfirst began in the 1970s. A large amount of research concerning fracture mechanicsparameters, which control the stress–strain field in front of the creep crack, hasbeen conducted using several types of specimens, including CT, CCT, SENand DEN. The C* parameter,1,2 based on the J integral,3 is generally used tocharacterize the stress-strain field in front of the crack tip and to correlate crackgrowth rate under extensive creep conditions. C(t),4,5 Ct

6 and Q*7 can also beused to evaluate creep crack growth properties. The standard test method formeasuring and evaluating high temperature creep crack growth for creep ductilematerials, ASTM E1457, was established in 19928 and revised to include creepbrittle materials in 2000,9 based on the results of VAMAS TWA11 and 19projects. Recently, the application of non-linear fracture mechanics to crackedcomponents has become an important subject.10

18.2 Non-linear fracture mechanics

Just as J characterizes the crack tip fields in an elastic or elastic–plasticmaterial, the stress and strain rate distribution in front of a crack tip under

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Fracture mechanics: understanding in microdimensions 505

steady-state creep conditions can be characterized by a path-independentenergy rate line integral, C* integral,1,2 which is defined by analogy to J.3

The definition of C* is the following line integral on a contour Γ surroundinga crack tip shown in Fig. 18.1:

C W y Tux

sii* = d – d∫ ∂

∂

Γ ˙ ˙[18.1]

˙ ˙˙

Wij

ij ij = d0

εσ ε∫ [18.2]

where W is strain energy rate density, σij and ε ij are the stress and strain ratetensors respectively, Ti is the outward traction vector, ui is the displacementrate vector and ds is the arc length along contour path Γ.

C* can be determined from a load–displacement rate diagram for a creepingmaterial. C* is defined from the shaded area of Fig. 18.2:

y

rθ

ds

x

Γ

18.1 Integration contour around the crack tip for the definition ofnon-linear fracture mechanics parameter C*.

18.2 Load–creep displacement rate responses for different crack sizesof non-linear material.

P

adU*

a + da

Creep displacement rate

Load

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CB

Ua

* = – 1 d *d

[18.3]

U P* = d0

˙˙

δδ∫ [18.4]

where U* is the potential energy rate, a is the crack length, δ is load linedisplacement rate and B is the thickness. For power-law creep, the C* parametercan be described as:

CP

B W aF* =

( – )δ

[18.5]

where W is the width of the plate and F is a non-dimensional factor whichdepends on geometry and creep exponent n of the power-law creep. For CTspecimens, the following calculation method for F is recommended in ASTME1457:8,9

F nn

W aW

= + 1

2 + 0.522 –

[18.6]

Usually, the relationship between creep crack growth rate (da/dt) and C* isgiven as:

dd

= *at

DC φ [18.7]

where D and φ are constants. Figure 18.3 shows the relationship betweencreep crack growth rates and C* obtained by Japanese research groups for1Cr–Mo–V turbine rotor steel using several sizes of CT specimen.11 C*correlates to creep crack growth rate better than net section stress (σnet) andstress intensity factor (K). In the relationship between da/dt and C* in Fig.18.3, transient behaviour (the tail part) is observed in the early stages ofcrack growth, where the values of C* and da/dt decrease. Only after thesteady-state creep crack growth stage can da/dt be expressed as Equation[18.7]. Transient crack growth behaviour is considered to occur through acombination of stress redistribution and primary creep.

ASTM E1457 recommends evaluating only the data for which time exceedstransition time, tT, by C*. The transition time is estimated as:

tKE n CT

2 2

= (1 – )

( + 1) *ν

[18.8]

where E is Young’s modulus, ν is Poisson’s ratio and K is the stress intensityfactor. The data where time is lower than tT are correlated using Ct.

6

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18.3 Effect of mechanical constraint

The effects of crack tip constraint arising from variations in specimen size,geometry and material ductility can influence the relationship between da/dtand C* in Equation [18.7].9 Figure 18.4 shows the effect of specimen size onda/dt versus C* for 1Cr–Mo–V steel using large CT (W = 254 mm) andstandard CT (W = 50.8 mm) specimens.12 Note that an increase in specimenthickness increases crack growth rate. The da/dt of the thickest specimen (B= 63.5 mm) with side grooves (SG) is about six times higher than that of thethinnest specimen (B = 6.35 mm) without SG. Figure 18.5 shows the fracturesurface of large CT specimens, 63.5 mm in thickness. Comparing non-SGspecimens with SG specimens, while the crack length at the thickest partwas nearly the same, the crack length at the surface was one-fifth that in SGspecimens. The reason why side grooves and thickness accelerate creepcrack growth rate is considered to be due to mechanical constraint ahead ofthe crack tip.

The stress and strain rate field under power-law creep, ε σ = A n , is givenas follows (HRR field):13,14

σ σ θijn

n

ijC

I Arn = * ( , )

1/ +1

˜ [18.9]

˙ ˙ε ε θijn

n n

ijA CI Ar

n = * ( , )/ +1

[18.10]

da /

dt

(mm

h–1

)100

10–1

10-2

10–3

10–1 100 101 102

C* (kJ m–2 h–1)

18.3 Relationship between creep crack growth rate versus C*parameter for a 1Cr–Mo–V turbine rotor steel.

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da /

dt

(mm

h–1

)

1

0.1

0.01

0.001

Plane

strain

254.0, 63.5, S.G.254.0, 63.5, Non S.G.254.0, 12.7, S.G.

W B

Plane

stres

s W B50.8, 25.4, S.G.50.8, 12.7, S.G.50.8, 6.35, S.G.50.8, 6.35, Non S.G.

0.01 0.1 1 10 100C* (kJ m–2 h–1)

18.4 Effect of specimen thickness, width and side groove on thecreep crack growth rate of 1Cr–Mo–V steel.

W = 254 mm, B = 63.5 mm without SG

W = 254 mm, B = 63.5 mm with SG

18.5 Fracture surface of large CT specimens with 63.5 mm inthickness after a creep crack growth test at 811 K for 1Cr–Mo–V steel.

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where σ ij and ε~

ij are non-dimensional functions of θ and n, and In is a non-dimensional function of n. From Equation [18.10], the following equation tocharacterize creep crack growth rate is derived:15,16

dd

= + 1 ( ) *

f* c

1/ +1/ +1a

tn Ar C

In

n

n n

ε

[18.11]

where ε f* is creep ductility and rc is the size of the creep process zone. Under

plane stress conditions, ε f* is the creep ductility obtained from smooth round

bar creep specimens. The creep crack growth rate calculated using Equation[18.11] is shown in Fig. 18.4 as the line that corresponds well with theexperimental data for the thinnest non-SG specimen (B = 6.35 mm). Whenthe specimen thickness increases and side grooves are introduced, ε f

* aheadof the crack tip decreases and the crack growth rate increases. The crackgrowth rate under plane strain conditions in Fig. 18.4 is obtained by assumingthat creep ductility under such conditions is 1/50 of uniaxial ductility.15 Theε f

* under a multiaxial stress state can be calculated as follows:17,18

εε

σσ

f*

f

m

e = sinh 2

3 – 1/2 + 1/2

sinh 2 – 1/2 + 1/2

nn

nn

[18.12]

where ε f* and εf are creep ductility under multiaxial and uniaxial condition,

and σm and σe are hydrostatic stress and equivalent von Mises stress,respectively. It is therefore necessary to keep the component dimensions inmind when applying the experimental data to large-scale structuralcomponents.9

18.4 Effect of microscopic fracture mechanisms

Creep crack growth properties are affected by microscopic fracture mechanisms,which change according to temperature and loading conditions. Over longoperating times, creep crack growth by grain boundary cavitation is oftenobserved. In most high-temperature, thick-section components operating undermultiaxial stress conditions, creep cracks grow accompanied by micro-cavities,voids and micro-cracks ahead of the crack tip. It is therefore important totake microscopic aspects into account when evaluating creep crack growthproperties with the aim of developing more accurate predictions of operatinglife.

Figure 18.6 shows the microscopic features of creep cracks observed inCT specimens of Alloy 800H.19 Three types of creep crack growth, brittleintergranular fractures caused by wedge-type cracks (W-type) at lowertemperature, ductile transgranular fractures (T-type) and void-type intergranular(V-type) fractures at higher temperature, were observed, depending on thetesting temperature and loading conditions. Large creep damage zones ahead

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of the crack tip were observed for V-type crack growth. The relationshipbetween da/dt and C* for Alloy 800H at various temperatures is shown inFig. 18.7. These relationships are dependent on the microscopic crack growthmechanisms; creep crack growth rate for W-type and V-type fracture washigher than that for T-type fracture. For 316 stainless steel, the same tendencywas reported.20

Figure 18.8 shows the da/dt at a constant C* value (1kJ/m2h) plottedagainst creep ductility in uniaxial creep tests for Alloy 800H and 316 stainless

18.6 Microscopic features of creep cracks of Alloy 800H tested at (a)873K, (b) 973K and (c) 1073K.

Void-type creep crack

Transgranular creep crack

Wedge-type creep crack

(a)

(b)

(c)

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steel,21 which is dependent on creep fracture mechanisms under theseconditions. For W-type and T-type fractures, da/dt is inversely proportionalto creep ductility, as shown by the solid curve in Fig. 18.8. Therefore, thecreep crack growth rate for W-type and T-type fractures can be predictedaccording to Equation [18.11]. On the other hand, for V-type fractures, da/dt is faster than that predicted from creep ductility. When creep damage isonly just in front of the crack tip, as in W-type and T-type fractures, da/dt canbe characterized by the creep ductility. Where there is a large creep damage

Cre

ep c

rack

gro

wth

rat

e, d

a /d

t (m

m h

–1)

1

0.1

0.01

0.001

0.0001Alloy 800H

0.01 0.1 1 10 100C* (kJ m–2 h–1)

Temp. Load Fracture(K) (kN) mode

873 13.69 Wedge-type

11.47 Wedge-type923 11.71 Transgranular

8.98 Transgranular973 7.62 Transgranular

4.36 Transgranular1073 3.11 Void-type

2.61 Void-type

18.7 Relationship between da/dt versus C* parameter for Alloy 800Hat various temperatures.

da/

dt

at C

* =

1 (k

J m

–2 h

–1)

(mm

h–1

)

0.08

0.06

0.04

0.02

0

W

V

W

W

V

V

V

T T T

0 20 40 60 80 100Reduction in area (%)

Wedge-typeTransgranularVoid-type

Alloy800H316 stainless

18.8 Relationship between da/dt at constant C* value 1 kJ m–2h–1

versus creep ductility for Alloy 800H and 316 stainless steel.

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zone formed ahead of the crack tip, as in voids and micro-cracks, da/dtaccelerates faster than predicted from creep ductility.

Analytical models for creep crack growth taking into account the effect ofcreep damage formed ahead of the crack tip have been proposed.5,16 Accordingto these analytical models, increase in creep damage (void density) in frontof the crack tip accelerates the creep crack growth rate. It can be consideredthat, under creep conditions, the diffusion of vacancies towards cracks andvoids would contribute to crack growth. The vacancy diffusion equationunder stress gradient is given as follows:22,23

∂∂

∇ ∇ ∇( )Ct

D C CRT

v = + [18.13]

ν = –σp∆V [18.14]

where C is the vacancy concentration, D is the diffusion coefficient, R is thegas constant, T is the absolute temperature, σp is the hydrostatic stress and∆V is the change to molar volume caused by vacancy diffusion. In thisequation, vacancy diffusion is controlled by the hydrostatic stress gradient.The modified equation to calculate realistic vacancy diffusion and concentrationusing weight coefficients α1 and α2 is as follows:24,25

∂∂

∇ ∇ ∇Ct

D C DRT

C = + 12

2α α ν [18.15]

The vacancy diffusion equation is solved for three dimensional CT specimenmodels using the hydrostatic stress gradient ∇σp computed by the finiteelement method (FEM).21 Figure 18.9 shows the computed examples forchanges in vacancy concentration ahead of the crack tip during creep. Thevacancy concentration increases faster and to higher levels in the centre of

18.9 Computed results for changes in vacancy concentration aheadof the crack tip during creep.

D = 1.5 × 10–9 m2s–1

D = 1.5 × 10–10 m2s–1

Vac

ancy

co

nce

ntr

atio

n C

/C0

3.0

2.5

2.0

1.5

1.0

0.50 100 200 300

Time (h)

Alloy800H

Surface

Surface

Centre

Centre

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specimen than at the surface. This is consistent with experimental resultswhere creep voids were often observed in the thickest part of the specimen,where stress multiaxiality is large. When D is low, at lower temperatures,cracks propagate before vacancies are accumulated, which corresponds totransgranular-type crack growth.

Creep crack growth properties can be characterized by creep ductility.Creep damage formed in front of the crack tip accelerates creep crack growth.Vacancy diffusion, void formation, and crack initiation and propagation areaccelerated under multiaxial stress fields at high temperatures.

18.5 Type IV creep crack growth in welded joints

In an attempt to reduce carbon dioxide emissions and save energy in thermalpower plants, the steam pressure and temperature conditions under whichboiler components are expected to operate is increasing. High Cr ferriticsteels (9–12%Cr steels), such as P91 (9Cr–1Mo–VNb), P92 (9Cr–0.5Mo–1.8W–VNb) and P122 (11Cr–0.4Mo–2W–CuVNb) steels, which have atempered martensite structure, are now used for boiler components in ultra-supercritical (USC) power plants operating at around 873 K because of theirhigh creep strength. The creep strength of weldment in these steels, however,decreases at above 873 K owing to type IV creep damage formed in the heat-affected zone (HAZ).26–30 It is important to predict initiation and growth oftype IV creep damage in the HAZ of the weldment.

The creep properties of HAZ are usually investigated using simulatedHAZ specimens, which are produced by rapid heating to peak temperaturesaround the AC1 to AC3 transformation temperature and cooling. Figure 18.10shows the relationship between creep rupture times and peak temperaturesduring simulated HAZ heat treatment for P122 steel at 923 K.29 The creeprupture time shows a minimum value for the specimens heated to around theAC3 transformation temperature. The microstructure of HAZ heated to AC3 ischaracterized by a fine-grained structure without lath martensite, with agrain size of about 5 µm. The creep rupture time for the simulated HAZspecimen heated to AC3 was about one-fifth that of base metal. Therefore,lower creep strengths in fine-grained HAZ structures are considered to bethe primary cause of type IV failure.

The width of the HAZ is very narrow in the welded joint. Therefore thecreep deformation of fine-grained HAZ (with low creep strength) ismechanically constrained by weld metal and base metal, which have highercreep strengths. Figure 18.11 shows the creep crack propagation profile in awelded joint of P92 steel. This test was conducted using a CT specimen witha notch tip located in the HAZ and fatigue pre-cracked. A very sharp creepcrack, which propagates in fine-grained HAZ, type IV creep crack, is observed.Creep voids were observed around the main cracks in welded joint specimens.

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The C* line integral of Equation [18.1] can be calculated by FEM usingpower-law creep obtained from uniaxial creep tests for simulated HAZ, basemetal and weld metal. Figure 18.12 shows the computed C* value for basemetal, welded joints and fine-grained HAZ for P122 steel plotted againstloading time.31 Here, the simulated HAZ specimen has a fine-grained structure

Tim

e to

ru

ptu

re (

h)

10 000

1000

100

Base Ac1 Ac3

120 MPa

P122 steel923 K

80 MPa

100 MPa

1100 1200 1300 1400 1500Peak temperature (K)

18.10 Changes in creep rupture time as a function of peaktemperature of simulated HAZ heat treatment for P122 steel at 923 K.

Creepcrack

Pre-crack

Basemetal

HAZ

Weldmetal

18.11 Creep crack propagation profile in welded joint of P92 steelusing CT specimen.

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and the width of fine-grained HAZ for welded joints is 1.2 mm. The computedC* value for the simulated HAZ is about one order higher than that of basemetal for the same loading conditions. The C* value of welded joints decreasesas the width of HAZ decreases.

Figure 18.13 shows the creep crack growth behaviour of welded joints,base metal and simulated fine-grained HAZ for P122 steel tested using CTspecimen at 923 K under the same loading conditions.31 For the weldedjoints, the early crack growth rate is nearly the same as that in base metal,but the crack grows rapidly in the accelerating stage after the incubationperiod. Consequently, the crack growth life of welded joints is about one-third that of base metals. The relationship between creep crack growth rateda/dt and C* in welded joints and base metal for P122 steel is shown in Fig.18.14. The creep crack growth rates of the welded joints are higher thanthose for base metal in higher C* regions, but the differences in da/dt betweenwelded joints and base metal are small in the lower C* region. From thisresult, the prediction of crack initiation time (incubation period), and whencrack growth accelerates, is clearly important for the evaluation of fracturelife of the welded joints.

The crack growth behaviour of welded joints mentioned above can becorrelated with stress multiaxiality in HAZ. Creep deformation for fine-grained HAZ with low creep strength is mechanically constrained by theweld metal and base metal, which have higher creep strength. Figure 18.15shows the computed stress triaxial factor TF in front of the crack tip forwelded joints and base metal, P92 steel, using FEM for three-dimensionalCT specimens with side grooves.32 TF is calculated as follows:

C*

(kJ

m–2

h–1

)

1

0.1

0.01

0.0011 10 100 1000

Time (h)

P122923 K

Simulated HAZ

Welded joint

Base metal

18.12 Computed C* value by FEM for base metal, welded joints andsimulated HAZ for P122 steel.

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TF = + + 1 2 3

eq

σ σ σσ [18.16]

where σ1, σ2, and σ3 are principal stress and σeq is equivalent von Misesstress. In Fig. 18.15, the value of TF for welded joints is larger than that forbase metal. The creep deformation in the HAZ of welds is small in the initial

Cre

ep c

rack

len

gth

(m

m)

8.0

6.0

4.0

2.0

0.00 500 1000 1500 2000

Time (h)

Simulated HAZ

Weldedjoint

Basemetal

P122 steel 923 KKin = 18 MPa m

18.13 Creep crack growth behavior of welded joints, base metal andsimulated fine-grained HAZ for P122 steel at 923 K for the sameloading condition.

Base metalWelded jointSimulated HAZ

P122 steel 923 K

0.001 0.01 0.1 1 10C* (kJ m–2 h–1)

da/

dt

(mm

h–1

)

1

0.1

0.01

0.001

0.0001

18.14 Relationship between creep crack growth rate da/dt versus C*parameter of welded joints and base metal for P122 steel.

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18.15 Stress triaxial factor TF ahead of the crack tip of three-dimensional CT specimen with S.G. for welded joint and base metalof P92 steel.

Base metalHeat affected zone

0 5 10 15Distance from the crack tip (mm)

Increase in creeptime up to 500 h

Increase in creeptime up to 500 h

3D elastic–plastic creep FEM analysisModel : CT specimen with SGTemp : 650°CLoad : 10000 N (Kin = 15 MPa m1/2)

Tria

xial

fac

tor

TF

12

10

8

6

4

2

0

stages owing to mechanical constraint under multiaxial stress conditions.However, void formation, growth and crack initiation are accelerated undermultiaxial stress conditions for welded joint specimens. Predicting the lifespanof welds therefore requires evaluating the initiation of creep voids and cracksunder multiaxial stress conditions.

18.6 References

1 K. Ohji, K. Ogura and S. Kubo, Proceedings 1974 Symposium on Mechanical Behaviorof Materials 1, The Society of Materials Science, Japan, 1974, 455–466.

2 J. D. Landes and J. A. Begley, ASTM STP 590, ASTM, 1976, 128–148.3 J. R. Rice, Trans. ASME, J. Appl. Mech., 1968, 35, 379–385.4 R. Ehlers and H. Riedel, Proceedings of ICF5, Volume 2, Pergamon Press, 1981,

691–698.5 H. Riedel, Fracture at High Temperatures, Springer-Verlag, Berlin, 1987.6 A. Saxena, ASTM STP 905, ASTM, 1986, 185–201.7 A. T. Yokobori, Jr. and T. Yokobori, Proceedings International Conference on Creep,

JSME, 1986, 135–140.8 ASTM E1457-92: Standard Test Method for Measurement of Creep Crack Growth

Rate in Metals, ASTM, 1992 1031–1043.9 ASTM E1457-00: Standard Test Method for Measurement of Creep Crack Growth

Rate in Metals, ASTM, 2000, 936–950.

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10 Code of Practice for Creep/Fatigue Testing of Cracked Components, ISO/TTA E2006(in press).

11 T. Yokobori, C. Tanaka, K. Yagi, M. Kikagawa, A. Fuji, A. T. Yokobori, Jr. and M.Tabuchi, Mater. High Temp., 1992, 10, 97–107.

12 M. Tabuchi, K. Kubo and K. Yagi, Eng. Fract. Mech., 1991, 40, 311–321.13 J. W. Huchinson, J. Mech. Phys. Solids, 1968, 16, 13–31.14 J. R. Rice and G. F. Rosengren, J. Mech. Phys. Solids, 1968, 16, 1–12.15 K. M. Nikbin, D. J. Smith and G. A. Webster, Trans. ASME J. Engng. Mater. Tech.,

1986, 108, 186–191.16 G. A. Webster and R. A. Ainsworth, High Temperature Component Life Assessment,

Chapman & Hall, London, 1994.17 A. C. F. Cocks and M. F. Ashby, Met. Sci., 1980, 14, 395–402.18 M. Yatomi, K.M. Nikbin and N.P. O’Dowd, Int. J. Pressure Vessels Piping, 2003, 80,

573–583.19 M. Tabuchi, K. Kubo and K. Yagi, Tetsu-to-Hagane, 1993, 79, 732–738 in Japanese.20 M. Tabuchi, K. Yagi and T. Ohba, ISIJ International, 1990, 30, 847–853.21 M. Tabuchi, J. Ha, H. Hongo, T. Watanabe and A. T. Yokobori Jr., Metall. Mater.

Trans. A, 2004, 35A, 1757–1764.22 H. P. Leeuwen, Eng. Fract. Mech., 1974, 6, 141–161.23 H. P. Leeuwen, Eng. Fract. Mech., 1977, 9, 951–974.24 A. T. Yokobori, Jr., T. Nemoto, K. Sato and T. Yamada, Eng. Fract. Mech., 1996, 55,

47–60.25 A. T. Yokobori, Jr., Y. Chinda, T. Nemoto, K. Sato and T. Yamada, Corrosion Sci.,

2002, 44, 407–424.26 G. Eggeler, A. Ramteke, M. Coleman, B. Chew, G. Peter, A. Burblies, J. Hald, C.

Jefferey, J. Rantala, M. deWitte and R. Mohrmann, Int. J. Pressure Vessels Piping,1994, 60, 237–257.

27 F. Masuyama, M. Matsui and N. Komai, Key Engng Mater., 171–174, 2000, 99–107.28 T. H. Hyde, W. Sun and A. A. Becker, Int. J. Pressure Vessels Piping, 2001, 78, 765–

77.29 M. Tabuchi, T. Watanabe, K. Kubo, M. Matsui, J. Kinugawa and F. Abe, Int. J.

Pressure Vessels Piping, 2001, 78, 779–784.30 Y. Hasegawa, T. Muraki and M. Ohgami, Tetsu-to-Hagane, 2004, 90, 609–617 in

Japanese.31 M. Tabuchi, H. Hongo, T. Watanabe and A. T. Yokobori, Jr., J. ASTM Inte, 2006, 3

(5), online.32 R. Sugiura, A. T. Yokobori, Jr., M. Tabuchi and T. Yokobori, Eng. Fract. Mech.,

2007, 74, 868–881.

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519

19Mechanisms of oxidation and the influenceof steam oxidation on service life of steam

power plant components

P. J. E N N I S and W. J. Q U A D A K K E R SForschungszentrum Juelich GmbH, Germany

19.1 Introduction

In the development of commercial 9% chromium steels, the principal aimhas been a significant increase in the creep rupture strength, to enableapplication temperatures of 600°C and even higher. Creep strength targetswere achieved, so that with the currently available commercial steels P91/T91, E911 and P92/T92, metal temperatures of up to 625°C can be considered,based on the creep rupture strength. Creep rupture testing to test durationsof many thousands of hours have been carried out on all these steels andit was found that a chromium content of 9mass% promoted the formationof highly protective oxide scales that provided excellent protection againstoxidation in air. However, the service environments of steam power plantcomponents are either steam or combustion gases, which also containsignificant levels of water vapour. Investigations into the effect of suchservice environments on oxide scale formation showed that in the presenceof water vapour, the oxidation rates of steels containing around 9mass% ofchromium were considerably accelerated. In contrast to thin protective scalesbased on haematite (Fe2O3) found after exposure in air, thick and highlydefective external and internal scales of mainly magnetite (Fe3O4) wereformed in steam-containing atmospheres. Figure 19.1 compares the oxidescales formed on the P92 steel in air and in an argon-50vol% steam atmosphereafter 10 000 h exposure at 650°C; the associated mass gains were 0.1 and45 mg cm–2, respectively.

The results shown in Fig. 19.1 are from specimens exposed in an argon-50vol% water vapour atmosphere not in pure steam. In laboratory testingthis is a convenient means of assessing the steam oxidation resistance ofsteels. Instead of a steam generator, argon is passed through a water bath ata constant temperature. By adjusting the temperature of the water bath, theconcentration of water vapour in the argon can be controlled. After passingthrough the water bath, the gas is led through heated pipes to the furnaceretort, to prevent the water condensing out. Comparison tests have shown

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that there is no significant difference between testing in pure steam and inargon-50 vol% H2O.

19.2 Mechanisms of enhanced steam oxidation

19.2.1 Stability of oxides

Figure 19.2 shows the thermodynamic stability of the oxides of iron andindicates which oxide, if any, is stable at a given temperature and oxygenpartial pressure of the environment for an iron activity of 1. In air (oxygenpartial pressure 0.2 bar) at 600°C, the stable oxide is haematite Fe2O3, whichhas a narrow homogeneity range with a low concentration of lattice defects(vacancies). Diffusion of ions through haematite is therefore relatively slowand as a result oxide growth rates are low. Oxide scales based on Fe2O3 willtherefore tend to be protective as evidenced by the results of long-termexposure of the chromium steels in air (Fig. 19.1).

Oxidation of iron by steam leads to the production of hydrogen, and thepresence of hydrogen in the atmosphere above the oxide scale reduces theoxygen partial pressure very rapidly. Haematite formation can no longer besustained and the next oxide, Fe3O4, magnetite, will form. This oxide has awide homogeneity range, that is, it can accommodate excess iron and oxygenions, maintaining electrical neutrality by an increasing concentration of lattice

19.1 Micrographs for P92 steel of the oxide scale formed after 10000 hat 600°C (a) in air and (b) in Ar-50 vol % steam (note difference inmagnification). (a) P92 steel, 10 000h/650°C in air, mass gain 0.1 mgcm–2; (b) P92 steel, 10000 h/650°C in Ar-50%H2O, mass gain 45 mgcm–2.

Ni layer to preserve scaleduring preparation

20 µm 200 µm

(a) (b)

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Mechanisms of oxidation and the influence of steam oxidation 521

vacancies. This allows very rapid ionic diffusion and leads to high oxidegrowth rates. Fe3O4 scales, if they form, are therefore thick and contain ahigh concentration of voids and gaps.

19.2.2 Stages of steam oxidation

Based on the results of short-term testing, Zurek (2004) proposed that theoxide scale formed on 9–12% chromium steels in steam-containingenvironments followed the sequence shown in Fig. 19.3.

• At time t1, a thin, protective scale based on (Fe,Cr)2O3 or (Fe,Cr)3O4

forms. The Cr/Fe ratio is greater than about 0.25 (Rahmel and Tobolski,1965).

• At time t2, the protective layer breaks down locally, perhaps owing tothe depletion of chromium beneath the scale so that the Cr/Fe ratiodecreases. Rapid growth of Fe3O4 then takes place and owing to the highoutward diffusion rates of the iron ions, scale grows at the gas/oxideinterface. The oxide surface is then not in equilibrium with the gas, theoxygen partial pressure at the oxide surface is lower than that in the gasand haematite cannot form. The rapid outward diffusion of ions to thescale/gas interface causes the formation of vacancies that condense intovoids and gaps. Beneath the magnetite scale, the oxygen partial pressureis still sufficiently high for internal oxides FeO and Cr2O3 to be formed.

Temperature (°C)900 800 700 600 500 400 300

Fe3O4 magnetite

Fe2O3 haematite

FeO wustite

Log

(o

xyg

en p

arti

al p

ress

ure

) (b

ar)

–5

–10

–15

–20

–25

–30

–358 10 12 14 16 18

Reciprocal temperature (K–1)

19.2 Thermodynamic stability of iron oxides at an iron activity of 1.

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522

19.3 Stages in the steam oxidation of 9–12% chromium steels; the time scale varies according to the exact compositionof the steel.

H2O H2 H2OH2

Alloysurface

Protective spinel

Alloy Alloy

FeO + Cr2O3 Fe3O4 + Cr2O3

Alloysurface

Protective spinel

Fe3O4

H2O H2

Alloy

Fe3O4 + (Fe, Cr)3O4

Fe3O4

Originalalloy

surface

Fe3O4 + Cr2O3

Fe3O4

FeO + Cr2O3

Alloy

FeO + Cr2O3

Fe2O3

Fe3O4

Alloy

Fe3O4 + (Fe, Cr)3O4

FeO + Cr2O3

Originalalloy

surface

t5t4t3

t1 t2

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• At time t3, thickening of the inner scale leads to partial transformation ofFeO to Fe3O4 and the Cr2O3 oxide particles become embedded in thescale. Further transformation leads to the formation of the spinel (Fe,Cr)3O4.Because Cr2O3 can only be formed internally, the place beneath whichthis oxide is observed coincides with the original alloy surface (Quadakkerset al., 2005).

• At time t4, the gap in the outer magnetite scale becomes more extensiveand the area available for diffusion paths between the gas/oxide andoxide/alloy interfaces is increasingly restricted.

• At time t5, the gap in the outer scale is now so large that outwarddiffusion of iron ions practically stops. The activity of iron in the outermostscale will then decrease and haematite will be formed. If haematite isobserved on the outermost surface of a steel oxidized in steam, thisindicates a high concentration of macroscopic defects in the scale (voidsand gaps) with associated poor scale adherence.

The actual times t1 – t5 vary widely depending on the steel composition,temperature, gas flow rates and other factors.

19.2.3 Development of defects and spalling of oxidescales

The presence of defects such as voids and gaps in the oxide scales formed insteam- containing atmospheres will affect the adherence of the scales, especiallyduring thermal cycling. The spalling (exfoliation) of scales may lead toenhanced oxidation rates as fresh material is exposed to the atmosphere.

Figure 19.4 shows the scale microstructure on 10Cr–Mo–W exposed at650°C in Ar-7% H2O for 5, 10, 50, 72 and 100 h. Formation of pores andeventually gaps in the individual stages of their evolution can be observed.After 5 h oxidation the scale is still mostly compact; owing to local vacancycondensation, some voids start to form on the original metal surface. After10 h oxidation the voids on the original metal surface become larger andthey begin to affect further oxide growth. The oxygen partial pressure in thegap probably stays constant and thus below the gap the gradient must becomelarger as the oxygen partial pressure at the scale/alloy interface equals thedissociation pressure of the oxide. Owing to the steeper oxygen partial pressuregradient, the transport of iron ions from the metal must increase. Thus‘secondary’ growth of oxide scale in the gap becomes possible and the gapmoves outward. Because of the enhanced iron ion outward diffusion, asecondary gap at the interface between the internal oxidation zone (FeO +Cr2O3) and the Fe, Cr spinel layer is formed. The appearance of haematite onthe top of the scale is a result of the large gap in the scale. During thefollowing hours of oxidation, the gradient of oxygen partial pressure remains

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unchanged, the gap becomes larger and the external haematite layer starts tospall. The completely spalled haematite layer allows an excess of molecularoxygen to enter the external voids and/or gap and healing of gaps formed inprevious stages of oxidation occurs. Figure 19.4 for 100 h clearly illustratesthe situation described above, where on the left-hand side a nearly compact

(a) (b)

(c) (d)

(e)

5 h 10 h

50 h

Gap

72 h

Gap

100 h

Compactoxide scale

Voids

Alloy50 µm

Alloy50 µm

Pores

Secondary gap

Secondary gap

Alloy50 µm

Alloy50 µm

Alloy50 µm

Gap

Secondary gapNearly compact

oxide scale

19.4 Metallographic cross-sections of 10Cr–Mo–W after exposure for5, 10, 50, 72 and 100 h at 650°C in Ar-7% H2O showing formation andevolution of the in-scale gap.

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oxide scale has formed (complete healing of the gaps) while on the right-hand side a very porous scale can be seen.

The considerations presented above concern the formation and evolutionof the ‘transient’ gap, which is formed after several minutes or hours ofoxidation. The voids and gaps observed after a few thousand hours oxidationare the result of other factors such as cooling during exposure, cracking,steel microstructure and formation of Cr-rich stringers in the inner part of theoxide scale. The first voids, as already shown, start to form after very shortoxidation times. The cracking and spallation of scales is correlated with thetype, morphology and growth of pores and voids in the scale and could beinfluenced by the steel microstructure. The formation of pores and gaps isthe dominating factor for the spalling characteristics of the oxide scalesformed in steam.

If isothermal exposure is continued, the gap may heal by growth of theinner oxide scale as gas molecules gain access through the outer layer. Ifisothermal exposure is continued for sufficiently long times, the gap mayeventually completely close which results in a more or less compact layer.Transport of Fe cations to the surface becomes possible again and thus theouter haematite transforms into magnetite after longer times (a few hundredhours) at 650°C.

If a thermal cycle is introduced in the early stages of the oxidation process,the poor adherence caused by the presence of the large gap results in spallationof the outer scale. This is illustrated in Fig. 19.5(a) for a 10CrMo–W steelexposed for 10 000 h at 625°C. Further exposure then results in growth ofthe freely exposed magnetite layer. Figure 19.5(b) shows the scalemicrostructure of the same steel after 10 000 h exposure at 650°C showingthat spalling has occurred within the inward growing oxide scale, owing tothe accumulation of voids there.

The question of whether spallation or healing of the outer layer above thegap occurs depends on the time at which for a given temperature the thermalcycle is introduced. This can be seen from the results in Fig. 19.4 for the10Cr–Mo–W–Si steel. If the specimen is cooled after 72 h exposure, spallationof the top layer occurs owing to the presence of the large in-scale gap. If thefirst temperature cycle is introduced after 250 h, sufficient time is availablefor (partial) healing of the gap and thus excellent oxide adherence is foundeven during exposure for several thousand hours.

19.3 Steam oxidation rates

19.3.1 Long-term exposure

The long-term steam oxidation behaviour of the ferritic Cr steels can berationalized on the basis of the schematic in Fig. 19.6. This graph should

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only be considered as a rough approximation to the real behaviour becausethe borderlines between the various oxidation regimes can be stronglyinfluenced by alloying additions, temperature and surface treatment. Thethree groups may be distinguished as follows:

(I) steels containing up to around 9–10%Cr form thick oxide scales, themain constituent of which is magnetite;

Initiation ofspallation

Alloy 100 µm

(a)

Initiation ofspallation

Alloy

100 µm

(b)

19.5 Different types of scale spallation during oxidation on 10Cr–Mo–W steel at (a) 625°C after 10000 h and (b) 650°C after 10 000 h in Ar-50% H2O.

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(II) steels containing 10–12%Cr exhibit highly variable steam oxidationbehaviour, the thickness and morphology of the oxide scales differingsubstantially as a function of test duration, temperature, minor alloyingadditions, grain size and surface treatment;

(III) steels with more than about 12.5%Cr possess excellent steam oxidationresistance and the thin scales consist mainly of Cr2O3, (Cr,Fe)2O3 orCr-rich (Cr,Mn,Fe)3O4 with an outermost layer of Fe2O3.

19.3.2 Influence of chromium content

The commercial 12%Cr–1%Mo steel should, on the basis of Fig. 19.6, exhibitgood steam oxidation resistance. However, the chromium content specifiedfor this steel is 10–12.5%, which means that if the chromium content is at thelow end of the specified range, the steam oxidation resistance will be similarto the low chromium steels. If the chromium content is at the high end of thespecification, then good steam oxidation resistance would be expected. Toensure adequate steam oxidation resistance, it is therefore important that theminimum chromium content is above about 12%.

19.3.3 Effect of minor alloying additions

Other minor alloying elements can significantly affect the steam oxidationresistance of the high chromium steels. Elements which promote the formation

Group I Group II Group IIIlo

g K

p

6 8 10 12 14 16Cr content (mass%)

19.6 Schematic illustration showing qualitative dependence ofoxidation rate on Cr content for ferritic/martensitic steels exposed inAr-50% H2O at 550–650°C; designations ‘Group I – III’ are explained inthe text.

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of spinel-type oxide scales or which enhance the diffusion rates may bebeneficial. There is some evidence that small additions of silicon and manganeselead to improved steam oxidation resistance (Quadakkers and Ennis, 2002)and the steels that show the highest steam oxidation resistance, such asVM12 (Gabrel et al., 2006), contain 11–12% chromium and around 0.5%silicon.

19.3.4 Anomalous temperature dependence

The formation of a protective oxide scale depends on the incorporation ofchromium into the external scale and therefore the diffusion rate of chromiumin the steel is an important factor. Because the diffusion rate of chromiumincreases with temperature, it is possible for the steam oxidation resistanceto increase with increasing temperature and this has been observed by Zureket al. (2004) for a number of chromium steels (Fig. 19.7). This effect shouldbe taken into account in the development of steels with improved steamoxidation resistance, as testing at the highest application temperature maygive a misleading assessment; the oxidation rate at lower temperatures couldwell be higher.

19.3.5 What is an acceptable oxidation rate?

Components are generally designed on the basis of the strength needed toensure that the planned service lifetime may be achieved. Oxidation effects

Rel

ativ

e m

ass

chan

ge

7

6

5

4

3

2

1

0

10Cr–Mo–W

10Cr–Mo–W–Si

11Cr–Mo–Co

HCM12

12Cr–Mo–V

540 560 580 600 620 640 660Temperature (°C)

19.7 Relative mass changes of different ferritic steels duringexposure for 1000 h in Ar-50% H2O showing different types oftemperature dependence of the oxidation rates.

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may lead to changes in the component dimensions and other properties andthese changes can reduce the service life. Because oxidation is a surface-related effect and the strength properties are a bulk effect, whether a givenoxidation rate will have a significant effect on the mechanical behaviour willdepend on the initial component dimensions. Considering the materialconsumed by the oxidation process, the net section loss that can be toleratedfor a component with a large surface to volume ratio, for instance a heatexchanger tube, will be much smaller than that of a component with a smallsurface to volume ratio, such as a turbine rotor.

An indication of what is an acceptable steam oxidation rate for a givencomponent may be obtained by looking at the long-term experience with lowalloy steels in power plant at lower temperatures. The 1Cr1/2Mo steel hasbeen successfully used at temperatures up to 550°C for durations of severalhundred thousand hours. Figure 19.8 compares the mass changes for thissteel at 550°C with the results obtained for P91 and P92 at 600 and 650°C.It is clear that the steam oxidation behaviour of the 9%Cr steels, with two toeight times higher mass gains in 10 000 h at 600–650°C than the 1Cr1/2Mosteel at 550°C will present a potential problem for the application of thesteels in power plant. The implications of these high oxidation rates on thecomponent service lifetime will now be considered.

Mas

s ch

ang

e (m

g c

m–2

)

50

40

30

20

10

00 2000 4000 6000 8000 10000 12000

Test duration (h)

P92 650°C

P91 650°C

P92 600°C

P91 600°C

1Cr0.5Mo550°C

19.8 Mass change curves for P91 and P92 at 600 and 650°C, withcomparison values for 1Cr1/2Mo at 550°C.

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19.4 Oxidation and service life

19.4.1 Effects of oxidation on service life

For discussion of the steam oxidation interaction with the stress rupturebehaviour, we shall take a mass gain of 30 mg cm–2 in 10 000 h as thebenchmark (see Fig. 19.8) which is equivalent to a total scale thickness of0.2 mm, of which about half is internal scale. Therefore a wall thicknessreduction of 0.1 mm from each side in 10 000 h would be seen and ifwe extrapolate linearly, this would mean a reduction in wall thickness of 1.5mm from each side in 15 years. A linear extrapolation can be justified as beingconservative, because of the likelihood of scale spallation discussed above.

The possible effects of the oxidation process on the mechanical behaviourare:

(1) Reduction of load-bearing cross-section:The thick oxide scales formed at 600–650°C in steam on the 9–10%Crsteels lead to a reduction in the wall thickness, which will in turn leadto increased stress (the load remains constant but the cross-sectionalarea is reduced). The increase in stress and the corresponding reductionin rupture life will, however, depend on the initial wall thickness of thecomponent.

(2) Thermal insulation effect of thick oxide scales:The thick external and internal oxide scales formed on the 9–10%Crsteels in steam will have a considerable influence on heat transferacross the tube wall. In addition to reducing the efficiency of the steamgeneration, the decrease in the heat transfer caused by thermallyinsulating, thick oxide scales could lead to overheating of the tubematerial, unless, of course, the scales spall.

(3) Spalling of oxide scales:The spalling of the thick oxide scales would be beneficial in terms ofreducing the above-mentioned thermal insulation effect. However, thespalled oxide itself can lead to tube overheating if it becomes entrappedin the system, thus reducing the flow rates inside the tubes. Erosion ofturbine components may also occur. The defect nature of the thickoxide scales formed in steam will promote spalling at some stage andtherefore it seems prudent to design systems so that the risk of blockagescaused by spalled oxide flakes can be reduced as far as possible.

19.4.2 Quantitative estimation of the effect of steamoxidation on service life

For P92, the stress rupture curves for temperatures in the range 600–700°Care shown in Fig. 19.9. From these curves it can be seen that an increase in

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the stress from 100 to 120 MPa at 600°C results in a decrease in rupture lifefrom 170 000 to 75 000 h, representing a life reduction of about 40%. Figure19.9 also shows that for a given stress level, say 100 MPa, increasing thetemperature from 612 to 625°C reduces the rupture life from 80 000 to33000 h, which is equivalent to a 65% reduction in life or a three-fold increasein the secondary creep rate for a 10 K rise in temperature.

Using the steam oxidation rate of 0.1 mm steel thickness loss in 10 000 hmentioned above, the service life of components can be estimated andcompared. Figure 19.10 shows the rupture life reductions for a thick-walledpipe and a thin-walled superheater tube at 600°C. The thick-walled pipe cantolerate a thickness loss of 1.5 mm, which could occur after 15 years exposureat 600°C, with only a slight reduction in service life. The thin-walled tubewith the same steam oxidation rate would experience a considerable reductionin service life from over 200 000 h to 60 000 h.

Laboratory tests have shown that the total scale thickness on 9–10%Crsteels exposed for 10 000 h at 600–650°C will be around 0.2 mm, whichcould lead to a temperature increase of 50 K and a very large associatedreduction in rupture life. The thermal insulation effect of the thick oxidescales would then become the most damaging effect with respect to thecomponent service life. Some confirmation of this has been demonstrated infield trials; the results of tests in a by-pass loop of a power station aredescribed and the creep failure of an E911 pipe was attributed to overheating

1000 10000 100000 1000000Time to rupture (h)

Str

ess

(MP

a)300

250

200

150

100

50

0

Mean stress rupture strength of P92

Temperature (°C)

550

600612625637650

700

19.9 Mean stress rupture curves for P92.

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produced by the thick oxide scale that had formed (Zabelt and Wachter,1995). The temperature increase was estimated to have been from 630 to675°C, which correlated roughly with the observed time to creep failure ofthe tube.

19.5 Development of steam oxidation-resistant

steels

19.5.1 Steel composition

It has been demonstrated that the steam oxidation resistance of chromiumsteels can be improved by increasing the chromium content to 12%. Toensure a fully martensitic microstructure, the increased chromium contenthas to be balanced by addition of elements that stabilize the austenite phasewithout reducing the ferrite/austenite transformation temperature. Cobaltand copper have been the favoured additions and the newer 12% chromiumsteels usually contain one of these elements.

Although the steam oxidation resistance of these steels is much betterthan that of the 9% chromium steels and in short-term tests the stress rupturestrength is at least as good, the stress rupture strength falls rapidly and fallsbelow that of the 9% chromium steels after long test durations (Gabrel et al.,2006). This has been attributed to the formation of the Z phase, CrNbN,which in turn causes the dissolution of the fine niobium and vanadium nitrides

19.10 Rupture life reduction for two P92 components, a pipe of 300mm diameter and wall thickness 40 mm and a tube of 40 mmdiameter, wall thickness 6 mm.

Ru

ptu

re l

ife

(h)

1000000

100000

10000

40 mm diameter, 6 mm wall thickness tube

300 mm diameter, 40 mm wall thickness pipe

P92 steel life reduction due to material lossby oxidation 600°C, pressure 300 bar

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6Material loss from each side (mm)

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and carbonitrides that provide the long-term strength. It is therefore concludedthat, based on current knowledge, it is not possible to obtain the requiredhigh stress rupture strength and at the same time good steam oxidationresistance in a single steel composition. Attention has therefore been turnedto coatings.

19.5.2 Coatings

The status of coatings development has been summarized by Aguero (2006).The aim has been to enrich the surface layer of the steel with either aluminiumor chromium, so that alumina or chromia oxide scales form. Silicon has alsobeen considered, but as silica reacts with steam to form volatile siliconhydrides, such coatings are not suitable. There are a number of constraintsconcerning the coating techniques employed. The temperatures used for thecoating processes should allow the basic microstructure of the steel to befully martensitic with the fine precipitates of niobium and vanadium nitridesand carbonitrides that are essential for the long-term creep strength. Theprocesses investigated are principally diffusion coatings and promising resultshave been obtained, especially for aluminium enrichment of the surface. Theup-scaling of the coating processes for real components still requiresinvestigation.

19.5.3 Surface modifications

It is well known that modification of the surface of austenitic stainless steelsby shot peening improves the oxidation resistance by enhancing the diffusionof chromium into the scale. Investigations of various mechanical treatments,including shot peening, grinding and polishing, on the steam oxidationresistance of martensitic steels has not revealed any significant improvementin the steam oxidation behaviour. This may be due to the fact that the initialtempered martensite microstructure consisting of fine martensite laths andthe high dislocation density arising from the martensite transformation duringcooling already provide the best achievable conditions for diffusion ofchromium, which are however not sufficient for the formation of a protectiveoxide scale.

19.6 Outlook

From the point of view of the creep strength, the 9% chromium steels aresuitable for applications at metal temperatures up to around 620°C and furtherdevelopments are in progress to enable temperatures of 650°C to be achieved.However, the steam oxidation resistance remains a significant problem.Increasing the chromium content to 12% or more leads to excellent steam

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oxidation resistance but at the cost of reduced creep strength. At the presenttime and in the light of the available data, it does not appear possible toobtain both steam oxidation resistance and sufficient creep strength in asingle steel composition. Coatings offer a potential solution and aluminizinghas been successfully applied on a laboratory scale.

Another line of development is to move to a fully ferritic steel with achromium content of, say, 14% or more. Because of the extremely lowsolubility of carbon and nitrogen in ferrite, it will not be possible to producea dispersion of strengthening carbide and nitride precipitates by conventionalsteel processing techniques. Instead, intermetallic phases, such as the Lavesphases, could be considered for creep strengthening, provided that suchprecipitates are sufficiently fine and resistant to coarsening at the applicationtemperatures.

19.7 Sources of further information

Kofstad P, High Temperature Corrosion, Elsevier Applied Science, Londonand New York 1988.Schuetze M, Protective Oxides and Their Breakdown, Wiley Series onCorrosion and Protection, Holmes D R (series ed.), J Wiley & Sons, Chichester,UK, 1997.

19.8 References

Aguero A (2006). ‘Coatings for protection of high temperature new generation steampower plant components; a review’, in Materials for Advanced Power Engineering2006, Lecomte-Beckers J, Carton M, Schubert F and Ennis P J (eds), Energy TechnologyVolume 53, Part II, Forschungszentrum Juelich, Germany, 949–964.

Gabrel J, Bendick W, Vandenberghe B and Lefebvre B (2006). ‘Status of development ofthe VM12 steel for tubular applications in advanced power plants’, in Materials forAdvanced Power Engineering 2006, Lecomte-Beckers J, Carton M, Schubert F andEnnis P J (eds), Energy Technology Volume 53, Part II, Forschungszentrum Juelich,Germany, 1065–1076.

Quadakkers W J and Ennis P J (2002). ‘The oxidation behaviour of chromium steels insupercritical steam power plant’, in Materials for Advanced Power Engineering 2002,Lecomte Beckers J, Carton M, Schubert F and Ennis P J, (eds), Energy TechnologySeries, Volume 21, Part II, Forschungszentrum Jülich, Germany, 1131–1142.

Quadakkers W J, Ennis P J, Zurek J and Michalik M (2005). ‘Steam oxidation of ferriticsteels: laboratory test kinetic data’, Materials at High Temperatures, 2005, 22 (1/2),37–47.

Rahmel J and Tobolski J (1965). ‘Einfluss von Wasserdampf und Kohlendioxid auf dieOxidation von Eisen in Sauerstoff bei hohen Temperaturen’, Corrosion Science, 1965,5, 333.

Zabelt K and Wachter O (1995). Ergebnisse von Feldversuchen in Kraftwerken mit 9 bis12-%-Chromstählen, 18 meeting of the Arbeitsgemeinschaft für warmfeste Stähle,

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Presentation 6. 1, Verein Deutscher Eisenhuettenleute (VDEh), Düsseldorf December1995.

Zurek, J (2004). Oxidation and Oxidation Protection of Ferrritic and Austenitic Steels inSimulated Steam Environments at Temperatures Between 550 and 650°C, DoctoralThesis, Rheinisch-Westfälische Hochschule, Aachen, September 2004.

Zurek J, Wessel E, Niewolak L, Schmitz F, Kern T U, Singheiser L and Quadakkers W J(2004). ‘Anomalous temperature dependence of oxidation kinetics during steam oxidationof ferritic steels in the temperature range 550 – 650°C’, Corrosion Science, 2004, 46/9, 2301–2317.

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Part III

Applications

537

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20Alloy design philosophy of creep-resistant

steels

M . I G A R A S H I, Sumitomo Metal Industries, Japan

20.1 Introduction

High Cr ferritic steels such as ASME P91 steel have successfully been usedfor large diameter and thick section boiler components such as main steampipes and headers in supercritical pressure boilers in fossil-fired power plants.1

Recent trends towards the utilization of clean energy leading to protection ofthe global environment have been accelerating the application of ultrasupercritical pressure (USC) boilers, which are operated with higher efficiencyin power generation than conventional boiler and thus release less carbondioxide.2,3 The USC boiler requires new steels with improved creep rupturestrength and steam oxidation resistance at elevated temperatures over 600°C,because of the increase in the operating temperature and pressure of the steam.Addition of W to the ferritic steel has been found to be effective in increasingcreep rupture strength at high temperatures and has been already used in somedeveloped steels such as T92/P92 and T122/P122 for the USC boilers.4

High strength austenitic stainless steels such as TP347HFG, SUPER304Hand HR3C have been used extensively as superheater and reheater tubes forthe latest USC boilers all over the world.5–8 In this chapter, the alloy designphilosophies for creep-resistant ferritic and austenitic stainless steels forvarious components in USC power plants are reviewed and demonstrated indetail and the future research will be discussed with regard to furtherimprovement of the steels for application in 700°C A-USC plants.9

20.2 Creep-resistant steels for particular

components in power plants and the

properties required

20.2.1 Water wall

Figure 20.1 shows a schematic illustration and photographs of various boilercomponents such as the water wall, superheater, reheater, header and main

539

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steam pipe and their typical materials used in a recent fossil fired powerplant.10 The water wall tubes are welded in a panel structure to achieveeffective heat exchange in order to produce pressurized steam from waterunder the super critical conditions in a furnace. Conventional steels andCrMo steels are used for the water wall tubes according to the operatingtemperatures and pressures, which are usually heated up to 500°C in thelatest USC plants.11 The materials requirements for the water wall are, therefore,strength and corrosion resistance at elevated temperatures as well as weldabilityand formability to construct the water wall panel. Recently, steels with higherstrength and good weldability have been developed and successfully used inthe latest USC plants described below.12,13

20.2.2 Superheater/reheater

The steams are superheated in a superheater (SH) up to the highest pressureat a designed temperature and reheated in reheater (RH) up to the highesttemperature that will achieve the designed thermal efficiency, for instance,42% at 600°C and 25 MPa with a SH and 605°C and 4.2 MPa with a RH inone of the latest USC plants.11 The materials requirements for SH and RH are,therefore, most severe for the components and high strength austenitic stainlesssteels are used for this purpose. They are required to have high creep strength

20.1 Schematic illustration and photographs of a fossil fired boilerand typical materials.

Header

Superheater

18 ~ 25%Craustenitic steels

9 ~ 12% Crferritic steels Main steam pipe

9 ~ 12%Crferritic steels

Furnace

Reheater

Main steam pipe

Water wall tubes

0.5 ~ 2% Cr steelsSeawater

Header

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Alloy design philosophy of creep-resistant steels 541

at a designed temperature and above and to have both steam oxidation resistancefor the inner surface and hot-corrosion resistance for the outer surface againstcoal ash containing sulfur, chlorine, vanadium and other corrosive salt-formingelements. Various types of austenitic stainless steels have been developed forSH/RH tubing application, which are described in detail later.

20.2.3 Header and main steam pipe

The superheated steams are gathered into a header pipe and transferred to aturbine system through main steam piping. The header and main steam pipingshould, therefore, be a heavy wall thickness pipe with a large diameter to keephuge amounts of steam pressurized. They are required to have high creepstrength with good ductility and toughness to prevent catastrophic failureduring operation and hydro-pressure testing. They are also required to havegood thermal fatigue resistance to the thermal stress imposed, because fossilfired plants are usually used to adjust the electricity supply during the day,which means they are used in a daily start and stop operation. For this reason,high Cr ferritic steels with high creep strength have been developed and havealready been used in the latest USC plants, because although ferritic steels arein principle inferior in strength to the austenitic stainless steels because theyhave a large diffusion constant at elevated temperatures caused by the differencein their crystal structures, they are superior in thermal expansion and conductivity.High Cr ferritic steels have been developed and successfully used in the latestUSC plants and will be described in detail later.

20.3 Alloy design philosophies of creep-resistant

steels

20.3.1 High strength low-Cr steels

Figure 20.2 shows a tree chart of developed ferritic steels.10 Strengtheningof ferritic steels is mainly achieved by dislocation strengthening using C andN for martensitic and bainitic transformation and the resultant microstructure,solid solution hardening by elements such as Mo and W, and precipitationhardening by carbo-nitrides containing Cr, V and Nb and Cu phase. Corrosionresistance of the steels is mainly achieved using Cr and Si. These ferriticsteels are used for various components in USC plants according to theirrespective strength and corrosion resistance.

Figure 20.3 shows the alloy design philosophy of 2.25Cr–1.6W–V–Nbsteel (HCM2S; KA-STBA24J1, T23/P23, ASME CC2199).14,15 2.25Cr–1.6W–V–Nb steel is used as a water wall, in SH and RH tubes, and in the headerand main steam pipe in fossil fired boilers and heat recovery boilers. Thesteel has been developed to improve the creep rupture strength of 2.25Cr–1Mo steel at elevated temperatures mainly by substituting Mo by W. High

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542C-steel ~ CrMo steels (2.25Cr) 5 ~ 9 Cr steels 12 Cr steels

11Cr–0.4Mo–2W–Cu–V–Nb(KA–SUS410J3TB, T122,

ASME CC2180)

+ W, Cu, B

12Cr–1Mo–1W–V–Nb(KA–SUS410J2TB)

–C, + W, Nb

X20CrMoV121

+ C, V

+ Cr

9Cr–1Mo–V–Nb(KA–STBA28, T91)

+ V, Nb

– C + Mo

9Cr–2Mo(KA–STAB27)

STBA 26STBA 25

KA-indicates that the steel isdesignated by the METI standard

Corrosion resistance

2.25Cr–1.6W–V–Nb(KA–STBA24J1, T23,

ASME CC2199)

+ W, V, Nb, B

+ Cr

STBA 24STBA 23STBA 22

KA-STBA21STBA20

KA-STBA10(CR1A)

+ Cr, Mo

+ Mo

STBA 13STBA 12

+Cr, Ni, Cu

+ Mn

+C

STB510

KA-STB480STB410STB340

Fe

Str

eng

th

20.2 Tree chart of developed ferritic boiler steels.

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strength is mainly achieved by the combination of a solid solution of W with(V, Nb)C dispersion hardening in a fully tempered bainitic matrix. Additionof B enhances the bainitic microstructure and is found to improve the toughnessof the steel to a great extent. Low C content has been chosen to improve theweldability of the steel and as a result no preheating and post weld heattreatment (PWHT) is potentially required for some applications. Figure 20.4shows the weldability and the crack sensitivity of 2.25Cr–1.6W–V–Nb steel.It is seen that 2.25Cr–1.6W–V–Nb steel exhibits superior weldability evenwithout preheating, while both T22 and T91 grade steels require preheatingabove 200°C and 150°C, respectively.

Figure 20.5 shows creep rupture data for 2.25Cr–1.6W–V–Nb steel pipes.16

It can be seen that the creep strength of this steel is estimated to be about 1.8times higher than those of T22/P22 steel. The longest creep rupture time of2.25Cr–1.6W–V–Nb steel pipes is about 90 000 h at 550°C. Their long-termcreep strengths are very stable at temperatures between 500 and 600°C. It is,however, noted that above 600°C the longer term rupture time tends todecline over the averaged curve, which has been found to show the effect ofoxidation when a small specimen was used and therefore no significantstrength degradation took place, like that observed for T22 grade steels.16

Figure 20.6 shows a calculated phase diagram for 2.25Cr–1.6W–V–Nbsteel at 600°C with changing Cr and C contents.17 The equilibrium phasediagram suggests that the final microstructure of 2.25Cr–1.6W–V–Nb steelconsists of ferrite (α) + MX ((V,Nb)C) + M6C and may include a smallamount of M23C6.

Figure 20.7 shows TEM micrographs of extraction replicas of 2.25Cr–1.6W–V–Nb steel (a) normalized and tempered and (b) crept for 12 567.6 hat 600°C. In the tempered specimen, M23C6 is formed along prior austenitic

Weldability Toughness Creep strength

max.Hv350

Low C-0.06 mass%

No preheating andPWHT after welding

Matching weldingfiller used without

PWHT

Hardenability :B addition

Fully temperedbainitic structure

0.06C-2.25Cr-1.6W-0.1Mo-0.25V-0.05Nb-B

Solutionstrengthening :

High W

Precipitationstrengthening :

V, Nb, B

20.3 Alloying philosophy of 2.25Cr–1.6W–V–Nb steel (HCM2S; KA-STBA24J1, T23/P23, ASME CC2199).

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grain boundaries and MX is formed along lath boundaries and in the bainiticmatrix. M7C3 is occasionally observed along lath boundaries. In the creptspecimen, on the other hand, no M23C6 and M7C3 are observed and insteadblocky M6C is formed along the prior austenitic grain boundaries and fineMX remains along lath boundaries and in the bainitic matrix as shown inFigure 20.7(b). It is noted that the pronounced lath structure is kept even

Cracking ratio C

C = h

H 100 (%)× T91

T22

hH

2. 25Cr-1.6W-V-Nb

0 50 100 150 200 250 300 350Preheating temperture (°C)

Cra

ckin

g r

atio

(%

)

100

90

80

70

60

50

40

30

20

10

0

20.4 Weldability of 2.25Cr–1.6W–V–Nb steel (HCM2S; KA-STBA24J1,T23/P23, ASME CC2199).

Str

ess

(MP

a)

500

400

300

200

10090807060

50

40101 102 103 104 105

Rupture time (h)

20.5 Creep rupture data for 2.25Cr–1.6W–V–Nb steel pipes (HCM2S;KA-STPA24J1, P23, ASME CC2199).

500°C

550°C

600°C

650°C

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C-c

on

ten

t (m

ass%

)

0.20

0.18

0.16

0.14

0.12

0.10

0.08

0.06

0.04

0.02

00 1 2 3 4 5

Cr-content (mass%)

α + M23C6 + M6C + MX

α + M X + M6C

20.6 Calculated phase diagram for 2.25Cr–1.6W–V–Nb steel at 600°C.

20.7 TEM micrographs of MX in 2.25Cr–1.6W–V–Nb steel(a) normalized and tempered, (b) crept for 12 567.6 h at 600°C.

1 µm 1 µm

M23C6 MX

MX M7C3

(b)

M6CMX

M6C

1 µm 1 µm

(a)

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after long-term creep deformation. Formation of M6C means that the steelloses W and/or Mo in solution, which is the key element for solid solutionstrengthening. However, W is superior to Mo in retarding the formation ofM6C during long-term creep deformation. Figure 20.8 shows the change inprecipitated W in 2.25Cr–1.6W–0.2Mo–V–Nb steel and Mo in 2.25Cr–1.0Mo–V–Nb steel during aging for up to 10 000 h at temperatures between 550°Cand 650°C.18 The lines indicate fitted curves, assuming that the growth rateis controlled by the Johnson–Mehrl–Abrami theory. It is seen that in 2.25Cr–1.6W–0.2Mo–V–Nb steel the growth rate of M6C is 10 to 100 times slowerthan that in 2.25Cr–1.0Mo–V–Nb steel. This is one of the major reasons forthe high creep strength of 2.25Cr–1.6W–V–Nb steel.

20.3.2 Martensitic high-Cr steels for heavy-wall thicknesspiping

11Cr–0.4Mo–2W–Cu–V–Nb steel (HCM12A; KA-SUS410J3TB/TP (T122/P122, ASME CC2180) and KA-SUS410J3DTB) has been developed to improvethe creep rupture strength and corrosion resistance of P91 type 9%Cr steelsabove 600°C, mainly achieved by higher Cr content and substitution of partof the Mo by W.19,20 In order to suppress δ-ferrite formation for thick wallpipes, Cu addition is chosen from the γ-forming elements shown in Fig.20.9. Cu, unlike Ni and Mn, is a γ-forming element which does not reducethe Ac1 temperature much and does not enhance coarsening of M23C6 carbide.Cu addition enables the combination of higher Cr content with high W andMo contents to be achieved.

Figure 20.10 gives a summary of the alloy design philosophy of 11Cr–0.4Mo–2W–Cu–V–Nb steel (HCM12A; KA-SUS410J3TB/TP (T122/P122,ASME CC2180) and KA-SUS410J3DTB).19 This steel was originally developed

Frac

tio

n o

f W

-pre

cip

itat

ed (

–)

1

0.8

0.6

0.4

0.2

0

Experiment value

650°C 600°C

550°C

1 10 100 1000 104 105

Aging time (h)(a)

Frac

tio

n o

f M

o-p

reci

pit

ated

(–)

1

0.8

0.6

0.4

0.2

0

650°C600°C

550°C

Experiment value

1 10 100 1000 104 105

Aging time (h)(b)

20.8 Changes in (a) precipitated W in 2.25Cr–1.6W–0.2Mo–V–Nb steeland (b) Mo in 2.25Cr–1.0Mo–V–N steel during aging for up to 10 000 hat temperatures between 550°C and 650°C.

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for two different purposes: one for large diameter, heavy wall thickness pipeand the other for a tubing application which often requires good corrosionresistance.19 The steel, therefore, has two different Cr versions: a steel withCr content lower than 11.5% having good toughness and one with Cr contenthigher than 11.5% having good corrosion resistance. The steel with lower Crcontent has an α′ single martensitic matrix microstructure, which is strengthenedby the combination of MX carbonitrides with M23C6 along the prior austeniticgrain boundaries. The steel with a higher Cr content has a δ-ferrite and α′martensite dual matrix microstructure. This difference in the microstructureimposes variation in the creep strength shown in Fig. 20.11.21 Figure 20.11shows creep rupture data for 11Cr–0.4Mo–2W–Cu–V–Nb steels with twodifferent Cr content levels, the α′ steels with Cr less than 11.5% and the theα′ + δ steels with Cr equal to and more than 11.5%. It is obvious from theseupdated rupture data that the creep rupture strength should be analyzedseparately for this steel according to the Cr content levels.

Figure 20.12 shows the stress dependence of the minimum creep rate andthe off-set strain on the acceleration creep of 11Cr–0.4Mo–2W–Cu–V–Nbsteels with two different Cr content levels, the α′ and α′ + δ steels shown inFig. 20.11.22 It can be seen that the minimum creep rate of both steelsdecreases with decreasing applied stress. The α′ steel exhibits a smoothchange in the minimum creep rate with stress. It is found that in the α′ steelthe power-law creep is relevant at a higher stress regime above 60 MPawhere the stress exponents are high enough at between 5 and 15, while theviscous creep is at lower stress where the stress exponent is as low as 1. Thisbehavior is very similar to that observed in P91 steel by Kloc and Sklenicka.23,24

The α′ + δ steel, on the other hand, exhibits a deviation from the expected

∆Ac 1

(°C

)+40

0

–40

–4 –2 0 +2 +4∆cr eq. (mass%)

γ-forming element

Creq = Cr+6Si+4Mo+1.5W+ 11V+5Nb+8Ti+12Al– 40C-30N-4Ni-2Mn– Cu-2Co (mass %)

0.05N

0.5Ni

2Cu

1Cr

2W

1Mo

Base; 0.1C–11Cr steelAddition (mass%)

α-forming element

20.9 Comparison of alloying elements with respect to changes in Ac1temperature and Creq for 0.1C–11Cr model steels.

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548

KA-SUS410J3J3TB/TP

KA-SUS410J3DTB

0.1C-10.5/12Cr-2W-0.4Mo-Cu-0.2V-0.05Nb-B-N

Corrosionresistance

Creepstrength

Weldability

Toughness

High Cr 12%

Superior to 9%Cr

VN, stable & effectiveNb(C, N), grain refinement

W, stable

Stable long-term strength

δ-ferrite5%(single phase)

δ-ferrite30%(dual phase)

10.0~12.5%Cr

V, Nb and Naddition

High W & Mo

B addition

Low Ni

Low C

Creq 9%

Cu addition

High temperaturetempering 770°C

20.10 Alloying philosophy of 11Cr–0.4Mo–2W–Cu–V–Nb steel (HCM12A; KA-SUS410J3TB/TP (T122/P122, ASME CC2180)and KA-SUS410J3DTB).

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smooth change in the minimum creep rate around the critical stress, wherethe rupture life of the α′ + δ steel starts to deviate from the straight line. Atlower stress a sudden drop in the minimum creep rate has been observed,indicating that the creep deformation mechanism has changed explicitlyfrom the high stress regime and viscous creep becomes dominant in thelower stress regime. Figure 20.12(b) shows that the change in the off-setstrain of the α′ + δ steel with stress is completely different from that of theα′ steel, in particular, at lower stress where a sudden drop in the minimumcreep rate has been observed. This suggests that at a particular stress level inthe α′ + δ steel the heterogeneous creep deformation takes place as aconsequence of the prohibition of ordinary homogeneous deformation andthe enhancement of localized deformation which gives rise to accelerationcreep even at a very small strain and a very small creep rate.

Figure 20.13 shows updated creep rupture data for 11Cr–0.4Mo–2W–Cu–V–Nb steel pipes with Cr < 11.5% (KA–SUS410J3TP, equivalent toP122, ASME CC2180).21 Three different fitted curves are based on the LMPmethod with original data and updated data, and the one determined by theregion splitting method of Kimura25 adopted by the SHC (SHC is the committeefor establishing allowable stress value, and soon, of high Cr ferritic steels forthe ministry of economy, Trade and Industry, set up in Japan since 2004).Using the region splitting method, the allowable tensile stresses obtained aremuch lower than with those originally proposed, as given in Table 20.1.

Open symbols; Cr<11.5%Solid symbols; Cr≥11.5% 700°C

550°C

600°C

650°C

101 102 103 104 105

Rupture time (h)

Str

ess

(MP

a)500

400

300

200

10090807060

50

40

30

20.11 Creep rupture data for 11Cr–0.4Mo–2W–Cu–V–Nb steels withtwo different Cr content levels, the α’ steels with Cr <11.5% and theα’ + δ steels with Cr ≥ 11.5%.

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α’ + δ

α’

5

1

1

20 30 40 50 60 70 80 90 100 200Stress (MPa)

(a)

Cre

ep r

ate

(h–1

)

10–3

10–4

10–5

10–6

10–7

α’ + δ

α’

20 30 40 50 60 70 80 90 100 200Stress (MPa)

(a)

Cre

ep r

ate

(h–1

)

10–1

10–2

10–3

20.12 Stress dependence of (a) the minimum creep rate and (b) theon-set strain on the acceleration creep of 11Cr–0.4Mo–2W–Cu–V–Nbsteels with two different Cr content levels, the α’ and α’ + δ steels.

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To achieve longer term creep strength in advanced ferritic steels at elevatedtemperatures over 600°C, the creep deformation mechanism of the steels andthe role of each precipitate in the respective creep deformation process havebeen examined using model steels with different initial microstructuresconsisting of an α′ single martensite matrix with M23C6 and with M23C6 andMX. Figures 20.14 and 20.15 are TEM micrographs showing themicrostructural evolution during creep deformation of the α′ + M23C6 andthe α′ + M23C6 + MX structures, crept and interrupted at 650°C with stressesof 70 and 100 MPa, respectively. It is seen that, in the α′ + M23C6 steelshown in Fig. 20.14, the excess dislocations are much reduced inside lathgrains in the transient creep region. Around the minimum creep rate themigrations of lath and block boundaries seem to start forming equi-axed

LMP method with original dataLMP method with updated dataRegion splitting method by SHC

Open symbols; original dataSolid symbols; updated data

700°C

550°C

600°C

650°C

101 102 103 104 105

Rupture time (h)

Str

ess

(MP

a)500

400

300

200

10090807060

50

40

30

20.13 Creep rupture data for 11Cr–0.4Mo–2W–Cu–V–Nb steel pipeswith Cr<11.5% (KA-SUS410J3TP, equivalent to P122, ASME CC2180).

Table 20.1 Allowable tensile stress values for 11Cr–0.4Mo–2W–Cu–V–Nb steelrevised according to the region splitting method adopted by the SHC in 2005.There are two steels designated according to different Cr content in thecorresponding METI specification

Temperature, (°C) 575 600 625 650

ASME CC2180-2(P122)* 107 83 61 (45)KA-SUS410J3TP pipe with 10.5/11.5 Cr 100 68 46 27KA-SUS410J3DTB tube with 11.5/12.5 Cr 94 52 25 16

* Allowable stress of ASME P122/T122 has been reassessed by an ASME subcommitteeand will soon be revised in Japan.

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20.14 TEM micrographs showing microstructural evolution duringcreep deformation of α’ + M23C6 steel, crept and interrupted at 650°Cwith 70 MPa. (a) As normalized and tempered (NT); (b) after 50 h,creep strain ε = 0.6%; (c) after 200 h, ε = 1.5%; (d) after 872 h, ε = 9.5%.

20.15 TEM micrographs showing microstructural evolution duringcreep deformation of the α’+M23C6+MX steel, crept and interrupted at650°C with 100 MPa. (a) As NT; (b) after 216 h, ε = 0.3%; (c) after1620 h, ε = 1.7% (d) after 5164 h, ε = 5.0%.

1µm

(a) As NT (b) After 50h, ε = 0.6%

(c) After 200h, ε = 1.35% (d) After 872h, ε = 9.5%

(a) As NT (b) After 216h, ε = 0.3%

(c) After 1620h, ε = 1.7% (d) After 516h, ε = 5.0%

1µm

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subgrains. In the acceleration creep region, subgrain formation propagates tothe whole specimen. In the α′ + M23C6 + MX steel shown in Fig. 20.15, themicrostructural change during the creep deformation is very similar to thatin the α′ + M23C6 steel but differs in the time taken to reach the samedeformation microstructure. This means that MX carbonitrides inside lathgrains serve as an effective obstacle against dislocation in motion and hencedelay the rearrangement of dislocations and formation of subgrains. It is,however, noted that this seems to enhance the heterogeneous creep deformationalong boundaries like the prior austenite grain boundary and packet, blockand lath boundaries, since the greatly strengthened matrix by MX does notdeform easily at low stress levels, although the regions around the boundarieswhich are softer than the matrix could deform much more easily.

Figure 20.16 shows a schematic representation of creep rate versus timecurves of the ferritic steels and the corresponding guiding principles for achievinglong-term creep strength, which is derived from microstructural evolution duringcreep deformation.26 To achieve a higher creep strength the creep rate in thetransient creep region can be reduced using fine dispersion of MX, α″, Cu-phaseand also M23C6 and the Laves phase. All the fine precipitates serve as obstaclesto dislocation in motion and hence delay the rearrangement of dislocations andformation of subgrains. In the acceleration creep region, however, subgrainformation proceeds to a great extent and homogeneous deformation becomesmore difficult. In such a deformation scheme, heterogeneous creep deformation

Cre

ep r

ate

Time

Reducingdisl. mobility

(1) Reducing mobility oflath-, block-boundaries

(2) Suppressing hetero-geneous deformation

Stabilization of martensite/bainite (C-free, Co)

Fine dispersion;MX, α”, Cu, M23C6, Laves

Solid solution (Mo, W etc)

Dispersion and stabilization ofboundary precipitates; M23C6,Laves (optimization of C, N, B)

Something else

20.16 Schematic illustration of creep rate versus time curvesrepresenting creep strengthening mechanisms for the ferritic steels.

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along the prior austenite grain boundary becomes significant, sometimes resultingin creep rupture in a short period. To increase the creep resistance in theacceleration creep region, stabilization of lath, block and packet boundaries ofmartensite and of the prior austenite grain boundary is useful. This can be doneby stabilization of the precipitates such as M23C6 and the Laves phase alongthese boundaries.

20.3.3 Austenitic stainless steels for superheater/reheatertubing

Figure 20.17 shows creep curves of Type304H austenitic stainless steel creptat 650°C with stresses of 147, 118 and 98 MPa, compared with those of a

20.17 Creep curves for Type304H austenitic steel crept at 650°C withstresses of 147 118 and 98 MPa. (a) Creep rate against time,(b) creep rate against creep strain, (c) minimum creep rate againststress.

Str

ain

0.8

0.7

0.6

0.5

0.4

0.3

0.2

0.1

0.0

147MPa

118MPa

98MPaα′+M23C6+MX

98MPa

0 2000 4000 6000 8000 1000 12000Time (h)

(a)

Cre

ep r

ate

(1/h

)10–1

10–2

10–3

10–4

10–5

10–6

10–7

147MPa

118MPa

98MPa

α′+M23C6+MX98MPa

0 0.05 0.1 0.15 0.2Strain

(b)

70 80 90 100 120 150Stress (MPa)

(c)

Min

imu

m c

reep

rat

e (h

–1)

10–3

10–4

10–5

10–6

10–7

304H13

13

4

0.12C–0.002N

0.08C–0.05N

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9Cr ferritic steel.10 It is seen, unlike for the ferritic steels, that in Type304Hsteel the creep rate decreases gradually in the transient creep region buthardly increases in the acceleration creep region, giving a large off-set strainat the minimum creep rate. This means that homogeneous creep deformationtakes place in Type304H steel.

Figure 20.18 are TEM micrographs showing microstructural evolutionduring creep deformation of Type304H steel crept and interrupted at 650°Cat 147 MPa.27 It is obvious that in the austenitic steel the initial dislocationsdensity is in principle negligible and substantial numbers of dislocation areintroduced homogeneously all over the grains during creep deformation inthe transient creep region so as to suppress localized hardening. Once thedislocation starts to rearrange into a subgrain structure in some portions, thetransition to acceleration creep takes place, as observed in the ferritic steels.This subgrain structure propagates to the other portions in the accelerationcreep region. Note that the dislocation density does not change very mucheven at the end of the acceleration creep. This gives rise to a homogeneousdistribution of dislocations and makes it difficult to achieve the deformationmode easily at a small strain as in the ferritic steels. This is confirmed by thespecimens crept at low stresses. At lower stress levels, M23C6 forms alongthe dislocations inside the grain and serves as an effective obstacle to dislocationin motion and hence delays the rearrangement of dislocations and formationof subgrains as shown in Fig. 20.19.27

1µm

(a) (b)

(c) (d) (e)

20.18 TEM micrographs of Type304H steel crept and interrupted at650°C and 147 MPa. (a) As ST; (b) after 50 h, ε = 2.6%; (c) after 100 h,ε = 4.8%; (d) after 246 h, ε = 13.6%; (e) after 429 h, ε = 20.0%.

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Figure 20.20 shows the creep deformation mechanism and guiding principlesfor further strengthening of the austenitic alloys proposed.10 According tothe process of microstructural evolution during creep deformation describedabove, all the fine precipitates such as MX, NbCrN, M23C6 , Cu, α-Cr, γ ′ andthe Laves phase can serve as obstacles to dislocation in motion in the transientcreep region and hence delay the rearrangement of dislocations and formationof subgrains. In the acceleration creep region, subgrain formation propagateseverywhere until an equilibrium subgrain microstructure has been achievedfor an easy deformation mode at a given applied stress. In such a deformationscheme, suppressing the migration of the subgrain boundaries and stabilizationof the grain boundary are necessary to increase the creep resistance of thealloys. This can in principle be done by stabilization of the fine particlessuch as M23C6 and the Laves phase inside grains and along boundaries.Solution hardening by elements like Mo, W and N is an essential fundamentalstrengthening mechanism to be used for the austenitic alloys, while phasestability against σ embrittlement is required for long-term creep ductilityand the resultant creep strength.28 These strengthening methods havesuccessfully been used in the developed steels such as SUPER304H withM23C6, Cu and NbN and HR3C with M23C6, MX and NbCrN as shown in thetree chart of the developed austenitic stainless steels in Fig. 20.21.

TP347HFG (fine-grained 18Cr–12Ni–Nb steel) is widely used as superheaterand reheater tubes in fossil fired boilers.28 The steel has been developed toimprove the steam oxidation resistance of conventional TP347H stainless

200nm

20.19 TEM micrograph of Type 304H steel crept at 650°C and 98 MPaand ruptured after 10558 h.

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steel by grain refinement through a specially established thermomechanicalprocess. The microstructure of the steel consists of an austenitic fine-grainedmatrix strengthened by M23C6 carbide mainly along the grain boundary andfinely dispersed NbC carbide in the matrix. NbC is fine and stable even afterlong-term creep exposure at high temperatures.

Figure 20.22 shows creep rupture data for TP347HFG tubes with averagecurves by the Larson–Miller parameter method.29 The longest creep rupturetime for TP347HFG tubes is about 60 000 h at 600°C. Their long term creepstrength is very stable and no degradation in creep strength is expected at thetemperatures up to 750°C.

Figure 20.23 shows the manufacturing process for establishing fine-grainedmicrostructures in 18Cr–9/12Ni–Nb steels, which is characterized by doublestage heat treatment in order to achieve a fine grain microstructure with avery fine dispersion of NbC in the matrix, compared with that for theconventional TP347H steel.30 In the developed process, NbC resolves intothe matrix during the presolution treatment at higher temperatures andreprecipitates finely in the matrix during the subsequent solution treatmentat lower temperatures. This gives rise to a fine grain microstructure with avery fine dispersion of NbC in matrix.

Figure 20.24 shows the initial microstructures of TP347HFG tubes andconventional TP347H steel. A homogeneous and very fine grain structurehas been achieved by using a double stage heat treatment in the developedprocess.30 Figure 20.25 shows the change in the amount of Nb precipitatedas NbC in TP347HFG after presolution treatment at various temperatures.30

NbC precipitated more when using the lower temperature solution treatmentafter a higher temperature presolution treatment.

Cre

ep r

ate Suppressing

(1) rearrangement of dislocations(2) migration of sub GBs(3) GB embrittlement

Hardening andintroduction of

dislocations

Solution hardening(H,Mo,W etc)Fine particles such asMX, NbCrN, Cu,M23C6, α-Cr, γ’, Laves

Solution hardening(Mo, W etc)Phase stability (N, Ni etc)Stabilization of boundaryprecipitates such asM23C6, LavesLong-term stabilizatin offine particles

20.20 Guiding principles for further strengthening the austeniticalloys.

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558

High-Cr austeniticsteels (22 ~ 25Cr)

18-8 type austeniticsteels (18Cr)

+Cu,Nb18Cr–9Ni–3Cu–Nb–N; SUPER304H

(KA-SUS304J1HTB, ASME CC2328)25Cr–20Ni–Nb–N; HR3C

(KA-SUS310J1TB, ASMECC2111, TP310HCBN)

Str

eng

th

Fine Grained 18Cr–12Ni–Nb; TP347HFG

KA-SUSTP347HTB

SUS321HTB SUS347HTB (Fine grain)

+ Ti + Nb

16Cr–12Ni–Mo(SUS316HT, AISI316H)

18Cr-8Ni(SUS304HTB, AISI304H)

18Cr-8Ni(AISI302)

+Cr, Ni

Corrosion resistance

22Cr-12Ni(SUS309TB, AISI309)

25Cr-20Ni(SUS310TB, AISI310)

+Cr, Ni

–C, + Mo, N

(KA-SUS309J2TB)

+ Nb, N

20.21 Tree chart of the developed austenitic steels.

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TP347HTP347HFG

SUPER304H

Developed process Conventional process

Fine NbC

precipitation

So

luti

on

Pre

-so

luti

on

Coldworking

CW 30% CW 20~30%

Fine NbC

dissolution

Fine grain& fine NbC

G.S. No.8

Coarse grain& coarse NbC

G.S.No.; 4~5

20.23 Manufacturing process to establish fine-grainedmicrostructures in 18Cr–9/12Ni–Nb steels compared with the processfor conventional TP347H steel.

Test Temp.

600°C650°C700°C750°C800°C

600°C650°C

101 102 103 104 105

Rupture time (h)

Str

ess

(MP

a)500400

300

200

1009080706050

40

30

20

700°C

750°C

800°C

20.22 Creep rupture data for TP347HFG (fine-grained 18Cr–12Ni–Nbsteel) tubes.

Figure 20.26 shows the effect of final-solution treatment temperature onthe grain size of TP347HFG tubes and conventional TP347H steel.30 Thegrain size of TP347HFG tubes is much smaller than that of the conventionalsteel even with an increase in the solution treatment temperature. This wasachieved by more precipitation of fine NbC carbide during the presolutiontreatment at lower temperatures.

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Figure 20.27 shows the oxidation behavior of TP347HFG with differentgrain size tested at 650°C and 700°C. Steam oxidation resistance improveswith reducing grain size of the steels. This steam oxidation resistance isachieved by the thin and tight protective Cr2O3 corundum-type inner oxidelayer formed in the fine-grained steel (see Figure 20.28)30.

SUPER304H (18Cr–9Ni–3Cu–Nb–N steel; KA-SUS304J1HTB, ASMECC2328) is used as superheater and reheater tubes in fossil fired boilers.31

The steel has been developed to substitute for conventional Type304H andType321H steels by the addition of copper and nitrogen to increase creepstrength at elevated temperatures and toughness after long-term exposure athigh temperatures. The microstructure of the steel consists of an austeniticmatrix strengthened by M23C6 carbide mainly along the grain boundary anda finely dispersed Cu-phase and NbCrN nitride in matrix. The Cu-phase is

100µ

(a) (b)

20.24 Optical micrographs of fine grain microstructure for(a) TP347HFG attained using a double stage heat treatmentcompared with (b) treatment for TP347H.

Ext

ract

ed N

b c

on

ten

t (%

)

1.0

0.8

0.6

0.4

0.2

0

Pre-solution treatment1250°Cx 10min

1300°Cx 10min

Precipitated Nb as NbCduring final-solution treatment

As pre-solution

1100 1150 1200 1250Solution treatment temperature (°C)

20.25 Change in the amount of Nb precipitated as NbC in TP347HFGafter presolution treatment at various temperatures.

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TP347HFGConventional TP347H

1100 1200 1300Solution treatment temperature (°C)

Gra

in s

ize

(AS

TM

No

.)

12

10

8

6

4

2

0

TP347H GS6)TP347HFG(GS8)TP347HFG(GS9)

700°C650°C

Inn

er-s

cale

th

ickn

ess

(µm

) 100

50

20

10 In steam

500 1000 3000Time (h)

20.26 Effect of solution treatment temperature on grain size ofTP347HFG and conventional TP347H steel.

Cr2O3(Fe, Cr)3O4

Fe3O4

(a)

Outer scale Inner scale

Metal

Fe3O4

Cr2O3

Outer scale Inner scale

Metal

(Fe, Cr)3O4

(b)

20.27 Steam oxidation resistance of TP347HFG tested at 650°C and700°C.

20.28 Schematic illustration of oxide layers formed in (a) fine grainedand (b) coarse grained 18Cr austenitic steels.

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fine and coherent is the matrix giving rise to a significant increase in creepstrength at elevated temperatures. NbCrN is also fine and stable even afterlong-term exposure at high temperatures. No significant sigma phase is expectedto form even after 10 000 h in the temperature range between 600 and800°C, mainly achieved by stabilization of the austenitic matrix with copperand nitrogen.

Figure 20.29 shows creep rupture data for SUPER304H (18Cr–9Ni–3Cu–Nb–N steel; KA-SUS304J1HTB, ASME CC2328) with average curves bythe Larson–Miller parameter method.31 The longest creep rupture datum ofSUPER304H tubes is over 85 000 h at 600°C. Their long-term creep strengthis very stable and no degradation in creep strength is expected at thetemperatures up to 750°C.

Figure 20.30 shows the allowable stress determined for SUPER304H(18Cr–9Ni–3Cu–Nb–N steel; KA-SUS304J1HTB, ASME CC2328) compared withthat for the conventional steel, TP347H and the corresponding strengtheningmechanisms used.32 The microstructural change in SUPER304H tubes afteraging for up to 10 000 h in the temperature range between 600°C and 750°Cis available in the literatures,32,33 There is no significant microstructural changeobserved even after aging for 30 000 h at 750°C. A detailed TEM observationof the specimens aged for 3000 h in the temperature range between 600°C and750°C has shown that a fine coherent Cu phase is dispersed in the matrix aswell as fine NbCrN nitrides and no harmful blocky precipitation such as asigma phase is found. This fine dispersion of the precipitates is the major

Test Temp.

600°C650°C700°C750°C800°C

600°C

650°C

101 102 103 104 105

Rupture time (h)

Str

ess

(MP

a)

500400

300

200

1009080706050

40

30

20

700°C

750°C

800°C

20.29 Creep rupture data for SUPER304H (18Cr–9Ni–3Cu–Nb–N steel;KA-SUS304J1HTB, ASME CC2328).

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strengthening mechanism of this steel in the creep region at higher temperatures,schematically depicted in Fig. 20.30.34

Figure 20.31 shows change in toughness of SUPER304H tubes after agingfor up to 30000 h in the temperature range between 500°C and 750°C.Toughness is reduced by aging for a short period but keeps at a high leveleven after aging for 30 000 h at 750°C.34

HR3C (25Cr–20Ni–Nb–N steel; KA-SUS310J1TBÅCTP310HCbN, ASMECC2115) is used for superheater and reheater tubes in fossil fired, blackliquor recovery and refuse fired boilers.35 An improved TP310 steel has beendeveloped by the addition of niobium and nitrogen with an increase is creepstrength at elevated temperatures. The microstructure of the steel consists ofan austenitic matrix strengthened by M23C6 carbide mainly along the grainboundary, and fine dispersed NbCrN nitride in matrix. NbCrN is fine andstable even after long-term creep exposure at high temperatures. No significantsigma phase has been found after aging for more than 30 000 h in thetemperature range between 600 and 800°C, achieved by optimizing the Ni-balance.

Figure 20.32 shows creep rupture data for HR3C tubes with average curvesby the Larson–Miller parameter method.36 The longest creep rupture datum ofHR3C tubes is for about 90 000 h at 700°C. Their long-term creep strength isvery stable and no degradation in creep strength is expected at the temperaturesup to 750°C.

The microstructural change in HR3C tubes after aging for up to 10 000 hin the temperature range between 600°C and 750°C is reported in the

0 100 200 300 400 500 600 700 800Temperature (°C)

Allo

wab

le t

ensi

le s

tres

s (M

Pa)

180

160

140

120

100

80

60

40

20

0

Tensile region Creep region

Solid solutionstrengthening (N)

18Cr–9Ni–3Cu–Nb–N

TP347H

Precipitationstrengthening

Nb (C,N)NbCrNM23C6

Cu phase

20.30 Allowable tensile stress determined for SUPER304H(18Cr–9Ni–3Cu–Nb–N steel; KA-SUS304J1HTB, ASME CC2328) according to theJapanese METI standard.

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literature.37,38 There is no significant microstructural change observed evenafter aging for 10000 h at 750°C. A detailed TEM observation of the specimensaged for 3000 h in the temperature range between 600°C and 750°C hasshown that fine NbCrN nitrides dispersion is identified in the matrix and noharmful blocky precipitation such as the sigma phase is found.

500°C550°C600°C

650°C700°C750°C

as sol. 1 10 100 1000 104 105

Aging duration (h)

Ch

arp

y im

pac

t va

lue

(J c

m–2

)

250

200

150

100

50

0

20.31 Change in toughness of SUPER304H(18Cr–9Ni–3Cu–Nb–N steel;KA-SUS304J1HTB, ASME CC2328) after aging for up to 30000 h at500°C and 750°C.

Test Temp.

600°C650°C700°C750°C800°C

600°C

650°C

101 102 103 104 105

Rupture time (h)

Str

ess

(MP

a)

500400

300

200

1009080706050

40

30

20

700°C

750°C

800°C

20.32 Creep rupture data for HR3C (25Cr–20Ni–Nb–N steel; KA-SUS310J1TB, TP310HCbN).

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20.3.4 Fe–Ni based austenitic alloys used for 700°C A-USC plant

In Europe and the USA, new research projects, THERMIE AD70039 andDOE USC project40 have been initiated to achieve much higher temperatureand pressure conditions such as over 700°C at 35 MPa and 760°C at 35 MPa,respectively. This is, in fact, a challenging program for the metallurgistssince none of the ferritic and austenitic steels developed seems to survive insuch a hostile environment, and there is no experience of using Ni–Co basealloys for large diameter and heavy wall thickness piping applications mainlyowing to their poor toughness, fatigue resistance and workability.

Figure 20.33 shows creep strength data of the candidate austenitic alloysfor 700°C A-USC boilers, compared with those of the advanced 9/12%Crferritic steels.27,40 It can be seen that only Ni–Co base alloys such as Alloys740 and 617 strengthened by γ ′ are applicable with respect to creep strengthat 700°C, while HR6W, Fe–Ni alloy is marginal and SUPER304H, austeniticstainless steel cannot be used at present.

Figure 20.34 shows creep rupture data for the Fe–Ni alloy, 23Cr–43Ni–7W at temperatures between 650 and 800°C. Unlike Alloys 740 and 617,this alloy is strengthened by a combination of the dispersion of precipitatessuch as M23C6, MX and the Laves phase and a solid solution of W withstabilization of the precipitates by B. This is confirmed by a TEM observationshowing that a fine dispersion of the Laves phase as well as M23C6 and MXwhich serve as effective obstacles to the dislocation in motion remainseven after long term-creep deformation for 58 798 h at 700°C with a stressof 98 MPa, as shown in Fig. 20.35.

Str

ess

(MP

a)

1000

100

1020 22 24 26 28

T(log t + 20) × 103

SUPER304HAlloy 800H

HR6W

Alloy 617

Alloy 617 modAlloy 740

Ni-Co base allosy

9.12Cr

600°C × 105h

650°C × 105h700°C × 105h

20.33 Creep strength of the candidate materials for 700°C USCboilers.40

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650°C700°C750°C800°C

101 102 103 104 105

Rupture time (h)

Str

ess

(MP

a)

400350

300

250

200

150

100

50

20.34 Creep rupture data for 23Cr–43Ni–7W alloy.

27.35 TEM micrographs of the specimen of 23Cr–43Ni–7W alloy creptand ruptured after 58798.4 h at 700°C.

Laves

MX

1µm

M23C6 250nm

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Note that the phase stability against σ phase formation is a key to achievingboth high strength and enough ductility in these austenitic alloys. This hasbeen demonstrated by 23Cr–43Ni based model alloys with Mo or W. Figure20.36 shows the Larson–Miller parameter plot of the creep rupture data for23Cr–43Ni based model alloys with 3%/5%Mo or 5%/7%W crept at 700,750 and 800°C. It can be seen that the alloys with W are superior to thosewith Mo for longer term creep strength at high temperatures. To analyze thisdifference between the alloys with Mo and those with W, the microstructuralevolution during creep deformation at high temperatures has been extensivelyexamined using a thermodynamic calculation and TEM observation.

Figure 20.37 shows a comparison of the equilibrium phase diagrams forthe alloys with 5%Mo and 7%W calculated by Thermo-Calc. In the alloywith 5%Mo, the σ phase is the dominant phase with M23C6 and a smallamount of MX at the temperatures between 700 and 800°C, while in thealloy with 7%W, the Laves phase is the major precipitate with M23C6 andMX. This characteristic difference in the phase equilibria of the alloys hasbeen confirmed by optical micrographs and extraction replicas from specimensaged for 3000 h at 750°C, shown in Fig. 20.38. In the alloys with Mo coarseblocky precipitates identified as σ phase form along grain boundaries, witha small amount of fine precipitates, M23C6 and coarse Laves phase inside thegrains. It is, however, noted that in the alloys with W, no blocky σ phase hasbeen identified and instead fine Laves phase forms homogeneously with fineprecipitates, M23C6 and MX both along the grain boundary and inside thegrains. It is thus considered that in the alloys with W long-term creep strength

M3M5W5W7

700°C × 105h

21.5 22.0 22.5 23.0 23.5 24.0 24.5 25.0T(log t + 20) × 103

Str

ess

(MP

a)

200

150

100

80

20.36 Larson–Miller parameter plot of creep data for 23Cr–43Ni alloyswith 3%/5%Mo or 5%/7%W.

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has been achieved by fine dispersion of the Laves phase along with M23C6

and MX, while in the alloys with Mo, coarse precipitates such as the σ phasealong the grain boundary and the Laves phase inside grains have beendetrimental to the long-term creep strength.

To explore further increase in creep strength of these alloys, the possibilitiesof increase in the Laves phase and of introduction of γ ′ and α-Cr phaseshave been examined using phase diagram calculations based on 0.08C–23Cr–43Ni–7W–0.1Ti–0.2Nb–0.003B, as shown in Fig. 20.39. The amountof Laves phase can be increased to a great extent with increasing W without

20.37 Equilibrium phase diagrams for (a) 23Cr–43Ni–5Mo and(b) 23Cr–43Ni–7W calculated by Thermo-Calc.

Mo

le f

ract

ion

of

pre

cip

itat

es

0.20

0.18

0.16

0.14

0.12

0.10

0.08

0.06

0.04

0.02

0

M5(5%Mo)

σ

αCr

M23C6

500 1000 1500Temperature (°C)

(a)

Mo

le f

ract

ion

of

pre

cip

itat

es

0.10

0.09

0.08

0.07

0.06

0.05

0.04

0.03

0.02

0.01

0500 1000 1500

Temperature (°C)(b)

W7(7%W)

αCr

Laves

C23C6

MXµ

MX

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20.38 Optical microstructures and extraction replicas from 23Cr–43Niwith 3/5Mo or with 5/7W alloys aged for 3000 h at 750°C.

M3 M5

σσ

W5 W7

20µm

Laves

M23C6

M5 M5

W7W7

Laves

1µm

5µm 1µm

5µm

M23C6

20.39 Calculated phase diagrams for stabilizing Laves, γ’ and α-Cr in23/30Cr-43/53Ni-7/10W-Ti-Al alloys and MX phases in Alloy120.

0.10

0.09

0.08

0.070.06

0.05

0.04

0.03

0.02

0.01

0

Mo

le f

ract

ion

of

pre

cip

itat

es

Temperature (°C)(a)

15001000500

Temperature (°C)(c)

15001000500

Temperature (°C)(b)

15001000500

Temperature (°C)(d)

15001000500

0.20

Mo

le f

ract

ion

of

pre

cip

itat

es 0.18

0.16

0.14

0.12

0.10

0.08

0.06

0.04

0.02

0

0.20

Mo

le f

ract

ion

of

pre

cip

itat

es 0.18

0.16

0.14

0.12

0.10

0.08

0.06

0.04

0.020

0.02Mo

le f

ract

ion

of

pre

cip

itat

es

0

0.04

0.06

0.08

0.10

0.12

0.14

0.16

0.18

0.20

23Cr-43Ni-10W

Laves

αCr µ

MX

M23C6

MXM23C6

αCr

µ

γ′

23Cr-53Ni-7W-2Ti-0.5Al

Alloy 120

MXLaves

αCr

σ

M23C6

MX

µ σ

αCr

30Cr-43Ni-7W

M23C6

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any other phase changes. γ ′ phase can be introduced by increasing Ti andNi, while, on the other hand, the Laves phase decreases and α-Cr phase andMX increase at higher temperatures. The α-Cr phase can be stabilized byincreasing Cr and Ni, while the Laves phase decreases to a great extent.MX is found to be used effectively in Alloy120, although the σ phase isdominant at the temperatures around 700°C. It is thus concluded from thesecalculations that there are several possibilities for further strengthening ofthe austenitic alloys. It is, however, noted that phase stability against σphase formation is a key to achieving both high strength and enough ductilityin these austenitic alloys to be used for 700°C A-USC boilers.

20.4 References

1 V. K. Sikka, C. T. Ward and K. C. Phomas, ASM International Conference Production,Fabrication, Properties and Application of Ferritic Steels for High-TemperatureApplications, ASM, Washington DC, 1981.

2 F. Masuyama, I. Ishihara, T. Yokoyama and M. Fujita, Thermal Nuclear Power,1995, 46, 498.

3 K. Muramatsu, Proceedings Advanced Heat Resistant Steels For Power Generation,R. Viswanathan and J. Nutting (eds), The University Press, Cambridge, 1998, 543.

4 F. Masuyama, Proceedings Advanced Heat Resistant Steels For Power Generation,R. Viswanathan and J. Nutting (eds), University Press, Cambridge, 1998, 33.

5 Sumitomo Seamless Tubes and Pipe Creep Data Sheets, Sumitomo Metal Industries,1993.

6 Y. Sawaragi, N. Otsuka, H. Senba and S. Yamamoto, Sumitomo Search, 1994, 56,34.

7 T. Kan, Y. Sawaragi, Y. Yamadera and H. Okada, Proceedings 6th InternationalConference on Materials for Advanced Power Engineering 1998, Liege,Forschungszentrum Julich GmbH, 1998, 60.

8 Y. Sawaragi, H. Teranishi, A. Iseda and K. Yoshikawa, Sumitomo Search, 1990, 44,146.

9 M. Igarashi, H. Okada and H. Semba, Proceedings 8th Workshop on the InnovativeStructural Materials for Infrastructure in 21st Century, NIMS, Tsukuba, 2004, 194.

10 M. Igarashi, M. Yoshizawa, H. Okada, H. Matsuo, Y. Yamadera and A. Iseda, CAMPISIJ, 2003, 17, 336 (in Japanese).

11 T. Otsuka, Y. Yamaji, S. Takenaka, M. Ichiryu, H. Momma and S. Takano, ThermalNuclear Power, 2006, 57, 734.

12 N. Komai, F. Masuyama, I. Ishihara, T. Yokoyama, Y. Yamadera, H. Okada, K.Miyata and Y. Sawaragi, Advanced Heat Resistant Steels For Power Generation,University Press, Cambridge, 1998, 96.

13 Y. Sawaragi, K. Miyata, S. Yamamoto, F. Masuyama, N. Komai and T. Yokoyama,Advanced Heat Resistant Steels For Power Generation, University Press, Cambridge1998, 144.

14 F. Masuyama, T. Yokoyama, Y. Sawaragi and A. Iseda, Materials for AdvancedPower Engineering, Part 1, Kluwer Academic Publishers 1994, 173.

15 Y. Sawaragi, T. Kan, Y. Yamadera, F. Masuyama, T. Yokoyama and N. Komai,Proceedings of the 6th International Conference on Materials for Advanced PowerEngineering 1998, Liege, Forschungszentrum Julich GmbH, 61.

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16 M. Yoshizawa, A. Iseda and M. Igarashi, CAMP ISIJ, 2005, 18, 1549. (in Japanese).17 K. Miyata, M. Igarashi and Y. Sawaragi, ISIJ Inte., 1999, 39, 947.18 K. Miyata and Y. Sawaragi, ISIJ Inter., 2001, 41, 281.19 A. Iseda, Y. Sawaragi, S. Kato and F. Masuyama, Proceedings of the Fifth International

Conference on Creep Materials, Florida, 1992, 389.20 Y. Sawaragi, A. Iseda, K. Ogawa and F. Masuyama, Materials for Advanced Power

Engineering, Part 1, Kluwer Academic Publishers 1994, 309.21 M. Yoshizawa and M. Igarashi, International Journal of Pressure Vessels and Piping,

2007, 84, 37.22 M. Igarashi, M. Yoshizawa, A. Iseda, H. Matsuo and T. Kan, Proceedings the 8th

Liege Conference on Materials for Advanced Power Engineering, J. Lecomte-Beckers,F. Schubert and P.J. Ennis (eds), Liege, Forshungszentrum, Jülich GmbH, 2006,Volume II, 1095.

23 L. Kloc and V. Sklencka, Proceedings 6th Liege Conference on Materials for AdvancedPower Engineering, J. Lecomte-Beckers, F. Schubert and P. J. Ennis (eds), Liege,Forshungszentrum, Jülich GmbH, 1998, Volume, I, 215–222.

24 L. Kloc and V. Sklencka, Mater. Sci. Eng. A, 1977, 234–236, 962.25 K. Kimura, Proceedings of ASME Pressure Vessel and Piping Conference, Denver,

CO, 17–21 July, ASME, 2005, Paper 71039.26 M. Igarashi, S. Muneki, H. Kutsumi, T. Itagaki, N. Fujitsuna and F. Abe, Proceedings

5th International Charles Parsons Turbine Conference, A. Strang, W.M. Banks,R.D. Conroy, G.M. McColvin, J.C. Neal and S. Simpson (eds), University Press,Cambridge, 2000, 334.

27 M. Igarashi, H. Okada and H. Semba, Proceedings 9th Workshop on the InnovativeStructural Materials for Infrastructure in 21st Century, NIMS, Tsukuba, 2005, 96.

28 Y. Sawaragi, N. Otsuka, H. Senba and S. Yamamoto, Sumitomo Search, 1994, 56,34.

29 Creep Properties of Heat Resistant Steels and Superalloys, Landolt-Bornstein NewSeries VIII-2B, Springer, 2004, 251–257.

30 H. Teranishi, K. Yoshikawa, H. Fujikawa, M. Kubota, K. Tokimasa and M. Miura,Proceedings of International Conference on Coatings and Bi-Metallics for EnergySystems Chemical Process Environment, ASM Conference, South Carolina 1984.


32 Y. Sawaragi, N. Otsuka, K. Ogawa, S. Kato and S. Hirano, Sumitomo Search, 48,1992, 50.

33 Y. Sawaragi, N. Otsuka, K. Ogawa, S. Kato and S. Hirano, Sumitomo Metals, 1991,43, 24 (in Japanese).

34 H. Senba, Y. Sawaragi, K. Ogawa, A. Natori and T. Kan Materia, 2002, 41, 120 (inJapanese).


36 Y. Sawaragi, Y. Teranishi, A. Iseda, K. Yoshikawa, Sumitomo Search, 1990, 44, 146.37 Y. Sawaragi, H. Teranishi, H. Makiura, M. Miura and M. Kubota, Sumitomo Metals,

1985, 37, 66.38 Y. Sawaragi, Y. Teranishi, A. Iseda and K. Yoshikawa, Sumitomo Metals, 1990, 42,

260 (in Japanese).39 R. Blum and R. W. Vanstone, Proceedings 6th International Charles Parsons Turbine

Conference, 16–18 September 2003, Dublin, Ireland, A. Strang, R. D. Conroy, W.

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M. Banks, M. Blackier, J. Leggett, G. M. McColvin, S. Simpson, M. Smith, F. Starrand R. W. Vanstone (eds), 2003, 487–510.

40 R. Viswanathan: Proceedings EPRI Conference on Materials and Corrosion Experiencefor Fossil Power Plants, US Program on Materials Technology for USC PowerPlants, South Carolina, USA 2003.

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573

21Using creep-resistant steels in turbines

T.- U. K E R N, Siemens AG Power Generation Group,Germany

21.1 Introduction

Energy is the source of general well being and the standard of living in eachcountry. Recent history has shown that a safe and sufficient energy supply isthe basis for the development of the whole population of a country andregion. The world situation today is characterised by a constantly increasingpopulation and the desire for improved living conditions. The industrialisationof a country is automatically connected to the availability of electricity. It isthe key to progress.

In the decades to come, there will continue to be heavy reliance on fossilfuels, such as coal, oil and natural gas, because nuclear technologies arealways under discussion. Research into fusion reaction is still under wayand could take a further 20 years to reach industrial importance. Becausereserves of oil and natural gas are unlikely to be sufficient to satisfy fullythe projected increased power demands and to fill the gap until the introductionof further advanced renewable energy sources, coal may be the only fuelavailable in substantial quantities around the world. This will force powerplant utilities to use the most advanced technologies available to increasethe efficiency of their power plants to meet the increasingly stringent emissionregulations for safeguarding health and preserving the environment for futuregenerations.

One of the best known and widely used standard technologies is the steampower plant using a steam generator (e.g. boiler), a steam turbine, condenser,generator and the thermodynamics of the steam–water cycle for the energyconversion of heat to electricity. A typical steam turbine turbo-set arrangementis shown in Fig. 21.1. The components are, from left to right: high pressureturbine HP with valves, intermediate pressure turbine IP with valves, twolow pressure turbines LP, the generator and, underneath, the steam condenser.This chapter deals with the application of creep-resistant steels in the steamturbine components in the past, today and in the future.

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21.2 Implications for industries using creep-resistant

steels

There are different turbine arrangements for a power plant possible dependingon the customer requirements and the technology of the turbine supplier. Butone thing remains: the extensive use of steels in the overall construction. Forexample, with a turbine arrangement for 800 MW as shown in Fig. 21.1, anoverall weight of steels of 1800 tonnes is required, comprising 130 tonnesHP, 350 tonnes IP and 1300 tonnes LP parts. For temperatures higher than400°C, mainly creep-resistant steels are applied because of their excellentmechanical, chemical and physical properties, and easy manufacturability.

Depending on the function, the requirements in service and the specificallyapplied design, the design criteria for different turbine parts and componentsin high pressure HP, intermediate pressure IP and low pressure LP turbinesor combined turbine variants are also different. The main critical componentsare covered in Table 21.1 from rotors and casings, to blades and bolts. Theproperty profile requested covers all areas of the material behaviour. It rangesfrom static strength, creep rupture strength, toughness, fatigue propertiesand crack growth to the influence of the environment characterized by corrosion,erosion and oxidation behaviour.

For a safe design, different strength and toughness criteria have to beapplied to ensure the safe operation of the turbo-set during its designedlifetime. Material properties have to be determined by standardized testswith specific specimens. Transfer of the results to the actual components ismade by models and rules based on official standards, whereas special aspects

21.1 Steam turbine arrangement, example for a 800 MW unit.

HP + valves600°C/300bar

IP + valves610°C/60bar

2 × LP350°C

Generator

Condenser

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Using creep-resistant steels in turbines

575

Table 21.1 Material data requirements for application in steam turbine design

Requirements HP- & LP-rotors, HP- & HP- & LP-blades HP- &IP-rotors LP-discs IP-casings IP-blades IP-bolts

Static strength:tensile strength

Creep rupture strength:creep behaviour

Toughness:fracture toughness

Fatigue properties:low cycle fatigue (LCF)high cycle fatigue (HCF)

Crack growth:static – creep CGalternating – fatigue CG

Corrosion:local corrosion

stress corrosion cracking

corrosion fatigue

Erosion behaviourOxidation behaviour

HP = high pressure turbine, IP = intermedium pressure turbine, LP = low pressure turbine.

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are investigated in research and development (R&D) projects and are broughtinto industrial practice. The summary in Table 21.2 shows the differentcriteria that play a key role for engineers. Depending on the service temperature,the material properties are sensitive to the time of application. Typical materialbehaviour for steels at T > 400°C is the phenomenon of creep. It is one of themost critical properties in application because the design life is increasingfrom 100000 h to 200000 h operation time owing to current customer requests.Long-term testing has to be performed and accurate extrapolation methodshave to be developed and validated. This point is discussed in more detail inother chapters of this book.

The description in Table 21.2 helps to clarify the different loading dataduring real operation and to mirror these with the possible material data. Theloading data are:

• the different types of stress σi

• stress changes ∆σi

• strain changes and differences ∆εi

• operation temperature Ti

• start-up and shut-down operation cycles Ni

• the overall operation time ti.

As an example, the different types of stress sources for rotors duringoperation, start-up and shut-down are schematically shown in Fig. 21.2.Centrifugal force and temperature gradients cause stresses which the componenthas to withstand during its operation time and start-up and shut-down cycles.Additional local loading occurs if there are natural or manufacturing defectsin the components volume.

Different kinds of analysis have to be performed to ensure reliable andsafe operation of the components. Starting with the stress analysis, fracturemechanic evaluation and fatigue evaluation have to be performed. Therespective material properties used are determined in the following materialtests:

• Tensile: 0.2 yield strength (YS), ultimate tensile strength (UTS), elongation• Creep rupture: creep rupture strength Ru for a time t at temperature T,

Ru/t/T• Creep elongation: creep elongation limit Rpε for a time t at temperature

T, Rpε/t/T

• Fracture toughness: linear–elastic static fracture toughness KIC

• Fatigue crack growth: amplitude of stress intensity ∆K, crack growthrate da/dN

• Creep crack growth: stress intensity K, crack growth rate da/dt• Creep crack initiation: stress intensity for static crack initiation KIid and

energy integral C*• Low cycle fatigue: amplitude of strain ∆ε, cycles to failure Nfi.

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577

Table 21.2 Summary of design criteria for steam turbines

Imposed Material data Limit Allowable Required materialload data values values properties

σi Stress evaluation0.2YS, UTS, elongation 0.2YS σi ≤ 0.2YS/S 0.2YS/UTS/Elong. = f (T)Creep rupture strength 0.8*Ru/t/T σi ≤ Ru/t/T/S Ru/t/T = f(T, t)Creep elongation limit 0.8*Rpε/t/T σi ≤ Rpε/t/T/S Rpε/t/T = f (T, t)

Fracture mechanic evaluation (a0 = start defect size)∆σi – short term safety KIC after

a0, KIC acrit a0 ≤ acrit,1/S Long term exposure

∆ε – long term safetycyclic : a0, da/dN ∆ai a0 + ∑ ∆ai ≤ acrit,2/S

da/dN = f (∆K, T)Ti static : a0, da/dt ∆ai at low stress rate

Fatigue evaluation Creep crack initiationNi – with defect (a0) tA = f (Klid)

cyclic : da/dN = f (∆K) Nfi = f (C*)ti static : da/dt = f (Klid) (tui) ∆ε = f (N, T)

– without defect (with notch) ∑ Ni /Nfi + ∑ t i/tui ≤ Eallow Ru/t/T = f (t, T)cyclic : Nfi = f (∆ε) Nfi Rpε/t/T

static : tui = f (Ru/t/T) tui

i, index for different service loading conditions; UTS, ultimate tensile strength; YS, yield strength; IC, mode I critical; Iid, mode I ideal elastic;u/t/T, ultimate strength for time t at temperature T.

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The initial crack sizes a0 have to be determined by appropriate non-destructivetesting during quality assurance in the material procurement process.

One of the most applied material classes for critical turbine componentsoperating at temperatures T > 400°C is the low alloyed CrMoV steel whichis also used for forgings and castings. The Cr content ranges from 0.5% upto 2.25%. The balance of properties described in Tables 21.1 and 21.2 can beachieved in practice. This is demonstrated for a typical 1CrMoNiV rotorforging in Fig. 21.3 in terms of basic properties.1 The strength was investigatedfor several parts with so-called radial cores going through the whole rotorbody in the radial direction. It was found to be stable whereas the toughnesscharacterized by the 50% fracture appearance transition temperature (FATT)shows clear changes from the surface at 20°C to the centre area at about80°C. The reasons for this are the cooling conditions during heat treatmentof large parts with thermal gradients and the material type itself which resultsin an annealed martensitic structure at the surface and an upper bainitestructure in the centre.

The tougher martensite shows disadvantages for long-term creep behaviour,as shown in Figs 21.4 and 21.5, with investigations simulating the differentpossible microstructures of 1CrMoV steels.2 The lowest creep rate was achievedwith the upper bainite structure. The results from a real rotor forging at thesurface and in the centre are also given. They demonstrate that the coolingconditions have to be optimized, for example by slowing down, to get ahomogenous upper bainite structure with the required best creep propertiesin the forging. The creep rupture strength achieved for times up to 90 000 hconfirms the advantages of upper bainite. The fatigue strength was alsodetermined, resulting again in a better behaviour for the upper bainite structure(Fig. 21.6).

TaTemperaturedistribution

Thermalstress

Centrifugalforce loading

Pressure loading

Ti

∆T

Stress caused bycentrifugal force

Stress causedby pressure

Thermal stress at thenut of turbine blade

fixture

σ∆Tαk · σ∆T

21.2 Loading conditions for a turbine rotor, schematically.

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Using creep-resistant steels in turbines 579

21.4 Influence of microstructure on creep behaviour of 1CrMoV rotorsteel at 530°C and 127 MPa.2

Ferrite – PearliteMartensiteLower bainiteUpper bainiteOutsideCentre

As deliveredtransverse

101

100

10–1

10–2

Str

ain

(%

)

101 102 103 104 105

Time (h)

Heat treatment contour

1.3

m

30 Mg

Radial core

Rm

Rp 0.2

Z

A

Surface Centre

FATT several rotors

Surface CentrePosition in the forging

800

700

600

500

Str

ess

(N m

m–2

)Te

mp

erat

ure

(°C

) 100

80

60

40

20

70

50

40

30

20

10

Du

ctili

ty (

%)

21.3 Balance of properties for a 1CrMoNiV rotor with respect tothrough-hardenability.1

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Summarizing these points, the martensitic structure at the surface of rotorforgings has to be avoided because the rim position later on holds the rotatingblades of the turbine shaft which are highly loaded during operation. Heattreatment and manufacture have to be performed in an accurate way toensure the right microstructure at the relevant points.

Ferrite – Pearlite ε (% min–1)Upper bainite 5.5–6Lower bainiteMartensiteWith hold time ± 20 minWithout hold time

.

Str

ain

∆ε t

(%

)

4

2

100

8

6

4

2

10–1

102 2 5 103 2 5 104 2 5 105

Cycles to crack initiation Ni

Time: 2 ·

RE tmt

creep = ˆ ∆ε

21.6 Influence of microstructure on fatigue behaviour of 1CrMoVrotor steel at 530°C.2

Ferrite – PearliteMartensiteUpper bainiteLower bainiteOutsideCentreNotchedSmooth

Str

ess

(MP

a)

1000

800

300

200

100

Hot tensiletest

Asdeliveredtransverse

Kt = 4.5

Scatterband of 1% CrMoNiV-steelsMean value ±20% (SEW555)

Specimen still running

Specimen withdrawn

101 102 103 104 105

Time (h)

21.5 Influence of microstructure on creep strength behaviour of1CrMoV rotor steel at 530°C.2

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Another important aspect is shown in Fig. 21.7. The investigations herewere made to obtain information about the notch weakening behaviour ofthis type of steel. Notch weakening means that notched creep specimens failearlier than the respective smooth specimens. The importance in practice isconfirmed by the fact that the surface of a rotor always has areas similar tonotches such as in the steam inlet area as well as in the blade attachmentgrooves. Using the wrong quality heat treatment, the steel can show a significantdrop in rupture ductility and therefore also notch weakening. Furthermore,the application of 1CrMoV steel in bolts is common, and here the notchesare clearly introduced by the thread itself. Cracks caused by the effect ofnotch weakening can form before the end of designed lifetime resulting infailure or damage. The delivery specification of the turbine makers has togive advice about the right heat treatment, also implying that they know theirmaterials very well.

Figure 21.8 demonstrates that not only does the creep rupture strengthplay an important role in the integrity of components but also the creepelongation and rupture ductility.1 Here the stresses of a stop and controlvalve in the cast design are calculated over the service lifetime. Starting withhigh elastic stresses at the inner wall (position A) and lower stresses at thesurface (Position B) caused by the internal pressure, the equivalent stress σvhas reduced after 100 000 h operation time at the inner surface but increased

E

lon

gat

ion

(%

)

Red

uct

ion

of

area

(%

)S

tres

s (N

mm

–2)

1000

500

200

100

50

100

80

60

40

20

0

900°C 1 h/oil + 750°C 2 h 1000°C 1 h/oil + 570°C 2 h

αk = 4.5

α = 4.5

0.1 1 10 102 103 104 105

Time to rupture (h)0.1 1 10 102 103 104 105

Time to rupture (h)

21.7 Creep rupture behaviour of 1CrMoV steels at 550°C for twodifferent heat treatments.1

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outside owing to relaxation and plastic deformation via creep. The tangentialcreep strain has significantly increased for the inner position A to 0.4% after100 000 h which is, compared to the outer surface position B with 0.1%, afactor of four. If the real creep elongation reaches the rupture ductility for theappropriate stress state at any position, cracks can occur and propagate bycreep crack growth. Cracks introduced by creep will propagate faster duringthe operational cycle via fatigue crack growth or creep crack growthmechanisms. Multiaxiality, reducing the toughness reserve of the material,can also strongly influence the creep and fatigue component behaviour. Designrules have to take into account the different operational, customer specificconditions and provide sufficient exact residual life tools to ensure the safeoperation of the components for the designed lifetime.

φ900φ500

σv Elastic σv after 105 hMises reference stress (N/mm2)

203050

3040

20

10 5060

70 110 50 60

20 40

30

Mean stress σ = 47 N/mm–2 according to design codes

B

A 250 bar540 °C

Location A

Location B

103 2 × 103 104 2 × 104 105 2 × 105

Time (h)

Tan

gen

tial

cre

ep s

trai

n (

%) 0.6

0.5

0.4

0.3

0.2

0.1

21.8 Creep stress and strain of a high pressure stop and controlvalve body from 1CrMoV cast steel.1

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21.3 Improving the performance and service life of

steel components

Up until the 1980s, based on the available materials and their property profile,especially the achievable creep strength, the thermodynamic parameters ofsteam turbines were restricted to main steam temperatures T of 540°C. Theoptimization in design and materials yielded an increase in T to 565°C in the1990s. The material classes used then were creep-resistant low and highalloyed steels of the 1-2CrMo(V) and 11-12CrMoV type available of thetime. Examples of typical creep-resistant steels for steam turbines in Europeare given in Table 21.3.

The increasing demand of the energy market for further improved efficienciesin power plants and political efforts for climate protection on Earth by reducingCO2 emissions have resulted in extensive R&D initiatives for materialdevelopment worldwide. The aim has been to increase the creep rupturestrength of materials by keeping the other properties stable.

At the beginning, single types of material were reviewed to identify themost promising ones. Different aspects of component application were takeninto consideration, for example that rotors for high temperatures can reachlarge diameters and weight, or casings have to be weldable for repair andconstruction purposes. Finally, work was started with the martensitic 9–12%CrMoV steels. The advantages are clearly visible:

• method of manufacture (forging, casting)• good through hardenability up to 1200 mm diameter• balance in strength and toughness (long and short term)• weldability• potential for high oxidation resistance• ease of fabrication• lower cost than austenites and Ni-base alloys

Table 21.3 Examples of creep-resistant steels for power plant application upto 565°C (in wt%)

Component type C Cr Mo V W N Nb

Rotor 0.30 1.20 1.10 0.300.22 2.10 0.85 0.30 0.65

Blade 0.19 11.0 0.60 0.20 0.10 0.300.22 12.0 1.00 0.300.20 13.0 1.00

Valve body 0.17 1.20 0.500.17 1.30 1.00 0.25

Casing 0.18 2.25 1.000.17 1.30 1.00 0.25

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In Europe efforts have been concentrated on the COST programme (Fig.21.9). This is a long established European programme aimed at coordinatingpre-competitive research activities in numerous areas of science and technology.Participants include turbine and boiler makers, manufacturers of forgings,castings, pipework and welding consumables, as well as testing institutesand utilities. The work has spanned the entire range, from basic research anddevelopment, through non-destructive and destructive testing of smallexperimental batches and large scale components, to standardization ofspecific steels and full-scale practical application in currently operating powerplants.

Action started with COST501 (1986–1997), was continued in COST522(1998–2003)3–5 and is currently run as COST536 (2004–2009).6,7 It hasestablished a strong trans-European network in this field. The effectivenessof past COST actions is exemplified by the introduction of newly developedmartensitic steels for forgings, castings and pipework. These improved steelsare in commercial operation in advanced European power stations and havemade it possible to increase the operating steam temperatures from 530–565°C to 580–600°C with a corresponding increase in thermal efficiency.

In Japan a R&D programme was initiated by the Electrical PowerDevelopment Company (EPDC) in 1981 to explore the possibilities ofdeveloping and applying new materials for high-temperature steam powerplants. The programme started with basic research into strengtheningmechanisms in 9–12%Cr steels and has led to the introduction of a numberof steels with much improved creep strength at temperatures of 600°C andabove.8 Full scale components, such as rotors, casings and boiler tubes andpipes have been manufactured and employed in a series of advanced steampower plants ordered by EPDC. Further basic research was started 1997 withthe NRIM-STX21 programme which aimed to find new alloy design conceptsfor boiler components. Current activities are concentrated on the transformationof the results from trial melts to real industrial applications.

In the USA, collaborative work has been performed within the EPRIprogramme RP1403, which began in 1978 with basic studies and startedpractical work in 1986. Activities were focused on the development of steelsfor applications such as thick-walled pipes. This programme was internationaland work was carried out by companies in the USA, Japan and Europe. Theproject was successful in that the materials P92 (NF616) and P122 (HCM12A)have been code approved by ASME. In addition, thick-walled pipes andheaders made from these steels have been manufactured and full-scale testedand are in operation in power plants in Europe.

The criterion for success in all of these developments for new 9–12CrMoVsteels was set as an improved creep rupture strength tested by conventionalcreep specimens. The other properties determined in standard tests havebeen accepted if they were not worse than those of 1CrMoV steels. A summary

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585

International projects of advanced power plants

Japan

R & D : EPDC

1981 – 1991

50 MW Pilot power plant

NRIM – STX 21 Project:USC 650°C/350 bar Boiler

1997 – 2012

Manufacturers, utilities, EPDC

316 bar 566/566/566°C314 bar 593/593/593°C343 bar 649/593/593°C

1989 – 19901991 – 19931994 – 2000 300 bar 630/630°C

Thick wall components

USA

R & D : EPRI

Manufacturers

Study 1978 – 1980

310 bar 566/566/566°C310 bar 593/593/593°C345 bar 649/649/649°CEPRI – RP 1403 – 15

300 – 900 MW

R & D : 1986 – 1993

Steels and components for Boiler + Turbine(USA, Japan, ALSTOM + MAN)

EPRI – RP 1403 –50–WO9000 –38

1990 – 1999

Thick wall pipes: P 92 + P 122(USA, Japan, UK + Denmark)

EUROPE

COST 50/501

1983 – 1997

Manufacturers, steelworks,utilities and R & D – institutes

300 bar 600/600/600°C300 bar 600/620°C

RotorsCast componentsBolting materialsPipes, tubes,Welds

COST 522 COST 536

2004–20091998 – 2003

300 bar 620/650°C

Rotors, cast components,bolts, pipes, tubes, welds

21.9 Development activities for the new 9–12CrMoV class steels.

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for the new European 600°C rotor materials is given as an example in Fig.21.10.4

In comparison with the known 1CrMoV and 12CrMoV steels, importantchanges were made in the chemical composition: the carbon content is reducedand the elements nitrogen, niobium and sometimes tungsten have been added.For the most promising alloy, 9Cr–MoCoVB steel, the elements boron andcobalt play an important role in highly increased creep strength.

Compared with conventional 1Cr steel, a 50–70 K increase in applicationtemperature without any significant design changes is possible. Theimprovement within the class of 12Cr steels is also demonstrated at a 30–50K higher temperature for the same 100 000 h creep rupture strength if weconsider the 9–10Cr steel data. The basis for the extrapolation to a relevantdesign life of 100 000 h is to perform real long-term tests in the laboratorywhich so far in Europe have reached up to 80 000 h for different melts fromreal components.9 This is of great importance for the reliability of theextrapolation because it is known that the new materials can show changesin the precipitation status after 30 000–50 000 h resulting in a strong decreasein the long-term strength. A recent example is the pipe steel P122 withstrongly reduced creep rupture data.10,11 The advantageous creep behaviourof the new steels enables the design of high temperature turbo-sets with thesame design rules as for the conventional machines.

appr. 70°C

appr. 30°C

Steel C Cr Mo W V Nb N B Co1CrMoV 0.28 1.0 0.9 – 0.30 – – – –12CrMoV 0.21 12.0 1.0 – 0.30 – – – –10CrMoV 0.12 10.0 1.5 – 0.20 0.05 0.05 – –10CrMoWV 0.12 10.0 1.0 1.0 0.20 0.05 0.05 – –9CrMoCoVB 0.12 9.0 1.5 – 0.25 0.05 0.02 0.010 1.0

(Weight%)200

100

0

100

000

h c

reep

ru

ptu

re s

tren

gth

500 550 600 650Temperature (°C)

1%CrMoV

12%CrMoV

10%CrMo(W)VNbN10-1(-1)

9%CrMoCoVNbB

21.10 Creep rupture strength at 100000 h for new 600°C steels inEurope.

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Similar to results for forgings, the development of cast materials for valvesand casings has been successful.7 It has resulted in alloys of 9Cr type withW, V, Nb, N and boron. Service temperatures can be increased up to 620°Cdepending on the material data. The advantage of having stationary partslike these loaded mainly by internal pressure lies in the possibility of reducingthe service stresses by increasing the wall thickness. For this, the thermalflexibility characterized by loading via low cycle fatigue during start-up andshut-down operation and thermally introduced secondary stresses has to bemaintained. In comparison, rotor stresses can only be handled when thematerial has a higher creep strength or the temperature is reduced by director indirect cooling. But this is not easily possible for casings and valves. Thecompositions of industrially applied cast materials for higher temperaturesare given in Table 21.4. Component examples for a high temperature rotorforging and an inner casing of a steam turbine with 600°C main steam areshown in Figs 21.11 and 21.12.

In Germany, the application of the new steels to real turbine componentsis accompanied by extensive research work in industry together with severalinstitutes to determine the relevant service behaviour. An overview of the

Table 21.4 High temperature cast materials applied in Europe (in wt%)

Component type C Cr Mo V W N Nb

Casing or 0.12 9.0 1.0 0.20 0.05 0.06valve 0.12 10.0 1.0 0.20 1.0 0.05 0.06

21.11 Rotor for high temperature application.

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main activities is given in Fig. 21.13. It shows that programmes have beenrun or are still ongoing to check:

• the influence of component size and technical ranges for the chemicalanalysis on long-term creep strength9

• the influence of the test location of a real power plant componenton long-term creep strength up to 100000 h for a safe lifetime extrapolation9

• the creep rupture strength and low cycle fatigue behaviour under service-relevant load changing conditions12,13

• the high-temperature crack growth behaviour on fracture mechanics14

• the influence of multiaxiality on material behaviour15

With the knowledge gained in these programmes the design can be optimizedfor new apparatus. In addition, lifetime evaluation during service with itscontinuously changing loading conditions is significantly improved.16,17

In the next part, two examples of material application for 565°C and600°C steam turbines of the Siemens type are discussed: high pressure (HP)and intermediate pressure (IP) turbine classes. The barrel type design of theHP turbine is a typical feature of Siemens steam turbines and was introducedabout 50 years ago (Fig. 21.14). The inner parts guide the stationary bladesin a near perfect axis-symmetric body. No flanges are required as this verticallysplit and bolted casing is compressed by high pressure steam, which iscontained in the inlet part of the outer casing and which has no horizontal orvertical split at all. Although over the years many different HP turbine designs

21.12 Inner casing in 10Cr steel for high temperature application.

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have been used, the principle of pressure containment in a barrel has neverchanged. The barrels of today are for different applications, either built usingan inlet and exhaust casing bolted together (for 600°C application, a materialcombination of 1CrMoV–10CrMoWV is used), or consist of one big barreltype outer casing of 1CrMoV cast material, closed with a bolted-on cover.The monoblock HP turbine rotor is made of either 1CrMoV steel or newlydeveloped 10CrMoWV steel.

21.13 R&D activities for qualification of 600°C steels in Europe.

Determination of component relevant material data

Standard Random Standard Long term LCF

Influence ofmanufacturing

Long term100000 h

As-servicerelevant

Single step

Creep rupturestrength

Low cyclefatigue

High temp.crack

behaviour

Influence ofmulti-axiality

Multi-axialUni-axial behaviour

1CrMoV cast steel (Divided if 600°C)565°C: 1CrMoV cast600°C: 10CrMoWV cast

565°C: 1CrMoV cast600°C: 10CrMoWV cast

565°C: 1CrMoV600°C 10CrMoWV

12CrMoV steel11–12CrMoV steel

21.14 HP turbine material application for a 565°C and 600°C design.

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In contrast to the single flow, barrel type HP turbine, the IP turbine consistsof a top and a bottom half for inner and outer casings and is either a singleor a double flow cylinder, depending on the power range (Fig. 21.15).

Nowadays, the design of IP turbines allows reheat temperatures of morethan 600°C. The double flow IP turbine shown in Fig. 21.15 was applied toa project with 610°C using a special internal cooling.18 With this coolingscheme, a temperature reduction of up to 16 K at the critical rotor surfacesection is achieved. The rotor is either made from 1CrMoV or 10CrMoWVsteel depending on the steam temperature and loading conditions. The innercasing material choice shows the same temperature dependence as the rotorsand is either made from 1CrMoV or 10CrMoWV cast steel. The blades in thehigh temperature region above 400°C are made from 10–12CrMoV(Nb)steel. Steels are not used in the highest inlet temperature areas. Ni-basedalloys have to be applied here depending on the service stresses. For lowertemperatures, a 13Cr steel is applied to take account of corrosion. The outercasing design requirements are well met with iron-based globular cast materialat lower costs than steel casts.

Other turbine manufacturers use a welded design for the rotors with 10Crmaterial in the steam inlet area and low alloyed steel in the area with T <500°C.19–22 The blades in the steam inlet area with highest temperatures areoften made with nickel base alloys or austenites.23,24

These examples show how deeply the design of turbo-sets is influencedby and connected to the class of creep-resistant steels. They allow flexibledesign and operation modes in service. Therefore continuous material

Globularcast iron

565°C: 1CrMoV cast600°C: 10CrMoWV cast

565°C: 1CrMoV600°C: 10CrMoWV

13CrMo

11CrMoVNb12CrMoV

21.15 IP turbine material application for a 600°C design.

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improvements and development of this material class are required and drivenby the industry.

21.4 Next steps into the future

In the 1980s and 1990s there were plans to design ultra super critical powerplants using austenites for very high stressed components.25,26 The idea wasto apply austenites up to temperatures of 650°C filling the gap between themartensitic materials (up to 600°C) and the Ni- or Co- based superalloys(>650°C). In Europe this concept was influenced by the experience alreadygained during the design, installation and service of small power plants withsteam temperatures of T = 600–650°C and an installed power output of 7–107 MW.27,28 The high temperature materials used here were austenites.These machines were designed in the early 1950s. Some of them are still inoperation, for example at Eddystone in the USA.29,30 They have been operatedas base load units with a low flexibility during operation, that is, fewer start-up and shut-down operations. For today’s applications, the main marketrequirements for power plants have radically changed towards a higher poweroutput of >300 MW, flexible operating conditions, short delivery times andcost-effective solutions.

Austenites do not fulfil the design requirements resulting from the newmarket conditions. The main disadvantages of austenites are the physicalproperties and lower strength resulting in high thermal stresses in largecomponents during start-up and shut-down. Therefore this material class isvery susceptible to strain-induced cracking and thermomechanical fatiguecharacterized by the R-value (resistance-to-crack), see Fig. 21.16.18 A highR-value will result in a lower probability of thermal induced crack initiationin thick walls. Prototype tests with an austenite rotor have shown a significantreduction in service life owing to start and stop operation at the 50 MWWakamatsu demonstration plant in Japan.31–33

The Ni-base alloys as well as the martensitic steels show higher R-valuesand are therefore better candidates for higher steam temperatures for thick walledcomponents. At the same time, the manufacturing of large components likerotors and casings is a very special task which requires much more R&D effortto obtain fully realised components with the required homogeneous properties.

As the success of the development of new martensitic creep-resistantsteels increases the applicability of these materials, the development ofnew austenitic alloys seems to be less attractive as it was in the 1990s,especially if the superior material properties of the martensites are takeninto account.34 They are very cost effective and offer a high flexibilityduring operation.

Turbine companies today are applying the 9–10Cr steels in current designsnot only for very high temperature machines but also in efforts to increase

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the performance of highly stressed components for steam parametersT < 600°C. On the other hand, steam power plants with temperatures of upto 600–610°C can be supplied for today’s customer by applying a specialdesign. The new generation of 9–11CrMoVCoB steels under developmentwill enable further advanced steam parameters up to 630°C and thereforeadditionally increased efficiency. A survey of applicable material classes forsteam turbine components is given in Table 21.5. It shows the expectationsthat the limit for creep-resistant steels is expected to be at about 650°C.

The next technological step is already on the way. There are now concreteplans to design and build 700°C steam power plants. The basic work wasstarted at the end of the 1990s in Europe.35,36 For 700°C steam power plants,steels no longer serve as candidate materials for the high temperature areas.Nickel- and cobalt-based superalloys play an important role.37,38,39

Unfortunately the manufacture of large components using them is restricted.40

Therefore new design concepts including multi-material components have tobe developed. Steels will serve for the parts where the temperatures are lowenough to design with their properties. The connection to the Ni-basedcomponents will be made by welding or other comparable technologies.This means that the higher the application temperature of new steels, thesmaller the amount of expensive Ni-base superalloys that have to be used.41,42

With this approach, the chances of realisation of 700°C steam power plantsare greatly increased.

The current plan of a large German utility is to build a 700°C demonstrationpower plant starting in 2010, to begin operation in 2014.43 This challenge

21.16 Susceptibility to crack initiation for austenites and Ni-basealloys at 650°C.

R = 0.2YS*K/(E*α)

K – thermal conductivityE – E-module at Tα – linear coeff. of expansion

High R-value → low probability ofthermal induced crack initiation

Temperature: 650°CAustenites

Type304

Type316

Type321

Type347 E

ssh

ete

1250

Type800

A 286

NF 709

X8C

rNi

Mo

VN

b 1

6–16

X8C

rNi

Mo

VN

b16

–13 X8C

rNi

Mo

BN

b 1

6–16

NiC

r20T

iAl

New

10C

r( 6

00° C

)

1Cr

(550

°C)

Martensites

Ni-base alloy

Mat

eria

l p

aram

eter

R8000

7000

6000

5000

4000

3000

2000

1000

0

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has been taken up by European boiler and turbine makers. Strong efforts arenow required to achieve this ambitious aim.

21.5 Summary

The importance of creep-resistant steels for steam turbine application is verygreat. Many of these steels are used in different compositions and for differentcomponents. The well-balanced property profile of these steels enables asafe and sound design of turbo-sets with lifetimes of >100 000 h.

The development of new creep-resistant steels in the last decade hasopened up the market for steam temperatures up to 620°C, resulting inhigher efficiencies for fossil-fueled steam power plants and therefore meetingefforts for climate protection by reducing CO2 emissions. The next generationof 700°C power plants is already being planned. Here also steels will keeptheir important role for design purposes and cost efficiency, enabling specialdesigns and optimized costs. In this way, an economically viable energysupply can be achieved despite the use of expensive nickel- and cobalt-basedsuperalloys.

Future R&D will concentrate on the development of new creep-resistantweldable steels for temperatures up to 650°C, improved superalloys for largeand thick-walled components in turbine and boiler and cost-optimizedsuperalloys to reduce overall costs for future steam turbines operating attemperatures of 700°C or above.

21.6 References

1 Muehle E E and Ewald J, ‘High-reliability steam turbine components – materialsand strength calculation aspects’, 4th International Conference High TemperatureMaterials for Power Engineering COST-501 and COST-505, Liege, Belgium, 1990,Research Centre Juelich, Germany.

Table 21.5 Materials for power plant components applied at hightemperatures

Temperature range

Component max. 565°C 600–610°C 620–650°C ≥ 700°C

Valve bodies 1–2CrMo steel 9–10Cr steel New Superalloys9–10Cr steel 9–11Cr steel

Rotor 1–2CrMo steel 9–10Cr steel New Superalloys9–11Cr steelNew

Turbine casing 1–2CrMo steel 9–10Cr steel 9–11Cr steel Superalloys

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2 Wiemann W, Ewald J, Niel K and Reiermann D, ‘Influence of microstructure of 1%CrMoV steel on creep and strain controlled fatigue behavior’, International ConferenceLCF of Materials, Munich, 1987, Deutscher Verband fuer Materialpruefung DVM,Germany.

3 Staubli M, Mayer K H, Kern T-U and Vanstone R W, ‘The European joint developmentprogram COST522’, 5th Proceedings International Charles Parsons Conference,Cambridge, UK, 2000, IoM Communication, UK.

4 Kern T-U, Staubli M, Mayer K H, Escher K and Zeiler G, ‘The European efforts indevelopment of new high temperature rotor materials up to 650°C–COST522’, 7thInternational Conference Materials for Advanced Power Engineering, Liege, Belgium,2002, Research Centre Juelich, Germany.

5 Scarlin B, Kern T-U and Staubli M, ‘The European efforts in material developmentfor 650°C USC power plant – COST522’, 4th International Conference Advances inMaterials Technology for Fossil Power Plants, Hilton Head Island, South Carolina,USA, 2004, EPRI, USA.

6 Kern T-U, Staubli M, Mayer K H, Donth B, and Zeiler G, ‘The European efforts indevelopment of new high temperature rotor materials–COST536’, 8th InternationalConference Materials for Advanced Power Engineering, Liege, Belgium, 2006,Research Centre Juelich, Germany.

7 Staubli M, Hanus R, Weber T, Mayer K H and Kern T-U, ‘The European efforts indevelopment of new high temperature casing materials - COST536’, 8th InternationalConference Materials for Advanced Power Engineering, Liege, Belgium, 2006,Research Centre Juelich, Germany.

8 Masuyama F, ‘Alloy development and material issues with increasing steamtemperature’, 4th International Conference Advances in Materials Technologyfor Fossil Power Plants, Hilton Head Island, South Carolina, USA, 2004, EPRI,USA.

9 Mayer K H, Blum R, Hillenbrand P, Kern T-U and Staubli M, ‘Development stepsof new steels for advanced steam power plants’, 7th International Conference Materialsfor Advanced Power Engineering, Liege, Belgium, 2002, Research Centre Juelich,Germany.

10 Igarashi M, Yoshizawa M, Iseda A, Matsuo H and Kan T, ‘Long term creep degradationin 12%Cr ferritic steel tubes and pipes’, International Conference ECCC Creep andFracture in High Temperature Components – Design and Life Assessment, London,UK, 2005, IoM Communication, UK.

11 Iseda A, Yoshizawa M, Igarashi M and Kan T, ‘Long term creep strength degradationin T122/P122 steels for USC power plants’, 8th International Conference Materialsfor Advanced Power Engineering, Liege, Belgium, 2006, Research Centre Juelich,Germany.

12 Schwienheer M, Haase H, Scholz A and Berger C, ‘Long term creep and creep-fatigue properties of the martensitic steels of type (G)X12CrMoWVNbN10-1-1’,7th International Conference Materials for Advanced Power Engineering, Liege,Belgium, 2002, Research Centre Juelich, Germany.

13 Berger C Schwienheer M and Scholz A, ‘Creep and fatigue properties of turbinesteels or application temperatures up to 625°C’, 8th International Conference Materialsfor Advanced Power Engineering, Liege, Belgium, 2006, Research Centre Juelich,Germany.

14 Mueller F, Scholz A and Berger C, ‘Crack behavior of 10Cr steels under creep andcreep-fatigue conditions’, International Conference ECCC Creep and Fracture in

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High Temperature Components – Design and Life Assessment, London, UK, 2005,IoM Communication, UK.

15 Ringel M, Roos E, Maile K and Klenk A, ‘Constitutive equations of adapted complexityfor high temperature loading’, International Conference ECCC Creep and Fracturein High Temperature Components – Design and Life Assessment, London, UK,2005, IoM Communication, UK.

16 Bagaviev A, ‘Steam turbine component integrity analysis based on high temperaturefracture mechanics’, International Conference ECCC Creep and Fracture in HighTemperature Components – Design and Life Assessment, London, UK, 2005, IoMCommunication, UK.

17 Bagaviev A and Sheng S, ‘Industrial application of creep/fatigue crack initiation andgrowth procedures for remaining life analysis of steam turbine components’,International Conference ECCC Creep and Fracture in High Temperature Components– Design and Life Assessment, London, UK, 2005, IoM Communication, UK.

18 Feldmueller A and Kern T-U, ‘Design and materials for modern steam power plants’,Proceedings 5th International Conference Charles Parsons Conference for AdvancedMaterials, Cambridge, UK, 2000, IoM Communication, UK.

19 Gerdes C and Bartsch H, ‘Steam turbine shafts of combined high pressure/lowpressure rotors’, 14th International Conference International Forgemasters Meeting,Wiesbaden, Germany, 2000, German Iron and Steel Institute.

20 Magoshi R, Nakano T, Konishi T, Shige T and Kondo Y, ‘Development and operatingexperience of welded rotors for high-temperature steam turbines’, InternationalConference Joint Power Generation Conference, Miami Beach, Florida, USA, 2000,EPRI, USA.

21 Nakano T, Tanaka K, Nakazawa T and Nishimoto S, ‘Development of large-capacitiysingle-casing reheat steam turbines for single-shaft combined cycle plant’, MitsubishiHeavy Industries Ltd, Technical Review, 2005, 42 (3), 23–29.

22 Magoshi R, Tanaka Y, Nakano T, Konishi T, Nishimoto S, Shige T and Kadoya Y,‘Development of welded rotors for high-temperature steam turbines’, InternationalConference Power Engineering (ICOPE), Chicago, Illinois, USA, 2005, EPRI,USA.

23 Scarlin B, ‘Advanced high-efficiency turbines utilizing improved materials’,International Conference Advanced Steam Plant, IMechE Conference Transactions,1997–2, Steam Plant Committee IMechE, UK.

24 Roberts B W and Vanstone R W, ‘Materials for today’s fossil and nuclear steamturbines’, International Conference Steam Turbine Retrofit Conference, Chicago,USA, 2006, EPRI, USA.

25 Zoerner W, ‘Steam turbines for power plants employing advanced steam conditions’,10th International Conference CEPSI, Christchurch, New Zealand, 1994.

26 Drosdziok A and Feldmueller A, ‘High-efficiency steam turbines for coal-fired powerplants’, International Conference Power Gen Asia, Singapore, 1995.

27 Haas H, Engelke W, Ewald J and Termuehlen H, ‘Turbines for advanced steamconditions’, International Conference American Power Conference, Chicago, Illinois,1982, EPRI, USA.

28 Haas H, Zimmermann A and Termuehlen H, ‘Turbines for advanced steam conditions– operational experience and development’, 1st International Conference ImprovedCoal-Fired Power Plants, Palo Alto, California, USA, 1986, EPRI, USA.

29 Campbell C B, Frank C C and Sphar J C, ‘The Eddystone superpressure unit’, ASMEpaper 56-A-156, 1957.

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30 Silvestri G J Jr, ‘Eddystone station, 325MW generating unit 1 – A brief history’,American Society of Mechanical Engineers, USA, 2003, special issure, 30 pp.

31 Yugami H et al., ‘Operating experience of Wakamatsu high temperature turbine,Step 1’, Mitsubishi Juko Giho, 1990, 27 (1), 3–16.

32 Miyashita K, ‘Overview of advanced steam plant development in Japan’, InternationalConference Advanced Steam Plant, IMechE Conference Transactions, 1997, SteamPlant Committee IMechE, UK.

33 Muramatsu K, ‘Development of ultra-super critical plant in Japan’, InternationalEPRI Conference Advanced Heat Resistant Steels for Power Generation, San Sebastian,Spain, 1998, EPRI, USA.

34 Fukuda M, Tsuda Y, Yamashita K, Shinozaki Y and Takahashi T, ‘Materials anddesign for advanced high temperature steam turbines’, 4th International ConferenceAdvances in Materials Technology for Fossil Power Plants, Hilton Head Island,South Carolina, USA, 2004, EPRI, USA.

35 Vanstone R W, ‘Advanced (700°C) pulverised fuel power plant’, 5th InternationalConference Charles Parsons Conference for Advanced Materials, Cambridge, UK,2000, IoM Communication, UK.

36 Blum R and Hald J, ‘Benefit of advanced steam power plants’, 7th InternationalConference Materials for Advanced Power Engineering, Liege, Belgium 2002,Research Centre Juelich, Germany.

37 Blum R and Vanstone R W, ‘Materials development for boilers and steam turbinesoperating at 700°C’, 6th International Conference Charles Parsons Conference forAdvanced Materials, Dublin, Ireland, 2003, IoM Communication, UK.

38 Blum R, Vanstone R W and Messelier-Gouze C, ‘Materials development for boilersand steam turbines operating at 700°C’, 4th International Conference AdvancedMaterials Technology for Fossil Power Plants, Hilton Head Island, South Carolina,USA, 2004. EPRI, USA.

39 Scarlin B, Vanstone R and Gerdes R, ‘Materials development for ultra-supercriticalsteam turbines’, 4th International Conference Advances in Materials Technology forFossil Power Plants, Hilton Head Island, South Carolina, USA, 2004, EPRI, USA.

40 Kern T-U, Wieghardt K and Kirchner H, ‘Materials and design solutions for advancedsteam power plants’, 4th International Conference Advanced Materials Technologyfor Fossil Power Plants, Hilton Head Island, South Carolina, USA, 2004, EPRI,USA.

41 Blum R and Vanstone R W, ‘Materials development for boilers and steam turbinesoperating at 700°C’, 8th International Conference Materials for Advanced PowerEngineering, Liege, Belgium, 2006, Research Centre Juelich, Germany.

42 Edelmann H, Effert M, Wieghardt K and Kirchner H, ‘The 700°C steam turbinepower plant – status of development and outlook’, International J Energy TechnolPolicy, 2007, in print.

43 Tschaffon H, ‘The European way to 700°C coal fired power plant’, 8th InternationalConference Materials for Advanced Power Engineering, Liege, Belgium, 2006,Research Centre Juelich, Germany.

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22Using creep-resistant steels in nuclear

reactors

S. K. A L B E R T, Indira Gandhi Centre forAtomic Research, India and S. S U N D A R E S A N,

Maharaja Sayajirao University, Baroda, India

22.1 Introduction

Nuclear power plants differ from fossil power plants mainly in the source ofheat for converting water into steam, which is subsequently used to run theturbine and produce electricity. In the former, the source of heat is nuclearfission (or fusion, in future fusion reactors), while in the latter, it is theburning of the fossil fuels like coal, oil or gas. In general, therefore, thestructural materials chosen for nuclear reactors should also meet therequirements of fossil power plants in terms of good creep resistance, oxidationresistance, low-cycle fatigue strength, thermal conductivity, and so on. Inaddition to these, the elements present in the structural materials should alsohave a low neutron absorption cross-section, that is the probability of neutronsproduced in the reactor being absorbed by these elements should be low.Further, the properties of these materials should not degrade under the highlevels of radiation that exist in nuclear reactors. Such degradation is generallyreferred to as radiation damage and includes irradiation embrittlement,irradiation creep, swelling, helium embrittlement, and so on, which aredescribed briefly later in the chapter.

In a nuclear reactor, heat is produced by fission of the heavy elements likeU or Pu, which is achieved by bombardment of the nuclei of these elementsby neutrons. A typical fission reaction is given below:

235U92 + 1n0 → 94Sr38 + 140Xe54 + 2 1n0 + ≈160 MeV

Fission reactions produce, on an average, two or three neutrons per fissionwhich can, in turn, take part in further fission, thus sustaining the reactionand generating energy continuously. However, not all neutrons would beavailable for the fission reaction: some of them would escape from thereactor and be permanently lost, others would be absorbed by the structuralmaterials present, and still others would be simply scattered away by nucleiof various elements present in the nuclear fuel and the structural materials.It is thus important that absorption of neutrons by structural materials should

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be minimized in order to sustain the fission reaction. It is in this context thatthe neutron absorption cross-section of various elements and their isotopespresent in the structural materials becomes important in material selectionfor reactor core application.1 The core is that part of the reactor where thenuclear reaction takes place; the structural materials that contain the fuel andfission products and facilitate the removal of heat from the fission productsare part of the reactor core. The unit of neutron absorption cross-section isthe barn (1 barn = 10–28m2/nucleus)1 and the higher the cross-section thegreater is the probability of neutron absorption. Minimizing the concentrationof elements with a high absorption cross-section is an important criterion inmaterial selection for the reactor core.

The absorption cross-section is a function of the energy of the neutrons.In thermal nuclear reactors which use low-energy thermal neutrons (≤0.025eV) to sustain the fission reactions, structural materials based on metals thathave a low absorption cross-section in this energy range only can be used forthe core applications. There are only a few metals like Zr, Al and Be thathave absorption low enough1 to make them suitable for core application inthe thermal reactors and among them only Zr-based alloys have the requiredmechanical properties. Average neutron absorption cross-section for iron inthis energy range is high (2.55 barn compared with 0.19 barn for Zr) andhence ferrous alloys are not used for core application in thermal reactors.However, in the case of fast breeder reactors (FBRs), the energy of neutronsused for both fission and breeding is significantly higher and in this energyrange the absorption cross-section for iron and most of the major alloyingelements present in the steels is quite low.2 In these reactors, therefore, alloysteels are the main structural materials even for the core applications. Further,steels are also actively considered for core application in future fusion reactors.

22.2 Radiation damage

Radiation damage is a general term employed for material deterioration in aradiation environment and includes radiation swelling, irradiation creep,irradiation embrittlement, helium embrittlement, and so on. It occurs whenhigh-energy particles displace atoms from their normal lattice sites to formFrenkel defects (vacancies and interstitials)3 and is usually expressed asdisplacements per atom or dpa, i.e. the number of times an atom is displacedduring the irradiation. Additionally, neutrons can cause transmutation reactionswith atoms of the irradiated material, resulting in solid or gaseous products.The former are generally considered harmless to material properties,4 butthey could be highly radioactive isotopes with long half-lives. This aspect isrelevant to the development of reduced-activation steels in which elementsthat can produce harmful radioactive isotopes in the reactor environment areeither removed or maintained at very low levels. The gases produced are

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Using creep-resistant steels in nuclear reactors 599

helium (through an (n,α) reaction) and hydrogen (through an (n,β) reaction).Small amounts of helium so produced can result in serious property deteriorationwhich is known as helium embrittlement.

22.2.1 Radiation swelling

Irradiation of some structural materials like austenitic stainless steels attemperatures in the range 0.3–0.5 Tm (Tm is the absolute melting temperature)and with intermediate to high neutron doses produces a significant densitydecrease and volume increase – a phenomenon known as radiation swelling.The point defects that form from displacement events are quite mobile atreactor temperatures and most vacancies and interstitials annihilate themselvesby recombination. The surviving defects migrate to sinks and get absorbed.However, dislocations which are the major sinks for these defects have aslightly higher preference for interstitials and hence over a period of timeexcess vacancies appear in the material that migrate to form clusters. Theseclusters are stabilized as three-dimensional voids by innate gases in thematerial and/or transmuted gas such as helium produced by (n,α) reaction.Voids are not formed below 0.3 Tm because of dominant mutual recombinationof the interstitials and the slow-moving vacancies and above 0.5 Tm becausethermal vacancy concentration then exceeds that induced by irradiation.5

The relationship between neutron dose and void swelling in materials isof much practical interest. It is characterized by an incubation dose belowwhich no swelling occurs. Above this, there is a linear increase in swellingwith increase in dose.6 The extensive experimental data for austenitic stainlesssteels have shown that the magnitude of swelling depends on major andminor elemental contents and the initial thermomechanical treatment. Thesefactors affect only the incubation or transient dose before swelling acceleratesto a constant rate (~ 1% per dpa in austenitic stainless steel) independent ofthe irradiation temperature in the peak swelling range.7

Austenitic stainless steels, presently employed in the fabrication of reactorcore components (clad tubes and wrapper) of FBRs, have poor resistance toradiation swelling. The composition of the conventional austenitic steels likeAISI 316 and AISI 304 and the temperature range and radiation dose inFBRs are such that these steels experience considerable swelling in theoperating environment of these reactors. Swelling in austenitic steelsbecomes unacceptable at dose rates higher than 50 dpa. However, a burn-upof 100 000 MW days/tonne corresponds to 85 dpa and the target burn-uprequired for FBRs to become economically viable is 200000 MW days/tonne, equivalent to fission of 20 atom% of the heavy metal present in thefuel. Hence the desired levels of burn-up cannot be achieved if clad tubesand wrapper made from conventional austenitic stainless steels are used inFBRs. This has led to the modification of both microstructure and composition

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of these alloys to improve the swelling resistance. Meanwhile, it was alsofound that steels with a ferritic/martensitic microstructure are much moreresistant to void swelling than those with an austenitic microstructure andthis resulted in the development of a large number of high-Cr ferritic steelsfor FBR core application.

The basis for the development of austenitic stainless steels with improvedresistance to void swelling has been the fact that void swelling can be reducedif the defects are pinned down by dislocations or solute atoms. As a result ofthese defect–dislocation or defect–solute atom interactions, coalescence ofvacancies to form voids is delayed and recombination of interstitials andvacancies formed owing to neutron bombardment is facilitated. Accordingly,it was found that cold-worked 316 stainless steel exhibits much higher resistanceto void swelling than the solution-annealed material. Further, increasing Niabove the levels present in AISI 316 steel and controlled additions of Ti andSi can also improve void swelling resistance in austenitic steels. Figure 22.1shows the variation with dose of the maximum deformation in the diameterof a fuel pin irradiated in the Phenix reactor for various austenitic claddingmaterials.8 Later it was found that controlled addition of P and trace amountsof B can improve the creep resistance of this class of alloys. These findingsled to the development of a new class of austenitic stainless steels with abase composition of 15Cr–15Ni with controlled additions of Ti and P in theUSA,9 Europe8 and Japan10,11 for use in core components of FBRs. Manycommercial and test reactors have clad tubes and wrapper materials made ofthis class of austenitic stainless steels subjected to an optimum level of 20%

CW 316CW 316 Ti

CW 15.15 Ti

CW Si-mod 15.15 Ti

Dose (dpa)120100806040

Dia

met

ral d

efo

rmat

ion

(%

)

8

7

6

5

4

3

2

1

0

22.1 Variation in the maximum diametral deformation of the fuelpins of different austenitic steels irradiated in Phenix with neutrondose.8

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cold work. It has been found that a burn-up of 100 000 MWd/t can be achievedsafely with core components of this class of alloys.

For burn-up levels higher than the above, one has to depend on ferriticalloys which are much more resistant to void swelling than the austeniticalloys. Several explanations have been put forward for the superior swellingresistance of ferritic steels relative to austenitic steels. They include mechanismsthat depend on solute trapping, the character of the dislocation loop structure,the lower relaxation volume of the body centred cubic (bcc) structure and thetempered martensitic structure.6 However, none of the proposed mechanismscompletely explains the superior swelling resistance of the ferritic alloys3,5,12

and it is felt that a combination of lower dislocation bias (for solute atoms)in bcc alloys, high self-diffusion, low helium generation rates and high subgrainboundary sink strength might be contributing to the higher swelling resistanceof ferritic steels in comparison with the austenitic alloys.

Kim et al.13 compared swelling in type 316LN stainless steel (SS), 9Cr–2WVTa steel and three oxide dispersion strengthened (ODS) steels usingdual beam (3.2 MeV Fe+, 330 keV He+) irradiation to 50 dpa and 260 appmof He at 650°C. The ODS alloys studied were: Fe–17Cr–0.25Y2O3 (17Y3),Fe–12Cr–0.25Y2O3 (12Y1) and Fe–12Cr–3W–0.4Ti–0.25Y2O3 (12YWT).The microstructures of the steels were quite varied, with the SS having veryclean grains and a low dislocation density (ρ<< 1013m m–3) and the 9Cr–2WTa steel having large M23C6 precipitates and a dislocation density of~1013 m m–3. The dislocation density in ODS alloys was typically in therange 1015–1016 m m–3 with 12YWT having a higher dislocation densitythan the others. Size and number densities of the oxide particles variedamong the three ODS alloys. They were 3–5 nm and 1–2 × 1023 m–3,respectively, in 12YWT, 10–40 nm and 1021–1022 m–3 for 12Y1 and a largerparticle size and lower number density in 17Y3 than in 12Y1.

The results showed that a bimodal distribution of bubbles (r < 0.5 nm)and voids (0.5 nm ≤ r < 10 nm) developed in the alloys in decreasing orderof 316LN SS, 9Cr–2WVTa, 17Y3 and 12Y1. In the 12YWT alloy, the bubblesformed were typically 1nm in size and associated with fine particles and ahigh density of dislocations with a number density of ~1023–1024m–3. Theseresults indicated that the more complicated the microstructure, the less is theswelling. The density of available sinks for vacancies and interstitials(dislocations, precipitates, bubble, etc.) determines the recombination of thesedefects and thus the void formation. A change in dislocation bias caused bya change in microstructure could also be important in void formation.Microstructure affects the length of the transient stage (shift from incubationto steady state) of the swelling with irradiation dose. Once sufficient voidsare nucleated, ferritic steels also swell and the steady state swelling rate is~ 0.2%/dpa, which is around one-fifth of that observed in austenitic stainlesssteels.14

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The results of neutron irradiation studies conducted in different FBRs ofvarious countries into swelling in ferritic steels have been reviewed by Kluehand Harries.6 Steels studied included 2.25Cr–1Mo, modified 9Cr–1Mo, HT9(12Cr–1MoVW), EM12(9Cr–2MoVNb), EM10 (9Cr–1Mo), DIN 1.4914 (12CrMoNiVNb), EP-450 (12Cr–1.5MoVNbB) and also some ODS alloys. Therange of temperatures used for irradiation varied from 400–650°C in differentreactors and these alloys were irradiated to different levels ranging from 30–200 dpa. Among the alloys studied, swelling was a minimum in 12Cr steelswith HT9 irradiated to 200 dpa at 420°C exhibiting a swelling of only 1.02%compared to 1.76% of 9Cr–1MoVNb steel. Steels with 9%Cr generallyexhibited more swelling than those with 12%Cr, which is in agreement withthe maximum swelling observed for 9%Cr binary alloys of iron and chromium.The degree of swelling was particularly high in the EM12 alloy, which wasattributed to its duplex microstructure and the presence of δ-ferrite. Further,the swelling resistance of ODS alloys was found to be comparable with thatof conventional alloys. Results also indicated that prior microstructure orheat treatment history also influenced the swelling behaviour of these alloys.

The fusion reactor irradiation environment differs from that of a fissionreactor in the high-energy component in the neutron spectrum. Neutronswith energy up to about 14.1 MeV are present in a fusion reactor, while theneutron energies in a fission reactor are lower than 2 MeV. The higherenergies in a fusion plant imply a higher cross-section for the (n,α) reactionwith elements like Fe, Cr and Ni. Significantly greater amounts of heliumand atomic hydrogen will form in steels exposed to the fusion environment,a consideration that is important in developing steels for the first wall andblanket structures in a fusion power plant.

22.2.2 Irradiation creep

Discovered in the 1950s, first in uranium and a few years later in stainlesssteel, irradiation creep refers to the strain undergone by a material whenirradiated under stress. However, in the absence of either stress or irradiationthe creep becomes insignificant.15 Thermal creep becomes prominent forirradiation at temperatures ≥ 0.5 Tm, while irradiation creep can be significantat much lower temperatures. Similar to the case of thermal creep, dislocationclimb and glide play an important role in the deformation processes thatoccur during irradiation creep.16

Much of the theoretical work on irradiation creep has been concernedwith steady-state point defect concentrations. More recently, it has beenrealised that large contributions to creep can occur as a result of severaltypes of transients in point defect populations.15 The point defects generatedby irradiation can be annihilated through stress-induced processes, for exampleby absorption on dislocations. If the absorption is asymmetric or preferential,

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it can cause the dislocations to climb, which can subsequently lead to glideof the dislocations. It must be noted that, if vacancies and interstitials wereabsorbed equally, annihilation would occur without climb, and there wouldbe no creep.16

The presence of an applied stress can change the capture efficiencies oforiented dislocations for the irradiation-induced point defects. In particular,there is a slight bias for self-interstitials to be absorbed by dislocationswhose Burgers vectors are more nearly aligned with the stress axis. Thiscauses dislocation climb and the different rates of climb for dislocationswith Burgers vectors at different orientations to the applied stress can resultin climb creep,17 which has been termed stress-induced preferential absorption(SIPA) creep.

When gliding dislocations become pinned at obstacles, the bias-driveninterstitials at dislocations can cause them to climb around the obstacles bythe SIPA process, after which the dislocations can glide again under theapplied stress until they are pinned by another obstacle. This results in anincrement of strain, referred to as preferred absorption glide (PAG) creep.

The climb-enabled glide mechanism can also operate as a result of swelling:as cavities grow, dislocations absorb excess interstitials and climb.17 Swelling-driven creep has also been called I-creep. The various mechanisms of irradiationcreep have been reviewed by Mathews and Finnis.18

Irradiation creep can help accommodate stresses induced by differentialswelling. However, excessive irradiation creep can lead to buckling and thestress relaxation can render devices such as clamps and bolts ineffective.15

The initial experimental studies on irradiation creep of ferritic/martensiticsteels showed that at temperatures of the order of 300°C, at which thermalcreep was negligible, considerable irradiation creep could occur. However,the steady-state irradiation creep rate for a 12Cr–MoVNb ferritic/martensiticsteel was nearly an order of magnitude lower than for austenitic stainlesssteels.16 A study of the behaviour of the 12Cr–MoVW steel (HT9) irradiatedat temperatures of 540°C and 595°C19 showed that this steel is superior toseveral austenitic stainless steels, although precipitation hardened Ni-basealloys were significantly more creep resistant. It should be remembered,however, that at these temperatures thermal creep could also occur in additionto irradiation creep.

The creep behaviour of an in-reactor tested specimen in the HT9 steel wascompared by Chin20 with that of a thermal control specimen of the samematerial, as illustrated in in Fig. 22.2. While the steady-state creep rates ofboth the specimens are nearly the same, the difference in total strain arisesfrom the existence of a primary creep stage in the thermal control specimen;irradiation appears to eliminate any detectable primary creep region. Figure22.320 show the variation in the creep coefficient B (which relates creepstrain to creep stress) as a function of temperature. The results illustrate that

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in-reactor creep is insensitive to temperature changes below about 500°C,that is, under conditions in which thermal creep makes an insignificantcontribution. Above 550°C, thermal creep becomes prominent and, as thecreep strengthening mechanisms rapidly deteriorate, large creep rates areobserved.20

While a transient creep stage was not found in Chin’s in-reactor creeptests, the results obtained with 12 Cr–MoVNb steel by Wassilew et al.21 didshow a transient stage. This difference has been attributed to the fact that thefluence range in the latter tests (using uniaxial specimens) was well belowthat in Chin’s pressurized-tube studies. Apparently, the first measurementsin the pressurized-tube tests were made after the primary stage was completed.16

Fluence (1022 neutrons cm2)1211109876543210

In-reactor

Thermal

Time (103 h)20181614121086420

HT-9540°C50 MPa

0.8

0.9

0.7

0.6

0.5

0.4

0.3

0.2

0.1

0.0

ε (%

)

22.2 Comparison of irradiation creep with thermal creep for HT9alloy.20

4.36.710.2

Fluence(1022 n/cm2)

HT-9 STA

Temperature (°C)750650550450350

10

8

6

4

2

0B (

10–2

8 cm

2 /neu

tro

n ×

MP

an)

22.3 Variation of creep coefficient B with temperature for HT9alloy.20

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Irradiation creep studies on several 9Cr–Mo ferritic/martensitic steelshave demonstrated their superiority to austenitic stainless steels such asType 316 and 15-15Ti, especially at temperatures up to about 500°C.22,23

Later work comparing the behaviour of HT9 and T91 showed that modified9Cr–1Mo steel underwent even lower irradiation creep than 12Cr–1MoVW.24

Klueh and Harries16 have estimated the irradiation creep coefficients (expressedin units of MPa–1dpa–1) of a number of Cr–Mo steels and the result showsthe essential similarity in behaviour of this class of ferritic/martensitic steels.The coefficient falls in the range of 1.25–10 × 10–7 MPa–1dpa–1 for thedifferent Cr–Mo steels that were examined.

It is of some interest to compare the performance of high-Cr ferritic/martensitic steels with that of ODS ferritic steels. The latter are candidatematerials for fuel cladding for fast reactors. Comparative tests under similarstress, temperature and irradiation conditions have shown that the creepresistance of the ODS steels was superior by a factor of 2–5 times. The Cr–Mo steels, however, performed 3–4 times better than Type 316 austeniticstainless steel.16 If the out-of-pile creep component is subtracted from the in-pile creep component, the true irradiation creep behaviour is revealed. Suchan analysis shows that, in conventional Cr–Mo steels, the total creep increasescontinuously with increasing temperature, since above 450°C thermal creepis dominant. With the ODS steels, on the other hand, irradiation creep ratedecreases above 450°C and, since thermal creep is also low in these steels,the total creep rate actually decreases with increasing temperature above450°C. The superior irradiation creep behaviour of the ODS steels might bebased on the presence of a large number of dispersed particle interfaceswhich act as sinks for the point defects and thus reduce the fraction going tothe dislocations to cause irradiation creep.16

22.2.3 Irradiation embrittlement

An area of important concern in the use of ferritic/martensitic steels in lightwater reactors, fast reactors and future fusion reactors is the adverse effect ofirradiation on fracture toughness.25 Irradiation is known to cause a largeincrease in the ductile-to-brittle transition temperature (DBTT) and a significantreduction in the upper shelf energy (USE). The practical implication is thatthe post-irradiation DBTT may rise above ambient temperature, although itmight have been well below room temperature prior to irradiation.

Irradiation embrittlement of ferritic steels is related to the lattice hardeningcaused by precipitation, induced or accelerated by the irradiation. Theseeffects arise from irradiation below about 0.4 Tm, where Tm is the absolutemelting temperature of the steel (Tm ≈ 1500°C). Hardening raises the flowstress and, if the fracture stress is assumed to be unaffected by irradiation,the DBTT is increased.25

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On account of their relatively superior void swelling resistance, the ferritic/martensitic steels are potential candidate materials for the first wall andblanket structures of future magnetic fusion reactors. The first wall of afusion reactor will experience displacement damage from the high-energyneutrons of the fusion reaction. Additionally, large amounts of transmutationhelium will form: for one year of irradiation of a martensitic steel, about 110appm (1 appm = 1 atom per million) of He is produced in a fusion reactorfirst wall compared with < 10 appm He in a fast reactor.26 Thus, the irradiationconditions expected in a fusion reactor cannot be simulated in a fast reactor.Investigators have therefore attempted to produce simultaneously displacementdamage and transmutational helium by adding nickel to the steel under studyand irradiating it in a mixed-spectrum reactor, such as the high-flux isotopereactor (HFIR). The fast neutrons in the spectrum cause displacement damage,while the thermal neutrons produce helium from 58Ni in a two-steptransmutation reaction 58Ni (n, γ) 59Ni followed by 59Ni (n, α) 56Fe. Thistechnique has been used to investigate the effect of helium on mechanicalproperties and swelling resistance.26

In the study of irradiation effects on impact behaviour, the limited irradiationspace available in most of the fission test reactors has dictated the use ofsmall-size Charpy specimens 1/2 and 1/3 the standard size.27 Such tests haveshown that the shifts in DBTT have been greatest at the lowest irradiationtemperature, which is usually below 400°C.28 Below 400–500°C, the magnitudeof the shift caused by displacement damage varies inversely with irradiationtemperature.25 Most irradiation embrittlement studies have been conductedin the temperature range 360–600°C.

Although the Charpy curves are shifted by irradiation hardening, the fracturemode is generally unaltered between the unirradiated and irradiated steels:cleavage or quasi-cleavage on the lower shelf and ductile void coalescenceon the upper shelf.25 However, in the situation where helium is produced bytransmutation, features associated with intergranular fracture have beenobserved. This aspect is discussed subsequently.

Investigations of high-chromium ferritic steels in fast-neutron environmentshave indicated that the shifts in DBTT tend to saturate with increase influence, as illustrated in Fig. 22.4.27 The hardening caused by irradiationalso tends to saturate in a similar manner. The saturation effect is, however,not observed in irradiation studies in mixed-spectrum reactors. This has beenattributed to the larger amount of helium generated by irradiation inenvironments containing thermal neutrons. Only small quantities of heliumare produced in fast reactors, whereas considerably larger amounts are formedin mixed-neutron environments.27 Similarly, a study using spallation neutronsources has also demonstrated the absence of the saturation effect in T91steel.29 The data showed that the DBTT increases continuously with dosagein the range investigated (0–9.4 dpa) as shown in Fig. 22.5.29 At the highest

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Normalized and tempered10 dpa, 365°C, FFTF17 dpa, 365°C, FFTF

12 Cr-1 MoVW Steel

Temperature (°C)3002001000–100–200

40

30

20

10

0

En

erg

y (J

)

22.4 Charpy curves for half-size specimens of 12Cr–1MoVW steel inthe normalized and tempered condition and after irradiation to 10and 17 dpa at 365°C in FFTF.27

DBTTSP USESP

Displacement (dpa)1086420

0.00

0.05

0.10

0.15

0.20

0.25U

SE

SP (

J)

DB

TT

SP (

°C)

–150

–190

–120

–90

–60

–30

22.5 Irradiation dose dependence of the DBTT from small punch (SP)test.29 USE is the upper-shelf energy.

dosage of 9.4 dpa, the increase in DBTT was as high as 295°C, when theshift was recalculated to standard Charpy conditions. The large increase inDBTT was also associated with intergranular fracture. These effects havealso been ascribed to the large quantities of helium produced during irradiation

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with spallation sources. In fact, the increase in DBTT showed a lineardependence on helium content, see Fig. 22.6.29

The microstructural changes caused by irradiation occur primarily overthe range 300–500°C and almost no change occurs at 600°C.30 Dependingupon the temperature of irradiation, some recovery of the as-tempered, cellularsubgrain structure might occur, as during long-term annealing. However, ofgreater relevance to the embrittlement process is the production of dislocationloops from the irradiation-induced point defects and, at temperatures of400°C and above, the formation of voids and micropores (which are presumablyhelium bubbles) along prior-austenite and lath boundaries.21,30 No newprecipitation was observed in T91 after irradiation, but there were changes inthe as-tempered carbide distribution and elemental concentration.30 However,in several other Cr–Mo steels, extensive precipitation has been observedover a range of irradiation temperatures, for example, the G phase observedin HT-9 and identified to be a nickel silicide, α′ and Mo2C.30 Such precipitationhas been considered to be related to the observed delays in the onset ofswelling in irradiated ferritic alloys, although it is also held to be responsiblefor irradiation hardening and embrittlement.31

These microstructural changes result in hardening and embrittlementdepending upon the temperature of irradiation. Investigations of several Cr–Mo steels have demonstrated that the degree of embrittlement is considerableand is a matter for serious concern. Irradiation of 12Cr–1MoVW steel in afast-flux test facility (FFTF) to 10 and 17 dpa at 365°C confirmed the saturationeffect under fast reactor irradiation. The same investigation also showed thatthe increase in DBTT in 2.25Cr–1Mo steel after irradiation under identicalconditions in the FFTF was similar in magnitude to the increase observed for

Slope = 0.135°C/appm He

Helium content (appm)800700600500400300200100

125

100

75

50

25

∆DB

TT

SP (

°C)

22.6 Helium content dependence of ∆DBTT on T91 steel obtained insmall punch test.29

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12Cr–1MoVW steel. However the final DBTT for the 2.25 Cr–1Mo steelwas lower than that for 12Cr–1MoVW because of the lower starting DBTTfor the 2.25 Cr–1Mo material.27

Investigations of 12Cr–1MoVW and 9Cr–1MoVNb irradiated in a fastreactor have revealed that, while the DBTT shifts tend to saturate withincreasing fluence, the hardening and embrittlement are significantly decreasedat irradiation temperatures of 450°C and above.32

Klueh and Alexander irradiated specimens of 9Cr–1MoVNb (T91) and12Cr–1MoVW (HT9) steels with and without 2% Ni addition in the mixed-spectrum high-flux isotope reactor (HFIR) at 300 and 400°C.33 While nosaturation was observed up to a displacement of ~ 42 dpa, the shifts in DBTTmeasured after irradiation at 400°C were the largest ever observed for thesesteels. The Ni-added steels exhibited a greater effect than those withoutnickel enhancement. Also, the shifts were much larger after irradiation in amixed-spectrum reactor than when the same steels were irradiated in a fastreactor, where little helium formed during irradiation.32

Many attempts have been made to study the possibility of improving pre-irradiation toughness by optimizing prior heat treatment, minor impurityconcentration and fabrication processes without significantly affecting thestrength, swelling resistance and creep resistance. Work on 9Cr–1MoVNbsteel has demonstrated that heat treatment can be optimized to reduce prior-austenite grain size and lower the pre-irradiation DBTT.34 The shift in DBTTowing to irradiation was, however, not affected by the prior heat treatment.This means that such optimized heat treatment can be used to ameliorate theeffect of irradiation on DBTT in 9Cr–1MoVNb steel. In the case of 12Cr–1MoVW steel, however, the same investigation showed that austenitizationtemperature and austenite grain size had little effect on the unirradiatedDBTT. The implication is that prior heat treatment cannot be used as ameans to offset the DBTT shift caused by irradiation in 12Cr–1MoVW. It isof interest to mention here that in the same steel the use of a relatively highaustenitization temperature (to dissolve all carbides and δ-ferrite present) inconjunction with a moderate temperature age at 700°C has been reported tocause a 50°C improvement in the pre-irradiation DBTT.35

The embrittlement caused by irradiation has also been studied in thecase of weldments. The results, however, have not been consistent. Hu andGelles28 found with 12Cr–1MoVW steel that the weld and heat-affectedzone samples exhibited the same or slightly lower DBTT and similar uppershelf energy (USE) as the base metal. This would mean that ferritic/martentisitic steels are not expected to suffer any penalty in performancecaused by the presence of welds. However, the same investigators found ina subsequent study on 9Cr–1Mo steel that the DBTT for the weldment wasabout 60°C higher than for the base metal exposed to the same irradiationconditions.32

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During the development of reduced-activation steels, systematic studiesof the effects of various elements were made. An important result was thatthe boron content seemed to be a controlling factor for DBTT. The 10Bisotope (20% of the natural composition) is a strong thermal neutron absorberand transmutes to He and Li even at moderate neutron fluxes. Helium bubbleshave been detected on the crack surfaces of irradiated material. The datashowed that the higher the boron concentration the steeper was the slope ofthe ∆DBTT versus irradiation dose curve. The boron effect overrode allother compositional factors.36 Another study with reduced-activation steelsshowed that, while the presence of boron generated He during irradiation inall the steels investigated, significant shifts in DBTT resulted only in somecases.36 In general, such studies demonstrate that at least with low neutronfluxes, the advanced reduced-activation alloys provide much improved impactproperties after irradiation.37

22.2.4 Helium embrittlement

In addition to the irradiation hardening and embrittlement taking place atirradiation temperatures below about 0.4 Tm , another form of embrittlementis exhibited at elevated temperatures higher than about 0.5 Tm followingirradiation at ambient (or elevated) temperature. Such high-temperatureembrittlement is manifested by austenitic stainless steels and some othermaterials and has been attributed to the presence of helium gas generatedduring irradiation. The subject has been comprehensively reviewed by Kluehand Harries.38

The production of helium is attributed to nuclear reactions involving Band Ni bombarded by thermal neutrons and Ni, Cr and Fe by fast neutronsin the reactor environment. Helium bubbles forming at grain boundaries ofthe irradiated austenitic stainless steels are believed to nucleate cavities thatgrow under stress and coalesce to form intergranular cracks in a manneranalogous to cavity growth during thermal creep. Although several solutionshave been suggested for mitigating the problem, such as grain refinement,trapping of helium at particles inside the grains and promoting grain boundaryprecipitation to reduce the production of vacancies, a fully satisfactory answerto the problem has yet to be found.38

In addition to the austenitic stainless steels, ductility reductions in elevated-temperature tensile testing have been observed in several nickel-base alloys,a niobium alloy and, to a lesser extent, in ferritic stainless steels. In contrast,however, the high-chromium Cr–Mo ferritic/martensitic steels have beenfound to be much more resistant to this form of embrittlement.

In a relatively early investigation,39 nickel-doped 9Cr–1MoVNb and 12Cr–1MoVW steels were irradiated in a mixed–spectrum reactor, which generatedup to 49 appm of helium, but subsequent tensile testing at 700°C showed no

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detrimental effects of irradiation or helium concentration on ductility. On theother hand, cold-worked Type 316 austenitic stainless steel exhibited asignificant reduction in ductility on similar irradiation and testing.

Helium implantation in a compact cyclotron with a beam of α-particleshas enabled helium concentrations of up to 3000 appm to be developed in ahigh-Cr ferritic/ martensitic steel corresponding to DIN 1.4914.40 However,tensile testing at temperatures up to 750°C revealed no helium embrittlement.In-beam creep tests at 600 and 700°C also showed no embrittlement, thefracture mode remaining transgranular except that, at the highest heliumconcentration and test temperature, portions of mixed trans- and intergranularfracture were observed. Helium implantation studies on 9Cr–0.5V and 9Cr–2W reduced-activation steels confirmed the resistance of high-Cr martensiticsteels to elevated temperature embrittlement, the tensile fracture always beingductile and transgranular.41

The relative immunity of ferritic/martensitic steels to such embrittlementis believed to be related to their lath microstructure containing fine carbideprecipitates in the ferrite matrix and on the low-angle lath boundaries. It ispostulated that such a structure enables the partitioning of helium atoms tothe lath and precipitate boundaries, so that the helium concentration on anyhigh-angle grain boundaries is relatively low.39 Microstructural evidencesupporting the above reasoning has been provided by subsequent investigationson high-Cr martensitic steels, which showed helium bubbles forming on thelath boundaries, sub-boundaries and dislocations in the laths, but no preferentialbubble growth at prior-austenite grain boundaries that could cause intergranularfracture.40,41

22.3 Embrittlement caused by ageing

An important limitation of high-Cr ferritic/martensitic steels is their marginalfracture toughness and their transition from ductile to brittle behaviour at atemperature not much below 0°C. A matter for further concern to that thetoughness tends to deteriorate during long-term exposure to elevatedtemperature and/or aggressive environments.

Since the presence of impurities can adversely affect toughness, it isnecessary to impose tight specifications to control impurity concentration.Investigations of the impact toughness of 9Cr–1Mo steels have shown thatphosphorus and silicon raise the DBTT and decrease the upper-shelf energy(USE) as a consequence of segregation effects and increase in the delta-ferrite content, respectively. The effect of sulphur is more pronounced thanthat of phosphorus and has been attributed to the formation of non-metallicinclusions which reduce toughness.42

Several studies have shown that in 9Cr–1Mo steels a loss of roomtemperature ductility and toughness and an increase in DBTT can occur as

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a result of thermal exposure in the range 500–600°C. Much of the mechanicalproperty deterioration that takes place during long-term thermal ageing canbe related to microstructural changes arising from such exposure. The pre-tempered material exhibits a tempered martensitic microstructure consistingof lath-shaped ferrite subgrains within prior-austenite grain boundaries, thesubgrains containing a concentration of dislocation networks. Precipitationis generally distributed along grain and subgrain boundaries, consisting ofcoarse M23C6-type and finer MC-type carbides.

Microstructural changes observed during long-term ageing (typically, upto 25 000 h) of 9Cr–1MoVNb steel can be separated distinctly into twotemperature regimes of quite different ageing effects.30,43 At lower ageingtemperatures in the range 482–593°C, little change occurs in lath/boundaryand carbide precipitate structures, except for a transient increase in dislocationdensity within the subgrains. However, the more important effect of lowertemperature ageing is the formation of significant amounts of the Fe2Mo-type Laves phase, which occurs to a great extent in steels containing a higheramount of silicon.30,43 The Laves phase forms preferentially at the lath andgrain boundaries. Abundant, fine vanadium carbide (VC) needles also format temperatures of 593°C and below. In contrast, during ageing at temperaturesin the range 593–704°C, there is detectable recovery and coarsening of thelath/subgrain boundary structure into equiaxed grains and coarsening of thecarbide precipitates, but new precipitation seems to be much reduced, especiallyin low-silicon steels.43

The DBTT increases sharply and USE decreases on ageing at temperaturesup to 593°C, reaching a maximum and a minimum, respectively, after25 000 h of thermal exposure. It is considered very likely that precipitation,especially of the intermetallic compound Laves phase, is responsible for theproperty deterioration, although the exact mechanism has not been identified.The amount of precipitates and relative amount of the Laves phase reach amaximum after ageing at 538°C, corresponding well with the maximumshift in DBTT. At ageing temperatures of 650°C and above, structuralcoarsening softens the material and the USE increases rapidly.

In low-alloy steels, a process of temper embrittlement is known to occurby the segregation of residual impurity elements P, Sn and Sb to the prior-austenite grain boundaries. High-Cr ferritic/martensitic steels also exhibitthis type of embrittlement on ageing, especially during short-term ageing.44,45

It has been shown that, in a high-phosphorus (>0.017%) 9Cr–1Mo steel,ageing at 550°C for 1000 h leads to significant reduction in ductility andincrease in DBTT, which is believed to be associated with the segregation ofphosphorus to the carbide–matrix interfaces at which de-cohesion occurs.For longer ageing durations of 5000 h and above, the formation of the Lavesphase becomes the dominant effect, leading to shifts in the DBTT of morethan 60°C.44 Even in such cases, it has been proposed that phosphorus could

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play a secondary role by being absorbed in the Laves phase and furtherembrittling it.45

Since conditions promoting ageing embrittlement and hydrogen embrittlementmay co-exist during the operation of a fusion reactor, the synergism betweenthe two has been studied in a 9Cr–1Mo steel.45 This investigation showed thata strong interaction existed between hydrogen embrittlement and temperembrittlement. Both acting together resulted in a greater degree of embrittlement,in which the dominant factor controlling the interaction was the precipitationof the Laves phase containing phosphorus in solution.

22.4 Use of heat-resistant steels in major reactor

types

Thermal reactors, which use low-energy thermal neutrons to sustain nuclearfission, are the most widely used reactors in nuclear energy production.There are different variants of the thermal reactors depending on fuel, moderator(used to slow down the high-energy neutrons produced in the reactor) andthe fluid used for the removal of the heat. They include boiling water reactors(BWR), pressurized water reactors (PWR), pressurized heavy water reactors(PHWR), Russian VVER type reactors. In all these reactors, the core structuralmaterials (used to contain fuel and extract heat from the fuel) are Zr-basedalloys. However, the reactor vessel, heat transport system, steam generators,and so on employ different classes of steels. As the steam temperature istypically below 300°C, creep-resistant steels are not extensively used in themajor components of these reactors. These are employed only for thosecomponents dealing with steam, like steam piping and turbines. Austeniticstainless steels find application as piping material, especially in BWRs,because of their corrosion resistance. In PHWRs, austenitic stainless steelhas replaced carbon steel in recent reactors as the material of constructionfor calendria, a major reactor component supporting the coolant tube, becauseof the embrittlement of the latter in service.

High temperature gas cooled reactors (HTGR) are a special case of thermalreactors that operate at relatively high temperatures. They use gases likehelium or carbon dioxide for heat transfer and graphite to moderate theneutron energy. In one HTGR type known as advanced gas cooled reactors(AGR), uranium oxide clad with stainless steel is used as fuel. As the coolanttemperature could be as high as 687°C, superalloys are used for the structuralcomponents of the reactor. Steam temperatures are lower at ~540°C.46

22.4.1 Fast breeder reactors

Fast breeder reactors use fast neutrons to sustain the fission reaction as wellas breeding. The non-fissile isotopes 238U and 232Th are converted to fissile

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isotopes of 239Pu and 233U, respectively, thus producing fresh fuel duringreactor operation. Use of such reactors as a source of energy is being activelypursued in India, Russia, France and Japan and these reactors are expectedto play a major role in the revival of the nuclear industry worldwide. Sincehydrogen nuclei present in water, the normally used medium to extract heat,absorb and slow down the neutrons, water or any hydrogen-containing fluidsis not suitable for extracting heat from these reactors. Hence, liquid metals(predominantly sodium) are used for extraction of heat produced from fission.Subsequently, in the secondary side of the reactor, liquid metal transfers theheat to water to produce steam which finally drives the turbine. Hence,compatibility with liquid metals is a major consideration in the selection ofthe structural materials for this class of reactors. A heat transport flow sheetfor a prototype fast breeder reactor (PFBR) that is under construction isshown in Fig. 22.7.47

Fast breeder reactors (FBRs) use heat-resistant steels extensively both inthe reactor core as well as in the conventional steam side of the reactor. Asalready mentioned, in the fast neutron spectrum, the neutron absorptioncross-sections for Fe and other alloying elements are very small and hence

2 Loops

Surge Tank613 K

753 K

3.2MPa

17.2MPa

795 K

Steamgenerator

Secondarysodiumpump

628

Air

Sodium

820 K

Core

PSP IHX

Reheater

To feed waterheaters

Turbine0.01MPa

HP IP LP 500 MWt

Generator TR

SEA

CEP 301 K– 305 K

HP heaters Deaerator

BFP LP heaters

22.7 Schematic heat transport flow sheet of India’s PFBR.PSP – Primary Sodium pump; IHX – Intermediate heat exchanger;HP – High pressure; LP – Low pressure; IP – Intermediate pressure;TR – Transmission; BFP – Boiler feed pump; CEP – condenserextraction pump.

4 loops6 LMW

Condenser

670 K

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neutron absorption is not a concern in the selection of the materials. Further,the normal operating temperature range for these reactors is typically~ 600°C – a range that is ideally suited to the use of heat-resistant (bothaustenitic stainless and high-Cr ferritic) steels.

Regarding the choice of materials, the major reactor components of FBRscan be classified into core components, structural components like the reactorvessel, safety vessels, piping, etc. and steam generators. Most of the FBRsfor which construction is completed and those which are under constructionuse austenitic stainless steels for the first two sets of components. However,for steam generators, ferritic steels are more appropriate. In the immediatefuture, ferritic steels might even be chosen even for core components.

22.4.2 Austenitic steels in FBRs

Core components

Although, for the reactors currently in operation and under construction,both core components and other reactor parts, except steam generators, aremade predominantly from austenitic stainless steel, the steels chosen forcore components differ from those chosen for other components. Clad tubeswhich contain the fuel and the sub-assembly wrapper which houses the cladtube are the major components of the reactor core. Figure 22.8 shows aschematic of a fuel sub-assembly and clad tube for India’s PFBR programme.These components are exposed to intense neutron irradiation which can inturn result in serious problems like void swelling, helium embrittlement andirradiation creep. Deformation of the sub-assembly components occurs owingto void swelling and creep induced by irradiation and internal sodium pressure.In fact, dimensional changes associated with void swelling limit the residenttime of these components in the reactor. The fuel clad tubes experiencetemperatures in the range of 400–700°C under steady-state conditions. Thewrapper tubes operate in the range 400–600°C. Besides adequate tensilestrength, ductility and creep strength, compatibility with liquid sodium whichis used as the coolant is also an important material property requirement.

The first generation materials for fast reactor core components belongedto Type 304 and Type 316 austenitic stainless steels, especially the latter.47

These steels were soon found to be inadequate because of unacceptableswelling at doses higher than 50 dpa and the first attempt was to introducecold work to improve the resistance to swelling. Subsequently, the alloycomposition was modified, as already reported, by increasing the nickelcontent, lowering the chromium content, raising the phosphorus level andadding titanium. It was found that there is a synergistic effect between titaniumand carbon and this has been related to the amount of free carbon availablefor trapping vacancies, the effect being smallest if titanium is totally absent

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E E

V V

H H

3071

010

300

CD

1000

(A

ctiv

e co

re)

2580

300

520

025

CD

Fuel subassembly

Cladtube

Section XX

Discriminator

Head

Adaptor

B4Cshielding

Steelshielding

Topblanket

Fuel

Bottomblanket

Coolantentry tube

608

4500

Wid

th A

/F 1

31.3

7 shield pins, OD 36±0.15, ID33.4,±0.15 pellet stack 03110.1

Section-MM

210 FUEL PMS

X X

Section-EE

Section-BB

Section-HH

22.8 Schematic of fuel sub-assembly and clad tube for a prototypefast breeder reactor.

Clad OD-LB10.00

0.4710.02PalletOD-LB10.00

0.1310.3

Spacer wire51.4510.00

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or if all the carbon atoms are bound by the titanium. On this basis, the needto maintain the optimum Ti/C ratio was identified.48 The beneficial effect ofphosphorus, which delays significantly the irradiation-induced swelling andcreep, is attributed to several possible mechanisms: strong binding betweenphosphorus in solution and vacancies, phosphide precipitates acting as trapsfor helium atoms, and so on.49 A small addition of boron is made in some ofthese grades to improve creep ductility.47 The composition of some of thecommercially available austenitic stainless steels developed for FBR corecomponents is given in Table 22.1.

It is necessary to know how the addition of elements such as phosphorusand titanium, done to increase the resistance to irradiation effects, influencesthe creep properties, both in the reactor environment and outside. Phosphorusadditions to Type 316 SS have been shown to increase the thermal creeprupture strength at 650°C, but not at 750°C.50,51 Similarly, titanium additionto a 15 Cr–15 Ni stainless steel was found to improve thermal creep behaviour,the effect reaching a maximum in the range 0.2–0.3%Ti. However, silicon isreported to exercise an adverse effect on the creep performance of cold-worked 15Cr–15Ni–Mo–Ti steel owing to accelerated precipitate coarseningeffects.

Another systematic study on the influence of Ti, Si, and P on thermalcreep properties has been carried out during the development of D9I in theUSA.49 This mainly involved adding controlled amounts of phosphorus andboron to a base D9 composition and optimizing their levels. The investigationclearly showed that both these elements increased rupture life at a testtemperature of 700°C. The creep ductility was also found to exceed 18% inall the compositions studied. Based on the results of this work, the chemicalcomposition of D9 was modified and the new composition was designated asD9I, as given in Table 22.1. Long-term biaxial creep rupture tests performedin the temperature range 650–760°C have confirmed the superior creepbehaviour of D9I in relation to D9. The in-reactor creep strength of D9 andD9I is compared in Fig. 22.9.52 It can be seen that the improvement in creeprupture life observed in out-of-reactor tests with controlled additions ofphosphorus and boron is also maintained during irradiation.

Structural components

Austenitic stainless steels with composition slightly modified from the standardAISI 304 or 316 steels are the major structural materials for most FBRcomponents (other than core components and steam generators). Thetemperatures during operation of these components range from 400°C to550°C and the environment is liquid sodium or argon with sodium vapour ornitrogen gas depending on the location of the component. Except for theregions near the core, irradiation is not a matter of serious concern. These

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618

Table 22.1 Nominal chemical compositions of some currently used austenitic steels for FBR core applications

Alloy Cr Ni Mo Mn Si C Ti P Others Fe

Type 316 18.0 13.0 2.6 1.9 0.80 0.05 0.05 Balance

D9 (USA) 13.8 15.2 1.5 1.7 0.9 0.052 0.23 0.003 – Balance

D9I(USA) 13.5 15.0 1.8 1.9 0.8 0.043 0.26 0.030 0.005 B Balance

15–15Ti (France) 14.7 14.7 1.2 1.6 0.43 0.096 0.43 0.007 – Balance

Mod. 15–15Ti (France) 14.9 14.8 1.5 1.5 0.9 0.085 0.50 0.007 – Balance

PNC 316 (Japan) 16.0 14.0 2.5 1.8 0.8 0.055 0.10 0.028 0.08 Nb Balance

PNC 1520 (Japan) 15.0 20.0 2.5 1.9 0.8 0.06 0.25 0.025 0.11 Nb Balance

DIN 1.4970 15.1 15.1 1.3 1.3 0.5 0.088 0.48 0.004 – Balance

D9 (India) 14.0 15.0 2.0 2.0 0.4 0.04 6×%C 0.02 0.05 Nb Balance

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steels are chosen because of their high-temperature properties, compatibilitywith liquid sodium, creep strength, weldability, availability of design dataand long experience in their use of these steels in the application temperaturerange in other industries. Low-cycle fatigue and creep-fatigue interactionsare important considerations in the selection of these materials.

Reactor designers prefer monometallic construction for liquid sodiumsystems for fear of migration of interstitial elements like carbon throughliquid sodium owing to differences in thermodynamic activity in bimetallicsystems. Austenitic stainless steels are thus used in the entire liquid sodiumsystem even if the temperatures of some components are low enough topermit the use of less expensive ferritic steels.47

Austenitic stainless steels like 304, 316, 321, 347, 304L and 316L havebeen used in various reactor components in the reactors built in differentcountries in the early stages of their FBR programmes, but the current trendis to use low-carbon, nitrogen-added 316 or 304 generally designated as316L(N) or 304L(N). Low-carbon varieties are chosen to minimize the riskof sensitization. Nitrogen is added to compensate for the loss of strengthcaused by reduction in carbon content. Compositions of different 316L(N)steels considered for the European Fast Reactor (EFR), the DemonstrationFast Breeder Reactor (DFBR), the Superphenix and Prototype Fast BreederReactor (PFBR) are given in Table 22.2. These steels have been extensivelystudied for their properties and compatibility with liquid sodium. Figures22.10 and 22.11 show a comparison of creep rupture strength and fatiguebehaviour, respectively, of various 316L(N) steels.47,53,54

575°C605°C670°C750°C

575°C630°C695°C775°C

D91D9 D91

D9

LMP, T(13.5 + log tp) × 10–3 (K-h)2018161412

1000

100

10

Ho

op

str

ess

(MP

a)

22.9 Comparison of in-reactor creep strength of D9 and D9I alloys.52

LMP is the Larson–Miller Parameter, temperature in Kelvin and timein h.

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Table 22.2 Chemical composition of major austenitic stainless steel structuralmaterials used in different FBRs (wt%)47

Element 316L(N) SS 316FR 316L(N ) SS 316L(N) SS(EFR) (DFBR) (Superphenix) PFBR

C 0.03 0.02 0.03 0.024–0.03Cr 17–18 16–18 17–18 17–18Ni 12–12.5 10–14 11.5–12.5 12–12.5Mo 2.3–2.7 2–3 2.3–2.7 2.3–2.7N 0.06–0.08 0.06–0.12 0.06–0.08 0.06–0.08Mn 1.6–2.0 2.0 1.6–2.0 1.6–2.0Si 0.5 1.0 0.5 0.5P 0.025 0.015–0.04 0.035 0.03S 0.005–0.01 0.03 0.025 0.01Ti NS* NS 0.05 0.05Nb NS NS 0.05 0.05Cu 0.3 NS 1.0 1.0Co 0.25 0.25 0.25 0.25B (ppm) 10–20 10 15–35 20Nb+Ta+Ti 0.15 NS NS NS

NS, not specified

316L(N) – Superphenix, France316FR – DFBR, Japan316L(N) – Germany316L(N) – PFBR, India316 – ORNL, USA316L(N) – RCC-MR design curve

Rupture time (h)105104103102101

500

400

300

200

100

Str

ess

(MP

a)

22.10 Comparison of creep rupture strengths of 316 and 316L(N)stainless steels at 873 K from various countries.47

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22.4.3 Ferritic steels for FBR core application

Though at present no commercial FBR is operating with a ferritic steel core,these steels have been extensively tested for this application in various reactorslike FFTF in the USA,55 Phenix in France,56,57 PFR in the UK, and BN-350and BN-600 in Russia.58,59 Extensive out-of-reactor irradiation studies havealso been carried out on ferritic steels. The incentive for the use of ferriticsteels is the high burn-up of fuel that can be achieved in reactors with suchcores, which would in turn reduce the fuel-cycle cost considerably. With theaustenitic stainless steels, which are the presently used clad and wrappermaterials in FBR, the burn-up that can be targeted is ~100 000 MW days/tonne, which corresponds to a neutron dose of 85 dpa and a life of ~ 2 yearsfor the fuel sub-assembly. The life is limited by the radiation swelling thatthe austenitic stainless steel undergoes at these dose rates. With the use offerritic steels, on the other hand, there is virtually no radiation swelling andhence the target burn-up can be at least doubled.

Though swelling resistance of ferritic steels is superior to that of austeniticstainless steels, the rise in the DBTT under irradiation, poor formability andlow high-temperature strength limit the choice of ferritic steels for FBR coreapplication. There is also concern with respect to transfer of carbon from theferritic steel side (high activity of carbon) to the austenitic steel (low activityof carbon) with flowing liquid sodium acting as transfer medium, which canfurther reduce the strength of the ferritic steel. However, the recent developmentof many advanced ferritic steels, some of them specifically for nuclear

Japan type 316FR, 1 × 10–3 s–1, 50 mm plateCRNL type 316 FR, 1 × 10–3 s–1 or higher, 50 mm plateUS type 316, 4 × 10–3 s–1, 538–566°C

Cycles to failure107106105104103102101

550°C

10–1

100

101To

tal s

trai

n r

ang

e (%

)

22.11 Comparison of fatigue behaviour of 316 and 316FR stainlesssteels.54

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applications, and extensive in-reactor testing of these steels have providedenough confidence for the use of these materials in the core components ofthe FBRs. Table 22.3 gives the chemical composition of various ferriticsteels considered for the core components of FBRs.

Swelling behaviour and irradiation embrittlement of different ferritic steelshave already been discussed in the previous sections. They indicate thatswelling is lowest for steels with 12%Cr, while shift in DBTT is minimumfor steels with 9%Cr. Figure 22.12 shows the influence of chromium contenton DBTT shift in ferritic/martenistic steels by neutron irradiation.60 However,the swelling resistance of both 9%Cr and 12%Cr ferritic steels is far lowerthan that of austenitic stainless steels. Hence, for wrapper application, 9%Crsteels seem to be more appropriate than 12%Cr steels. Since the DBTT of9%Cr steels is below ambient temperature even after irradiation, it willensure that, during handling of the fuel sub-assembly removed from thereactor at ambient temperature, the risk of fuel pins coming out owing tofailure of the wrapper will be minimum.

Material selection for the clad tube is more challenging than for the wrapper.The operating temperature of the clad tube is high (600–700°C), with sometransient that can go even higher, though the neutron dose experienced iscomparable to that in the wrapper. Ferritic steels made through the conventionalprocessing route of melting, casting and forming do not possess adequatecreep resistance in this temperature range and hence are not ideal replacementsfor the presently used austenitic alloys. This has led to the development of anew class of ODS ferritic/martensitic steels for use as clad material. Y2O3 isthe most widely used type of oxide particles added in these alloys and theyreact with Ti present in the steel to form complex oxides of yttrium andtitanium. Ukai et al.61 studied the creep properties of ODS alloys containingY2O3 (12Ce–2WTiY2O3) and have demonstrated that at 700°C, the creepproperties of these alloys are superior to those of conventional ferritic steelslike HT9 and PNC-FMS, austenitic stainless steel SUS316 and another ODSalloy MA957 (14Cr-0.9Ti–0.3MoY2O3) developed in the USA.62 However,fabrication of thin clad tubes from such steels, which are produced throughthe powder metallurgical route, is a major technological challenge that has tobe overcome before these alloys are used in FBRs.

At present, mixed oxides (a mixture of uranium and plutonium oxides)are used as fuels for FBRs and a ferritic wrapper with austenitic clad seemsto be the appropriate choice for these fuels. However, there is a strongincentive to switch from oxide fuels to metallic fuels to achieve higher burn-up levels. The alloys that are under active consideration are based on the U–Pu–Zr system. Use of metallic fuel necessitates shifting clad material fromaustenitic to ferritic steels. This is because in addition to reduced swelling inthe ferritic steels, the damage caused by fuel-clad interaction (FCI) is less forferritic steels than the austenitic steels. FCI can be either mechanical or

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Using creep-resistant steels in nuclear reactors

623

Table 22.3 Chemical composition of major ferritic steels developed and studied in different countries for FBR core applications(wt%)

ALLOY C Cr Ni Mo Si Mn V Nb Ti P S N B Others

UKFI 0.15 13.0 0.47 – 0.30 0.45 – – – – – – – –FV607 0.13 11.1 0.59 0.93 0.53 0.80 0.27 – – – – – – –CRM–12 0.19 11.8 0.42 0.96 0.45 0.54 0.30 – – – – – – –FV448 0.10 10.7 0.64 0.64 0.38 0.86 0.16 0.30 – – – – – –

FranceF17 0.05 17.0 0.10 – 0.30 0.40 – – – ≤0.008 ≤0.008 0.020 – –EM10 0.10 9.0 0.20 1.0 0.30 0.50 – – – ≤0.008 – – – –EM12 0.10 9.0 0.30 2.0 0.40 1.00 0.40 0.50 – ≤0.008 ≤0.008 – – –T91 0.10 9.0 <0.40 0.95 0.35 0.45 0.22 0.08 – ≤0.008 ≤0.008 0.050 – –

Germany1.4923 0.21 11.2 0.42 0.83 0.37 0.50 0.21 – – – – – – –1.4914 0.14 11.3 0.70 0.50 0.45 0.35 0.30 0.25 – – – 0.029 0.007 –1.4914 0.16– 10.2– 0.75– 0.45– 0.25– 0.60– 0.20– 0.10– – – 0.010 0.0015 –mod 0.18 10.7 0.95 0.65 0.35 0.80 0.30 0.25 max. max.

USAHT9 0.20 11.9 0.62 0.91 0.38 0.59 0.30 – – – – – – 0.52 (W)403 0.12 12.0 0.15 – 0.35 0.48 – – – – – – – –

JapanPNC–FMS 0.2 11 0.4 0.5 – – 0.2 0.05 – – – 0.05 – –RussiaEP450 0.2 11.0– 0.3– 1.2– – – 0.1– 0.3– – – – – 0.004 –

13.5 0.5 1.8 0.3 0.6

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chemical. Fuel clad mechanical interaction (FCMI) arises from the appliedstress when element design restrains fuel swelling and results in plasticdeformation of the clad. Fuel clad chemical interaction (FCCI) is a complexmulti-component diffusion problem; U and Pu in the fuel and rare-earthfission products generated during irradiation have a propensity to interactmetallurgically with the clad material at high temperature.63–67 FCCI is animportant consideration in the choice of the clad material, as the compositionof the clad material has a significant effect on the above phenomenon. Ni isthe major element that contributes to FCCI; it forms low-melting eutecticswith U and enhances diffusion of U and Pu into the clad, thus increasing thediffusion layer thickness. As Ni is an important alloying element in austeniticstainless steel, FCCI is understandably greater in this class of steel than inferritic steels. Further, loss of Ni from the clad surface can also change thestructure at the surface from austenitic to ferritic, introducing additionaluncertainty with respect to its performance. As a result, when using metallicfuel in FBRs, it is essential that both clad and wrapper should be of ferriticsteels. However, presently available commercial ferritic steels produced viathe ingot route do not possess the required high-temperature strength.This again brings to the fore the importance of ODS ferritic alloys as a cladtube material for FBRs. An alternative option is to design the reactors withlower temperatures of operation than in the present design with oxide fuelsso that ferritic steels chosen for the wrapper application can themselves beused.

22.12 Variation in ∆DBTT with Cr content in ferritic steels.60

7dpa; 365°C10dpa; 365°C38dpa; 410°C

Chromium content (wt%)121086420

250

200

150

100

50

0

DB

TT

sh

ift

(°C

)2.25 CrV

2.25 Cr1WV10 dpa; 365°C

JLF-42Cr-1.5V

2.25Cr-2W

2.25Cr-2WV5Cr2WV

7dpa; 365°C

36dpa; 410°C

JLF-3

F82H

9Cr-2WV

12Cr2WV

12Cr-6MN-1V

JLF-612Cr-6MN-1W

Cr-1VJLF-1

9Cr-1W

9Cr-2WVTs

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22.4.4 Fusion reactors68

In fusion reactors, energy is produced from the fusion of nuclei of lightelements and the most suitable fusion reaction occurs between the nuclei oftwo heavy isotopes of hydrogen, viz., deuterium and tritium, to form a heliumnucleus, accompanied by the release of a neutron and energy:

2D1 + 3T1 → 4He2 (13.5 MeV) + 1n0 (14.1 MeV)

The temperature required for this reaction is of the order of millions ofdegrees Celsius and at such temperatures fuel changes from the gaseousstate to plasma. The hot plasma is magnetically contained in a vacuum vesseland isolated from the walls of the vessel. The most promising magneticconfinement systems are toroidal (ring-shaped) and the most advanced ofthese is the Tokomak reactor. Inertially confined fusion systems in whichenergy is produced by repeated ignition of D-T pellets by focused laserbeams are also being examined.

The deuterium fuel is abundant and is easily extracted from water. Tritiumcan be produced by bombarding Li with neutrons. Lithium is kept as ablanket surrounding the vacuum vessel for breeding tritium and a breedingratio of more than one can be achieved with this blanket. The overall fusionreaction can thus be represented as follows:

2D1 + 6Li3 → 2 4He2 + 22.4 MeV

The inner wall of the fusion reactor vessel, which faces the plasma and isexposed to nuclear radiation and fusion products, is often referred to as thefirst wall. Selection of materials for this component is an important aspect inthe development of fusion reactors. The first wall is subjected to the followingsevere conditions during service:

• mechanical and electromagnetic loading, and alternating thermal stressesinduced by surface and volumetric heating owing to the pulsed nature ofthe operation;

• irradiation with high-energy (14.1 MeV) fusion neutrons producingdisplaced atoms and helium, hydrogen and transmutation products, leadingto changes in bulk properties; and

• bombardment with ions and energetic neutral atoms from the plasma,resulting in surface and near-surface damage.

Potential structural materials considered for this application includeaustenitic steels based on the Fe–Cr–Ni and Fe–Cr–Mn systems, Cr–Moferritic/martensitic steels, alloys based on vanadium, niobium, molybdenum,titanium, and so on and SiC/SiC composites. The Cr–Mo and high-Cr ferriticsteels initially considered for this application are the same as those envisagedfor clad and wrapper applications of fast breeder reactors.

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Unlike in fission reactors, there are no highly radioactive products resultingfrom nuclear fusion. However, the high-energy neutrons can be absorbed bynuclei of various elements present in the structural material and these nucleiundergo transmutation producing radioactive isotopes of various elements.There is thus a specific interest in eliminating or minimizing those elementsthat can produce radioactive products as a result of neutron absorption andsubsequent transmutation.

A major limitation in evaluating structural materials for fusion reactorapplication is the inability to conduct in-reactor irradiation studies on thecandidate materials. One has to rely heavily on the data on neutron cross-section, radioactive isotopes, emission from these isotopes, half-life, and soon. Accordingly, inventory codes and cross-section and decay libraries havebeen developed in various countries to predict the radionuclide inventoriesof materials exposed in fusion reactors.69,70 Using these codes and libraries,the type of radiation, activity levels after stipulated cooling time, decay heat,and so on are predicted for the reference neutron spectrum to which variousreactor components like first wall, blanket, shield and magnetic coils areexposed.

In the case of ferritic steels considered for fusion reactor applications, theeffects of various alloying elements on contact γ-dose rate, induced activityand decay heat have been evaluated using the codes and libraries for aneutron loading of 2 MW m–2 for 2.5 years. These results show that Cr (anyconcentration), V (≤8%), Mn (≤1%), Ta (~1%) and Si (<0.4%) are acceptable,while Mo (>100 ppm), Nb (>1 ppb) and Ni (>50 ppm) are unacceptablealloying additions. Further, C, B and Ti in the concentration that is generallypresent in these steels do not detrimentally affect the activation parametersand W is not an ideal replacement for Mo in these steels.71 In addition tomajor and minor alloying elements, the effect of impurity elements in thesesteels on the γ-dose rate contribution has been evaluated using these librariesand codes and the results show that impurities like Pd, Ag, Bi, Hf, and so onshould be maintained below 1 ppm.72

These results clearly show that compositions of the various ferritic steelspresently considered for fast breeder reactor core application will not satisfythe stringent requirement on contact γ-dose rate (≤25 µSv h–1 after 100 ycooling), induced activity (<103Bq kg–1 for unrestricted release) and decayheat (<1 W m–3 for low level waste after 50 years of interim storage).73–75

This has resulted in the development of reduced activation ferritic steelsspecifically for fusion reactor applications. The principal approaches adoptedin the development are the replacement of radiologically undesirable Mo,Nb and Ni in the commercial steels by elements such as W, V, Mn, Ta and Tiwhich have equivalent or similar effects on the constitution and structuresand the removal of impurities that adversely affect the induced activities anddose rates, even when present in low concentrations in steels. Typical examples

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of reduced activation steels are EUROFER (Europe), F82H and JLF-1 (Japan)and 9Cr–2WVTa (USA). Compositions of these steels are given in Table22.4.68

22.4.5 Ferritic steels for steam generators and turbines innuclear plants

As mentioned in the introduction, in most nuclear reactors, the steam producedusing the heat generated from nuclear fission is used to drive the turbine.As in the case of fossil power plants, the materials for construction ofturbines are the ferritic steels. For rotors, steels like Cr–Mo–V steel areused and for turbine blades and shrouds 12%Cr martensitic stainless steelsare used.

The steam generators for PWRs and PHWRs do not use creep-resistantsteels as the steam temperatures are low (<300°C) in these reactors. However,in the case of FBRs, steam temperatures are comparable to those in fossilpower plant and creep-resistant steels are employed. FBR steam generatorshave liquid sodium in the shell side and high temperature water and steam inthe tube side. Since any leak in the tube would bring water or steam intodirect contact with liquid sodium, resulting in a very aggressive reaction, thechoice of the material and the design of the steam generator must be suchthat the possibility of tube rupture is eliminated. Hence, though austeniticstainless steel is chosen as the main structural material for FBRs, ferriticsteels are considered for steam generators. The ferritic steels possess betterresistance to stress corrosion cracking in the presence of high-temperaturesteam or water, good thermal conductivity (required for transfer of heat fromliquid Na to water) and a low thermal expansion coefficient and have sufficientcreep strength at the operating temperature of ~500°C.47 One matter of concern,however, is the possibility of carbon transfer from ferritic steel to austeniticsteel through liquid sodium owing to the difference in carbon activitiesbetween the two steels. This can result in carburization of the latter makingit brittle and decarburization of the former making it soft. In fact, it isinteresting to note that modified 9Cr–1Mo steel was originally developed byOak Ridge National Laboratory specifically for use in steam generators ofFBRs.76 The attempt was to reduce the carbon activity in the steel by addingNb and V and introducing N to form carbonitrides instead of just carbides asin conventional 2.25Cr–1Mo and ordinary 9Cr–1Mo steels. As the FBRprogramme in the USA came to a standstill, this alloy did not find immediateapplication in FBRs, but its superiority over the conventional Cr–Mo steelswas recognized by power industries worldwide and this alloy emerged as thepreferred material for fossil power plant replacing the 2.25Cr–1Mo and12Cr–1MoVNb steels that were in wide use in the industry at that time.Modified 9Cr–1Mo steel led the new generation of advanced ferritic steels

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628

Table 22.4 Typical/nominal compositions of reduced activation martensitic steels with a favourable combination of properties (wt%)68

Programme Name C Si Mn Cr W V Ta N B

CEC LA12TALC 0.09 0.03 1.0 8.9 0.8 0.40 0.10 0.02 –

EUROFER 0.10– 0.05 0.4– 8.0– 1.0– 0.20– 0.06– 0.02– 0.004–

0.12 (max) 0.6 9.0 1.2 0.30 0.10 0.04 0.006

Japan F82H 0.10 0.20 0.50 8.0 2.0 0.20 0.04 <0.01 0.003

JLF-1 0.10 0.08 0.45 9.0 2.0 0.20 0.07 0.05

USA 9Cr–2WVTa 0.10 0.30 0.40 9.0 2.0 0.25 0.07

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pursued actively by the fossil power industries to increase the efficiency ofpower plants and reduce CO2 emissions.

At present, modified 9Cr–1Mo steel is the first choice as structural materialfor steam generators of FBRs; for India’s PFBR, steam generators madefrom it are under fabrication. The Demonstration Fast Breeder Reactor (DFBR)and European Fast Breeder Reactor (EFR) also chose this steel for steamgenerators, although the reactors themselves were not constructed. However,FBRs constructed or conceived in earlier years like BN-600 and BN-800(Russia), SNR 300 and MONJU ( Japan) and Phenix (France) had either2.25Cr–1Mo steel or a combination of 2.25Cr–1Mo steel (for evaporator)and austenitic stainless steel (for superheater) for their steam generator. ForSuperphenix of France, Alloy 800 was the material of construction for thesteam generator.47

Though many advanced ferritic steels with better high-temperature propertiesthan modified 9Cr–1Mo steel are currently available, it is unlikely that any ofthe new steels would be considered for FBR steam generators in the immediatefuture. This is because the steam temperature in FBRs is limited by the maximumtemperature of liquid sodium, which is in turn depends on the maximumtemperature of the reactor core. The reactor core temperature is limited by theproperties of the structural material, and with both the currently used austeniticstainless and the ferritic steels of the immediate future not much increase in thereactor temperature is expected. Further, at ~ 500°C, type IV cracking, resultingfrom the poor creep strength of the weld joint, is not a serious problem inmodified 9Cr–1Mo steel. However, once metallic fuel and ODS clad tubes areintroduced in FBRs, it would be possible to increase the steam temperature andhence new alloys could be considered.

22.5 Fabrication and joining considerations

From the above discussion, it is clear that austenitic stainless steels andadvanced ferritic steels are the main creep-resistant steels considered fornuclear applications. Among these, austenitic stainless steels are widely usedin FBRs and advanced ferritic steels are the future structural material forcore components of FBRs and fusion reactors. Components of FBRs varywidely in size and dimensions; for Indian Prototype Fast Breeder Reactor(PFBR) clad tubes have a wall thickness <0.5 mm with a diameter of ~ 6 mmand a length of a few metres, while the grid plate, which supports the fuelsub-assemblies, is a massive plate that is 60–80 mm in thickness and ~6 min diameter. Fabrication of components from plates, pipes, tubes, and so onmade from austenitic stainless steels have not been a serious technologicalchallenge because of their good formability and weldability.

However, there have been specific cases of reactor component fabricationthat needed extensive development even with the use of austenitic stainless

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steels. One such case is the end cap welding of fuel clad tubes made from the15Cr–15Ni–Ti class of steels. The hot cracking susceptibility of D9 alloy(15Cr–15Ni–Ti steel) chosen for clad and wrapper and 316L(N) steel chosenfor all the major structural components of PFBR has been studied in detailusing the Varestraint test. The results have demonstrated the much greatersusceptibility of D9 alloy to solidification cracking than 316L(N) steel.77–79

The hot cracking problem in stainless steel welding is generally avoided byensuring some minimum volume fraction of δ-ferrite in the weld metal,which in turn is achieved by modification in the composition of the weldingconsumables employed. However, end cap welding of the clad tube is doneusing autogenous gas tungsten-arc welding with no filler addition and theD9 alloy solidifies in the fully austenitic mode with no possibility of δ-ferrite formation. Elements like Ti, P, S and N present in the steel alsocontribute to cracking by forming low-melting eutectics. It has been foundthat cracking in the end cap welding of the clad tube can be avoided only bythe use of an end plug made of a different material like 316L(N) steel.Further, welding parameters need to be carefully controlled in a narrowrange to ensure defect-free welds.

Another weld joint in FBRs that has been subjected to extensive study isthe dissimilar joint involving austenitic stainless steel piping and the steamgenerator header made from ferritic steel. Operating experience from manyfossil power plants has shown that these joints fail prematurely in the ferriticsteel side of the joint80 and the reasons attributed to the failure are thesecondary stresses generated by the difference in thermal expansion coefficientsof the two steels, carbon migration from the ferritic steel side to the austeniticsteel owing to difference in the carbon activities, strain accumulation at theferritic steel owing to the difference in creep strengths of the two steels,oxide notching in the ferritic steel side, and so on. A trimetallic transitionmetal joint involving the use of an intermediate Alloy 800 piece betweenthe austenitic and ferritic steels has been suggested for improving the life ofthe joint. Austenitic stainless steel is welded to Alloy 800 using ER 16-8-2filler wire (an austenitic stainless steel consumable), and Alloy 800 is weldedto the ferritic steel using ER NiCr–2/ENiCrFe–3 (Ni-base) consumables.81,82

The properties of these joints have been studied extensively and the resultshave shown that the life of the trimetallic joints is many times longer thanthat of the direct joint between austenitic stainless and ferritic steels.81–86 Inaddition to Alloy 800, another Ni-base alloy Alloy 600 has also been used asa transition piece in the dissimilar joints between austenitic and ferritic steelsdeveloped for fusion reactor applications.87,88 A recent trend in the transitionjoints involving these two classes of steels is the use of inserts of gradedcomposition manufactured by powder metallurgy using the hot isostatic pressing(PM-HIP) process.89

As the formability and weldability aspects of ferritic steels are considerably

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different from those of austenitic stainless steels, changing the material ofconstruction for core components from the latter to the former in the nearfuture would introduce additional issues of fabrication and joining. Drawingthin clad tubes and forming a hexagonal wrapper from ferritic steels aremuch more difficult than producing them from austenitic stainless steels.Similarly, welding ferritic steels calls for preheating and postweld heat treatmentand this would necessitate localized heat treatment of the weld joint withouthaving any adverse effect on the rest of the component.

At present, ODS ferritic steels developed specifically for core componentsof FBRs and fusion reactors suffer from anisotropy in properties and difficultiesin forming and joining. Processing them into thin clad tubes for FBR applicationhas proved to be a difficult task. Hamilton et al.62 describe the technologicalefforts taken up to develop successfully the processing route for an ODSalloy MA 957 from rods to clad tubes. Tubes were made from bar stockusing a combination of processes involving gun drilling, rod drawing, reroddingand final plug drawing. The product was also subjected to annealing in thedifferent stages of the forming operation. It was required to control theannealing and forming temperatures and reduction in each stage of formingvery carefully to complete the forming operation satisfactorily. Later, endcap welding of the tube was also successfully demonstrated. In spite of thesedevelopments and the significant R&D activities currently taking place, it isunlikely, however, that ODS alloys will be used in commercial reactors inthe near future. High-Cr ferritic steels produced by the conventional processingroute and their low-activation counterparts are more likely to be the materialsthat may find application in the nuclear reactors in the near future.90

22.6 Summary

Austenitic stainless steels and high-Cr ferritic steels are the major classes ofheat-resistant steels considered for nuclear applications. Among these, theaustenitic steels are already in use as major structural materials for FBRs.These steels are also used as a major piping material in BWRs and are beingconsidered for fusion reactor application. However, the high susceptibilityof austenitic steels to radiation swelling, helium embrittlement, irradiationcreep, and so on have necessitated replacement of these alloys for the reactorcore components with ferritic steels which are much more resistant to radiationswelling. Extensive in-reactor studies on advanced ferritic steels have shownthat these alloys can tolerate neutron doses as high 200 dpa and energyproduction as high as 200 MW days/tonne can be achieved with the use offerritic wrapper and clad material, compared to <100 MW d ton–1 currentlyachieved with austenitic clad and wrapper in FBRs. However, the creepstrength of these steels is inferior to that of austenitic stainless steels andthey also exhibit irradiation embrittlement (reflected as a large shift in DBTT

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under irradiation). Creep strength is important for ferritic steels consideredfor clad tubes, for which the operating temperature is high, and a relativelylow shift in DBTT is important for ferritic steels considered for wrapperfabrication. ODS alloys with improved creep resistance have been developedfor clad application using the powder metallurgy route and thin clad tubeshave been produced successfully from them. For wrapper application, thechoice is between 12%Cr steels like HT9, which exhibit good resistance toradiation swelling and 9%Cr steels like EM10 or 9Cr-1Mo steels whichshow minimum shift in DBTT; between them, the preference is for the latter.

Ferritic steels considered for fusion reactor application (first wall of thefusion chamber) have been derived essentially from those studied for FBRcore application. Emphasis is also being given to minimizing or avoidingthose elements in the steel that produce radioactive isotopes with long lifewhen exposed to high-energy neutrons present in fusion reactors. This hasled to the development of reduced-activation steels specifically for use infusion reactors. Elements like Ni, Mo, Nb, which produce radioactive isotopeswhen exposed to the reactor environment, are replaced by W, V and Ta.Further, trace elements that can produce radioactive isotopes are identifiedand upper limits are specified in these alloys. The purpose is to ensure safetyof the personnel during maintenance, waste disposal and recycling of thecomponents.

There is renewed interest worldwide in nuclear energy and FBRs aregoing to play a major role in the revival of nuclear energy as an alternativeto fossil fuel. Commercial reactors operating on nuclear fusion are alsolikely to be a reality in the future. It may therefore be expected that theuse of creep-resistant steels, especially advanced ferritic steels, in nuclearindustries is going to increase in a manner similar to their use in fossil powerplants.

22.7 References

1 S. Glasstone and A. Sesonskei, Nuclear Reactor Engineering, Vol. 1, CBS Publishersand Distributors, New Delhi, 2001, 72.

2 V. McLane, C.L. Dunford and P.F. Rose, Neutron Cross Section, Vol. 2, NeutronCross Section Curves, Academic Press, 1988, 184.

3 L.K. Mansur, in Kinetics of Non-Homogeneous Processes, G.R. Freeman (ed.), JohnWiley & Son, New York, 377.

4 R.L. Klueh and D.R. Harries, High Chromium Ferritic and Martensitic Steels forNuclear Applications, Chapter 8, ASTM, Philadelphia, 2001, 81.

5 E.A. Little, J. Nucl. Mater., 1993, 206, 324.6 R.L. Klueh and D.R. Harries, High Chromium Ferritic and Martensitic Steels for

Nuclear Applications, Chapter 9, ASTM, Philadelphia, 2001, 90.7 F.A. Garner, J.F. Bates and M.A. Mitchell, J. Nucl. Mater., 1992, 189, 201.8 J.L. Seran, V. Levy, P. Dubuisson, D. Gilbon, A. Maillard, A. Fissolo, H. Touron, R.

Cauvin, A. Chalony and E. Le Boublbin, in Effect of Radiation in Materials: 15th

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International Symposium ASTM STP 1125, R.E. Stoller, A.S. Kumar and D.S. Gelles(eds), ASTM, Philadelphia, 1992, 1209.

9 M.I. Hamilton, C.D. Johnsion, R.J. Puigh, F.A. Garner, P.J. Maziasz, W.J.S. Yangand N. Abraham in Residual and Unspecified Elements in Steel, ASTM STP 1042,A.S. Melilli and E.G. Nisbett (eds), ASTM, Philadelphia, 1969, 124.

10 Y. Yateishi, J. Nucl. Sci. Tech., 1989, 26, 132.11 I. Shibahara, N. Akasaka, S. Onose, H. Okada and S. Ukai, J. Nucl Mater., 212–215

1994, 487.12 G. R. Odette; J. Nucl. Mater., 1988, 155–157, 921.13 I.S. Kim, J.D. Hunn, N. Hashimoto, D.L. Larson, P.J. Maziasz, K. Miyahara and

E.H. Lee, J. Nucl. Mater., 2000, 280, 264.14 F.A. Garner, M.B. Toloczko and B.H. Sence, J. Nucl. Mater., 2000, 276, 123.15 L.K. Mansur, ‘Theory and experimental background on dimensional changes in

irradiated alloys’, J. Nucl. Mater., 1994, 216, 97–123.16 R.L. Klueh and D.R. Harries, ‘High chromium ferritic and martensitic steels for

nuclear applications Chapter 9, ASTM, Philadelphia, 2001 p. 113.17 J.O. Stiegler and L.K. Mansur, Ann. Rev. Mater. Sci. 1979, 9, 405.18 J.R. Mathews and M.W. Finnis, J. Nucl. Mater. 1988, 159, 257.19 E.R. Gilbert and B.A. Chin, Nucl. Tech., 1981, 52, 273.20 B.A. Chin, ‘An analysis of the creep properties of a 12Cr–1MoWV steel’, Topical

Conference on Ferritic Steels for Use in Nuclear Technologies, J.W. Davis and D.J.Michel (eds), AIME, Warrendale, PA, USA, 1984, 593.

21 C. Wassilew, K. Herschbach, E. Materna-Morris and K. Ehrlich, in Topical Conferenceon Ferritic Steels for Use in Nuclear Technologies, J.W. Davis and D.J. Michel(eds), AIME, Warrendale, PA, USA, 1984, 607.

22 R.J. Puigh and G.L. Wire, in Topical Conference on Ferritic Steels for Use in NuclearTechnologies, Eds. J.W. Davis and D.J. Michel, AIME, Warrendale, PA, USA, 1984,601.

23 J.M. Dupuoy, Y. Carteret, H. Aubert and J.L. Boutard in Topical Conference onFerritic Steels for Use in Nuclear Technologies, Eds. J.W. Davis and D.J. Michel,AIME, Warrendale, PA, USA, 1984., 125.

24 M.B. Toloczko, F.A. Garner and C.R. Eiholzer, J. Nucl. Mater. 1994, 212–215, 604.25 R.L. Klueh and D.R. Harries, High chromium ferritic and martensitic steels for

nuclear applications, Chapter 9, ASTM, Philadelphia, 2001, 139.26 R.L. Klueh and D.J. Alexander, J. Nucl. Mater, 1992, 187, 60.27 R.L. Klueh and D.J. Alexander, in Effect of Radiation in Materials: 15th International

Symposium ASTM STP 1125, R.E. Stoller, A.S. Kumar and D.S. Gelles (eds), ASTM,Philadelphia, 1992, 1256.

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Symposium ASTM STP 1125, R.E. Stoller, A.S. Kuma and D.S. Gelles (eds), ASTM,Philadelphia, 1992, 1135.

31 D.S. Gelles and L.E. Thomas, Topical Conference on Ferritic Steels for Use inNuclear Technologies, J.W. Davis and D.J. Michel (eds), AIME, Warrendale, PA,USA, 1984, 559.

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Symposium Part II, ASTM STP 956, F.A. Garner, C. H. Heneger Jr. and N. Igata(eds), ASTM, Philadelphia, 1987, 83.

33 R.L. Klueh and D.J. Alexander, J. Nucl. Mater., 1992, 187, 60.34 R.L. Klueh and D.J. Alexander, J. Nucl. Mater., 1992, 191–194, 896.35 B.A. Chin and R.C. Wilcox, Topical Conference on Ferritic Steels for Use in Nuclear

Technologies, (eds), J.W. Davis and D.J. Michel, AIME, Warrendale, PA, USA,1984, 347.

36 H.-C. Schneider, B. Dafferner and J. Akata, J. Nucl. Mater, 2001, 295, 26.37 R.L. Klueh, M.A. Sokolv, K. Shiba, Y. Miwa, J.P. Robertson, J. Nucl. Mater., 2000,

283–287, 478.38 R.L. Klueh and D.R. Harries, High-Chromium Ferritic and Martensitic Steels for

Nuclear Applications, American Society for Testing and Materials, Philadelphia,2001, 135.

39 R.L. Klueh and J.M. Vitek, ‘The resistance of 9Cr–1MoVNb and 12Cr–1MoVWsteels to helium embrittlement’, J. Nucl. Mater., 1983, 117, 295–302.

40 U. Stamm and H. Schroeder, ‘The influence of helium on the high-temperaturemechanical properties of DIN 1.4914 martensitic steel’, J. Nucl. Mater., 1988, 155–157, 1059–1063.

41 Hasegawa and H. Shiraishi, ‘Helium implantation effects on low activation 9Crmartensitic steels’, J. Nucl. Mater., 1992, 191–194, 910–914.

42 K.J. Harrison, S.H. Rou and R.C. Wilcox, J. Nucl. Mater., 1986, 141–143, 508.43 D.J. Alexander, P.J. Maziasz and C.R. Brinkman, in Microstructures and Mechanical

Properties of Ageing Materials, P.K. Liaw, R. Viswanathan, K.L. Murthy, E.P. Simonenand D. Frear (eds), The Minerals, Metals and Materials Society, Warendale, PA,1993, 343.

44 F.W. Noble, B.A. Senior and B.L. Eyre, Acta. Met., 35(5), 709.45 C.A. Hippsley and N.P. Howarth, Mater. Sci. Tech., 1988, 4, 791.46 S. Glasstone and A. Sesonske, Nuclear Reactor Engineering, Volume 2, CBS Publishers

& Distributors, New Delhi, 2001, 820.47 S.L. Mannan, S.C. Chetal, Baldev Raj and S.B. Bhoje, in Materials R&D for PFBR,

S.L. Mannan and M.D. Mathew (eds), IGCAR, Kalpakkam, 2003, 9.48 K. Herschbach, W. Schneider and K. Ehrlich, J. Nucl. Mater., 1993, 203, 233–245.49 M.I. Hamilton, C.D. Johnson, R.J. Puigh, F.A. Garner, P.J. Maziasz, W.J.S. Yang and

N. Abraham, ‘The effect of phosphorus and boron on the behaviour of titaniumstabilised austenitic stainless steel developed for fast breeder reactor service’ ReportNo. PNL-SA\15303 1987.

50 M. Fujiwara, H. Uchida, S. Ohta, S. Yuhaara, S. Tani and Y. Sato, in RadiationInduced Changes in Microstructure 13th International Symposium Part 1, ASTMSTP 955, F.A. Garner, N.H. Packan and A. S. Kumar (eds), ASTM, Pa, 1987, 127.

51 M. Fujiwara, H. Uchida and S. Ohta, J. Mater. Sci., 1994, 13, 908.52 M.L. Hamilton, G.D. Johnson, R.J. Puigh, F.A. Garner, P.J. Maziasz, W.J.S. Yang

and N. Abraham, ‘The effect of phosphorus and boron on the behaviour of titanium-stabilized austenitic stainless steel developed for fast breeder reactor service’ ReportNo. PNL-SA-15303 1987.

53 Schirra M, Creep and Fracture of Engineering Materials, Parker, J.D. (ed.), TheInstitution of Materials, London, 612.

54 Brinkman C.R. ‘Elevated temperature mechanical properties of an advanced type316 stainless steel’, Report No. ORNL/CP-101053, National Technology InformationScience (NTIS), USA 1999.

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55 A.J. Lovell, D.R. Wilson, D.F. Leibnitz and W.H. Sutherland, Topical Conference onFerritic Steels for Use in Nuclear Technologies, eds. J.W. Davis and D.J. Michel,AIME, Warrendale, PA, USA, 1984, 135.

56 C. Brown, V. Levy, J.L. Serin, K. Ehrlich, R.J.C. Rojer and H. Bergmann, ProceedingsFast Reactors and Related Fuel Cycles, FR-91, Vol. 1, Atomic Energy Society,Tokyo, 1991, 7.5.

57 C. Brown, R.J. Lilley and G.C. Crittenden, Nucl. Eng., 1994, 35, 122.58 V.S. Khabarov, A.M. Dvorisahin and S.I Porollo, Proceedings of Technical Committee

Meeting on Influence of High Dose Irradiation on Advanced Reactor Core Structuraland Fuel Materials, IAEA-TECH DOC 1039, IAEA, Vienna, 1998, 139.

59 V.M. Poplavsky and L.M. Zabudko, Proceedings of Technical Committee Meetingon Influence of High Dose Irradiation on Advanced Reactor Core Structural andFuel Materials, IAEA-TECH DOC 1039, IAEA, Vienna, 1998, 7.

60 A. Kohyama, A. Hishinuma, D.S. Gelles, R.L. Klueh, W. Diets and K. Ehrlich, J.Nucl. Mater., 1996, 231–237, 138.

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65 D.D. Keiser Jr., ‘Review of compatibility of FR fuel and austenitic stainless steel’,Report ANLIFE/PP-83153, Illinois, NIIS No DE 97000 549, 1994.

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Dispersion and non-oxide advanced and alternative fuels for power and researchreactors, IAEA TECDOC 1374, Ch. 3, Sept. 2003.

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applications’, International Conference at Warren, PA, 1981, 65.77 V. Shankar, T.P.S. Gill, S.L. Mannan and S. Sundaresan, Sci. Tech. Welding Joining,

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Mater. Sci. Tech., 1988, 4 (11), 1020.82 A.K. Bhaduri, S. Venkadesan, P. Rodriguez and P.G. Mukunda, Int. J. Pressure

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(1), 65–76.84 Mopati Sireesha, Shaju K. Albert and Subramania Sundaresan, Steel Res., 2002, 73

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1995.90 R.L. Klueh and D.R. Harries, High-Chromium Ferritic and Martensitic Steels for

Nuclear Applications, American Society for Testing and Materials, Philadelphia,2001, 208.

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637

23Creep damage – industry needs and future

research and development

R . V I S WA N AT H A N and R . T I L L E YElectric Power Research Institute, USA

23.1 Introduction

The phenomena of creep and creep fatigue have assumed great industrialsignificance worldwide in recent years. There are a variety of reasons forthis. First and foremost, industrial plants operating at high temperatures havebeen subject to aging and the fleet of these aging plants which have beenoperating well past their design life of 30 to 40 years is growing. Reducingthe cost of production is paramount for staying competitive. Reducing capitalcosts by deferring replacement of expensive components, and reducingoperating and maintenance costs by optimizing operation and maintenanceprocedures will both be the key strategic objectives of plant owners. Inaddition, operation at higher than design temperatures and cycling caused bythe de-regulated energy market have exacerbated the failure process byinteraction of creep and fatigue. A lot of research has therefore focused onquantifying damage occurring owing to creep and creep fatigue mechanisms.This chapter will address industry needs, present selected results from studiesfocusing on these needs and identify future needs in research and development(R&D).

Creep damage in components can take many forms. Excessive deformationcaused by creep can lead to failure by wall thinning, loss of clearance andother dimensional changes. A second form of damage can consist of theformation of cavities at the grain boundaries, which can then link up to formlarge cracks at high temperatures. In another failure scenario creep damagemay be initiated at high temperatures, but final fracture may occur in a brittlemode at lower temperatures owing to increased thermal stresses arising fromtransient conditions during shut-down or start-up. While dimensional changescan be readily seen and measured, in the other two failure modes, however,the failure is insidious and damage initiation can be observed only usingspecialized techniques. While most studies have focused on initiation ofdamage, more recently the crack growth regime has also received attention.Literature pertaining to crack initiation by progressive damage and subsequent

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growth of cracks is too vast to be described here. Any review of this literaturehas therefore to be selective. The emphasis here is on industry perspectiveon creep damage assessment and future R&D needs. Since much of theindustry interest centers on characterizing and quantifying creep damage andrelating the results to life consumption, this paper will deal selectively withissues relating to three categories of damage measurement techniques, thatis, calculational, non-destructive and destructive analysis.

23.2 Calculational methods for estimating damage

The very first step taken in assessing creep damage is using calculationalmethods. Although it is relatively easy to quantify damage in laboratorycreep tests conducted at constant temperature and stress (load), componentsin service hardly ever operate under constant conditions. Start–stop cycles,reduced power operation, thermal gradients and other factors result in variationsin stresses and temperatures. Procedures are needed that will permit estimationof the cumulative damage under changing exposure conditions.

23.2.1 Damage rules for creep

The most common approach to calculation of cumulative creep damage is tocompute the amount of life expended by using time or strain fractions asmeasures of damage. When the fractional damages add up to unity, thenfailure is postulated to occur. The most prominent rules are as follows:

Life-fraction rule (LFR)1:

Σ = 1i

ri

tt

[23.1]

Strain-fraction rule2:

Σ = 1i

ri

εε [23.2]

where ti and ε i are the time spent and strain accrued at condition i, and triand ε ri are the rupture life and rupture strain under the same conditions.Several authors have shown that the LFR is valid for temperature changesbut not for stress changes. It has also been shown that the ductility andfailure mode are also important. In an industrial context the LFR is oftenused to calculate the cumulative damage leading up to crack initiation. Toevaluate total life of the component, crack growth analysis is needed.

To calculate cumulative damage, the plant records and the time–temperaturehistory of the component are reviewed. The creep or creep–fatigue life fractionconsumed is calculated using the history, material properties and damage

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Creep damage – industry needs and future research 639

rules. This procedure is usually inaccurate owing to errors in assumed history,in the material properties and in the damage rules. Temperature historyinformation may be refined by supplemental non-destructive or destructiveexaminations such as microstructural studies, hardness measurements andoxide-scale measurements. The uncertainties in material properties can bereduced by building a database, particularly with respect to service-retrievedcomponents, taking into account environmental effects and heat-to-heatvariations

23.2.2 Linear damage rule for creep–fatigue

A variety of rules have been enunciated for calculating total damage whenboth creep and fatigue damage are present in a component. The most popularamong these is the linear damage rule, in which the life fractions consumedin creep (t/tr) and fatigue (N/Nf) are added as follows:

Σ Σ + = f r

NN

tt

D [23.3]

Differing approaches to calculating N/Nf, t/tr and the value of D at failurehave led to alternate rules which have been discussed by Viswanathan.3

A critique of damage rules as they apply to fossil plant components hasbeen published by Viswanathan.4,5 A detailed review of literature shows thatthere are divergent opinions regarding which damage approach provides thebest basis for life prediction. It is quite clear that a number of variables, suchas test temperature, strain range, frequency, time and type of hold, waveform,ductility of the material and damage characteristics, affect the creep–fatiguelife. The conclusions drawn in any investigation may therefore apply only tothe envelope of material and test conditions used in that study. The validityof any damage approach has to be examined with reference to the materialand service conditions relevant to a specific application. Broad generalizationsbased on laboratory tests, which often may have no relevance to actualcomponent conditions, do not appear to be productive. Thus, one should usea tailored, case-specific approach for any given situation.

23.2.3 Ductility exhaustion

Another approach that is widely used is the ductility–exhaustion approach.The ductility–exhaustion approach is simply a strain-based life-fraction rulein which the fatigue damage and creep damage are summed up in terms ofthe fractional strain damage for each category, as follows:6

1 = + f

p

p

c

cN D D∆ ∆ε ε

[23.4]

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where ∆ερ is the plastic strain-range component at half life, Dρ is the fatigueductility obtained from pure fatigue tests, ∆εc is the true tensile creep-straincomponent and Dc is the lower-boundary creep-rupture ductility of the material.The first term in Equation [23.4] denotes the fatigue-damage component andthe second term denotes the creep-damage component. Although the firstterm is fairly easy to understand and obtain from test data, the second term,especially the definitions of ∆εc and Dc, needs to be clarified. The problemarises from two issues. First, not all creep strain is viewed as damaging andonly that strain which accumulates below a critical strain rate necessary tocause constrained cavity growth is viewed as damaging.7 Second, the ruptureductility of a material is not a constant but decreases with decreasing strainrate. Hence, in defining a failure criterion, an appropriate lower boundaryvalue has to be defined for Dc.

23.2.4 Lacunae in calculational methods

One of the major problems in evaluating the applicability of calculationalmethods is that in many cases it is necessary to use all the available data inderiving the damage rule and thus it is possible to examine only the accuracywith which a given method describes the data. There also is a scarcity ofinstances in which service experience has been compared with results fromspecific life-prediction methods. In general, the available methods are utilizedonly to predict the lives of samples tested under laboratory conditions.Validation against test data in the laboratory and in-service data on actualequipment would lead to more confidence in the use of the various rules.

Results from most studies show that even the best of the available methodscan predict life only to within a factor of 2 to 3. Some of the cited reasonsfor these inaccuracies have already been discussed. Some additional reasonsare: failure of the methods of modeling changing stress relaxation and creepcharacteristics caused by strain softening or hardening, use of monotoniccreep data instead of cycle creep data, and lack of sufficiently extendedduration test data. None of the damage rules available today is entirely basedon sound mechanistic principles. They are all phenomenological in nature,involving empirical constants that are material dependent and difficult toevaluate. Extrapolation of the rules to materials and conditions outside theenvelope covered by the specific investigation often results in unsuccessfullife predictions. Material behavior under isothermal LCF conditions in thelaboratory often turns out to be totally irrelevant to material behavior underthermomechanical fatigue cycles involving in-phase or out-of-phase thermalcycles in the field. For application to service components, the stress–strainvariation for each type of transient and its time dependence must be knownwith accuracy. Such calculations are difficult and expensive to perform.Because of these limitations and the simplicity of the linear damage summation

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using the life-fraction rule, the latter approach continues to enjoy popularityin engineering applications.

Despite the inaccuracy of calculation procedures to estimate cumulativedamage, they will continue in use because of their simplicity, inexpensivenessand non-invasive character. They lend themselves to on-line monitoring sincedamage can be calculated in real time. Improved accuracy can be achievedby keeping better plant records, better definition of material properties bystoring and aging samples from the original heat of material and treatment ofthe various uncertainties by probabilistic techniques. Furthermore, thecalculation procedure will remain a mainstay of on-line monitoring techniques.

23.2.5 Non-relevance of isothermal low cycle data tothermomechanical fatigue

It is extremely important that in calculating creep–fatigue damage undercyclic conditions, the appropriate data obtained using a strain cycle simulatingthe actual service strain cycle is used. In many instances of fatigue, thetemperature varies along with the strain, giving rise to what is known asthermomechanical fatigue (TMF). Two simple waveforms in TMF testingare shown and compared with the LCF test in Fig. 23.1. If maximumtemperature corresponds to peak compression, as in the center diagram inthe figure, it is known as the out-of-phase cycle (OP). If the maximumtensile stress occurs at the peak temperature, it is known as the in-phase (IP)cycle. Depending further on when the hold time is superimposed, variouscycle shapes are possible.

In the past, thermal fatigue traditionally has been treated as beingsynonymous with isothermal LCF at the maximum temperature of the thermalcycle. Consequently, life prediction techniques have evolved from iso-thermalLCF literature. More recently, advances in finite element analysis and inservohydraulic test systems have made it possible to analyze complex thermalcycles and to conduct TMF tests under controlled conditions that simulatethese complex cycles. The assumed equivalence of isothermal LCF tests andTMF tests has been brought into question as a result of number of studiesand the pertinent issues have been reviewed by Viswanathan and Bernstein.8

High tensile strains at high temperature (IP) would favor creep, whereashigh tensile strains at low temperature (OP) would favor cracking of theoxide scale at the low temperature followed by environmentally inducedaccelerated creep damage during subsequent high-temperature exposure. Hence,Kuwabara et al. rationalized that in case of materials where damage is drivenby creep, IP cycles would be more damaging than OP cycles, for a givenstrain range.9 In other materials, where the environmental contribution issignificant, OP cycles may be more damaging than IP or isothermal LCFcycles. In addition to environmental effects, differences also arise between

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642

Isothermal fatigueThermal fatigue

Out-of-phase In-phaseTe

mp

. T T = Tmax

Time t0

+

0

–+

0

–Str

ess.

σS

trai

n.

ε

∆ε t

t

∆εσ

σ0

Tem

p.

T

Tmax

Time t0

+

0

–+

0

–

Str

ess.

σS

trai

n.

ε

∆ε t

t

∆ε

σ

ε0

Tmax

Tm

Tmin

TmaxTmTmin Tem

p.

T

0

+

0

–+

0

–

Str

ess.

σS

trai

n.

ε

Tm

TminTmaxTmTmin

Time t

∆ε

t

t

∆ε

σ

ε

23.1 Schematic diagrams showing waveforms of temperature, strain, and stress in thermal and isothermal fatigue tests.

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cycles in terms of the relaxed mean stresses. The relative severity of thedifferent cycles can also change with material ductility, maximum temperatureand hold time. Consequently, a simple classification of material behavior isnot possible.

A case in point is the cracking encountered in the ligaments betweentube penetrations in CrMo steel piping in power plants illustrated in Fig.23.2(a). The figure shows a deep crack, filled with oxides, indicating a verystrong environmental contribution. The cracking mode has been identifiedas creep fatigue. The creep–fatigue damage summation approach was foundto be inconsistent with the early initiation of cracks observed in the pipes.The actual failure involved crack initiation by repeated cracking of oxidescales at low temperatures during shutdown transients and subsequent creepdamage at the high operating temperature. The metallography of the crackedregion showed numerous oxide spikes (see Fig. 23.2) confirming that theoxide cracking is a crack initiation mechanism. This example clearly illustratesthe need to use appropriate thermomechanical fatigue data simulatingof actual component cycles in predicting the crack initiation life ofcomponents.9

23.3 Non-destructive evaluation methods

Conventional non-destructive evaluation (NDE) methods such as ultrasonictesting (UT), dye penetrant examination and eddy current examination reveallarge flaws and crack-like defects. More advanced techniques are needed todetect fine-scale creep damage prior to formation of cracks and to be able torelate it to the remaining life. These techniques have specific limitations andhave been reviewed in detail elsewhere. High-resolution NDE techniquessuch as positron annihilation, ultrasonic velocity measurements, phased andfocused beam ultrasonic techniques and mechanistic property changemeasurements have been demonstrated to be capable of detecting creepcavities in steels and are continuing to evolve as field deployable, inspectionoptions. In particular, linear phased array UT has gained significant applicationexperience in creep damage detection and characterization in steam pipingweld areas. The availability of portable equipment supporting 16-channeltransducers is a key enabling development for this application.

The key drivers for these advancements include the following: the abilityto sweep the ultrasonic beam through multiple angles in fine increments; theability to focus ultrasonic energy into a very localized region and the abilityto capture inspection data and perform more elaborate signal processingusing computer-based techniques to display and characterize the signal featuresbetter and to correlate such features to the damage present. Results are,however affected by microstrucual and grain size variations owing to variationsin manufacturing and specific correlations are therefore needed.

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(a)

(b)

23.2 Ligament cracking in boiler headers, (b) oxide spike in aligament crack.

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23.3.1 Phased array ultrasonic testing

The use of phased array technology to focus the ultrasonic beam has beensuccessfully applied to detecting pre-crack damage.10 In this approach, aspecially designed probe with an annular array of elements produces a focusedbeam in which the depth of focus can be varied by changing the time delaysfor pulsing and receiving. The enhanced signal-to-noise performance of thistechnique provides the basis for refined detection and characterization. Theinspection process with focusing is, however, inherently slow and generallyrequires a very localized application. Accordingly, other techniques, such asconventional ultrasonics or analytical modeling, are used to establish theareas for detailed inspection.

The advanced transducer and conventional UT technologies benefit stronglyby the opportunity for refined computer processing to the acquired ultrasonicresponse. Techniques involving waveform analysis, spectrum processing,noise analysis and others offer opportunities to improve the correlation betweeninspection results and the actual level of damage present in the component.In general, these areas are evolving through an extensive process of dataacquisition for plant components followed by component removal anddestructive verification of the actual level of damage. To date, these approacheshave been offered as a means of improving defect characterization and avoidingfalse calls. Continued improvement can be expected in this area which willgreatly benefit all forms of non-destructive testing.

23.3.2 Magnetic testing

The magnetic property response by materials has long been identified asaffected by early stage microstructural change (such as produced by creepevolution). Accordingly, considerable research has focused on the use of thiseffect as a technique to detect creep damage. In general, the significantchallenge for these types of NDE tools for creep damage has been the needto separate microstructural changes caused by to creep damage from thosecaused by material processing. As noted previously, the creep damage processproduces changes in material microstructure as cavitation evolves alonggrain boundaries and develops into cracking. Non-destructive testing basedupon measuring changes in magnetic properties such as hysteresis, remanenceand coercivity by applying a magnetic field to the component of interestwould be sensitive to the development of creep damage. Specifically, it ispostulated that the presence of voids along grain boundaries has an impacton domain wall activity and demagnetization response. Research has validatedthese qualitative relationships as shown in Fig. 23.3.12

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23.3.3 Magnetic acoustic emission

In addition to the measurement of the magnetic property response, it ispossible to use an applied magnetic field to induce acoustic emissions in thematerial. This process, called magnetoacoustic emission (MAE), occursgenerally when domain walls, in a changing magnetic field, move in discretejumps from one pinned position to another, producing acoustic waves, whichcan then be detected by a transducer. Research, in modeling and experiment,indicates a proportionate reduction in MAE owing to microstructural creepdamage.11 It is noted that this process relies on the high sensitivity of acousticemission testing but is an active rather than passive approach to generatingthe emissions. Further, the domain wall excitation produced by the appliedmagnetic field is not identical to the stress variations during plant operationswhich induce defect areas to emit acoustically, as detected under the acousticemission monitoring approach discussed previously. Research to refine thebasic approach for magnetic based NDE is on-going. A recent effort hasaddressed the issue of creep damage development in austenitic materials.The interest here is that the microstructural changes associated with agingproduce an increased magnetic response over time.

It has been found that long-term creep of certain steels results in a decreasein the electrical resistivity. Hardness decreases have been correlated withexpended life for softening-type ductile creep damage in rotor-type steels inJapan, although lack of knowledge of initial hardness makes direct correlationinaccurate. Density decreases have been correlated with the degree of cavitation

Seamweld

with µHAZ , µ

Weld with creep damage

Seamweld

with µHAZ , µ

Weld but no creep damage

Red

uct

ion

in

EM

F fr

om

bas

e m

etal

pla

te (

%)

60

50

40

30

20

10

1 2 3 4Current (A)

23.3 Magnetic response variation owing to the presence of creepdamage. µHAZ and µweld are the magnetic permeability of the heataffected zone of the weld, respectively.

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damage. Use of microstructural catalogs, interparticle spacing measurementsand carbide/ferrite chemistry changes have been investigated in steels andsuperalloys but the results have not been consistent in providing quantitativeanswers regarding life consumed. Microstructural features are used primarilyas supplementary tools to estimate the local temperature. Strain measurementsusing capacitance-type strain gauges have been shown to be capable ofmonitoring long-term, localized strain accumulation and are likely to findwidespread use.

23.3.4 Strain monitoring

Strain measurements are often employed to detect and measure creep damage.Gross changes such as component swelling and other dimensional changeshave been monitored in the past. Owing to unknown variations in the originaldimensions, changes in dimensions cannot be determined with confidence.Dimensional measurements fail to provide indications of highly damagingand localized creep strains such as those in the heat-affected zones of weldsand regions of stress concentrations in the base metal. Cracking can frequentlyoccur without manifest overall strain. Furthermore, the critical strainaccumulation preceding fracture can vary widely with a variety of operationaland material parameters, and with stress state. To enable measurement oflocalized strains, an ‘off-line’ condition surveillance system has been developed.The system uses the replication principle to evaluate localized strains andlife consumption.13 A surface grid is scribed at the region of interest andpreserved by means of an oxidation-resistant coating. A hard replica of highstability is used to duplicate the grid. Biaxial strain assessment is then madeby high-resolution measurement of the replicas taken at successive plantinspection shutdowns. A predictive strain-rate lifetime model approach isused to establish ‘fitness for service’. No field experience of this techniquehas been reported. More recently capacitance strain gauges have come intowide use. They have been found to be capable of monitoring long term,localized strain accumulation and are likely to find widespread use.

23.3.5 Acoustic emissions monitoring

Defects in a material structure are known to produce detectable acousticemissions under conditions of applied stress. In this fashion, acoustic emission(AE) has potential use as an on-line monitoring technology for creep andother damage. The localized sources of acoustic emission can originate froma number of causes. The resultant transient sound waves propagate radiallythroughout the structure, attenuating with distance from the source as afunction of frequency. The wave direction and wave mode types can changeas they are reflected and refracted at the boundaries of the structure. The

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waves are detected by a piezoelectric sensor, which transforms their mechanicaldisplacement into an electrical signal. These signals are very small at thesensor (microvolts to millivolts) and must be filtered and amplified near thesensor before transmission via cable to electronic data acquisition andprocessing equipment.

In AE inspection, the sound emitted by the material during changes instress states is detected and correlated to its source. Recent advancements inboth equipment and signal interpretation are providing a basis for movingthe detection of damage from the cracking stage to the stage of pre-crackcavitation. In this case, the measurement is indirect since the acoustic emissionsare from groups of cavities and are detected without absolute location withinthe thickness of the material.

In recent years, specific application of AE monitoring has been made tocreep damage detection in high energy steam piping.14 The key benefit of anacoustic emission monitoring approach is the economic leverage gained byinspection during actual plant operation and without the full removal of thepiping insulation that would be required for conventional inspection techniques.Additionally, piping systems typically undergo variations in temperature andpressure that provide a variable stress environment for activating damagewhen emitting acoustically. Key improvements have been made in order todifferentiate signals of no interest from those of critical interest via a rangeof pre- and post-data acquisition filtering of signal characteristics. On-goingwork is now examining the potential for quantitative (extent of damage)determinations from the AE processes.

23.3.6 Evolution of creep cavitation

One of the most widely used non-destructive techniques is surface replication.In this technique, the damage surface features are completely replicated onan acetate tape that can then be examined under a microscope at highmagnification to reveal the extent of creep cavitation. Portable microscopeshave also been used for field examination of the components. Alternatively,extremely small slices of the component have also been removed and usedfor metallographic observation. These approaches were popularized whenNeubauer and Wedel15 characterized cavity evolution in steels at four stages– isolated cavities, oriented cavities, linked cavities (microcracks) and macro-crack and formulated recommendations corresponding to the four stages ofcavitation, as shown in Fig. 23.4.

To provide a theoretical and quantitative basis for cavity evolution, Caneand Shammas16 used a constrained cavity growth model and proposed arelationship between the number fraction of cavitated boundaries (A parameter)and the life fraction consumed, t/tr, using heat-specific constants. Values ofthese constants either had to be assumed or determined experimentally for

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each heat, thus diminishing the usefulness of the model. Based on interruptedcreep tests on simulated heat-affected zone in 1Cr–0.5 steels, it wassubsequently concluded that the data contained too much scatter to verifythe life prediction model proposed by Cane and Shammas. The data couldnevertheless be used empirically, by plotting all the data together in the formof a scatterband whose lower limits are defined by the equation:

A = 0.517(t /tr) – 0.816 [23.5]

Direct correlations between cavity classification and expended life fractionhave also been established for 1.25Cr–0.5Mo steels, as shown in Fig. 23.5.The cavity classification system has been further refined and enhanced toinclude additional subcategories. While these models apply to cavitation inthe coarse-grained heat-affected zone (CGHAZ) of welds, cavitation in thefine-grained heat-affected zone (FGHAZ) known as type IV cracking is yetto be quantified. The literature pertaining to life prediction of type IV damagedwelds has been reviewed Ellis and Viswanathan.17 Replication is also notuseful when creep cavitation occurs below the surface. Catastrophic failuresof longitudinal seam welded piping have been reported in the USA, wherecracks started at the midwall region owing to stress concentration effects atweld cusps and resulted in catastrophic failure of a hot reheat pipe at theMojave Boiler Station as shown in Fig. 23.6.18

The field replication data forming the basis of the Wedel–Neubauerrecommendations have shown no consistent trends in cavitation evolutionwith operating time or with calculated creep life fraction consumed. Morefield data are needed before clear correlation can be established betweenreplica results and cumulative creep damage.

In using the A parameter method, the specific procedure used to measureA is crucial. The A parameter is defined as the number fraction of cavitatinggrain boundaries encountered in a line parallel to the direction of maximum

Cre

ep s

trai

n

Exposure time

IIA

B

C

III

D

Fracture

23.4 Creep life assessment based on cavity classification.

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principal stress. To measure A reproducibly, the procedure needs to bestandardized. The density of cavitation has also been used as a measure ofcreep damage. There is a need for standardized measurement of the ‘A parameteras well as the cavity/density since there is no way of comparing/correlating

Dam

age

clas

sifi

cati

on

D

C

B

A

1

00 0.2 0.4 0.6 0.8 1.0

Expended life fraction, (t/tr)

Ratings1 UndamagedA Isolated cavitiesB Oriented cavitiesC Linked cavitiesD Macrocracks

23.5 Correlation between damage classification and expended creeplife fraction for 1.25Cr–0.5Mo steels. (1) Undamaged, (A) isolateddamage, (B) oriented cavities, (C) linked cavities, (D) macrocracks.

23.6 Rupture in Monroe No. 1 north hot reheat line.

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results from different studies. There are a number of other limitations of thismetallographic technique. Quantitative correlations vary with steel composition,steel ductility and other properties. Consequently, a large database is neededto define the degree of scatter in the quantitative relations. More field validationof the replication technique is also required.

23.3.7 Analysis of carbides

Several investigators have attempted to correlate microstructural change withthe extent of creep damage. Carbide evolution in 2.25Cr–1Mo steel has beeninvestigated by Stevens and co-workers19,20 and Munson.21 In both studies,the amount of M6C carbides as a percentage of the total weight percent ofcarbides increased with time and temperature. Plotted in terms of a Larson–Miller time–temperature parameter, the results (Fig. 23.7) show the familyof curves from the Stevens and Flewitt study to be shifted laterally by asignificant amount compared to that of Munson.

These studies suggest that although the evolution of M6C may be used asa qualitative index of service temperature, wide variations may occur owingto differences in initial microstructure and composition. Data in Fig. 23.7also show that the kinetics of M6C formation are considerably accelerated byphosphrous. Improved procedures for carbide extraction, as well as a largerdatabase on samples with various initial compositions and heat treatmentsare needed before the microstructure technique can be used assessment ofplant life.

0.04 %P0.02 %P0.005 %PAged under stressMunson

Stevens and Flewitt

M6C

(%

)

60

50

40

30

20

10

034 35 36 37 38 39 40 41

LMP = (T + 460) (20 + logt)

23.7 Evolution of M6C in 2.25 Cr–1Mo steel as a function of agingtime and temperature.

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Stevens and Flewitt also found that the evolution of the M6C phase wasunaffected by stress. Nakatani et al. have developed correlations betweenthe percent M6C and rupture-life reduction so that the percent M6C in aservice-exposed component could be used to deduce the creep-life fractionconsumed.22

23.3.8 Hardness-based techniques for creep damageassessment

The changes in hardness in low-alloy steels as a function of time andtemperature have been extensively quantified, so that hardness changes canbe used to estimate the operating temperature.23,24 Correlations have alsobeen established between tensile (hence hardness) and the Larson–Millerrupture relationships for low-alloy steels. These correlations enable selectionof the appropriate Larson–Miller plot corresponding to a given hardnesslevel, which can be used to calculate the remaining life,25 see Fig. 23.8.

Goto26 and Kadoya et al.27 have proposed using hardness as a stressindicator in the remaining life assessment of CrMoV rotors26,27 used in steamturbines. They have quantified the effect of stress on the aging process so

60–70 ksi70–75 ksi75–80 ksi80–90 ksi90–100 ksi100–110 ksi110–115 ksi115–120 ksi120–130 ksi> 130 ksi

60 ksi100 ksi140 ksi

LMP = 40 975 + 57 (UTS) –5225log σ–2450 (log σ)2

25000 30000 35000 40000 45 000 50000LMP = T (20 + log tr)

Log

str

ess

(ksi

)

2.2

2.0

1.8

1.6

1.4

1.2

1.0

0.8

0.6

0.4

0.2

0

23.8 2.25 Cr–1Mo rupture data showing UTS dependence. ksi is1000 psi (pounds per square inch).

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that by comparing the kinetics of hardness change in the rotor with that ofthermally aged samples, the local stress can be determined. This value ofstress and the known value of service temperature are used in conjunctionwith the Larson–Miller rupture plot to estimate tr for the material in itscurrent damage state.

In contrast, McGuire and Gooch show that the magnitude of the stress isunimportant28 once a ‘threshold stress level’ has been exceeded. For allstress levels in the range 70–240 MPa, the hardness change in the stressedcondition was found to be 21% higher compared to the change in the unstressedcondition.

A creep model incorporating structural degradation as monitored by hardnesschanges has been proposed by Cane and Bissall.13 By equating the kineticsof hardness change to the kinetics of interparticle spacing changes, the decreasein the threshold back stress for creep was modeled. Substitution of the thresholdback stress in the Norton creep law yields an expression for secondary creeprate ε in terms of hardness changes. By integrating the ε between t = 0 andt = tr, where tr is the time to failure defined in terms of the time to reach anarbitrarily chosen critical strain, the remaining rupture life is predicted. Themodel is currently based on limited data and involves numerous assumptionsthat can only be justified by further research.

23.3.9 Hardness and low-cycle fatigue life

Considerable work has been carried out (26-23) in applying hardness tocalculation of fatigue-life consumption in the groove regions of a CrMoVsteel rotor used in steam turbines. It has been observed that low-cycle fatiguedamage results in strain softening and can be measured as a hardness decrease.The premise, therefore, is that if the fatigue curve corresponding to thecurrent hardness (in service) could be defined, the fatigue-life fraction consumedcould be calculated by entering the appropriate total strain range ∆ε t versusnumber of cycles to crack initiation (Nf) curve, as shown in Fig. 23.9.26

These relationships have been quantified.26

23.4 Accelerated destructive tests

One of the techniques widely used for life assessment of components involvesremoval of samples from a service component and conducting creep tests orstress rupture tests under accelerated conditions. Acceleration is achieved byincreasing the stress or temperature or both. The data are extrapolated toservice conditions using various parametric techniques reviewed in detail byViswanathan.3 An example of commonly used procedures is briefly describedhere.

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23.4.1 Parametric extrapolation techniques

Larson and Miller first introduced the concept of a time–temperature groupingin the form T(K + log t), where generally t is the time to rupture for a givenmaterial. A plot of stress versus the above parameter resulted in a single plot,within the limits of scatter, for any combination of time and temperature. Avalue of K = 20 is commonly used, although the constant is now known tohave a range of material-specific values. For the purposes of the remaininglife evaluation, specimens from the aged component are tested at higherstresses and temperature and the results are extrapolated to service conditionsby plotting the stress versus the Larson–Miller parametric combination oftemperature and time to rupture.

Monkman and Grant29 found that, for many alloy systems, the relationshipbetween the minimum creep rate, ε , and time to rupture, tr, can be expressedby:

log tr + m log ε = constant [23.6]

where m is a constant. For most materials the evaluated m has valuesapproaching unity, so that Equation [23.6] can be rewritten as:

ε tr = constant [23.7]

This relationship enables an estimation of tr if the minimum creep rate canbe determined either from accelerated tests or by creep or stress relaxationtests under service conditions. In general, the relationship between the minimumcreep rate and stress is given by:

ε σ = A n [23.8]

Hv = 252

Hv = 252

Hv = 203

102 103 104 105

Number of cycles to crack initiation

Tota

l st

rain

ran

ge,

∆ε t

(%

)

23.9 Estimation of low-cycle fatigue properties by hardness for aCr–Mo–V rotor forging at 500°C. Hv = Vickers hardness number.

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where A and n are stress-independent constants. Because creep is a thermallyactivated process, its temperature sensitivity would be expected to obey anArrhenius-type expression, with a characteristic activation energy, Q, for therate-controlling mechanism. Equation [23.8] can therefore be rewritten as:

ε σ = exp0– /A n Q RT [23.9]

Since ε and tr are correlated through Equation [23.8], Equation [23.9] isusually applicable to express tr instead of ε with the signs reversed for n andQ. This is the basis for the parameters defined by Orr et al.30 and is often usedto extrapolate ε or tr from one set of stress–temperature conditions to another.

A very general description of the creep curve under constant stress conditionsis given by the θ projection concept put forward by Evans et al.,31 in whichcreep strain, ε , is considered to be the sum of two competing processesusing the equation:

ε = θ1[1 – exp(–θ2t)] + θ3[exp(θ4t) – 1] [23.10]

where, θ1, θ2, θ3, and θ4 are all experimentally determined constants whichare functions of stress and temperature: θ1 and θ2 define the primary ordecaying strain rate component, and θ3 and θ4 describe the tertiary oraccelerating strain rate component. The absence of a steady second-stagecreep rate is implied by the model. A wide range of creep curve shapes canbe modeled with various combinations of the constants. Analysis of extensivecreep data on ferritic steels has shown that the log θ values vary systematicallyand linearly with stress. Hence, for a given material, if the θ functions can bedefined on the basis of short-time tests at high stresses, then the values atlower stresses (longer times) can be obtained by extrapolation and the long-time creep curves under low-stress conditions can be readily predicted bysubstituting the θ values in Equation [23.10]. From the shape of the creepcurve, the remaining life is estimated.

23.4.2 Isostress rupture tests

This technique, widely used in Europe and the USA, is based on the conclusionsstated earlier that the LFR is valid for temperature-based extrapolations andnot for stress-based extrapolations. Rupture tests are conducted at elevatedtemperatures at a constant stress close to the service stress and the tr versusT data are extrapolated linearly to the service temperature to estimate theremaining life.

Refinements of this technique include the use of miniature samples,consideration of oxidation and specimen size effects, verifications of theLFR, the applicability of uniaxial results to predict component behaviorunder multiaxial stresses and the effects of cycling. The use of miniaturespecimens 2 mm in diameter and 10 mm long instead of the conventional

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large-size specimens has been demonstrated. Based on the fact that largespecimens (base metal) last two to three times as long as laboratory-sizecreep specimens tested in air, correction factors have been determined forapplying laboratory data to predict the life of heavy section components. Theapplicability of the LFR under isotress test conditions has been verified. Ithas been shown that uniaxial test data can be used to make conservativeestimates of the remaining life of tubes operating under biaxial loading.Selection of time–temperature combinations must be carried out judiciously.It has also been observed that the isotress extrapolation procedure using trversus T yields the most conservative estimate of remaining life comparedwith tr versus 1/T and the Larson–Miller-type extrapolations, although contraryresults have also been reported in the literature.

23.4.3 Small punch testing

The small punch or disk bend test has particular value in life prediction ofoperating equipment since the test requires very small amounts of material(a typical specimen disk is 0.25 inches (6.35 mm) diameter × 0.020 inches(0.5 mm thick) and often the required volume of material can be acquired byoperating equipment in a virtually non-destructive manner (see for exampleFoulds and Viswanathan).32

The application of the small punch (SP) test for creep has gained significantinterest in the last decade, primarily as a result of research in Europe (see forexample, a 2003 summary of the European round-robin testing).33 Mostrecently, the CEN (one of three European standardization organizationsrecognized by the EC) has been working to develop a code of practice for thesmall punch test. The code of practice, intended to achieve a practical levelof uniformity in implementation of the test method, includes material creeptesting in addition to the more mature application of the test for tensile andtoughness properties. The practice documents are being developed as aworkshop agreement and are nearing completion.34

The focus on development of the SP test for creep has been on the use ofthe small punch test to predict the time to rupture of a uniaxially loaded testspecimen under specific conditions of applied load and temperature. This isbecause the creep life prediction for operating equipment has been based onextrapolation of conventional, uniaxially loaded test specimen rupture times tothe field application in question (see for example Foulds and Viswanathan).35

In essence then, the current challenge of the use and interpretation of smallpunch test data boils down to answering the question of what is the load thatshould be applied to the small punch test specimen to produce a rupture timeequal to that which would be produced in a uniaxial test specimen under theload or initial applied stress condition of interest. For example, to estimatethe remaining life of a pressure vessel operating at a particular stress and

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temperature, we need to small punch test the as-removed vessel samplematerial at some load that would give us rupture times equal to those that auniaxial test specimen would give us when subjected to the same stress asthe pressure vessel. However, since the stress state in the SP test (biaxial)differs from that in a uniaxial test (uniaxial) and since the stress field in thesmall punch test varies through the test, interpretation is not straightforward.

Research over the past decade has led to an encouraging semi-empiricalinterpretation of the test that could lead to field use. A summary of availablecorrelations between the small punch test load and the ‘equivalent’32 uniaxialtest specimen stress has been summarized,36 all of which ignore bending inthe punch specimen, assuming the rupture time is predominantly driven bymembrane stresses. The correlations have been reviewed and the Europeanround-robin data regressed over a simplified general form to suggest thefollowing usable equation.33

F k R rt/ = 3.33 SP–0.2 1.2σ [23.11]

where F is small punch test load, σ is the ‘equivalent’ uniaxial specimenstress, R is the small punch receiving die opening radius, r is the punchradius and t is the punch specimen thickness, and kSP is an empirical correlationfactor that should be determined for given material and test specimen geometry/dimensions. Application of the SP test requires initial testing to determinethis correlation factor.

The basic approach to using the small punch test for creep life predictioncan follow the temperature-accelerated method popular for uniaxial testing(e.g., Foulds and Viswanathan)32 that is, following establishment of kSP andwith knowledge of the stress and temperature of the operating equipment inquestion, a series of small punch tests at temperatures elevated above theequipment operating temperature may be run, each at a constant load determinedfrom the above equation. The small punch rupture times may then be extrapolatedto the operating temperature of interest on a temperature–log (rupture time)basis (common with low alloy steels) or the 1/temperature–log (rupture time)basis (sometimes used with austenitic stainless steels). Alternative accelerationmethods using a combination of elevated stress and temperature may also beused wherein extrapolation is performed using the Larson–Miller parameter,although the temperature acceleration method is considered more reliable. Asthe SP test sees more field use and as the specimen and test configurationsachieve better uniformity, we can expect that its application to creep life predictionwill increase.

23.4.4 Stress relaxation testing

The best short term (less than one week of testing) test is the stress relaxationtest (SRT), with a constant displacement rate (CDR) test providing an

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independent measure of embrittlement.37 Three or four 20 h test runs canprovide a plot of creep rate versus a time–temperature parameter. Using theMonkman–Grant relation, the rupture life at the desired stress and temperaturecan be estimated. Alternatively, the data may be used to compute the stressfor a creep strain of 1% as a function of temperature. This may be convertedto rupture life under the desired test conditions using the Gill–Goldhoffcorrelation. This approach is the one advocated generally, and is described indetail elsewhere.37

23.5 High temperature crack growth

All of the techniques described so far relate to life prediction from a crackinitiation point of view. For heavy wall components, the initiation criteriahave to be combined with crack growth data to perform a fracture mechanicsanalysis of the remaining life. Fatigue crack growth analysis procedures arewell established. For creep crack growth and creep–fatigue crack growth,however, methodologies and data needed for analysis have been gatheredonly during the last few years.

23.5.1 Creep crack growth

Extensive creep crack growth data pertaining to CrMo piping steels andCrMoV rotor steels have been collected, analyzed and consolidated.38,39 Ithas been observed that a crack-tip driving force parameter termed Ct, whichtakes time-dependent creep deformation into account, correlates much betterwith crack growth rates (da/dt) than the traditionally used elastic stressintensity factor K. The relation between da/dt and Ct can be expressedas:38–40

dd

= at

bC tm [23.12]

and

Ct = σ ε (A, n) aH (geometry, n) [23.13]

where σ is the stress far from the crack tip, obtained by stress analysis, ε isthe strain rate far from the crack tip, which is a function of the constants Aand n in the Norton relation, a is the crack depth obtained from NDEmeasurements and H is a tabulated function of geometry and the creepexponent n. The values of A and n are either assumed from prior data orgenerated by creep testing of samples. By assembling all the constants needed,the value of Ct can be calculated.

Once Ct is known, it can be correlated to the crack growth rate through theconstants b and m in Equation 23.12. Combining Equations 23.12 and 23.13

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provides a first order differential equation for crack depth (a) as a functionof time (t). Theoretically, this equation can be solved by separating variablesand integrating. However, the procedure is complicated by the time dependencyof Ct and the crack size dependence of the term H. To circumvent this, crackgrowth calculations are performed with the current values of da/dt (or a), todetermine the time increment required for incrementing the crack size by asmall amount ∆a (i.e. ∆t = ∆a/a). This provides new values of a, t and Ct, andthe process is then repeated. When the value of a reaches the critical size acas defined by KIC, JIC, wall thickness, remaining ligament thickness, or anyother appropriate failure parameter, failure is deemed to have occurred.

A number of variables affect the crack growth rate by modifying b, Ct orm and a large database is needed in order to apply the methodology. Amongthe variables to be considered are service degradation, presence of inclusions,impurities, ductility, test temperature, crack tip constraint and primary creep.

23.5.2 Creep–fatigue crack growth

Major advances have been made in developing the methodologies and dataneeded to treat crack growth under the combined effects of creep and fatigueat elevated temperatures. The loading conditions in elevated temperaturepower generation components can often be simply represented by a trapezoidalwave shape consisting of a loading period, a hold time and an unloadingperiod.

For creep–fatigue crack growth, the total crack growth rate is then givenby:

dd

= dd

+dd

total fatigue hold

aN

aN

aN

[23.14]

dd

= ( ) + [( ) ]total

1 t ave haN

C K C C tn q

∆ [23.15]

where the first term denotes the Paris law for the fatigue crack growthcomponent and the second term combines the crack growth owing to creep,including that owing to stress relaxation.

The average da/dt and Ct are obtained as follows:

dd

= 1 dd

avg h hold

at t

aN

[23.16]

and

( ) = 1 dt avgh 0

t

h

Ct

C tt

∫ [23.17]

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The (da/dN) hold is the crack growth during the hold period and is obtained bysubtracting the cycle-dependent crack growth rate from the total crack growth rate.

Figure 23.10(b) shows a plot of (da/dt)avg against (Ct)avg for 1.25Cr–0.5Mo steels at 538°C. The data include test results for a 90-s hold time,10-min hold time and also the creep crack growth rate data. When plotted asa function of (Ct)avg, the time rates of crack growth for these very differentconditions fall on the same trend curve. In contrast, when the crack growthper cycle is plotted as a function of ∆K (see Fig. 23.10(a)) different curvesare obtained. The significance of the above trend with regard to predictingthe hold-time effect in engineering components is obvious since creep crackgrowth data can be used to estimate creep–fatigue crack growth data andvice versa. In order to use this approach to predict crack growth during holdtime it is necessary to estimate the magnitude of (Ct)avg for components. Anequation has been proposed for estimating the (Ct) of any geometry for amaterial deforming by elastic–cyclic plasticity and power-law creep.40 Thedetails of this equation are given by Yoon et al.40 and are outside the scopeof this overview. The applicability of this approach to steels other than Cr–Mo steels is yet to be explored.

The potential of time-dependent fracture mechanics (TDFM) in establishingthe design life of new components or safe inspection intervals for componentsin service, or for performing risk assessments is obvious. The technology hascome a long way in the past 20 years but much still remains to be done todevelop total confidence in the approach. A majority of tests and analysesperformed assume isothermal conditions in which the influence of environmentis not explicitly included. More research into understanding creep–fatigueenvironment interactions is necessary for accurate life predictions. Considerableresearch is needed into analytical methods for treating crack growth underthermalmechanical loading and new test methods are needed that providecrack growth data under temperature gradients. The limitations of parameterssuch as ∆J under thermal gradients should be explored.

An area that has not been explored much is that of load interactionsduring crack growth at elevated temperatures. In the presence of transientthermal stresses, it becomes quite important to treat the effects of overloadon the crack growth rate during the subsequent hold time. There are significantopportunities for developing standard methods for creep–fatigue crack growthtesting. These tests are highly specialized and require very precise controlsand measurements. The data analysis is also complex so that forcing someuniformity into how data are treated will also help the overall goal of developinga well accepted life prediction methodology.

Extension of these methods to directionally solidified alloys single crystalalloys and to intermetallics is needed. These materials can exhibit a range ofbehavior not seen in Cr–Mo ferritic and austenitic stainless steels. For example,depending on the loading conditions and orientation, the same alloy may

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behave as a creep-ductile or a creep-brittle alloy. Solving this problem willrequire good numerical simulations that are now well within the capabilityof the current technology. Monitoring of service experience is very importantin determining which aspects of the problem deserve priority over others.Service experience is also important to validate the models after they aredeveloped and implemented.

1/10/1 (sec)1/98/11/600/11/900/11/24h/1

Trapezoidal waveshape

1.25Cr–0.5 Mo538°C (1000°F)

10 20 30 40

∆K .(ksi inch)(a)

da/

dN

(in

ch c

ycle

–1)

10–2

10–3

10–4

10–5

5 × 10–6

10–1

10–2

10–3

(b)

mm

h–1

MPa in.20 30 40

1/10/11/95/11/600/1

1/900/11/24/1CCG

Trapezoidal waveshape

1.25Cr–0.5 Mo538°C (1000°F)

10–5 10–4 10–3 10–2 10–1 1(Ct)avg (kips inch h–1)

(b)

(a)

(da/

dt)

avg (

inch

h–1

)

10–1

10–2

10–3

10–4

10–5

10–6

KJ m–2 h–1

10–2 10–1 1 101 102

1

10–1

10–2

10–3

10–4

mm

h–1

23.10 Comparison of creep–fatigue crack growth rates with(a) fatigue crack growth rate plotted against ∆K, (b) creepcrack growth data against (Ct)avg.

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23.5.3 Unique issues to weldments

There are many issues unique to weldments that are often not taken intoaccount in life prediction procedures using calculations or accelerated testing.High-temperature steam pipe seam weldment failures have occurred underoperating conditions believed to be well within design, and at accumulatedservice hours far short of any lifetime predicted by the base metal rupturedatabase (i.e. at very low expended life fractions). Premature failures couldbe attributed to inherently weaker weldment properties, strain concentrationsarising from inhomogeneous creep properties, or to some combination of thetwo effects. Design databases for high-temperature components are generallybased on data for homogeneous base metal tested as small specimens inuniaxial tests. None of these conditions apply to the behavior of heavysection weldments.

The application of isostress rupture testing to a high-temperature weldedstructure typically involves consideration of weld configuration, time-dependentstress distribution, inhomogeneous creep deformation, creep ductility variationacross the weldment, and so on. To add to the uncertainty of these effects,very little data are available on the stress rupture properties of weldmentsand limited data show that certain types of weldments can be inherentlyweaker compared to base metal owing to segregation of impurities andinclusions (see Viswanathan and Foulds).35 Typical weldments in operatingcomponents are subject to multiaxial stresses and this introduces yet anotherlevel of complexity to application of the method.

Specific concerns with respect to welds are:

• Welds are composite materials that contain many zones and pose manyproblems for the stress analyst.

• Failure can occur by crack growth at any of the interfaces, and it is hardto predict/simulate/accelerate the failures.

• Properties of welds vary widely with process.• Multi-pass welds lead to through-thickness variability in properties.• Sub-surface (type IV) cavities are hard to inspect.• Strength mismatch, cusp angle and roof angle influence failure

profoundly.• Interfaces in the weldment are subject to carbide growth, depletion,

embrittlement, and so on. It is hard to predict/model these phenomena.• Crack growth data are limited and fracture mechanics methodology is

underdeveloped.

23.6 Future trends

It is clear from the discussion in this chapter that there are many discrepanciesin current technology for detecting, characterizing and quantifying creep and

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creep damage in components operating at high temperatures. The demandson creep damage assessment and its effect on the remaining life of componentsunder creep will lead to increased use of on-line monitoring and decisionmaking tools, in situ replication, microscopy, chemical analysis, and so on,which will be used on a more quantitative basis. The initiation and evolutionof type IV cracking will be of great interest. Remote, rapid, wide coverageNDE techniques will be developed to perform quick and large scale surveysof damage. NDE techniques not requiring preparatory work (scaffolding,insulation removal, blade removal) will be developed. Improved signalprocessing and pattern recognition techniques will be developed better tocharacterize defects. Methods for non-destructive evaluation of in-servicetoughness, creep and fatigue damage will become more commonplace.Better tracking of operating history, improved databases on material propertiesand use of relevant TMF data will be utilized leading to improved calculationalprocedures and understanding of creep–fatigue. More realistic test proceduressimulating service conditions need to be developed. Treatment of uncertaintiesin data will increasingly be treated by probabilistic methods. Miniaturespecimen (e.g. small punch test), stress relaxation and other improveddestructive test techniques will be of greater interest. The effects of weldprocedures, weld geometry of PWHT, optimized fillers, mismatch effects,and other variables in weld performance in creep will be more thoroughlyinvestigated. Weld strength reduction factors will be defined and the resultsapplied to design as well as to life prediction. Creep and creep–fatiguecrack growth methodologies, especially with focus on weldment will bedeveloped.

It is difficult to prioritise among these many areas competing for R&Ddollars. In general, going from calculational to NDE to destructive assessmentinvolves increasing cost but also increasing accuracy. Which ones of thisthree-phase approach a utility company might choose to implement woulddepend on the age of the plant, budget considerations, operation andmaintenance policies and many other factors. It is often the case that acompany has only a given budget and a decision has to be made about whichactions will provide maximum ‘bang for the buck’. A probabilistic approach,combined with a sensitivity analysis might help determine where the maximumvalue for money could be obtained. In conclusion, it is reasonable to say thatsubstantial progress has been made during the last decade in conditionassessment of in service components. Some more work remains to be done,as reviewed in this chapter.

23.7 References

1 Robinson E L, ‘Effect of temperature variation on the creep strength of steels’,Trans. ASME, 1938, 160, 253–259.

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2 Lieberman Y, ‘Relaxation, tensile strength and failure of E1 510 and Kh1 F-L steels’,Metalloyed Term Obrabodke Metal, 1962, 4, 6–13.

3 Viswanathan R, Damage Mechanisms and Life Assessment of High TemperatureComponents, ASM International, Metals Park, 1989.

4 Viswanathan R, ‘Low cycle fatigue life prediction in LCF 3’, Low Cycle Fatigueand Elastic Behavior of Materials, K T Rie, Gounling H W, Konig G, Neumann P,Nowak H, Schwalbe K H and Seeger T, Elsevier, UK, 1992, 695–721.

5 Viswanathan R, ‘Creep fatigue life prediction of fossil plant components in creep’,Fatigue, Flaw Evaluation and Leak Before Break Assessment, Y S Garud (ed.),ASME, PVP, New York, 1993, volume 266, 33–51.

6 Priest R H, Beauchamp D J and Ellison E G, ‘Damage during creep-fatigue’, inAdvances in Life Prediction Methods, ASME Conference, Albany, American Societyof Mechanical Engineers, 1983, 115–222.

7 Miller, D A, Priest R H and Ellison E G, ‘A review of material response and lifeprediction techniques under fatigue–creep loading conditions’, High Temp. Mater.Proc., 1984, 6 (3 and 4), 115–194.

8 Viswanathan R and Bernstein H, J. ‘Some issues in creep–fatigue life predictionof fossil power plant components’, Trans. Indian Inst. Metals, 2000, 59 (3), 185–202.

9 Kuwabara K, Nitta A and Kitamura T, Advances in Life Prediction, Ed., D A Woodfordand Whitehead R (eds), ASME, New York, 1985, 141–152.

10 Bisbee H and Nottingham, L, December ‘Longitudinal seam weld characterizationby focused ultrasonics’, SPIE Proceedings, 1996, 2947, 88–99.

11 Sablik M J and Augustyniak B, 2000, ‘Modeling the magnetic field dependence ofmagnetoacoustic emission and its dependence on creep damage, Review of Progressin Quantitative Non-destructive Evaluation, Thompson D O and Chimenti D E(eds), American Institute of Physics, New York, Volume 19B, 1557–1564.

12 Govindaraju M R, Kaminski D A, Devine M K, Biner S B and Jiles D C, ‘Nondestructiveevaluation of creep damage in power plant steam generators and piping by magneticmeasurements’, NDT & E Inte., 1997, 30, 11–17.

13 Cane B J and Bissall A M, ‘Predictive assessment of damage in elevated temperatureweldment’, paper presented at the EPRI Plant Maintenance Technology Conference,Houston, November 14–16, 1986.

14 EPRI, Acoustic Emission Monitoring of High-Energy Steam Piping, Volume 1: AcousticEmission Guidelines for Hot Reheat Piping, EPRI Report TR-105265-V1, November1995.

15 Neubauer B and Wedel V, ‘Restlife Estimation of Creeping Component By Means ofReplication’, D A Woodford and R Whitehead (eds), Advances in Life Prediction,ASME, New York, 1983, 307–314.

16 Cane B J and Shamma M, A Method for Remanent Life Estimation By QuantitativeAssessment of Creep Cavitation on Plant, Report TPRD/L2645/N84,?, UK. As citedby Viswanathan et al. 1994, ‘Life Assessment Of Superheater/Reheater Tubes InFossil Boilers’, J. Pressure Vessel Technol., 1994, 1, 59–75.

17 Ellis F and Viswanathan R, ‘Review of Type IV Cracking’, Fitness for ServiceEvaluations, Volume 380, ASME PVP, 1998, 59–75.

18 Wells C and Viswanathan R, ‘Life Assessment of High Energy Piping’, in Technologyfor the 90s, Au Yang M K (ed.) ASME Pressure Vessel and Piping Division, NewYork, 179–216.

19 Stevens R A and Lonsdale D, Isolation and Quantification of Various Carbide Phases

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in 2.25Cr–1Mo Steel, SER/SSD/-84-0046/N, Central Electricity Generating Board,England, June 1984.

20 Stevens R A and Flewitt P E J, The Effect of Phosphorus on the Microstructure andCreep Strength of 2.25Cr–1Mo Steel, SER/SSD/84-0020/R, Central ElectricityGenerating Board, England, March 1985.

21 Munson R, Radian Corporation, Austin, TX, 1990, private communication.22 Nakatani H, et al., ‘Metallurgical Damage Detection and Life Evaluation System for

Boiler Pressure Parts’, Paper presentated at the EPRI Conference on PredictiveMaintenance of Fossil Plant Components, Boston, MA, October 1990.

23 Askins M C, Remaining Life Estimation of Boiler Pressure Parts, Vol. 4, MetallographicModels for Weld Heat Affected Zones, Report CS-5588, Electric Power ResearchInstitute, Palo Alto, CA, November 1989.

24 Ellis F V, Robert B W and Henry J F, Remaining Life Estimation of Boiler PressureParts, Vol. 4, Metallographic Models for Weld Heat Affected Zones, Report CS-5588, Electric Power Research Institute, Palo Alto, CA, November 1989.

25 Grunloh H and Ryder R H, Life Assessment of Boiler Pressure Parts, Vol. 7, Superheater,Reheater Tubes, Report TR-103377, Vol. 7, Electric Power Research Institute, PaloAlto, CA, November 1993.

26 Goto T, Study on Residual Creep Life Estimation Using Nondestructive MaterialProperty Tests, ‘Mitsubishi Technical Bulletin, No. 169, Mitsubishi Heavy Industries,Takasago, April 1985.

27 Kadoya Y, Goto T, Uaeke M and Fujii H, 1985 ‘Material Characteristics NDESystem from High Temperature Rotors’, Paper No. 85-JPGC, PWR-10 presented atthe ASME/IEEE Joint Power Generation Conference.

28 McGuire J and Gooch D J, ‘Metallographic Techniques for Residual Life Assessmentof 1CrMoV Rotor Forgings’, Proceedings of the International Conference on LifeAssessment and Extension, Nederlands Instituut Voor Lastichniek, The Hague, 1989,Volume 2, p 116.

29 Monkman F C and Grant N J, ‘An Empirical Relation Between Rupture Life andMinimum Creep Rate in Creep Rupture Tests’, Proc. ASTM 56, 1938, 1956, 593–620.

30 Orr R L, Sherby O D and Dorn J E, Trans. ASM, 1954, 46, 113.31 Evans R W, Parker J D and Wilshire B, ‘An extrapolative procedure for long-term

creep strain and creep life prediction’, In Recent Advances in Creep and Fracture ofEngineering Materials and Structures, Pineridge Press, Swangear, 135–184.

32 Foulds J and Viswanathan R, ‘Nondisruptive Material Sampling and MechanicalTesting’, J. Nondestructive Evaluation, 1996, 15 (3), 151–162.

33 Bicego V, Di Persio F, Hurst R and Ranta J H, ‘Small Punch Creep Test Method:Results from A Round Robin carried out within EPERC TTF5’, 29th MPA Seminar,Stuttgart, 9–10 October 2003.

34 CEN Workshop Agreement WS21, Small Punch Test Method for Metallic Materials,CEN, Brussels Belgium, 2004.

35 Foulds J and Viswanathan R, ‘Accelerated Stress Rupture Testing for Creep LifePrediction – Its Value and Limitations’, J. Pressure Vessel Technol., May 1998, 120,105–115.

36 Li Y and Sturm R, ‘Small Punch Tests for Welded Heat Affected Zones’, InternationalConference on Welds 2005, Geesthacht, September 2005.

37 Woodford D.A. ‘Creep strength evaluation of serviced and rejuvenated T91 usingthe stress relaxation method’, in Advances in Materials Technology For Fossil Power

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Plants, Chicago, American Society For Metals ASM, Metals Park OH, 2004, 1101–1115.

38 Saxena A, Han J and Banerji K, Creep Crack Growth in Boiler and Steam PipeSteels, Report CS-5583, Electric Power Research Institute, Palo Alto, CA, January1988.

39 Saxena A, Creep Crack Growth in CrMoV Rotor Steels, EPRI RP2481-5 Report,Palo Alto, CA.

40 Yoon K B and Saxena A ‘Characterization of Creep-Fatigue Crack Growth BehaviorUnder Trapezoidal Waveshape Using Ct Parameter’, Int. J. Fracture, 1993, 59, 95–102.

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Index

A-USC plants 565–70accelerated destructive tests 653–8acceleration of cyclic creep 375, 384acoustic emissions (AE)

monitoring 647–8activation energies 241–2advanced turbines 174age hardening 158ageing embrittlement 611–13alloy design 539–70

austenitic steels 45–8, 554–70Fe-Ni based austenitic alloys 565–70for header pipes 541for heavy-wall thickness

piping 546–54high strength low-Cr steels 541–6for main steam pipes 541martensitic high-Cr steels 546–54and oxidation 527–8for power plant components 539for reheaters 540–1, 554–64for superheaters 540–1, 554–64for water walls 539–40

analytical models of crack formation 512anelastic deformation 371, 372anomalous temperature dependence 528arc welding 476argon-oxygen decarburization (AOD) 66ASME Code 158, 169ASME P91 steel 30, 539athermal yield stress 265, 270–5, 277,

359–61atomic bonding 224–5, 226–7austenitic steels 23, 42–64, 68–70

15%Cr-15%Ni 51–218%Cr-8%Ni 43, 44, 51

20-25% Cr 44, 52–3age hardening 158alloy design 45–8, 554–70austenitising temperature 23boiler tube applications 48–57chemical plant applications 62–4cold working 286constant stress creep 368–9cooling from the annealing

temperature 158Discaloy 60, 61electrical conductivity/resistivity

236–7electro slag remelting (ESR)

process 61, 66grain boundaries 341–5heat exchanger applications 48–57low initial dislocation density 379–80microstructure 366nickel alloyed 42–3in nuclear reactors 615–21and radiation swelling 599–600strengthening mechanisms 292–5Tempaloy A-1 steel 44thermal conductivity 231for thick-section pipe 57–62TP 316 steel 57, 59for turbine components 57–62, 591under-stabilising technique 47void swelling 599–600Young’s moduli 224

austenitising temperature 23

back-stress concept 322–4bainitic low Cr steels 287–9boiler tube applications 48–57

667

WPNL2204

Index668

bolting materials 142Boltzmann constant 345, 346bond energy 226–7bond length 227bond strength 227boron 38, 329–30, 336–7, 342, 345Brite-Euram-Projects 80brittle failure 25

calculational methods of damageassessment 638–43

cumulative damage 638–9ductility exhaustion 639–40lacunae 640–1life-fraction rule 638, 641–3linear damage rule 639strain-fraction rule 638

carbide analysis 651–2carbon steels 19casting process 174, 175–82

12Cr steels 199–200electric arc furnaces (EAF) 175,

178–80electroslag remelting (ESR) 61, 66,

175, 180–1hot topping process 181–2ladle refining furnaces (LRF) 66, 175,

178–80open hearth furnaces 64, 66, 175vacuum arc remelting (VAR) 182vacuum degassing 66, 175vacuum induction melting

(VIM) 175, 182cavity growth 352–5, 648–51cavity surface diffusion 354–5CB8 alloy 306–12, 315–20, 324–5CCT (continuous cooling transformation)

diagrams 188–9CEN standards 81, 95–150

bolting materials 142component creep tests 147design standards 148general principles 95material specifications 95relaxation testing 147–8residual life assessment 149–50testing programmes 147testing techniques 144–6user specifications 149

welding consumables 142welding procedure

qualifications 142, 144chemical composition 155chemical plant applications 62–4chromium

10-12% Cr steels 32–611% Cr steels 30, 36–812Cr steels 199–200, 20115%Cr-15%Ni system steels 51–218%Cr-8%Ni austenitic

steels 43, 44, 5120-25% Cr steels 44, 52–39-12% Cr steels 26–32, 45, 67–8,

290–2, 295–301and grain boundaries 330CrMoV steels 21, 25, 26–7, 192–5and oxidation 527and turbine components 192–204

classical nucleation theory (CNT) 313–14coatings for oxidation-resistant steels 533cobalt diffusion 250Coble creep 10, 275, 276cold rolling 284–5, 286component creep tests 147composite model 388–9COMTES projects 211conductivity

electrical 234–8thermal 230–4

constant stress creep 368–70constituent atoms 244–5constitutive equations 403–17

data distribution 412defect assessment 416end-of-life criteria 414–15future trends 416grain boundaries 345–6and material characteristics 412model fitting effectiveness 405–11model selection 412multi-axial stress rupture 415service life predictions 412–16

continuous cooling transformation (CCT)diagrams 188–9

cooling from the annealing temperature155, 158

core components 615–17, 621–4COST programmes 79–80, 306–12, 584

WPNL2204

Index 669

CB8 alloy 306–12, 315–20, 324–5costs to industry 637crack formation 459–68, 484–95,

504, 643analytical models 512creep crack growth 658–9creep-fatigue crack growth 659–61detectability 494–5in dissimilar welds 488–90growth rate equation 509high temperature crack growth

658–62microscopic features 509–13reheat cracking 486–8role of constraint 493–4type I cracks 484, 485type II cracks 484, 485type III cracks 484–5, 486–8type IV cracks 485, 491–5,

513–17, 649V-type cracks 509–11W-type cracks 509–11in welded joints 484–95, 513–17see also fracture mechanics

creep crack growth 658–9creep curve shape variations 429–30creep deformation 637–63

athermal yield stress 265, 270–5, 277calculational methods of

assessment 638–43cavity evolution 352–5, 648–51Coble creep 10, 275, 276costs to industry 637damage evolution 5definition 3deformation mechanism

map 9–11, 275–7fracture mechanism map 11–13,

350–64homologous temperature 3–4logarithmic creep 6microstructure evolution 5minimum creep rate 5, 243modes of deformation 265, 637Nabarro-Herring creep 10, 275, 276necking of specimens 4–5NIMS Creep Data Sheets 9non-destructive evaluation

643–53, 663

NRIM Creep Data Sheets 9and oxidation 7, 519–34rate curves 6–7, 265rupture strength 7–9, 16–17, 279stages of 3–4stress-strain responses of

materials 265–7temperature and strain rate

dependence 267–9threshold stress 78in welded joints 483–4yield stress 265, 267–9, 270–5, 277

creep ductility 298, 351–2, 431–3,639–40

creep strain curves 403creep-fatigue behaviour 446–69

crack behaviour 459–68, 659–61experimental procedures 447–9fatigue stresses 217–18life estimation 449–56multiaxial behaviour 456–9Nelder-Mead method 459stress-strain behaviour 449

creep-induced strain 422–7CrMoV steel 192–5cross gliding 248crystallographic slip 365cumulative damage calculation 638–9cyclic creep 374–5, 382

acceleration 375deceleration 374–5

cyclic variation of stress 384

damage assessment calculations 638–43damage evolution 5damage rules 638–9damage tolerance values 430–1data distribution in constitutive

equations 412data rationalization 425–6DBTT (ductile to brittle transition

temperature) 337–8, 605–10deceleration of cyclic creep 374–5, 384defect assessment 416defect development and oxidation 523–5deformation mechanism map

9–11, 275–7degradation of particle hardening 383–4delta ferrite 200, 478

WPNL2204

Index670

design standards 148detectability of crack formation 494–5deuterium 625development of creep-resistant

steels 15–70austenitic steels 23, 42–64, 68–70ferritic steels 19–42purity of heat-resistant steels 64–7requirements for heat-resistant

steels 18–19steel melting 64–7

differential heat treatment 205, 207diffusion behaviour 241–63

activation energies 241–2constituent atoms 244–5cross gliding 248data application and searches 260–3dislocation climb 248dislocation glide 249and failure 250glide velocity 249grain boundary diffusion 252–3interdiffusion reactions 245interdiffusion values 258in iron and iron-base alloys 255–60lattice diffusion 245–6magnetic transformation 250–2matrix deformation 248–9microstructure change 249–50oxygen diffusion 250segregation of minor elements 253–4short-circuit diffusion 243, 246–8solute atoms 249stacking-fault energy 254–5time-dependent deformation 242–3tracer diffusion values 257vacancies 243–4, 260

DIN standards 79Discaloy 60, 61dislocation

climb 248glide 249hardening 284–6, 541models 385–9precipitate-dislocation

interaction 320–1strengthening 541

dispersion hardening 281–4dissolution of fine carbonitrides 296–8

ductile-to-brittle transition temperature(DBTT) 337–8, 605–10

ductility exhaustion 639–40ductility of materials 351–2durability strength 15–16, 17DVM creep rate limit test 15–16, 19dynamic recrystallisation 330, 334–6

ECCC see European Creep CollaborativeCommittee (ECCC)

18%Cr-8%Ni austenitic steels 43, 44, 51elastic behaviour 219–25

modulus of elasticity 221–5stress and strain 219–21

elastic deformation 371electric arc furnaces (EAF) 175,

178–80electrical conductivity/resistivity 234–8electro slag remelting (ESR) process 61,

66, 175, 180–111% Cr steel 30, 36–8embrittlement caused by ageing 611–13end-of-life criteria 414–15engineering grain boundaries 329enhanced steam oxidation 520–5EPERC see European Pressure

Equipment Research Council(EPERC)

estimation of damage 638–43Eurofer ODS 339Euronorms standards 80European Carbon and Steel

Collaboration 79European Commission 81–2European Creep Collaborative Committee

(ECCC) 79, 80, 85–92contribution to standardisation 90–2future of 92Memorandum of Understanding

(MoU) 85motivation and history 85organisation 90Working Groups 90, 91–2

European Pressure Equipment ResearchCouncil (EPERC) 80, 92–5

contribution to standardisation 94–5history 92, 94objectives 92, 94organisation 94

WPNL2204

Index 671

European specifications see specificationsand standards (Europe)

fabrication and joining 629–31failure and diffusion behaviour 250fast breeder reactors (FBRs) 598,

613–24welded joints 630

fatigue behaviour see creep-fatiguebehaviour

FATT (fracture appearance transitiontemperature) 207

Fe-Ni based austenitic alloys 565–70FEM (finite element method)

analysis 185ferrite effects on strengthening

mechanisms 300–1ferritic steels 19–42

10-12% Cr steels 32–611% Cr steel 30, 36–856T5 steel 279-12% Cr steels 26–32, 45, 67–8,

290–2, 295–301boron additions 38, 329–30, 336–7,

342, 345carbon steels 19constant stress creep 369–70cooling from the annealing

temperature 155CrMoV steel 21, 25, 26–7, 192–5dislocation strengthening 541FV448 steel 27grain boundaries 330–40H46 steel 27HCM2S steel 25HCM 12 steel 30–1, 32high initial dislocation density 380–1laths 331low alloy steels 19–26microstructure 366–7, 474–82molybdenum steels 20in nuclear reactors 621–4, 626–9, 631oxide dispersion strengthened (ODS)

292, 338–40, 631P91 steel 30, 539rotor steels 30, 31STX 21 research project 38–42TAF steel 27thermal conductivity 231

for thick-section boilercomponents 38–42

TMK 1 and 2 steels 31ferromagnetic state 250ferromagnetic transformation 25215%Cr-15%Ni system steels 51–256T5 steel 27finite element method (FEM)

analysis 185fission reactions 597–8forging process 174, 183–6

12Cr steels 200fracture appearance transition

temperature (FATT) 207fracture mechanics 504–17

effect of mechanical constraint 507–9microscopic 509–13non-linear 504–7see also crack formation

fracture mechanism map 11–13, 350–64athermal yield stress 359–61cavity growth 352–5ductility of materials 351–2modes of fracture 351multi region analysis 361–2rupture strengthchanges 356–8, 359–61and microstructural

degradation 358–9stress and temperature

dependence 352–5fusion reactors 625–7fusion welding 474–5FV448 steel 27

German specifications andstandards 78–9

Gibbs equilibrium segregation 333–4Gibbs free energy 312, 313glide velocity 249Gr.91 steel 433–6Gr.92 steel 436–40grain boundaries 329–47

austenitic steels 341–5and boron 329–30, 336–7, 342, 345cavity growth controlled by 352–4and chromium behaviour 330constitutive equations 345–6diffusion 252–3

WPNL2204

Index672

ductile to brittle transitiontemperature 337–8

dynamic recrystallisation 330, 334–6engineering 329ferritic steels 330–40future trends 348–9and hafnium 338, 342, 345, 347intergranular cracking 330, 353molecular dynamics (MD)

modelling 347precipitate-grain boundary interaction

322precipitate-subgrain boundary

interaction 321–2precipitation 329, 331–3, 341properties 345–6segregation effects 329–30, 333–4,

341–2subgrain boundary migration 393triple junction mobility 347vacancy concentration 354and zirconium 342, 345

grain growth zone 479grain refined zone 479grain refining behaviour 188

H46 steel 27hafnium 338, 342, 345, 347Hall-Petch relationship 330hardening see strengthening mechanismshardness-based testing techniques 652–3HCM2S steel 25HCM 12 steel 30–1, 32header pipes 541heat affected zone (HAZ) 472, 478–80,

497–8, 513–17HAZ simulation 480–2

heat exchanger applications 48–57heat indication test (HIT) 192heat transfer mechanisms 231heat treatment process 174, 186–90

12Cr steels 201continuous cooling transformation

(CCT) diagrams 188–9differential heat treatment 205, 207grain refining behaviour 188normalizing heat treatments 187–8pearlite transformation 188preliminary heat treatments 187–8

quality heat treatment 188–90quenching 188, 189–90, 201stress relief treatment 190tempering 190

heavy-wall thickness piping 546–54helium embrittlement 610–11high initial dislocation density 380–1high pressure-low pressure combination

(HLP) rotors 204–7high strength low-Cr steels 541–6high temperature crack growth 658–62history of creep-resistant steels see

development of creep-resistantsteels

homologous temperature 3–4Hooke’s Law 221hot topping process 181–2Hull-Rimmer equation 345

in situ TEM observations 389–93ingot making process 174, 182–3interdiffusion reactions 245interdiffusion values 258intergranular cracking 330, 353iron and iron-base alloys diffusion

behaviour 255–60irradiation creep 602–5irradiation embrittlement 605–10isostress rupture tests 655–6Italian specifications and standards 79

Japanese specifications see specificationsand standards (Japan)

JIS Code 158–69JSME Code 158

kinetics of dislocation glide 385–7kinetics model 312–13Kirkendall effect 245

lacunae in calculational methods 640–1ladle refining furnaces (LRF) 66, 175,

178–80Larson-Miller parameters 361, 567laths 331lattice diffusion 245–6life assessment see service lifelife extrapolation in Gr.91 steel 433–6life extrapolation in Gr.92 steel 436–40

WPNL2204

Index 673

life-fraction rule 638, 641–3linear damage rule 639lithium 625logarithmic creep 6, 405longitudinal seam welds 472Lorenz number 236low alloy steels 19–26low initial dislocation density 379–80low-cycle fatigue life 653low-temperature tempering 300

machining 190magnetic acoustic emissions 646–7magnetic testing 645magnetic transformation 250–2main steam pipes 541martensitic steels 155, 224

high-Cr 546–54tempered 9-12Cr steels 290–2, 381for turbine components 591

material characteristics 412material specifications 95matrix deformation 248–9mechanical constraint 507–9mechanical properties 155, 169, 173mechanical tests 190–2mechanisms of creep 365–402

austenitic steel microstructure 366composite model 388–9constant stress creep 368–70crystallographic slip 365cyclic creep 374–5, 382cyclic variation of stress 384degradation of particle

hardening 383–4dislocation models 385–9equations 375–6evolution of dislocation structure

model 387ferritic steel microstructure 366–7in situ TEM observations 389–93kinetics of dislocation glide 385–7microstructural interpretation of creep

rate 375–84microstructural model 401–2particle hardening 387–8, 393–4primary creep and loading strain

379–81solid solution hardening 376–8

stress change responses 370–3, 381–2subgrain coarsening 382–3tertiary creep 382–3transmission electron microscope

(TEM) observations 389–93velocity of glide 394

Metallographic Atlas 9metallurgical tests 190–2METI Code 158–69microscopic fracture mechanics 509–13microstructural model 401–2microstructure

of austenitic steels 366changes 249–50degradation 358–9evolution 5of ferritic steels 366–7, 474–82interpretation of creep rate 375–84and precipitation hardening

306–12, 320–2Microstructure Data Sheets 9Miner rule 451minimum creep rate 5, 243modelling

bolt relaxation testing 148in complex systems 312–15composite model 388–9crack formation 512dislocation 385–9evolution of dislocation structure

model 387fitting effectiveness 405–11microstructural model 401–2molecular dynamics (MD)

modelling 347precipitation kinetics model 312–13selection 412

modes of deformation 265, 637modes of fracture 351modulus of elasticity 221–5molecular dynamics (MD)

modelling 347molybdenum steels 20Monkman-Grant relationship 6, 346monotonic creep 382multi region analysis 361–2multiaxial behaviour 456–9multiaxial stress rupture 415multilayer welding 480

WPNL2204

Index674

MX carbonitrides 286, 553

Nabarro-Herring creep 10, 275, 276national standardisation bodies 81, 84–5necking of specimens 4–5Nelder-Mead method 459neutron absorption cross-section 598nickel steel 42–3, 232NIMS Creep Data Sheets 99-12% Cr steels 26–32, 45, 67–8, 290–2,

295–301non-destructive evaluation 643–53, 663non-equilibrium segregation 334non-linear fracture mechanics 504–7normalizing heat treatments 187–8Norton’s law 6notch weakening 17, 581NRIM Creep Data Sheets 9nuclear reactors 597–632

and austenitic steels 615–21core components 615–17, 621–4embrittlement caused by ageing

611–13fabrication and joining 629–31fast breeder reactors (FBRs) 598,

613–24and ferritic steels 621–4, 626–9, 631fission reactions 597–8fusion reactors 625–7helium embrittlement 610–11irradiation creep 602–5irradiation embrittlement 605–10neutron absorption cross-section 598radiation damage 598–611radiation swelling 599–602specifications and standards 79steam generators 627–9thermal nuclear reactors 598turbines 627–9types of reactors 613

nucleation of precipitates 313–14

ODS steels 292, 338–40, 631oil quenching 190, 192open hearth furnaces 64, 66, 175Orowan stress 272, 281–4, 388Orr-Sherby-Dorn (OSD) method 361–2Ostwald ripening 383–4over-tempered region 480

overlay welding 201–2oxidation 7, 519–34

acceptable rates of oxidation 528–9and alloying additions 527–8anomalous temperature

dependence 528and chromium content 527defect development 523–5hydrogen production 520mechanisms of enhanced steam

oxidation 520–5rates of steam oxidation 525–9and service life 530–2spalling of oxide scales 523–5, 530stability of oxides 520–1stages of steam oxidation 521–3testing steam oxidation resistance

519–20voids and gap formation 525

oxidation-resistant steels 532–3coatings 533composition 532–3surface modifications 533

oxide dispersion strengthened (ODS)292, 338–40, 631

oxygen diffusion 250

P91 steel 30, 539parametric extrapolation techniques

654–5and strain analysis 423–5

partially transformed zone 479particle hardening 383–4, 387–8, 393–4

degradation 383–4patterns of strain accumulation 427–33pearlite transformation 188PED (Pressure Equipment Directive) 80,

82–3phased array ultrasonic testing (UT) 645physical and elastic behaviour 217–23

elastic behaviour 219–25electrical resistivity and

conductivity 234–8fatigue stresses 217–18thermal properties 225–34thermal stress parameter 217–18

piping and tubing steels 158plastic deformation 265, 266, 267, 278,

371

WPNL2204

Index 675

Poisson’s ratio 220post weld heat treatment (PWHT) 543power law behaviour 426–7power plant components 539

see also nuclear reactors; turbinespower-law creep 507precipitation at the grain boundary 329,

331–3, 341precipitation hardening 281–4, 305–26,

476–7back-stress concept 322–4effects of precipitates 305kinetics model 312–13loss of strengthening 324–5microstructure analysis 306–12microstructure-property relationships

320–2modelling in complex systems 312–15nucleation of precipitates 313–14precipitate evolution 307–12, 315–20precipitate-dislocation interaction

320–1precipitate-grain boundary interaction

322precipitate-subgrain boundary

interaction 321–2thermodynamic equilibrium analysis

315–16see also Z-phase precipitation

precipitation of new phases 296–8preferential recovery of microstructure

298preferred absorption glide (PAG) creep

603preliminary heat treatments 187–8Pressure Equipment Directive

(PED) 80, 82–3primary creep 4, 405

and loading strain 379–81producer certifications 83production of steels for turbines 174–214

12Cr steels 195–204casting process 174, 175–82CrMoV steel 192–5forging process 174, 183–6future trends 208–12heat treatment process 174, 186–90high pressure-low pressure combination

(HLP) rotors 204–7

ingot making process 174, 182–3machining 190testing and non-destructive

examination 190–2progressive coarsening 298purity of heat-resistant steels 64–7

quality heat treatment 188–90quenching 188, 189–90, 201radiation damage 598–611

helium embrittlement 610–11irradiation creep 602–5irradiation embrittlement 605–10radiation swelling 599–602

rate of creep curves 6–7, 265rates of steam oxidation 525–9reheat cracking 486–8reheaters 540–1, 554–64relaxation testing 147–8requirements for heat-resistant steels

18–19residual life assessment 149–50resistivity 234–8role of constraint 493–4rotor steels 30, 31rotors 22, 578

welded design 590rupture life 352–5rupture strength 7–9, 16–17, 279, 356–61

and microstructuraldegradation 358–9

rupture tests 16–17, 145

secondary creep 4, 5, 405segregation effects 329–30, 333–4, 341–2segregation of minor elements 253–4service life

constitutive equations 412–16estimation 449–56, 662and oxidation 530–2rupture life 352–5of turbine components 583–91

shear stress 220–1short-circuit diffusion 243, 246–8small punch testing 656–7solid solution hardening 279–81, 376–8solute atoms 249spalling of oxide scales 523–5, 530

WPNL2204

Index676

specifications and standards (Europe)78–151

Brite-Euram-Projects 80CEN standards 81, 95–150COST programmes 79–80, 306–12,

584DIN standards 79Euronorms standards 80European Carbon and Steel

Collaboration 79European Commission 81–2European Creep Collaborative

Committee (ECCC) 79, 80,85–92

European Pressure EquipmentResearch Council (EPERC) 80,92–5

future trends 150–1in Germany 78–9in Italy 79national standardisation

bodies 81, 84–5non PED creep applications 83–4and nuclear power 79Pressure Equipment Directive (PED)

80, 82–3producer certifications 83Thermie-Projects 80, 208, 211and turbine manufacturers 84in the UK 79

specifications and standards (Japan)155–73

ASME Code 158, 169JIS Code 158–69JSME Code 158METI Code 158–69piping and tubing steels 158steam turbine steels 169super alloy steels 169types of Japanese steels 155, 158

Srolovitz mechanism 283, 284stability of oxides 520–1stacking-fault energy 254–5stages of creep deformation 3–4stages of steam oxidation 521–3steam generators 627–9

see also turbinessteam oxidation see oxidationsteel melting 64–7

strain analysis 421–42appraisal of data analysis 440–1creep curve shape variations 429–30creep ductility 431–3creep strain curves 403creep-induced strain 422–7damage tolerance values 430–1data rationalization 425–6future trends 441–2life extrapolation in Gr.91 steel 433–6life extrapolation in Gr.92

steel 436–40parametric approaches 423–5patterns of strain accumulation

427–33power law behaviour 426–7practical implications 433–41see also stress-strain behaviour

strain monitoring 647strain-fraction rule 638strengthening mechanisms 279–301

austenitic steels 292–5bainitic low Cr steels 287–9dislocation hardening 284–6, 541dispersion hardening 281–4dissolution of fine

carbonitrides 296–8ferrite effects 300–1loss of creep ductility 298loss of strengthening mechanisms

295–301, 324–5low-temperature tempering 300oxide dispersion strengthened (ODS)

338–40precipitation hardening 281–4,

305–26, 476–7precipitation of new phases 296–8preferential recovery of microstructure

298progressive coarsening 298recovery of excess dislocations 300solid solution hardening 279–81,

376–8sub-boundary hardening 286–7subgrain strengthening 382–3tempered martensitic 9-12Cr steels

290–2, 301stress relaxation testing 657–8stress-strain behaviour 219–21, 449

WPNL2204

Index 677

responses of materials 265–7, 370–3,381–2

stress relief cracking 486–8stress relief treatment 190temperature dependence of

rupture 352–5see also strain analysis

STX 21 research project 38–42sub-boundary hardening 286–7subgrain boundary

migration 393precipitate interaction 321–2

subgrain coarsening 382–3subgrain strengthening 382–3super alloy steels 169superheaters 540–1, 554–64surface modifications 533surface replication 648–51

TAF steel 27Tempaloy A-1 steel 44temperature

austenitising temperature 23ductile-to-brittle transition 337–8,

605–10fracture appearance transition 207high temperature crack

growth 658–62low-temperature tempering 300stability in testing 146strain rate dependence 267–9time-temperature parameter

analysis 361–2tempering 190, 290–2, 301, 38110-12% Cr steels 32–6tertiary creep 4, 382–3, 405testing

accelerated destructive tests 653–8acoustic emissions (AE)

monitoring 647–8at constant load 3at constant stress 7carbide analysis 651–2CEN standards 144–7component creep tests 147data assessment 146deformation upon loading 269DVM creep rate limit test 15–16, 19hardness-based techniques 652–3

heat indication test (HIT) 192isostress rupture tests 655–6magnetic acoustic emissions 646–7magnetic testing 645mechanical tests 190–2metallurgical tests 190–2non-destructive evaluation

643–53, 663parametric extrapolation

techniques 654–5relaxation testing 147–8rupture tests 16–17, 145small punch testing 656–7steam oxidation resistance 519–20strain monitoring 647stress relaxation testing 657–8stress and temperature ranges 265surface replication 648–51temperature stability 146thermocouple calibrations 146thermomechanical fatigue (TMF)

testing 641of turbine steel 190–2ultrasonic testing (UT) 643, 645

thermal fatigue 217–19, 238thermal nuclear reactors 598thermal properties 225–34

conductivity 230–4expansion 225–9

thermal stress parameter 217–18THERMIE project 80, 208, 211thermocouple calibrations 146thermodynamic equilibrium

analysis 315–16thermomechanical fatigue (TMF)

testing 641thick-section components 38–42, 57–62Thomas process 64, 66threshold stress 78time-dependent deformation 242–3, 350time-temperature parameter (TTP)

analysis 361–2TMK 1 and 2 steels 31TP 316 steel 57, 59tracer diffusion values 257transmission electron microscope (TEM)

observations 389–93triple junction mobility 347tritium 625

WPNL2204

Index678

turbines 169, 573–93, 627–9advanced turbines 174and austenitic steels 57–62, 591critical components 574manufacturers specifications and

standards 84material properties 576notch weakening 581rotors 22, 578welded design 590service life of components 583–91see also production of steels for

turbines12Cr steels 195–20420-25% Cr steels 44, 52–3type I cracks 484, 485type II cracks 484, 485type III cracks 484–5, 486–8type IV cracks 485, 491–5, 513–17, 649

UK specifications and standards 79ultimate tensile strength (UTS) 266ultrasonic testing (UT) 643, 645under-stabilising technique 47uniaxial relaxation testing 148user specifications 149

V-type cracks 509–11vacancies 243–4, 260, 354vacuum arc remelting (VAR) 66, 182vacuum carbon deoxidation (VCD) 66,

175vacuum degassing 66, 175vacuum induction melting (VIM) 66,

175, 182vacuum oxygen decarburization

(VOD) 66velocity of glide 394void swelling 599–600voids and gap formation 525

W-type cracks 509–11water quenching 189–90water spray quenching 189–90water walls 539–40

Wedel-Neubauer recommendations 649weld strength factor (WSF) 495–6weld strength reduction factor (SRF) 495welded joints 472–98

arc welding 476CEN standards 142, 144crack formation 484–95, 513–17in dissimilar welds 488–90creep behaviour 483–4delta ferrite residuals 478in fast breeder reactors (FBRs) 630and ferritic steel

microstructure 474–82fusion welding 474–5future trends 496–8grain growth zone 479grain refined zone 479heat affected zone (HAZ) 472,

478–80, 497–8, 513–17HAZ simulation 480–2implications for industry 495–6and life prediction procedures 662longitudinal seam welds 472metal development 482–3multilayer welding 480over-tempered region 480partially transformed zone 479post weld heat treatment (PWHT) 543and precipitation strengthening 476–7weld metal development 482–3zone of unchanged base material 480

welding consumables 142welding procedure

qualifications 142, 144Wiedemann-Franz Law 236

X22CrMoV steel 26–7

yield stress 265, 267–9, 270–5, 277Young’s modulus 218, 221, 224, 238

Z-phase precipitation 68, 239, 320Zener stress 394zirconium 342, 345zone of unchanged base material 480

WPNL2204

WPNL2204

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Creep-resistant steels - allaboutmetallurgy.comallaboutmetallurgy.com/.../uploads/2016/12/Creep-Resistant-Steels.pdf · Creep-resistant steels Edited by Fujio Abe, Torsten-Ulf Kern