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STP 1259

Steel Forgings: Second Volume

Edward G. Nisbett and Albert S. Melilli, editors

ASTM Publication Code Number (PCN):04-012590-02

ASTM100 Barr Harbor DriveWest Conshohocken, PA 19428-2959

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LIbrary of Congress CatalogIng-In-PublIcatIon Data

Steel forgings. Second volume / [edited by] Edward G. Nisbett andAlbert S. Melilli.

p. em."Contains papers presented at the Second Symposium on Steel

Forgings held in ... New Orleans, Louisiana, on November 20-21, 1996... sponsored by ASTM Committee A-1 on Steel, Stainless Steel, andRelated Alloys"uP. iii.

"STP 1259.""ASTM publication code number (PCN) 04-012530-02."Includes bibliographical references and index.ISBN 0-8031-2423-61. Steel forgings--Congresses. I. Nisbett, Edward G.

II. Melilli, Albert S. III. American Society for Testing andMaterials. Committee A-1 on Steel, Stainless Steel, and RelatedAlloys.TS320.S744 1997672.3'32--dc21 97-22639

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To make technical information available as quickly as possible, the peer-reviewed papers inthis publication were prepared "camera-ready" as submitted by the authors.

The quality of the papers in this publication reflects not only the obvious efforts of theauthors and the technical editor(s), but also the work of these peer reviewers. The ASTMCommittee on Publications acknowledges with appreciation their dedication and contributionof time and effort on behalf of ASTM .

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Foreword

This publication, Steel Forgings: Second Volume, contains papers presented at the SecondSymposium on Steel Forgings in Hyatt Regency New Orleans, New Orleans, Louisiana, onNovember 20-21, 1996. The symposium was sponsored by ASTM Committee A-Ion Steel,,Stainless Steel, and Related Alloys. The symposium was chaired by E. G. Nisbett, NationalForge Company; A. S. Melilli, Consultant, Winchester. They also served as editors of thispublication.

ContentsOverview VII

PRESSURE VESSEL AND NUCLEAR FORGINGS

New Materials and Forgings Used for Pressure Vessels Operating inHydrogen Environment-p. BOCQUET,A. CHEVIET,L. COUDREUSE,ANDR. DUMONT 3

Improved Mechanical Properties of the A 508 Class 3 Steel for NuclearPressure Vessel Through Steelmaking-JEONG-TAE KIM, HEE-KYUNGKWON,KOOK-CHULKIM, ANDJOONG-MYOUNGKIM 18

Effects of Composition and Heat Treatment on the Toughness of ASTM A508Grade 3 Class 1Material for Pressure Vessels-MINFA UN.STEVENS. HANSEN,TODDD. NELSON,AND ROBERTB. FOCHT 33

Current Forgings and Their Properties for Steam Generator of NuclearPlant-HISASHI TSUKADA,KOMEISUZUKI,MIKIOKUSUHASHI,AND IKUOSATO 56

Forging Technology Adapted to the Manufacture of Nuclear PWR AusteniticPrimary Piping-FRAN<;:OISE MORIN,PIERREBOCQUET,AND ALAINCHEVIET 65

The Optimization of Mechanical Properties for Nuclear Transportation Casksin ASTM A350 LF5-STEPHEN PRICEAND GRAHAMA. HONEYMAN 79

GENERAL INDUSTRIALFORGINGS

Developments in Forging Ingot Production at Bethforge Inc.-JOHNE. FIELDING,ROBERTB. FOCHT,KENNETHF. REPPERT,ANDEUGENEL. TIHANSKY 93

Application of Nitrogen-Alloyed Martensh.ic Stainless Steels in the AviationIndustrY-GERALD STEIN,WALTERKIRSCHNER,AND JOACHIMLUEG 104

Process Model Development for Optimization of Forged Disk ManufacturingProcesses-cHRISTIAN E. FISCHER,JAYS. GUNASEKERA,AND JAMESC. MALAS 116

Manufacturing and Properties of Continuous Grain Flow Crankshafts forLocomotive and Power Generation Diesel Engines-DANIEL J. ANTOSANDEDWARDG. NISBETT 129

Reducing Stress Related Problems in Steel Forgings Using Sub-HarmonicVibrational EnergY-THoMAS E. HEBEL 148

Development of a Process for Toughening Grain-Refined, High-StrengthSteels-MIcHAEL J. LEAP, JAMESC. WINGERT,AND CHARLESA. MOZDEN 160

The Chemistry Modifications to ASTM A707 for Offshore StructuralIntegritY-MARTIN A. WALSHAND STEPHENPRICE 196

TEST METHODS

Ultrasonic Signal Processing Using Indication Sets for Detection andCharacterization-JOHN R. M. VIERTL 213

Application of Fern-Based Modeling of Open-Die Forging to Product andProcess Development at Bethforge, Inc.-wILLIAM REESEII.DAVIDC. RONEMUS,KEYANGHUANG,EMORYW. ZIMMERSJR., ANDTHIRUVEERABADRAN 224

A Rare Type of Flake-Like Forging Burst in Heavy ForgingS-JOHN E. STEINERAND EDWARDL. MURPHY 241

A New Model for Calculating Maximum Blow Force of Die-ForgingHammer- YUNRUI LI, KEYIYANG,ANDZHENLINNI 249

TURBINE ANDGENERATOR FORGINGS

Martensitic 11 % CrMoNiNb Steel for Turbine Rotors in Geothermal PowerStations-KARL H. SCHONFELD,RALFLEVACHER,MICHAELP. MANNING,ANDPAULF. MURLEY 259

Development of High-Strength 12% Cr Ferritic Steel for Turbine RotorOperating above 600°C-YOICHI TSUDA, MASAYUKIYAMADA,RYUICHIISHII,YASUHIKOTANAKA,TSUKASAAZUMA,AND YASUMIIKEDA 267

Historical Overview of Improving Cleanliness of Rotor Steels for ElectricUtility Applications-RAMASWAMY VISWANATHAN 280

Prediction and Control of Segregations in CrMoV Steel Ingot for MonoblockHLP Rotor Forgings Using Experimental Results Obtained from 8 TonSand Mold Ingots-AKIHIRO ITOH, HITOHISAYAMADA,ANDTOMOOTAKENOUCHI 305

High Strength 12 % Cr Heat Resisting Steel for High Temperature SteamTurbine Blade-RYUICHI ISHII, YOICHITSUDA,AND MASAYUKIYAMADA 317

Manufacturing and Properties of Newly Developed 9% CrMoVNiNbN High-Pressure Low-Pressure Rotor Shaft Forging-TsuKASA AZUMA,YASUHIKOTANAKA,TOHRUISHIGURO,HAJIMEYOSHIDA,AND YASUMIIKEDA 330

Hydrogen and Flaking after 40 Years of Vacuum Degassing-JOHN E. STEINER,EDWARDL. MURPHY,AND ROBERTD.WILLIAMS 344

International Business, Codes, and Material Specifications-MIcHAEL GOLD 353Indexes 363

Overview

Steel is supplied in many product forms, most of which are produced in terms of basicdimensions such as width and thickness, or diameter and with length describing quantity.These products may be used by the foot for example as concrete reinforcing bar, or railroadrails or may be fabricated by bending, and welding into products such as storage tanks.Often they essentially lose their identity in the process. Forgings and castings by contrast arediverse in shape and form and are individually made for a specific purpose, either as selfcontained units such as crankshafts, valve bodies or turbine rotors, or as discrete compo-nents to be fabricated into a larger assembly, as for example a nozzle for a pressure vessel.The specification and testing of forgings is therefore more varied, complex, and demandingthan is the case for other product forms. This is augmented by the fact that forgings are oftenexpected to give better reliability and service performance than can be expected when thesame part is fabricated from sections of other steel product forms, if this were in factpractical. Given these unique circumstances the exchange of ideas on forging manufacturingtechniques and experience, materials data and service experience has been an essentialdriving force in developing forging techniques and applications in every industrial field. Inturn these user driven needs and producer developments for manufacture have promoted thedevelopment of product specific standards that ASTM, by virtue of its organization capabili-ties and goals, is able to supply promptly and efficiently. This then was the underlyingpurpose for both this symposium, and its predecessor held in Williamsburg Virginia inNovember 1984.

The symposium was sponsored by ASTM Committee AOI on Steel, Stainless Steel, andRelated Alloys, and was organized by Subcommittee AO1.06 on Steel Forgings and Billets.The symposium was international both in terms of the papers presented and the attendance.The format of the symposium was similar to that of Williamsburg, focusing on the scope ofthe subcommittee in the areas of pressure vessel and nuclear forgings, turbine and generatorforgings, general industrial forgings, and test methods for forgings. Several of these authorswho contributed to the first symposium also submitted papers for this the second symposiumand so demonstrated an expansion of the developments in their organizations. This wasgratifying because time and financial restraints on travel have had a tendency to reduce theexchange of experience and data between those making steel forgings and those who usethem-to the detr.iment of both. Although the maximum benefit will be gained by those whoboth attended the symposium and obtain this record of the proceedings, it is hoped that thispublication extending as it does the published work of the Williamsburg conference willserve as a valuable reference volume for future forging applications.

The keynote address, developed by Mike Gold at very short notice but with keen insightinto the current way in which business is being done in the international market, shows thatthe traditional way of manufacturing equipment in the established industrialized countriesand exporting it to the underdeveloped nations is changing to the point that the equipmenttends to be built in the destination country itself under a cooperative arrangement. Howeverthere is still a niche where critical components, that may possibly include forgings, are madeby the more experienced producers.

Although forgings for the domestic commercial nuclear applications are limited to thereplacement of items such as steam generators for existing power generating stations, it will

vii

viii STEEL FORGINGS

be seen that the development of new manufacturing techniques, such as the forged stainlesssteel reactor piping units in France that will reduce in service inspection demands andimprove component reliability, and the steam generator forging developments in Japanindicate that the nuclear technology continues to progress. Developments intended to im-prove the mechanical properties of the ASTM A 508 Grade 3 steel, used for many nuclearand other pressure vessel applications, have been described both from domestic and Koreanproducers, and these may result in revisions to that material grade in A 508, a potentialexample of specification development through technical exchange. Developments in pres-sure vessel materials for forgings to be used in high pressure hydrogen environments in thepetrochemical industry, and for the manufacture of spent nuclear fuel transportation casksalso show how progress is being made in other sections of the pressure vessel industry. Thedemand for very large and complex components for high temperature catalytic crackervessels again for the petrochemical industry has spurred material development with conse-quent material specification revisions.

A potpourri of forging information was included in the General Industrial forging session.This included process model development for the optimization of forging disks, and finiteelement modeling for open die forging. Both of these papers were from domestic sources,and illustrate the drive to improve forging techniques. A third paper on the forging processthis time from China discussed forging hammer force calculation. Sub harmonic treatmentof forgings to relieve thermally induced residual stresses and the latest developments in theunique nitrogen alloyed stainless martensitic steels produced in Germany by the pressurizedESR melting process increased the diversity. Other papers in this session also looked atcurrent forging ingot production for the sole remaining domestic producer of very largeopen die forgings. The manufacture of continuous grain flow crankshafts for medium speeddiesel engines is described together with the required materials and properties. The demandfor this product has continued to increase, in part because of the use of natural gas for fueland the potential for high thermal efficiencies when waste heat recovery is included in theinstallation. Improved toughness grain refined high strength steels for forgings are describedin the paper by Leap.

The information given is of a very practical nature and could prove to be useful inspecifying heat treatments. The often used sequence of normalizing, quenching, and tem-pering possibly owes its success to the mechanisms described in that paper. One last area ofinterest here that could lead to specification revision also was the paper from England on thecopper bearing age hardening steels for offshore tension leg platforms. An area of forgingproblems-all too rarely written about, but none-the-Iess real was discussed by two veryexperienced and long time members of the subcommittee. This and their other paper onhydrogen flaking problems-or the apparent lack of them-in forgings gave rise to somespirited discussion which although it does not appear in this account, gave food for thoughtfor those present, and deserves close attention to readers of this volume. The germ of an ideafor future papers on failure analysis in forgings came out of these discussions.

Always a source of information on the extremes of forging application the turbine andgenerator forgings session discussed developments in the martensitic stainless steels forturbine rotors and blades, as well as the combined high pressure-low pressure rotor shaftsin a modified 9CrlMo high temperature steel. A study in the control of segregation inCrMoV steel ingots for the combined high pressure-low pressure rotors was also pre-sented, both papers coming from Japan. The reader's attention is drawn to the excellentreview of the superclean steel forging technology for rotor manufacture that has beenspearheaded by EPRI. This steel making practice, made possible by great strides in steelmaking technology was in its early days at the time of the Williamsburg meeting. IUs being

OVERVIEW ix

extended to the high temperature pressure vessel field as a way to reduce in service embrit-tlement.

The steel forgings subcommittee has developed several widely used standards for special-ized test methods for forgings, and this subject was covered by two papers on the ultrasonicexamination of rotor forgings, one of which is included in this volume. The advantages ofbeing able to record ultrasonic examinations for base reference purposes will spur furtheractivity in this area.

Although forging is an ancient production process long predating the industrial revolu-tion, the development of steel forgings shows no sign of being exhausted, new forgingmachines continue to appear to make better use of the starting material and reduce cost, andnew applications are put forward to meet the expanding needs of industry. Symposia such asthis one will assist in obtaining the best from our resources.

Edward G. NisbettNational Forge Company,

symposium co-chairman and STP editor

Pressure Vessel and Nuclear Forgings

Pierre Bocquetl, Alain Cheviet2, Lionel Coudreusel and Rene Dumont3

NEW MATERIALS AND FORGINGS USED FOR PRESSURE VESSELS OPERATING INHYDROGEN ENVIRONMENT

REFERENCE: Bocquet P., Cheviet A., Coudreuse L. and Dumont R., "New Mate-rials and Forgings Used for Pressure Vessels Operating in Hydrogen Environ-ment" Steel FOf!~infiS: Second Volume. ASTM STP 1259, E.O. Nisbett and A.S.Melilli, Eds, American Society for Testing Materials, 1997.

ABSTRACT : To improve the in-service behaviour of Cr Mo (V) steel grades usedfor the pressure vessels operating in hydrogen environment at high temperature forthe oil industry, the manufacture of heavy forgings needs a high quality. Improve-ment of the standard and enhanced strength (ASME Case 1960) 2 1/1.Cr 1 Mo steelgrades may be achieved by reducing drastically the impurities (S, P, X, etc ...) to ex-tra low level and avoiding segregates at the inner surface of the shells. For high tem-perature operation, new V modified steel grades are proposed (ASME case 1961,Code Case 1973 and Code Case 2098). Their conventional mechanical propertiesare similar to those of enhanced strength 2 1/4 Cr 1 Mo but they offer higher creepproperties and improved resistance to hydrogen damage.

KEYWORDS : hydrogen damage, pressure vessel, Cr Mo (V) steels, high tempera-ture, embrittlement.

The various high pressure vessels commonly used in the refining oil industry(hydrotreaters, hydrodesulfurisers, hydrocrackers) operating usually at high tempe-rature, high pressure as well as high partial hydrogen pressure are exposed, to seve-ral types of damages during their service life.

Made usually from the Cr Mo materials (1.25 Cr - 0.5 Mo, 21/4 Cr 1 Mo) tho-se pressure vessels are susceptible to temper embrittlement at operating tempera-tures (up to 450°C). To estimate this susceptibility to temper embrittlement the me-chanical properties of the material are tested in the as heat treated and in post weldheat treated conditions as well as after the step cooling treatment which is supposedto simulate the e.mbrittlement of the material after long term exposure to high tem-perature.

Furthermore, if locally the temperature happens to rise to a high level, the ma-terial can become creep embrittled.

In conversion processes using high hydrogen partial pressure, the materials areexposed to several types of hydrogen damage.

1Research Scientists, Centre de Recherche des Materiaux du Creusot, Creusot-Loire Industrie, BP 56, Le Creusot, France 71200.

2Product Manager, Creusot-Loire Industrie, BP 68, Rive de Oier, France42800.

3Sales Department, Creusot-Loire Industrie, BP 56, Le Creusot, France 71200.

3

4 STEEL FORGINGS: SECOND VOLUME

The most well known damaging mechanism is the hot hydrogen attack whichoccurs at high service temperature.

Another damage from hydrogen may occur when cooling down the reactor toroom tempemture, mainly for change of catalyst, due to residual tensile stresses andhydrogen embrittlement. Hydrogen embrittlement has very detrimental effect on themechanical properties of the material, especially the fracture toughness in the basemetal and in the weld metal.

And last, because most of the hydrogen conversion processes contain sulfuracids (H2S), the material must be protected against corrosion with a stainless steelcladding. Also, during the cooling down operation the hydrogen content grows at theinterface between base metal and the overlay and is a potential source of crackswhich creates the disbonding, decohesion of overlay from base metal.

The evolution of in-service temperature and pressure W made it necessary todevelop improved materials for a better safety of these pressure vessels.

IMPROVEMENTSTO STANDARDANDENHANCEDSTRENGTH2 1/4 Cr 1 Mo MATE-RIALS

Temoerature and time effects

The temper embrittlement susceptibility of the 2 1/4 Cr 1 Mo steel grade hasfor a long time been related to the high detrimental effect of the impurities, Phospho-rus, Tin, Antimony, Arsenium, of which P plays the most important role.X ppm [(lOP + 5Sb + 4Sn + As).1O-2] and J factor (%) [(P+Sn) (Si+Mn) 104] arethe most commonly used criteria to evaluate this susceptibility for each heat of steel.

At CLI, we estimate the parameter (P+Sn) as the most significant for the eva-luation of long term embrittlement L2."]. No significant shift in transition temperaturemay be expected when P+Sn is lower than 0.010 %.

BOCQUET ET AL. ON NEW MATERIALS AND FORGINGS 5

The scatterband of results of standard 2 1/4 Cr - I Mo materials representativeof products manufactured at the beginning of the eighties, is compared to data obtai-ned more recently on enhanced 2 1/4 Cr - 1 Mo materials.

It can be concluded that enhanced material does not present higher sensitivityto temper embrittlement than the standard one.

Temper embrittlement is known to be enhanced in coarse grain area and mar-tensitic microstructure, so it can be expected that in some areas of the heat affectedzone (HAZ) of welds, the embrittlement effect is higher.

Tests have been made to evaluate the toughness properties of the HAZ of thestainless steel weld overlay near the fusion line where the grain size number is about4 (from ASTM E 112). The base material had P = 0.006 % and Sn = 0.006 % (P+Sn= 0.012 %).

The results of Charpy test in PWHT condition and after step-cooling are pre-sented in Table 1.

TABLE 1 -- Effect of step cooling and PWHT on Chm:py V properties

Low PWHT Standard PWHT

Temperatures 660°C 690°C(Enhanced Material) (Standard Material)

in °C

Base HAZ CO Base HAZ CO

TK54J -90 -90 -87 -90

as PWHT

TK54J -80 -56 -69 -79

after S.c.

~ TK54J +10 +34 +18 +11

S.c. - PWHT

TK + 3 ~T -60 +12 -33 -57

TK = 54 J transition temperature

It is clear that the PWHT temperature has higher effect on the susceptibility totemper embrittlement of the coarse grain of HAZ than that of base material.

Considering the HAZ of weld seams, we can note that the coarse grain areasare very small and the global properties of these HAZ are not so different from thoseof base material.

In such welds, it is generally the weld material which presents the lower tough-ness.

Hydrogen disbonding of the stainless steel weld overlayConcerning the disbonding phenomenon, extensive investigations have been

conducted at CLI's research center (CRMC). Disbonding occurs when the conjunc-tion of hydrogen peak at the interface and of a sensitive microstructure in that area.It has been demonstrated that an increase of carbon content of the interface is a ma-jor factor of disbonding sensitivity [II. SO, it is important to control the carboncontent at the material surface to be clad as shown in Fig. 2.

From Fig. 2 it can be noted that, for a carbon content of 0.15 % at the weld interface,the disbonding tests give 10 % and 30 % respectively of cracked area for ESW andSAW weld overlay. In comparison, for areas with C 0.18 % (20 % of segregation ofcarbon) the disbonding tests give 20 % and 40 % of cracked area respectively forESW and SAW cladding.

To increase safety margin against that risk of cracking, CLI manufactures itsforging shells from hollow ingots which present no segregated areas on the inner sur-face to be clad by stainless steel weld overlay. [1]

CLI also has developed special forging sequences and techniques to manufac-ture, special shapes (such as conical, hemispherical shells etc) from hollow ingot inaddition to simple shells, to control the carbon content of the inner surface. The pho-tograph (Fig. 3) shows a such element.


RECENT EVOLUTION OF MATERIALS

Enhanced strength 2 1/4 Cr 1 Mo material (ASME Code Case 1960) is nowbeing used for the reduction of thickness and weight of big pressure vessels operatingup to 454°C (850°F). However, to increase the design temperature, it was necessary todevelop materials with increased creep resistance compared to the standard or enhan-ced 2 1/4 Cr 1 Mo and with at least equivalent resistance to hydrogen damages.

The V modified steels developed during the eighties are now being used for thenew operating conditions.

The main characteristics of the materials already being used in the industry aregiven in Table 2.

TABLE 2 -- Steels proposed for pressure vessels in high temperature hydrogen ser-vice

Chemical Minimum MechanicalSteel grade composition PWHT properties (MPa)

temperature YS UTSA336 Cl F22 2.25 Cr 1 Mo > 675°C > 310 515/690

Code Case 1960-3 > 650°C > 380 585n60A336 Cl. F21B 3 Cr 1 Mo 0.25 V > 675°C >415 585n60Code Case 1961 TiB

Code Case 2098.1 2.25 Cr 1 Mo 0.25 V > 675°C > 415 585n60A336 Cl F91 9 Cr 1 Mo V Nb N > 730°C > 415 585n60

Code Case 1973


Development of V Modified Cr Mo steels

The 2 1/4 Cr 1 Mo 1/4 V (Code Case 2098) and 3 Cr 1 Mo 1/4 Ti B (Code Case1961) are new materials with only a short industrial experience and may be conside-red still under development. At the Vienna Conference in 1994 121 about ten papersgave a complete view of the situation of these materials which provide allowablestresses similar to the enhanced material but their creep resistance properties allowtheir use at temperature up to 482°C (900°F).

The 9 Cr 1 Mo V Nb N modified steel (called "grade 91 ") initially developpedfor nuclear applications in fast breeder reactors, is largely used in the power plantsand petroleum refineries as pipes and tubes. Research work by CLI has shown the in-terest in such material for future generation of pressure vessels in the oil industry. [QJ

CLI has produced industrial heats (up to 190 tons for 3 Cr 1 Mo 1/4 V Ti B ma-terial) of all these type of new materials and made an extensive characterization oflarge components whose results have been published [IJ [1i]. They confirm the possi-bility to manufacture these new materials successfully for thicknesses at least up to300 mm.

Hereafter we shall discuss the two main areas of interest to the fabricator ofpressure vessels and the enduser, the weldability and the hydrogen resistance, of thesenew materials as compared to standard ones.

WELDABILITY ASPECTS OF NEW MATERIALS

The fabricators of pressure vessels need to control two types of problems asso-ciated to the material properties :- first, they must select welding products and optimize the PWHT conditions to ob-tain the required mechanical properties of the weld metal and base metal;- second, they have to avoid the development of cracks in the weld (HAZ and/orWeld Metal) during welding (cold cracking) and during the PWHT (reheat cracking).

For the first question, the choice of PWHT conditions results from a commonapproach of all concerned parties (fabricator, purchaser of base material, purchaser ofwelding products). The reason is that the main mechanical properties are directly rela-ted to the PWHT conditions (tempering parameter, including time and temperatureeffects). As an example, the curves of evolution of the tensile properties for V modi-fied 2 1/4 Cr and 3 Cr 1 Mo steels with the tempering parameter TP = T (20+logt) areshown in Fig. 4 and Fig. 5 for both base and weld materials.

To achieve creep properties equivalent to base material, the weld material mustcontain Nb addition, but that element decreases significantly the toughness properties.[2.]. It is clear from Fig. 6 that the properties of submerged arc welds (SAW) are high-ly dependant of the PWHT condition and the P content. The margin with the requiredCharpy V-notch level'is increased when reducing UTS for a given chemistry. So it isdesirable for these materials to optimize the PWHT temperature in the upper range inaccordance with the minimum requirement for UTS of base material.

For the V modified 9 Cr I Mo steel (grade 91) a similar approach has to be ma-de for optimizing the materials properties [lQ].

The cracking sensitivity of steels may be evaluated by different testing me-thods :

The cold cracking sensitivity has been determined by the implant test method inview to define the minimum preheating temperature to avoid cracking. Fig. 7 showsthe relation ship between the minimum preheating temperature when applying astress of 500 MPa to the implant sample and the total hydrogen content of weld me-tal. The V addition appears to have no significant effect on the risk of cold cracking.



Concerning the steel Grade 91, the same type of tests show that a minimum preheatingtemperature of 200°C and a low hydrogen content of weld metal are required.

The reheat cracking sensitivity has been evaluated by the measurement of HAZductility at the PWHT temperature. The HAZ coarse grain area is simulated on a tensilespecimen with the Gleeble thermomechanical simulator. The tensile test is performed atPWHT temperature with a relative loading rate of 10-1 %.s-l. The main results presen-ted on Fig. 8 show that the effect of V is detrimental (compare 2.25 Cr 1 Mo and 2.25Cr 1 Mo V), but limited because the high purity of materials, and an increase of Crcontent provides a significant improvement of the ductility (compare 2.25 Cr Mo V and3 Cr 1 Mo V Ti B). The 9 Cr 1 Mo Nb N steel grade is not sensitive to reheat crackingdespite the addition of Nb, another detrimental element.

HYDROGEN RESISTANCE

Hydro~en embrittlement and hydro~en attack

The hydrogen introduced by the processing conditions may induce two types ofdamages to base material and weld seams.

At high temperature (in operation) hydrogen may combines with the carbon of thesteel to form methane (CH4) bubbles. The steel properties may be affected by this de-carburization effect and a loss of ductility and even of strength may be observed.


The high stability of complex carbides formed in these Cr Mo V material gives a veryhigh resistance to this phenomenon.

Another phenomenon due to hydrogen is observed after cooling down. The hy-drogen introduced in the material cannot escape quickly and an embrittlement effectmay be obtained at room temperature (ductility loss).

Both types of hydrogen damages may be characterized by the ductility loss mea-sured by tensile test on specimens containing hydrogen compared to reference speci-mens without hydrogen with the following formula:

DL (Ductility Loss %) = [(RAUC - RAHC) / RAUel x 100RA : Reduction of AreaUC : UnchargedHC : Hydrogen Charged.

Hydrogen attack

To characterize the sensitivity to hydrogen attack, the charging conditions in au-toclave are performed at higher temperature than normal operating temperature inview to accelerate the phenomenon.

Typical charging conditions are 720 to 1000 hrs at 600°C and hydrogen pressureof 150 to 230 bars.

A reference specimen is also exposed to the same temperature and time, but un-der Helium pressure (neutral gas) to evaluate separately the possible softening effectof temperature and time on mechanical properties.

The test results are presented in Fig. 9 where the charging conditions in autocla-ve are noted in the upper right window and the PWHT of materials is noted under thecharts.

It can be observed that all new V modified materials are highly resistant to hy-drogen attack in comparison with standard and enhanced 2 1/4 Cr 1 Mo steels.

Hydrogen embrittlement

To evaluate the sensitivity to hydrogen embrittlement, the charging conditionsare chosen more closely to the operating conditions. In fact, in that type of test, the ob-jective is to introduce a controlled quantity of hydrogen in the specimen and to test itquickly to measure the direct effect of hydrogen content on the steel ductility ortoughness.

Typical charging conditions were selected as follows:- In autoclav.e under 150 bar H2 at 450°C for a few days.- Cathodic charging in an acidic solution (HC1/H2S04)'- H2S charging in saturated NACE solution.

This permitted to cover a range of hydrogen content from 1 to 10 ppm.Fig. 10 shows the ductility loss (noted here F %) versus the hydrogen content

measured on the specimen just after testing.All new V modified materials appear to have an improved resistance to hydrogen

embrittlement when compared with the reference standard 2 1/4 Cr 1 Mo steel grade.It has been shown that Vanadium containing steels have a higher trapping capa-

city due to Vanadium carbide distribution lli]. PWHT which modifies the carbidedistribution and size is a major parameter of material sensitivity.


Other tests performed on specimens taken in weld metals show generally a higher em-brittlement resistance than base materials. This behaviour is attributed to the additivetrapping effect of numerous and homogeneously distributed micro non metallic sphericinclusions on hydrogen atoms. So a a small quantity of hydrogen is trapped on each par-ticle reducing the local embrittling effect.

Disbondin~ of overlay

As described previously for standard and enhanced materials disbonding of overlayoccurs when a high hydrogen peak is present at the interface whose hydrogen sensitivityhas been increased by carbon dilution and diffusion from base materials. V modified CrMo steels are expected to have a higher resistance to disbonding than standard and en-hanced materials for at least two reasons :- one concerns the higher trapping effect of hydrogen in the base material whose conse-quence may be that the hydrogen peak intensity at the overlay interface will be reduced,- the second is related to the higher stability of the V and complex carbides formed du-ring PWHT whose consequence is the carbon diffusion to the interface is lower than forstandard material.

Calculations have been made U2J for similar operating and shut down sequencesto evaluate the hydrogen peak intensity at the interface : in comparison with standard 21/4 Cr 1 Mo, the hydrogen peak intensity has been found to be reduced by 2 times for 9Cr 1 Mo V Nb N steel grade and more than 10 times for 3 Cr 1 Mo 1/4 V steel grade.

The best way to evaluate the disbonding resistance is to perform high severity testswhich consist in hydrogen charging of representative specimens and then cool downquickly to maximalise the hydrogen peak at the interface of the overlay. Then, ultrasonicdetection is practiced on the specimen to measure the percentage of disbonded interface.[12]

Ten cladding conditions on both 2 1/4 Cr and 3 Cr 1 Mo V modified steels havebeen tested (Tables 3 and 4) with more severe hydrogen charging and cooling conditionsthan the most severe applied to the reference standard material : temperature 482°C,pressure 150 bars, time 48 hours and rapid cooling (675°C/hr). On V modified mate-rials, no disbonding was observed on any sample. For similar severe testing conditions,except temperature at 454°C, the best results obtained on standard material were of 20% of disbonding Ll.3J and it was shown that only a deembrittling heat treatment of theinterface gave no disbonding for such test condition.

So, these tests results confirm the very good resistance of the overlay deposited onV modified steels to the disbonding phenomenon. [ll]


TABLE 3 -- Overlay welding conditions on 3 Cr I Mo 114 Ti B material for disbanding tests

Welding Condition ESW 309 LNB SAW 309L SAW 347 SAW 347

+ ESW 316L + ESW 347

PWHT 690°C-25 h 690°C-25 h 690°C-20 h/40 h 690°C-20 h/40 h

C 0.045 0.030 0.050 0.037

Mn 1.54 1.37 1.29 1.46b' Si 0.52 0.35 0.86 0.40'"'§

Cr 20.5 18.5 18.2 19.0.cu

~ Mo 0.18 2.24 0.29 0.24"5> Ni 10.7 12.0 9.7 10.20

Nb 0.64 - 0.51 0.40

V 0.10 0.06 - -

TABLE 4 -- Overlay welding conditions on 2114 Cr 1 Mo 114material for disbanding tests

Welding Condition SAW 347 SAW 347 SAW 309L Nb ESW 309L Nb ESW 308L Nt ESW 308L Nb

60mm +ESW 347 + ESW 347 90mm 60mm

PWHT 690°C-20 h/40 h J90°C-20 h/40 h 705°C-30 h 705°C-30 h 705°C-30 h 705°C-30 h

C 0.038 0.036 0.030 0.032 0.037 0.032

~ Mn 1.28 1.22 1.21 1.36 1.20 1.08'"'s Si 0.82 0.80 0.35 0.27 0.25 0.29~u Cr 19.0 19.1 19.2 19.8 18.4 18.6

i Mo 0.16 0.14 - - - ->0 Ni 10.3 10.4 10.3 10.6 9.4 9.5

Nb 0.54 0.52 0.41 0.45 0.47 0.41


CONCLUSIONS

The steel grade 2 1/4 Cr 1 Mo has a large background and it will remain for a longtime the main reference for materials used in hydrogen environment in the petrochemicalindustry. A lot of improvements have been made to make this material safer than in thepast and ensure a long life time to reactors built recently with a such material. The en-hanced steel grade (ASME Code Case 1960) has now also been largely used in the re-cent years for the construction of vessels of large size to save weight and constructioncosts.

The new V modified materials, in particular 2 1/4 and 3 Cr 1 Mo V steel grades arebeginning to be used for the construction of new units, and because their exceptionnalproperties of resistance to hydrogen they are of major interest for the increase of serviceconditions (temperature and hydrogen pressure). They require only a more drasticcontrol of the PWHT conditions than standard material, all other manufacturing condi-tions (welding precautions) been very similar to those of standard materials. They are thematerials of to-day and to-morrow for advanced heavy petrochemical reactors.

Grade 91 steel appears as an alternative material for longer future because its possi-bilities for increasing the temperature to higher step.

REFERENCES

W Humphries MJ., "Performance requirements for hydroprocessing reactormaterials", Creusot-Loire Industrie Seminar: Two-day workshop on newgeneration on Cr Mo and Cr Mo V steels for use in refining reactors in Europe,Abbaye de Collonges, France, December 9-10, 1992.

[2.] Blondeau R.P., Berthet J.A., Duranseaud J.M. and Roux J.H., API Technicalsession on Temper Embrittlement, Chicago, Illinois, May 1981.

[3.] Blondeau R.P., Berthet J.A., Cheviet A., Duranseaud J.M. and Pressouyre G.M."Contribution to a solution to the disbonding problem in 2 1/4 Cr 1 Mo heavy wallreactors" Proceedin~s of 1st Int. Conference on Current Solutions to HydroL!enProblems in Steels, e.G. Interrante and G.M. Pressouyre, Eds, A.S.M., p. 356,Washington DC, 1982.

[±J Poitrault I., Coet 1., Dumont R., Bocquet P. and Blondeau R. "Hollow ingot, thebetter way to high quality forgings for pressure vessels" 32nd Mechanical Workingand Steel Processing Conference, Cincinnati, Ohio, October 1990.

[5] Second International Conference on Interaction of Steels with Hydrogen inPetroleum Industry Pressure Vessel on Pipeline Service, Vienna, Austria, October19-21, 1994. Conference organizer: The Materials Properties Council.

L6J Blondeau R., Bocquet P. and Cheviet A. "Development of a Modified 9 Cr 1 MoSteel for Use in Hydrotreating Reactors" The 1990 Pressure Vessels and PipingConference, Nashville, Tenessee, June 17-21, 1990. ASME PVP, Vol. 201 - MPCVol. 31.


[1] Bocquet P., Bourges Ph., Burlat J and Cheviet A. "Manufacturing experience ofthick plates for pressure vessels" Proceedings of ECSC Information Day - TheManufacture and Properties of Steel 91 for the Power Plant and Process Industries,Diisseldorf, Germany, November 5th, 1992.

I]] Bocquet P., Cheviet A. and Dumont R. "Recent Evolution in High Pressure Vesselsfor the Petrochemical Industry: Materials and Products". 12th Int. ForgemastersConference, Chicago, illinois, September 11-17, 1994.

[2J Hauck G., Demuzere R., Fuchs R. and Bocquet P. "Welding and strip cladding ofthe new 2 1/4 Cr and 3 Cr 1 Mo 1/4 V steels for hydrocrakers" AWS Convention,Chicago, illinois, April 1996.

[1QJ Bocquet P., Bourges Ph. and Cheviet A. "Properties of heavy components of steelgrade 91 and their welds" Nuclear Engineering and Design 144 (1993) 149-154.

[11] Coudreuse L., Bocquet P. and Cheviet A. "Hydrogen trapping in Cr-Mo steels forhydroprocessing reactors" The 1992 Pressure Vessels and Piping Conference,New Orleans, Louisiana, June 21-25, 1992. ASME PVP Vol. 239 - MPC - Vol. 33 -p.I13.

Ul] Pressouyre G.M., Chaillet J.M. and Valette G. "Parameters affecting the hydrogendisbonding of austenitic stainless cladded steels" Proceedings of 1st Int. Conferenceon Current Solutions to Hydrogen Problems in Steels, C.G. Interrante and G.M.Pressouyre,Eds, A.S.M., p. 349, Washington DC, 1982.

LUI Vignes A., Palengat R. and Bocquet P. "Disbonding mechanism and its prevention"Proceedings of 1st Int. Conference on Interaction of Steels with Hydrogen inPetroleum Pressure Vessel Service, Paris, France, March 28-30, 1989. ConferenceOrganizer : The Materials Properties Council.

Ll.4] Shimomura 1., Sugie E., Nakano Y., Nakano S., Oka Y. and Ueda S. "Hydrogenattack and overlay disbonding in 2 1/4 Cr 1 Mo steels" Same conference.

Jeong- Tae Kiml, Hee-Kyung Kwon2, Kook-Chul Kim2, and Joong-Myoung Kim3

Improved Mechanical Properties of the A 508 Class 3 Steel for Nuclear PressureVessel Through Steelmaking

REFERENCE: Kim J T., Kwon H. K, Kim K c., and Kim J M., "ImprovedMechanical Properties of the A 508 Class 3 Steel for Nuclear Pressure VesselThrough Steelmaking" , Steel Forgings: Second Yolume ASTM STP 1259, E.G.Nisbett and A. S. Melilli, Eds., American Society for Testing and Materials, 1997.

ABSTRACT: The present work is concerned with the steelmaking practices whichimprove the mechanical properties of the A 508 class 3 steel for reactor pressure vessel.Three kinds of steelmaking practices were applied to manufacture the forged heavy wallshell for reactor pressure vessel, that is, the vacuum carbon deoxidation(YCD), modifiedYCD containing aluminum and silicon-killing. The segregation of the chemical elementsthrough the thickness was quite small so that the variations of the tensile properties atroom temperature were small and the anisotropy of the impact properties was hardlyobserved regardless of the steelmaking practices. The Charpy Y-notch impact propertiesand the reference nil-ductile transition temperature by drop weight test were significantlyimproved by the modified YCD and silicon-killing as compared with those of the steel byYCD. Moreover, the plane strain fracture toughness values of the materials by modifiedYCD and silicon-killing practices was much higher than those of the steel by YCD. Thesewere resulted from the fining of austenite grain size. It was observed that the grain sizewas below 20~m (ASTM No. 8.5) when using the modified YCD and silicon-killing ,compared to 50~m (ASTM No. 7.0)when using YCD.

KEY WORDS: A508 class 3 steel, reactor pressure vessel(RPY), mechanical properties,transition temperatures, steelmaking practice, YCD, modified YCD, silicon-killing, grainsize, fracture toughness

The reactor pressure vessel(RPY) for the nuclear power plants has large dimensions andheavy wall thickness. Strict requirements are imposed on the materials to provide adequatesafety margins against the possibility of failure. The forgings and plates in manganese-nickel-molybdenum low alloy ferritic steels designated according to the ASTM A 508class 3 and A 533 Type B class 1 are the predominant materials used in nuclear pressurevessels This steel is able to be manufactured by YCD and the silicon-killing practices.

IProfessional Researcher and 2Researcher, Respectively, Research and DevelopmentCenter,3Yice President, Castings and Forgings Business Division, Korea Heavy Industries andConstruction Co., LTD. 555 Guygok-Dong, Changwon, 641-792, Korea

18

KIM ET AL. ON IMPROVED MECHANICAL PROPERTIES 19

Steelmakers have been constantly developing the steelmaking techniques to improveand adjust the quality of their products to increase the safety of the nuclear power plants.Many investigators[l, 2] have studied the effects of the steelmaking practices on themechanical and metallurgical properties of the A 508 Class 3 steel. P. Bernabei et al[l]recommended VCD process for the A 508 class 3 steel, for a given grain size, because theCharpy impact energy after welding was always higher than that in quenched andtempered condition. Comparing with the steels manufactured by various steelmakingpractices, they observed that the VCD steel showed better value of fracture appearancetransition temperature(F ATT) and absorbed energy than the conventionally killed melts.When the heavy wall thick steel was manufactured by VCD, however, the grain size wascoarsened and the secure toughness properties were not obtainable. Therefore, manyforgemasters[2 - Q] have preferred to add aluminium or deoxidize by silicon to producethe better forging shell for RPV.

Since the influence of the steelmaking practices on the mechanical properties of theRPV steel had not been systematically investigated, we studied to improve the mechanicalproperties of the A 508 class 3 steel through the steelmaking practices. The conventionalhot pierced ingots for shell forgings were manufactured by each steelmaking practice, thatis, VCD, modified VCD with aluminium, and silicon-killing practices. In case of themodified VCD practice, the aluminium was added after the vacuum treatment. In thesilicon-killing practice, the silicon was added during the vacuum treatment, and thealuminium was also treated after the vacuum treatment.

The improved results of metallurgical quality and mechanical properties of the RPV steelthrough the steelmaking practices are presented in this paper. Especially, the improvedresults of the impact toughness with the applied steelmaking practice were described withthe size of effective grain within prior austenite grain.

EXPERIMENTAL PROCEDURES

Several shell forgings for RPV of 1 000 MW nuclear power plants were manufacturedby using ingots of 270-ton in weight, which were made by the vacuum stream degassingmethod. The geometry of the ingot is 3 090mm in mean diameter and 3 865mm in height.The chemical compositions at ~ -thickness location of each forged shell are described inTable 1. It should be noted that the contents of copper and phosphorus, which aredeleterious elements for the resistance for the neutron irradiation embrittlement[l, ~],were suffici~ntly low. The copper and phosphorus contents in the present materials variedfrom 0.02 to 0.06 weight% and from 0.005 to 0.008 weight%, respectively. These levelsare remarkably small compared with the ASTM A 508 class 3 requirements.The dimensions of the examined shells were 300mm in thickness, 4 100mm in inner

diameter, and 4 700mm in length. The forged shells were normalized at 900°C, quenchedfrom 880 - 900 °C in water, and then tempered at 650°C. The cooling rates duringquenching in water were measured with the thermocouples attached at various thicknesslocations of the shell. The measured cooling curves are depicted in Fig. 1. The simulatedpost weld heat treatment for the test blocks was executed for 32h at 620 - 630°C. Thetest coupons were taken out from the shells as per ASTM A 370.

FIG. 1--The cooling curves during quenching in water at various thickness of the shell.

RESULTS AND DISCUSSION

FIG. 8 -- Anisotropy of the 68Joule transition temperature(vTr68) and fracture appearancetransition temperature(FATT) of the A 508 class 3 steel produced with each steelmakingpractices. The specimens were taken out at the 1/4 thickness of the shell.

tests were relatively small and the scatter of data was also small. Figure 8 describes ananisotropy of the vTr68 and FATT measured in the both transverse and longitudinaldirections. The differences of the transition temperatures between transverse andlongitudinal directions were scarcely observed in each material produced with three steel-making practices. These homogeneity of mechanical properties by the Charpy V-notchimpact and tensile tests could be achieved by homogeneous chemical compositionsthrough the thickness of the large shell and by accurate control of quenching andtempering temperatures.

The mechanical properties of the materials by the silicon-killing and modified VCDtreated aluminium were generally superior to those of the material by the VCD. These areresulted from the reduction of sulphur content as shown in Table 1 and the refining of theaustenite grain size. The non-metallic inclusion of manganese-sulphide in the steel bysilicon-killing practice, whose sulphur content of 0.002%, was hardly observed in thefractured surfaces of the Charpy impact specimen. Figure 9 depicts the opticalmicrographs of each material manufactured by three steelmaking practices. The austenitegrain size of the aluminium treated materials was much finer than that of the VCD steel. Itshould be noted that the austenite grain size of the modified VCD and silicon-killed steelswas below 20f.lm, and that of the VCD steel was nearly 50f.lm.Figure 10 shows the effects of prior austenite grain size and mean effective grain size on

the vTr68 and FATT of the materials produced by three steelmaking practices. Thedependence of the impact transition temperatures on the cooling rate and steelmakingpractices was great in the ferritic steel. When the cooling rate from quenching temperature

Fig. 9 --Optical Micrographies of the A 508 class 3 steel produced with steelmakingpractices. VCD(a), modified VCD(b) and silicon-killing(c).

between 800 and 500 'c was above 100 'c /min, the FATT was below -70 'C. However,when the cooling rate at the center location was near 20'C/min, the FATT was nearo 'c . The prior austenite grain size at surface and inner the ~ -thickness location wasalmost same. The .effect of the cooling rate on transition temperatures was much greaterthan the effect of the prior austenite grain size on those. By the effect of a prior austenitegrain size, it was difficult to explain the reasons why the secure impact properties at thesurfaces of the each shell were obtained. Therefore, the effective grain size within the prioraustenite grain was introduced to explain the variations of the impact properties throughthe thickness. Figure 11 shows the SEM fractography of the impact specimen fracturednear lower shelf region. The space between the two neighboring tear lines indicated byarrows in Fig. 11 is the effective grain size. The effective grain size is defined as a regionin which cleavage cracks by impact fracture go through in a nearly straight fashion[lO]. Alinear relationship was obtained between the impact transition temperatures(vTr68 andFATT) and <d>-V" where d is the mean effective ferrite grain size. As the cooling ratefrom quenching temperature increased, the austenite grain size decreased, and the effectivegrain size also decreased and the impact properties was significantly improved.

FIG. 11 --SEM fractographies of the A 508 class 3 steel fractured at lower shelf energyregion tear lines are indicated by arrows. (a) Cooling rate of 100aC/min from 800aC to500aC,(b) Cooling rate of20aC/min from 800aC to 500ae.

FIG. 12 --Dependence ofK1Cand KQ(J)values on the normalized temperature of the A 508Class 3 steels with steelmaking practices.

Figure 12 shows the fracture toughness values of the materials manufactured by threesteelmaking practices and also depicts the minimum K1Ccurve of the RPV steel requiredby the ASME section XI[llJ. The K1Cvalues of the steels produced by three steelmakingpractices were higher than the K1Cspecified in the ASME section XI. At the above uppershelf energy region by the Charpy V-notch impact test, as the test temperature increased,K1Cvalues decreased. The K1Cvalues at near room temperature were great in the order ofthe steel manufactured by the silicon-killing, modified VCD, and VCD practice. The KICvalues of the materials by the silicon-killing and modified VCD were no difference with thesteelmaking practices, however, these results were better than those of the VCD at 288·C .It is noted that this K1Ccurve could be expressed with the function of the normalizedtemperature by subtracting RTNDT from the test temperature. The minimum specified K1Ccurve[lll is useful to evaluate the safety of the materials containing the pre-crack and/or apotential crack. As the fracture toughness value increases, the safety margin for the RPVcontaining a potential crack increases, and long residual life of RPV after detection ofcrack can be estimated under cyclic loading conditions. Considering the tensile and CharpyV-notch impact properties, it seems that the higher yield strength and upper shelf energy,the higher fracture toughness. The fracture toughness values of the steels by the modifiedVCD and silicon-killing practices were much securer than those of the steel by the VCD.Therefore, the mechanical properties of A 508 class 3 steel for RPV could be improvedthrough steelmaking practices which added aluminium, which resulted in the refining ofaustenite grain size.

KIM ET AL. ON IMPROVED MECHANICAL PROPERTIES 31

CONCLUDING REMARKS

The present work examines the improvement of the mechanical properties of the A 508class 3 steel for reactor pressure vessel through the steelmaking practices. The appliedsteelmaking practices to manufacture the forged heavy wall shell were VCD, modifiedVCD containing aluminium and silicon-killing. The segregation of the chemical elementsthrough the thickness was quite small so that the variations of the tensile properties atroom temperature were small and the anisotropy of the impact properties was hardlyobserved regardless of the steelmaking practices. However, the impact properties atsurfaces of the shell were superior to those at the ~ -thickness and center locations. Thesewere resulted from the difference of cooling rate between the surface and inner location.These results could be explained by the effective grain size within the prior austenite grain,that is, as the cooling rate increased, the effective grain size was refined. The Charpy V-notch impact properties and the reference nil-ductile transition temperature by drop weighttest were significantly improved by the modified VCD and silicon-killing as compared withthose of the steel by VCD. Moreover, the fracture toughness values of the materials by themodified VCD and silicon-killing practices were much higher than those of the steel byVCD. These were resulted from the fine austenite grain size. It was observed that thegrain size was below 20J..lmwhen using the modified VCD and silicon-killing, compared towas 50J..lmwhen using VCD.

REFERENCES

[l] Bernabei, P, Callegari, L., Scepi, M., and Salinetti, T.,"Seamless Shell CourseForgings for Heavy Wall Reactor Vessels: A Forgemaster's Critical Review", SteelForging, ASTM STP 903, E. G. Nisbett and A. S. Melilli, Eds., American Societyfor Testing and Materials, Philadelphia, 1986, pp 275-300.

[2] Erve, M., Papouscheck, F, Fischer, K, and Maidorn, Ch., "State of the Art in theManufacture of Heavy Forgings for Reactor Components in the Federal Republic ofGermany", Nuclear Engineering and Design, Vol. 108, 1988, pp 487-495.m Haverkamp, K D., Forch, K, Piehl, K H., and Witte, w., "Effect of Heat Treatmentand Precipitation State on Toughness of Heavy Section Mn-Mo-Ni Steel for NuclearPower Plants Components", Nuclear Engineering and Design, Vol. 81, 1984, pp207-217.

[1] Sakai, Y., Kikutake, T., Kosaki, A., Takahashi, L, Miya, K, and Ando, Y, "A Studyon the Fracture Toughness of Heavy Section Steel Plates and Forgings for NuclearPressure Vessels Produced in Japan", Journal of High Pressure Institute of JapanVol. 22, No.4, 1984, pp 210-220.m Onodera, S, Kawaguchi, S., Tsukada, H., Moritani, H., Suzuki, K, and Sato, L,"Manufacturing of Ultra-Large Diameter 20 MnMoNi55 Steel Forgings for ReactorPressure Vessels and Their Properties", Nuclear Engineering and Design, Vol. 84,1985, pp 261-272.

[~] Kawaguchi, S., Tsukada, H., Suzuki, K, Sato, L, and Onodera, S., "Manufacturingof Large and Intergral-Type Steel Forgings for Nuclear Steam Supply System


Components", Steel Forgings, ASTM STP 903, E. G. Nisbett and A. S. MeliIIi, Eds.,American Society for Testing and Materials, Philadelphia, 1986, pp 398-409.

[1] Steele, L. E., Neutron Irradiation Embrittlement of Reactor Pressure Vessel Steels,International Atomic Energy Agency, Vienna, 1975.

[Ji] US Nuclear Regulatory Commission Nuclear Regulatory Guide 1.99 Revision 1,1977.

[2] US Nuclear Regulatory Commission: lOCFR 50, Appendix G, Fracture ToughnessRequirements, 1985.

[ill] Matsuda, S., Inoue, T., Mimura, H., and Okamura, Y., "Toughness and EffectiveGrain Size in Heat-Treated Low-Alloy High-Strength Steels", in Toward ImproveDuctility and Toughness, Climax Modybdenum Company Publication, 1971, pp 45-67.

[il] ASME Boiler and Pressure Vessel Code Section XI, Division 1, Appendix A, 1992.

Minfa Lin', Steven S. Hansen', Todd D. Nelson2, and Robert B. Fochf

EFFECTS OF COMPOSITION AND HEAT TREATMENT ON THETOUGHNESS OF ASTM A508 GRADE 3 CLASS 1 MATERIAL FORPRESSURE VESSELS

REFERENCE: Lin, M., Hansen, S. S., Nelson, T. D., and Focht, R. B., "Effects ofComposition and Heat Treatment on the Toughness of ASTM A508 Grade 3 Class 1Material for Pressure Vessels," Steel Forgings: Second Volume, ASTM STP 1259, E.G. Nisbett and A. S. Melilli, Eds., American Society for Testing and Materials,Philadelphia, 1997.

ABSTRACT: A laboratory study was conducted to evaluate the effects ofcomposition and heat treatment on the toughness of ASTM A508 Grade 3 Class 1material for pressure vessels. Five steels were vacuum induction melted and cast asingots in the laboratory. These heats included a base steel representing thespecification mid-range analysis, a steel containing higher levels of Si, Ni, and Cr(high-side composition) as compared to the base steel, and three steels derived fromthe high-side composition by adding AI, AIIN, and Nb, respectively. The ingots wererolled to plate, heat treated, and evaluated. Among these steels, the high-sidecomposition with additions of AI and N displays the best strength/toughnesscombination. For example, a 75 mm-thick plate of this steel has acceptable strengthand a reference nil ductility transition temperature (RTNDT) of::; -29°C afteraustenitizing at 875°C, air cooling, and tempering at 660°C for up to 20 hours.Upper-nose temper embrittlement (UNTE) occurs in all these steels. This UNTE isattributed to the precipitation of needle-like Mo-rich carbides during tempering, and issignificantly reduced by increasing the cooling rate after austenitizing.

,Research Engineer of Hot Rolled Product Development and Division Manager

of Cold Rolled & Coated Sheet, respectively, Research Department, Bethlehem SteelCorporation, Bethlehem, PA 18016.

2 Managers of Quality Control & Development and Primary Operations,respectively, BethForge, Inc., Bethlehem Steel Corporation, Bethlehem, PA 18016.

33


KEYWORDS: A50X Grade 3 Class 1, toughness, Al/N additions, austenitizing,tempering, upper-nose temper embrittlement, pressure vessels, reference nil ductilitytransition temperature

INTRODUCTION

The Westinghouse advanced nuclear reactor program has selected ASTM A508Grade 3 Class 1 as the material of choice for pressure vessels [1]. For thisapplication, a reference nil ductility transition temperature (RTNDT) of -29°C or loweris required. A laboratory study was conducted to determine if the RTNDT requirementcould be met within the compositional limits of the ASTM A508 Grade 3specification. In this study, five laboratory heats were prepared, rolled to plate, heattreated, and evaluated.

EXPERIMENT AL PROCEDURE

Steel Design

The compositions of the five laboratory-melted steels are listed in Table 1,compared with the chemistry requirement for ASTM A508 Grade 3 material. In allthe steels, the C level is kept at 0.19% to ensure a reasonable strength/toughnessbalance; increasing the C content increases the strength, but decreases toughness [2].The Mo content is maintained at nominally 0.50% in all the heats to provide sufficienthardenability, while minimizing the formation of MozC type carbides which are alsodetrimental to toughness 131. Among these heats, the "Base" composition representsthe mid-range analysis of the A50X Grade 3 specification. The steel designated"High" contains higher levels of Si, Ni, and Cr as compared to the Base steel, and wasprepared to evaluate the effect of increased hardenability on toughness. The calculatedhardenability factors (01) for the Base and High steels are 65 and 88 mm, respectively,based on ASTM A255. The last three steels in the table, "High+Al", "High+Al/N",and "High+Nb" are variations on the High steel obtained by adding AI, Al/N, and Nb,respectively. These alloy additions were made in an attempt to refine the austenitegrain size obtained on reheating, and thereby improve the toughness [2,3].

The steels were prepared as 190 kg, vacuum-induction melted ingots. Eachingot was initially heated to l230°C, held for 5 hours, hot rolled to 50 mm- and75 mm-thick plates, and then air cooled to room temperature.

Preliminary Heat Treatment

The as-rolled plates were austenitized at 900°C for 6 hours and air cooled(4°C/min) to room temperature. The plates were then tempered at 650°C for 10 hoursand air cooled to room temperature to obtain a spheroidized microstructure.

UN ET AL. ON EFFECTS OF COMPOSITION 35


Final Heat Treatment

After preliminary heat treatment, the plates were austenitized at 875°C for 20hours. Following austenitizing, the 75 mm-thick plates were air cooled (9°C/min),while the 50 mm-thick plates were either fan cooled (22°C/min) or quenched into apolymer/water solution (2.5% aqueous solution of Aqua Quench 1103, cooling rate of51}5°C/min). These different woling media were used to simulate the cooling rates (atthe mid-thickness location) of water-quenched production forgings of 560, 300, and 50mm thick, respectively, and to evaluate the effect of cooling rate after austenitizing onthe mechanical properties of the steels. After austenitizing, the plates were temperedat ()()O°Cfor various times ranging from 2 to 50 hours, and air cooled to roomtemperature. The minimum tempering temperature and time given in the specificationare 050°C and 0.5 hour /25 mm of thickness, respectively. Accordingly, theminimum tempering time required for the thickest forging simulated (560 mm) is 11hours. Thus, the range of tempering times (up to 50 hours) evaluated in this study hasadequately covered the effects of retempering and repeated stress relief and post-weldheat treatments of production forgings.

Dilatometry

Cylindrical pins measuring 4 mm (diameter) x 8 mm (length) were machinedfor the determination of the critical temperatures on heating, and the transformationtemperatures on cooling, using a Materials Measuring Corporation computer-controlleddilatometer. The critical temperatures on heating were measured using a heating rateof 55°C/hour. The transformation temperatures on cooling (such as, bainite-start (Bs)

and bainite-finish (Br) temperatures), from an austenitizing temperature of 900°C, weredetermined for cooling rates of 30 and 600°C/min, which approximate the coolingrates of fan cooled and aqua-quenched 50 tnm-thick plates, respectively.

Mechanical Testing

Two transverse round-bar tensile (13 mm diameter) and ten full-size transverseCharpy V-notch (CVN) specimens were machined from the plate mid-thicknesslocation. The tensile sped mens were tested at room temperature, while the Charpyspecimens were initially broken in triplicate at -29 and +4°C. The remaining Charpysped mens were then used to fill in the energy, fracture appearance, and lateralexpansion transition curves. The selection of -29 and +4°C for the initial Charpy testswas based on the required RTNilI'temperature of -29°C. The nil ductility transitiontemperature (TNI>T)was determined in accordance with the provisions of ASTM E208using standard P-3 drop weight specimens (16 x 51 x 127 mm) and a "single-pass"crack-starter weld bead. The drop weight specimens were also machined from theplate mid-thkkness location. When needed, additional Charpy specimens wereprepared and tested at TNIlT+ 33°C to determine the RTNDT'as described in the

3 Aqua-Quench II()is a tradename of E. F. Houghton & Co.


American Society of Mechanical Engineers Boiler and Pressure Vessel Code [4]. Atthis temperature, each Charpy specimen should have a lateral expansion of at leastn.'} mm, and an absorbed energy of 68 J or more. If these requirements are met, theTNIlTbecomes the reference temperature RTNDT.However, if the Charpy values atTNllT+ 33°C are below these requirements, then additional Charpy tests must beconducted to determine the temperature (TcJ at which the 0.9. mm/68 J requirementsare met. In the latter case, the RTNDTis equal to Tcv - 33°C. Thus, the referencetemperature RTNDTis the higher of TNDTor Tcv - 33°C.

Metallography

Specimens for optical microscopy were prepared according to standard pro-cedures, and etched with picral/nital. Two-stage carbon extraction replicas wereprepared and examined in a Philips 300 transmission electron microscope (TEM) and ava Scientific model HB 501 dedicated scanning transmission electron microscope(STEM), both operated at lOOkv. Scanning electron microscopy (SEM) was used toassess the Charpy fracture surface morphology of selected samples.

RESULTS AND DISCUSSION

Critical and Austenite Transformation Temperatures

The AC1 temperatures of the experimental steels are determined to be about710°C, which is above the tempering temperature of 660°C. This confirms that tem-pering was conducted in the single-phase ferrite region. The AC3 temperatures areabout 8600C, which is 15°C below the austenitizing temperature of 875°C, indicatingthat the steels were also fully austenitized. The Bs temperatures of the Base steel areabout 600 and 560°C for cooling rates of 30 and 600°C/min, respectively. Asexpected, the Bs temperatures for the other four steels are lower due to the increasedalloy levels. For example, the Bs temperatures of the High+AVN steel are about 590and 520°C for cooling rates of 30 and 600°C/min, respectively.

As-Quenched Hardness/Microstructure

As-quenched through-thickness hardness profiles were determined for all theas-cooled plates. As an example, the profiles for the 75 mm air-cooled plates (insimulation of 560 mm water-quenched forgings) are shown in Figure 1. The hardnessprofiles of the 50 mm fan-cooled plates (in simulation of 300 mm water-quenchedforgings) are very similar to these profiles, while the hardness levels of the 50 mmaqua-quenched plates (in simulation of 50 mm water-quenched forgings) are about 13Rc higher for a given steel. For each simulation, the Base steel, due to its loweralloying level, exhibits lower hardness than the other four steels. As shown inFigure I, these higher-alloy steels all have similar hardness levels. The microstruc-tures of 75 mm air-cooled and 50 mm aqua-quenched plates of the Base andHigh+Al/N steels are shown in Figure 2. In general, the microstructures of the

relatively slowly-cooled (75 mm air-cool and 50 mm fan-cool) plates are mixtures offerrite, pearlite, and bainite. Consistent with its lower hardness level, the base steelcontains more ferrite, and pearlite, and less bainite in these conditions than thecomparable High+AI!N steel. The microstructures of the 50 mm aqua-quenched platesare mainly mixtures of bainite and martensite. In this condition, the prior austenitegrain sizes were determined to be about 18 flm (ASTM No. 8.30) for the Base, High,and High+AI steels, and 15 flm (ASTM No. 8.82) for the High+AI/N and High+Nbsteels. The modest austenite grain refinement observed in the High+AI/N andHigh+Nb steels is presumably attributed to the presence of AIN and Nb(C,N) particles,respectively.

Effect of Tempering Time on Mechanical Properties of the 75 mm Air-Cooled Plates

Upper-nose temper embrittlement (UNTE)-- The 75 mm air-cooled plates (insimulation of 5(i() mm water-quenched forgings) were tempered at 660°C for times


Figure 2--0ptical micrographs (at mid-thickness location) of 75 mm air-cooledand 50 mm aqua-quenched plates of the Base and High+AI/N steels (in the as-cooledcondition).

ranging from 5 to 50 hours to determine the effect of tempering time on mechanicalproperties. The as-cooled microstructures of these plates are a mixture offerrite/pearlite/bainite, as shown in Figures 2(a) and 2(b). Figure 3 illustrates thechanges in tensile strength as a function of tempering time at 660°C. For all thesteels, the tensile strength decreases continuously with increasing tempering time.Similar trends were observed in yield strength (0.2% offset) as a function of temperingtime. For a given tempering time, the Base steel has the lowest strength (due to its.lower alloying level), and the High+Nb the highest strength (due to its higher alloyinglevel and finer grain size). The High, High+AI and High+AI/N steels have anintermediate strength level. Over the range of tempering times evaluated, the High,High+AI, High+AI/N, and High+Nb steels exceed the tensile strength requirement of


Figure 3--Effed of tempering time at MO°C on tensile strength of 75 mm air-cooled plates. The minimum tensile strength requirement of A508 Grade 3 Class 1 isindicated.

A50X Grade 3 Class I. However, the Base steel does not meet the requirement whentempered over 20 hours. Figure 3 also suggests that a leaner alloy composition (suchas, lower end of the specified range) may not meet the tensile strength requirement inthick sections of forgings after tempering. On the other hand, all the steels exceed theyield strength (>345 MPa), tensile elongation (>18%), and reduction of area (>38%)requirements of the A50X Grade 3 Class lover the range of tempering timesevaluated.

Figure 4 shows the Charpy energy at both +4°C and -29°C as a function oftempering time. For all the steels, the Charpy energy decreases continuously withincreasing tempering time, even though the strength also decreases with increasingtempering time (Figure 3). This result indicates that the steels are being embrittled ontempering. Since both the strength and toughness decrease with increasing temperingtime, it is helpful to plot tensile strength/toughness combinations for a given temperingtime as shown in Figure 5. In this figure, the strength/toughness data of the five steelsare grouped according to their tempering times in order to assess the effect of


Figure 4--Effect of tempering time (at 660°C) on Charpy energy of 75 mm air-cooled plates, (a) CVN energy at +4°C and (b) CVN energy at -29°C. The minimumCVN energy requirement of A508 Grade 3 Class 1 is shown in (a).

composition on the strength/toughness balance of the steels (to be discussed later). Itis apparent that the strength and toughness of all the steels decrease simultaneously asthe tempering time is increased. Figure 6 shows the fracture appearance of two


Figure 5--Tensile strength and toughness combinations of 75 mm air-cooledplates after tempering at 660°C for 2 to 50 hours, (a) CVN energy at +4°C and (b)CVN energy at -29°C. The data points are grouped according to their tempering timesas indicated. The hatched region in (a) indicates the minimum strength and toughnessrequirements for A50X Grade 3 Class 1 material.

Charpy specimens (tested at -29°C) from the Base steel for tempering times of 5 and50 hours. As the tempering time is increased, and the Charpy energy is reduced


Figure 6--Effect of tempering time (at 660°C) on the fracture appearance ofCharpy specimens (tested at -2YoC) from the 75 mm air-cooled plate of the Base steel,(a) 5 and (b) 50 hour tempering.

(Figure 4), the size of the shear zone next to the notch decreases and the area of brittle(unstable) fracture increases. At both tempering times, the brittle fracture regionsexhibit cleavage (Figure 7). Similar fracture characteristics were observed for theother four steels. There is no evidence of intergranular fracture (indicative of temperembrittlement) in any of these specimens. Considering the tempering temperature andthe fracture characteristics, this reduction in toughness with increasing tempering timeis attributed to upper-nose temper embrittlement (UNTE) [5]. Upper-nose temperembrittlement has also been observed in other steels [6-13]. Important mechanismsassociated with UNTE have been reported to be grain boundary carbide coarsening19,10, II], recrystallization/grain growth [9], and the formation of a dual-phase ferrite-martensite microstructure [9,121, on tempering.

Optical microscopy on samples taken from the Charpy specimens did notreveal any evidence of recrystallization/grain growth, or noticeable changes in carbidedistribution with increasing tempering times. However, electron microscopy of carbonextraction replicas from these samples shows the presence of needle-like carbides, anddemonstrates that the size of these carbides increases with increasing tempering time(Figure 8). Using EDX in a STEM, these needle-like carbides were identified as Ma-rich, and contain RlUl±O.R%Mo, 5.0±O.6% Fe, 5.0±O.5% Mn, and 2.0±o.3% Cr(average over ten analyses on different carbides). An analysis of the electron


Figure 7--Fractographs of the brittle fracture region of the Charpy specimens(Figure 0) showing cleavage fracture characteristics for both tempering times of (a) 5and (b) 50 hours.

diffraction patterns obtained from these carbides suggests that they are probably M2C.Based on these observations, the upper-nose temper embrittlement observed in thesesteels is attributed to the precipitation and coarsening of needle-like Mo-rich(M2C-type) carbides on tempering.

Effect of Composition on Mechanical Properties of the 75 mm Air-Cooled Plates

Effect of alloy level--As shown in Figures 4 and 5, for a given tempering time,the Base steel exhibits slightly lower strength, but slightly better toughness than theHigh steel. The higher levels of Si, Ni, and Cr lead to increased strength, but reducedtoughness (classic strength-toughness tradeoff) for a given tempering time.

Effect of a Nb addition-- The addition of Nb to the High steel generally in-creases both the strength and toughness at a given tempering time, resulting in animproved strength/toughness balance (Figures 4 and 5). This improvement in strengthand toughness is presumably due to the presence of Nb(C,N) particles which refine thegrain size on austenitizing. In fact, the strength/toughness balance of the High+Nbsteel is better than that of the Base steel at the longer tempering times.


Figure 8--Carbon extraction replicas prepared from the 75 mm air-cooled plateof the Base steel tempered at li60°C for (a) 5, (b) 20, and (c) 50 hours, showing theeffect of tempering time on the size of the needle-like carbides precipitated duringtempering. The needle-like carbides in the Base steel were found to contain 88.0±o.8%Mo, 5.0±O.6% Fe, 5.0±O.5% Mn, and 2.0±0.3% Cr, using EDX in a STEM.

Effect of AI and N additions-- The addition of Al alone to the High steelprovides a modest improvement in toughness and a somewhat reduced strength level(Figure 5). The ad.dition of both Al and N to the High steel also results in a decreasein strength, but provides a larger increase in toughness than obtained with Al only.Consistent with results reported previously [2,3], this toughness improvement due tothe combined addition of Al and N is attributed to the austenite grain refinement fromAIN particles.

As shown in Figure 4, the Charpy energy of the Base steel decreases morerapidly with increasing tempering time than do the Charpy energies of the High+AI,High+AI/N, and High+Nb steels. For tempering times longer than 10 hours, thetoughness of the Base steel is lower than the toughness of the High+AI, High+AI/N,and High+Nb steels (Figure 4(b). TEM micrographs (Figure 9) show that when the




excess of about 20 hours (Figure 3). In contrast, the tensile strength of the High+Al/Nsteel is still above 550 MPa, even after tempering for 50 hours at 660°e. Note thatthe tensile strength of the High+Al/N steel is within the required strength range of 550to no MPa for the entire range of tempering times evaluated. Overall, while both theBase and High+Al/N steels have acceptable toughness levels, the High+Al/N steelgenerally exhibits better strength/toughness combinations.

RL,DT. temperature--Based on the Charpy toughness data presented above, the75 mm air-cooled and tempered plates (in simulation of 560 mm water-quenchedforgings) of the Base and High+Al/N steels were selected for drop-weight testing todetermine their TNDT and RT NOT temperatures. The RT NOT temperatures measured forplates tempered for various times are plotted in Figure 10 as a function of tensilestrength. For the Base steel, as the tempering time is increased from 2 to 10 hours,the RTNDT temperature increases from -35 to -29°C, while the tensile strengthdecreases from 585 to 560 MPa. As a result, after tempering for 10 hours, the Basesteel marginally meets the strength and RTNOT requirements. In contrast, the RTNOT

temperature of the High+Al/N steel remains at -35°C until the tempering time isincreased to 20 hours, when the RTNDT increases to -29°C. In addition, at comparabletempering times, the tensile strength of the High+Al/N steel is higher than that of theBase steel. For the tempering time of 20 hours, the Base steel does not meet thestrength and RTN/lT requirements (Figures 3 and 10). Overall, the High+Al/N steelexhibits a better combination of RTNDT and tensile strength compared to the Base steel.

Effect of Cooling Rate after Austenitizing on Mechanical Properties

Effect of cooling rate after austenitizing on UNTE--To evaluate the effect ofcooling rate on the extent of UNTE, a 50 mm plate of the Base steel was quenchedinto a water/polymer solution (cooling rate = 595°C/min, in simulation of 50 mmwater-quenched forgings) after austenitizing. As shown in Figure 2(c), the as-quenched microstructure is predominantly martensitic. The as-quenched plate wasthen tempered at 660°C for times ranging from 2 to 50 hours, and its toughness wascompared with the 75 mm air-cooled plate (9°C/min) of the Base steel (Figure 4). Themicrostructure of the 75 mm air-cooled plate is a mixture of ferrite, pearlite, andbainite (Figure 2(a». As shown in Figure 11, the Charpy energy of the 50 mm aqua-quenched plate remains relatively constant, while that of the 75 mm air-cooled platedecreases rapidly as the tempering time is increased from 2 to 50 hours. This indi-cates that a predominantly martensitic microstructure is less prone to UNTE than theferrite/pearlite/bainite microstructures produced at slower cooling rates; this result isconsistent with the conclusions of reference [13]. Figure 12 shows carbon extractionreplicas from these samples for tempering times of 5, 20, and 50 hours. Compared toFigure X, the needle-like carbides in this mostly martensite microstructure do notcoarsen as the tempering time is increased from 5 to 50 hours. This explains therelatively constant toughness level with increasing tempering time. In summary, theextent of UNTE is significantly reduced by increasing the cooling rate afteraustenitizing to produce a predominantly martensitic microstructure. In a productionforging, higher cooling rates can be achieved by increasing the severity of quenching,


Figure Il--Effects of cooling rate (after austenitizing) and tempering time (at66()°C) on the Cl1arpy energy of the 75 mm air-cooled (9°C/min) and 50 mm aqua-quenched (595"C/min) plates of the Base steel, (a) CVN energy at +4°C and (b) CVNenergy at -29°C.

such as, using high pressure spray quenching, brine quenching, or energeticmechanical agitation, which should lead to a reduced UNTE.

Effect of cooling rate after austenitizing on toughnessnTo evaluate the effect ofcooling rate after austenitizing on toughness, the tempering time for a given steel waskept the same, while the cooling rate after austenitizing was varied. Based on


Figure 12--Carbon extraction replicas prepared from the 50 mm aqua-quenchedplate of the Base steel tempered at 660°C for (a) 5, (b) 20, and (c) 50 hours, showingthe effect of tempering time on the size of the needle-like carbides precipitated duringtempering.

Figure 3, to obtain a tensile strength of about 585 MPa for the 75 mm air-cooledplates, the tempering times for the Base, High, High+AI, High+Al/N, and High+Nbsteels should be about 2, 20, 20, 20, and 50 hours, respectively. These temperingtimes were also used to temper 50 mm fan-cooled plates (in simulation of 300 mmwater-quenched forgings) and 50 mm aqua-quenched plates (in simulation of 50 mmwater-quenched forgings) for the evaluation of the effect of cooling rate on toughness.

The tensile properties of the steels which were quenched and tempered aresummarized in Table 2, compared with ASTM A508 Grade 3 Class 1 tensilerequirements. In all cases, the tensile load versus elongation curves displayeddiscontinuous yielding behavior. For both the 75 mm air-cooled plates and the 50 mmfan-cooled plates, a tensile strength level in the range of about 570 to 590 MPa isobtained for all the steels. However, for the 50 mm aqua-quenched plates, a slightlyhigher tensile strength level (range of 605 to 630 MPa) is obtained. In all the cases,the tensile properties after tempering meet the tensile requirements of ASTM A508Grade 3 Class 1.


Figure 13--Charpy transition \:urves of 75 mm air-cooled plates tempered to atensile strength of 5X5 MPa, (a) energy, (b) lateral expansion, and (c) % ductile

fracture.



Although full Charpy transition curves were determined for all three coolingconditions, only those for the 75 mm air-cooled plates are shown in Figure 13 toillustrate trends. Based on these transition curves, the absorbed energies at -29°C andat +4°C, and the 50% fracture appearance (50%FATI) transition temperature for eachsteel were determined. These data are plotted in Figure 14 as a function of coolingrate. The Charpy impact toughness of all the steels generally increases as the coolingrate is increased; note that the tensile strength of each steel also increases withim;reasing cooling rate (Table 2). This increased toughness is attributed to theim:reased amount of bainite/martensite and the less-severe UNTE present in the faster-cooled microstructures.

For the 75 mm air-cooled (9°C/min) and 50 mm fan-cooled (22°C/min) plates,the Base steel exhibits better toughness than the other four steels. This is because theBase steel was tempered for only 2 hours, while the other four steels were temperedfor either 20 or 50 hours. Consequently, the better toughness of the Base steel isattributed to less UNTE. However, when the cooling rate is increased (50 mm aqua-quenched (595"C/min) plates), the Base steel exhibits poorer toughness than the otherfour steels. Note that in the 50 mm aqua-quenched plates, the microstructures areprimarily martensitic, and therefore are less prone to UNTE. In this type ofmicrostructure, the difference in tempering time among the steels no longer has asignificant effect on toughness (via UNTE). Instead, increasing the hardenability(through alloy additions) actually increases the toughness, presumably by increasingthe volume fraction of the martensitic constituent.

As described earlier, the addition of Nb to the High steel does provide someaustenite grain refinement. Nevertheless, the High+Nb steel still exhibits the lowesttoughness level among all the steels for the three cooling conditions evaluated. This isbecause the High+Nb steel was tempered for 50 hours, compared to 20 hours for theHigh, High+AI, and High+AIIN steels, and 2 hours for the Base steel. Consequently,the High+Nb steel experiences more UNTE (especially at the slower cooling rates),and hence has lower toughness levels than the other four steels.

At the tempering time of 20 hours used for the High, High+AI, and High+Al/Nsteels, the High+AI!N steel exhibits the best toughness for all the cooling rates evalu-ated. This is presumably due to the austenite grain refinement provided by AINpartides, consistent with results reported previously [2,3]. On the other hand, eventhough the Nb addition provides similar grain refinement, overall the Nb addition hasa negative effect on toughness. Again, this result is rationalized on the basis ofdifferent tempering times (50 hours for the High+Nb steel versus 20 hours for theHigh+AI/N steel), and different levels of UNTE.

CONCLUSIONS

I. Upper-nose temper embrittlement (UNTE) is observed in all the experimentalA508 Grade 3 steels on tempering at 660°C. In all the Charpy specimensexamined, deavage is the dominant fracture mode, and there is no evidence ofintergranular fracture (indicative of temper embrittlement). This UNTE isattributed to the precipitation and coarsening of needle-like, Mo-rich carbides. The


extent of UNTE is significantly reduced by increasing the cooling rate afteraustenitizing to produce more acicular (martensite/bainite) microstructures. Tominimize the negative effects of UNTE on toughness, the tempering treatment forthese steels should be as short as possible.

2. At a constant tempering time, increasing the levels of Si, Ni, and Cr in the Basesteel leads to increased strength and reduced toughness. A further addition of Nbincreases both the strength and toughness, a consequence of austenite grain refine-ment. The addition of AI (only) reduces the strength and improves the toughnessslightly. In comparison, the addition of both Al and N results in a decrease instrength, but a relatively large increase in toughness. This toughness improvementis attributed to austenite grain refinement by AIN particles.

3. At a given tempering time, the High+AVN steel exhibits the best strength!toughness combination among the five steels evaluated. For example, for atempering time of 20 hours, the tensile strength and RTNDT temperature of theHigh+AI!N steel meet the minimum 550 MPa and maximum -29°C requirementsfor the Westinghouse advanced reactor program, while those of the Base steel failto meet these requirements.

4. At a strength level of about 585 MPa, the Base steel exhibits the best toughnesswhen the microstructure is a ferrite/pearlite/bainite mixture (slower cooling rates).However, when the microstructure is martensite/bainite (faster cooling rates), theBase steel has the lowest toughness. This is because the Base steel was temperedfor only 2 hours to reach this strength level, while the other four steels requiredtempering times of 20 hours or more. Consequently, the Base steel experiencesless UNTE and has better toughness. However, as the cooling rate is increased toproduce martensite/bainite microstructures, the extent of UNTE is significantlyreduced and the effect of tempering time becomes less important. In this case, theBase steel has the lowest toughness level because it has the lowest hardenability(less martensite) among all the steels.

ACKNOWLEDGMENTS

Special thanks go to R. L. Bodnar for reviewing the manuscript and W.Furdanowicz for conducting STEM analysis of carbides. The assistance ofR. 1. August, K. 'E. Downey, L. L. Hahn, T. R. Knauss, S. J. Lawrence, R. R. Lichty,F. J. Marsh, C. Santos, R. E. Steigerwalt, and G. E. Weiss in this work is alsogratefully acknowledged.

REFERENCES

r II D. E. Ekeroth, "Forging Design for Westinghouse Advanced Nuclear PlantsComponents", 12th International For~emasters Meetin~, Chicago, Illinois, USA,September II-I n, IYY4.


[2] H. Takashima et aI., "Properties of Nuclear Reactor Pressure Vessel Steel Plate(ASTM A533 Class 1", Nippon Steel Technical Report, No. 22, December 1983,pp. 19-34.

[3] K. D. Haverkamp et aI., "Effect of Heat Treatment and Precipitation State onToughness of Heavy Section Mn-Mo-Ni Steel for Nuclear Power PlantsComponents", Nuclear Engineering and Design, Volume 81, 1984, pp. 207-217.

14] American Society of Mechanical Engineers Standard: Boiler and Pressure VesselCode-Section III, Article NB 2330, 1992.

15] L. D. Jaffe and D. C. Buffum, "Upper Nose Temper Embrittlement of a Ni-CrSteel", Journal of Metals, AIME, January, 1957, pp. 8-16.

[6] R. L. Bodnar and K. A. Taylor, "Structure!Property Relationships in Medium-Carbon Bainitic Steels for Thick Sections", Iron and Steelmaker, Volume 17, no.8, August 1990, pp. 47-63.

[7] T. D. Nelson, R. L. Bodnar, and J. E. Fielding, "A Critical Assessment of ASTMA508 Class 2 Steel for Pressure Vessel Applications", 32nd Mechanical Workingand Steel Processing Conference Proceedings, ISS-AIME, Volume XXVIII, 1991,pp. 323-341.

[8] R. L. Bodnar, T. D. Nelson, and R. F. Cappellini, "Enhancements in the Integrityof Heavy Forgings", Iron and Steelmaker, Volume 18, no.7, July 1991, pp. 26-36.

[9] H. J. DeKlerk and W. M. DeVilliers, "Microstructural Evolution and Embrittle-ment During Tempering of Low-Alloy Mn, Ni, Mo, Mn-Mo and Mn-Mo-NiSteels", 36th Mechanical Working and Steel Processing Conference Proceedings,ISS-AIME, Vol. XXXII, 1995, pp. 453-470.

[10] 1. P. Naylor and M. Guttmann, "Mechanism of Upper-Nose TemperEmbrittlement in Mn-Ni-Mo A533 Grade B Steel", Metal Science, Vol. 15,October, 1981, pp. 433-441.

[11] G. Pienaar, "Effects of Vanadium on Upper-Nose Temper Embrittlement andOther Mechanical Properties of Cr-Ni-Mo Low-Alloy Steels", Materials Scienceand Technology, Vol. 2, October, 1986, pp. 1051-1061.

[12] W. M. DeVilliers, H. J. DeKlerk, and G. Pienaar, "Upper-Nose TemperEmbrittlement Phenomena in a Mn-Mo-Ni Pressure Vessel Steel", InternationalJournal of Pressure Vessels and Piping, Vol. 50, 1992, pp. 215-230.

[13] P. E. Martin, C. Rogues, and P. Bastien, Revue de Metallurgie, Vol. 59, 1962,pp. 829-843.

Hisashi Tsukada', Komei Suzuki', Mikio Kusuhashi3 and Ikuo sato'

CURRENT FORGINGS AND THBIR PROPBRTIBSFOR STEAM GBNERATOR OF NUCLEAR PLANT

REFERENCE: Tsukada, H., Suzuki, K., Kusuhashi, M., and Sato, I. ,"Current Forgings and Their Properties for Steam Generator of NuclearPlant" Steel Forqinqs: Second Volume. ASTM STP1259. E.G. Nisbett and A.S.Melilli, Eds., American Society for Testing and Materials, 1997.

ABSTRACT: Current steel forgings for steam generator (SG) of PWR plantare reviewed in the aspect of design and material improvement. Thefollowing three items are introduced.

1. The use of integral type steel forgings for the fabrication of steamgenerator enhances the structural integrity and makes easier fabricationand inspection including in-service inspection [1] • The followingexamples of current integral type forgings developed by the Japan SteelWorks, Ltd. (JSW) are introduced

(1) Primary head integrated with nozzles, manways and supports(2) Steam drum head integrated with nozzle and handhole(3) Conical shell integrated with cylindrical sections and handholes

2. In order to decrease the weight of steam generator, the highstrength materials such as SA50B, CI.3a steel have been adopted in somecases. The properties of this steel are introduced and the chemistry andheat treatment condition are discussed.

3. As one of the methods to minimize the macro- and micro-segregations, the use of vacuum carbon deoxidization (VCD), Le.deoxidization of steel by gaseous CO reaction, with addition of Al forgrain refining was investigated. The properties of SA50B, C1.3 steelswith Low Si content are compared with those of conventional one.

KEYWORDS: forging, steam generator, nuclear plant, integrated forging,SA50B, CI.3 steel, high strength, low si steel

1) Managing Director, Research Develc.p1BI1tHeadquarters, '!beJapan steel Works Ltd.2) General Manager, Steel Business Headquarters, '!beJapan steel Works Ltd.3) Manager, Chief Research Engineer, Muroran Research Iaboratory4) General Manager, Atonic Energy Dept. Muroran Plant,'!beJapan Steel W:>rksLtd.

4 chatsu-machi Muroran Japan

56

TSUKADA ET AL. ON CURRENT FORGINGS AND PROPERTIES 57

1. INTBGRAL TIPB FORGINGS OF SGFig. 1 shows a possible material layout in

the typical design of SG. The currentmaterial layout consists of the followingintegral type forgings.

(1) Forged heads integrated with nozzlesfor steam drum head and primary head

(2) Forged shell, in some cases integratedwith handholes

(3) Forged conical shell integrated withcylindrical section and in some casesintegrated with handholes

With the application of above integraltype forgings, the weld seams can beextensively reduced. These forgings havebeen realized by the development inmanufacturing technologies.

1.1 Forged HeadsThe stearn drum and primary heads for

SG had been made by forming of plates andwelding fabrication, or made of casting.To eliminate weld seams and to makeinspection easier, the forged head integratedwith nozzles has been developed with thespecial die forg1ng process[2].This process is rotating die forgingtechnique using partial male die and femaledie as shown in Fig. 2. By this process,an ingot is directly forged into a shape ofa head integrated with a nozzle afterpreliminary solid forging.

Several type and diameter of forged headscan be manufactured by the combination ofseveral dies, as one of advantages of thisprocess.


(1) Steam Drum headPhotos 1 and 2 show several type of steam drum heads of SG. Photo 1 is

one piece forged elliptical head with 4>4. 5m (177") integrated with steamoutlet nozzle. Photo 2 is one piece forged spherical head with 4>4. 5m

(177") integrated with steam outlet nozzle and manway. A small skirt isalso integrated in inside surface of head in order to make joint weldingof steam dryer easy.

(2) primary headPhotos 3 to 5 show several type of primary heads. Photo 3 is one piece

forged primary head with tp3.4m (134") integrated with two (2) nozzles,two (2) manways and four (4) supports. photo 4 is also one piece forgedprimary head with 4>3.8m (ISO") integrated with three (3) nozzles and two(2) manways. photo 5 is one piece forged primary head with 4>4. 6m (181")integrated with support ring.

1.2 Forged Shell and Conical ShellThe forged shell ring have been adopted to eliminate the longitudinal

weld seams. Forged rings of conical shell with cylindrical sections havebeen developed by the contour shape forging process. The enlarging ismade by the special designed upper die and mandrel.

TSUKADA ET AL. ON CURRENT FORGINGS AND PROPERTIES 59

(1) Forged shell with nozzlePhoto 6 shows the forged shell ring with 4. 5m (178") high integrated

with handhole.(2) Forged conical shell

Photo 7 shows the forged conical shell with 2.6m (102") high, 200mm(8")cylindrical section for upper side, 420mm(16.5") for lower side andintegrated with handholes.

1.3 Mechanical propertiesOne of the advantages in the use of the forged head is in lessdirectionality of the material properties. Fig. 3 shows the charpyimpact properties at head and nozzle portions of the forged primary head.Nosignificant difference at both location is observed. This homogeneityis acquired by the proper steel making, forging process and heattreatment.


4. CONCLUSIONAt the time of SG replacement and new construction, the use of integral

type forgings are increased from the advantage of reduction offabrication period and extent of ISI. And the high strength steel hasalso adopted in some cases to reduce the thickness of component. Thefollowing points are reviewed in this paper.

(1) Several integral type forgings have been developed and currentforgings for SG are introduced.

(2) High strength steel, SA508 CI.3a can be manufactured by theproper chemical composition and tempering condition.

(3) Low si content steel with addition of Al is effective to improveimpact properties and better homogeneity.

Further efforts will be continued on developments for the demands of theintegral type forgings. And a close cooperation with designers andfabricators of SG is very important to assure a higher quality and lowercost of component from material layout point of view.

REFERENCE:[1] H.Doner., E.Michel., rAdvanced Technique of the Primary circuit

of KWU-PWRs with regard to Safety and ReliabilityJ lAEA 21-25 Mar.1983 Stuttgart

[2] H.Tsukada., K.Suzuki., M.Shimazaki. and I.Sato., rRecent Progressof Large and Integrated Forgings for Nuclear Power PlantJ The 3rdJSME/ASME Joint International Conference on Nuclear Engineering,April 23-27,1995 Kyoto,Japan

[3] S.Kawaguchi., H.Tsukada., K.Suzuki., M.Kusuhashi., rApplication of20MnMoNi55 steel with lowered Si content to Heavy Thick SteamGenerator Tube Sheet ForgingsJ 10 MPA-Seminar, 10-13 Oct. 1984Stuttgart

Fran~oise Morin 1, Pierre Bocquet2 and Alain Cheviet3

FORGING TECHNOLOGY ADAPTED TO THE MANUFACTURE OF NUCLEAR PWRAUSTENITIC PRIMARY PIPING

REFERENCE: Morin F., Bocquet P. and Cheviet A., "Forging Technology Adaptedto the Manufacture of Nuclear PWR Austenitic Primary Piping" Steel Forgings:Second Volume. ASTM STP 1259, E.G. Nisbett and A.S. Melilli, Eds, American Societyfor Testing Materials, 1997.

ABSTRACT: In order to reduce the number of welds (in service inspection costs), to im-prove the thermal ageing resistance and the internal soundness of components, Electricite deFrance decided to develop a solution for forged austenitic stainless steel in place of castingsto manufacture the straight parts and the elbows of the primary piping of PWR power plants.The paper presents the interest of such solution and the manufacture conditions of such for-gings applied, first, to a cold leg for the Civaux I unit and, second, to a complete loop forthe Civaux 2 unit.

KEYWORDS: austenitic stainless steel, piping, forging, ultrasonic testing.

To increase the safety and reduce the operating costs of the primary piping of electro-nuclear power plants of the Pressurized Water Reactor type, Electricite de France has encou-raged its French industrial partners to change the design and manufacture of these compo-nents. That evolution resulted in the manufacturing of large forged parts in austenitic stain-less steel.

The main objective of a such optimization is to improve the life of new power stations,increasing their Safety margin, in addition to a reduction of the in service inspection andmaintenance costs.

The paper presents the interest of using austenitic stainless steel forgings in place ofcastings and describes the manufacturing conditions of such forgings for the French Civaux1 and 2 1400 MW power plants.

lQuality Inspection Department, Products Division, Electricite de France, Evry,France 91000.

2Research Department, Centre de Recherche des Materiaux du Creusot, Creusot-LoireIndustrie, BP 56, Le Creusot, France 71200.

3product Manager, Creusot-Loire Industrie, BP 68, Rive de Gier, France 42800.

65


EXPECTED IMPROVEMENTS DUE TO THE NEW DESIGN

Safety I Ouality improvementsThe steel casting used up to now for manufacturing these components contain a signi-

ficant amount of delta ferrite (to avoid solidification cracking and to reach a convenient levelof tensile properties). Ferrite induces a high sensitivity to the embrittlement associated tothermal ageing in the temperature range concerned (290 - 325°C), in particular when the ma-terial is enriched in Si, Cr and Mo [1] [2].

To lower the sensitivity to this type of embrittlement it has been proposed to use a pu-rely austenitic grade (AISI 304L type) made by forging. In addition, this manufacturing tech-nique is expected to improve the internal soundness of the material by the reduction of volu-metric flaws (cavities, bubbles ...) inherent in cast products.Inspection conditions improvements

The testability of a component consists in being able to examine it in respect of volumeby non destructive means. The ultrasonic (UT) method cannot be used on large castings inaustenitic steel because the large grain size and grain orientation makes the material no per-meable to the ultrasound waves. Here again, the use of a forging technology with control ofthe grain size enables better ultrasound permeability to be achieved.

Taking into account the risk of defect connected with the welding operations, eachweld must be inspected with regard to volume in order to certify that it has been done cor-rectly. Testability of the component therefore also involves the associated welds being capa-ble of inspection. Having fewer welds is therefore of advantage to the operator.

DESCRIPTION OF THE PREVIOUS DESIGN

The Primary Piping of pressurised water reactors is used to carry the primary water,heated in the vessel by nuclear fuel, to the steam generators. The power stations currentlybuilt (of the N4 type) have 4 steam generators connected to the vessel by 4 loops, each incor-porating a motorised pump unit which circulates the heat-conveying fluid at a speed of 10m/s.

Figure I shows a diagram of a loop and the following 3 lines can be seen there.- The hot leg which connects the reactor vessel directly to the steam generator, composed ofa straight part and a divergent 50° elbow, giving it a total length of more than 7 metres. Onloop No.1 this line has a straight stub pipe which joins the expansion line.The grade of the straight tubing was originally stainless steel Z2CND 18-12 with controllednitrogen (316L). The tubes were then extruded and therefore they could not be longer than 4metres.Later, in order to reduce the number of welds, 8 metres long cast tubes made by centrifugingwere procured. The corresponding grade: Z3CN20-09M steel (derived from 304L) is anausteno-ferritic stainless steel.- The cross over leg which connects the steam generator to the primary pump, including two90° elbows and one 40° elbow.The elbows were generally conventionnal gravity castings. Their original grade wasZ3CNDI9-IOM (derived from 316L) but was abondoned in favour ofZ3CD20-09M steel(derived from 304L) so as to reduce their sensitivity to embrittlement by thermal ageing.The next change only concerned the 50° elbow of the hot lines which, from the second N4section (CHOOZ B2) is a forging of grade Z2CN 19-10 with controlled nitrogen (304L). It isextruded and then hot formed.


- The cold leg which connects the primary pump to the reactor, consisting of a straight partand a 22° elbow, giving it a length of 8 m. Each of them has a stub pipe at an angle of 45°which takes the safety injection system.

The 4 stub pipes at an angle of 45° have always been gravity castings and are the samegrade as the elbows. Therefore, they have changed from grade Z3CND 19-1OM to Z3CN20-09M. The size of them at the baseplate to be welded to the Primary Piping (TP) is llOO mmlengthwise and 650 mm widthwise (see Fig. 2).

MANUFACTURING OF THE FORGED AND BENT COLD LINE FOR CIVAUX UNIT

The choice of making a prototype cold line was made to overview all the potential dif-ficulties that way arise during the manufacturing. By mutual agreement between EDF andFramatome, the longest line was choosen, which in addition had the most bulky stub pipe: itis the cold line known as BF4.

In view of the size of the product to be made, Framatome turned to Creusot-Loire In-dustrie which has the industrial tool and the know-how necessary for taking up this challen-ge. It should be noted that since 1988 CLl has held the qualification for induction bendingsuccessfully implemented in Mannesmann's workshops on geometry corresponding to a 50°elbow.

Manufacturing sequence :- Steelmaking: the austenitic stainless steel grade Z2CNI9-lO (AISI 304L) with

controlled nitrogen was processed in an electric arc furnace by CLl. Its refining was comple-ted in a vacuum ladle with silicon killing. The solid 87 tons ingot was bottom poured.

- Forging: CLI's forging programme is broken down into blooming of the whole in-got and then drawing on either side of the area reserved for the stub pipe. The volume reser-ved for the stub pipe is forged prior to drawing to the final forging dimensions. In this blankthe inclined geometry of the stub pipe was premachined as shown in Fig. 3.A special forging sequence has been developed in view to control the grain size of the mate-rial.

- Machining: Making the tubular product is done in a machining operation (drillingand boring) in the metallurgical centreline of the workpiece.

- Induction heating for bending: The induction bending operation was done under thesame conditions as the qualification pipe and was therefore preceded by specific machiningof the area to be bent.The bending as such, done at Mannesmann's in Mulheim, was followed by quality heattreatment (solution heat treatment). The bending plant is shown diagrammatically in Fig. 4.

Cold Line No.4 of Civaux 1 is shown, in the final delivery condition, in Fig. 5.

MORIN ET AL. ON FORGING TECHNOLOGY 69



Metallurgical and mechanical propertiesChemical analysis

A synthesis of the results found recapitulated in Table 1 - both at the pouring and invarious locations on the workpiece, shows a good uniformity of the chemical composition.

Metallographic examinationMetallographic examinations were performed after chromic acid electrolytic etching:

- on specimens taken from test coupons- on replicas at different locations on the forging itself.

Figure 6 shows the location of replicas and the results obtained for the grain size num-ber G between 0 and 2.

That prove a'very good uniformity in spite the large size of the component.No significant amount of delta ferrite was noticed at any location, even on the test

rings. « 1 %)The mechanical properties were established on overlengths for acceptance tests at the

head and foot sides, and at the stub pipe end as shown in Figure 7.All of the results are given in Table 2, they are uniform and are significantly higher

than required.



74 STEELFORGINGS:SECONDVOLUME

Ultrasonic inspections and material soundness

The non destructive inspections on the cold line in the final state were carried out ultra-sonically in straight waves and transverse waves. The equivalent diameter method as per NFStandard A 04.308 was applicable. The tracing was carried out through the outer surface ofBF4 and completed through the inside surface for the stub pipe and for the bend. No noticia-ble indication was revealed.

The ultrasound permeability was checked and carried out with the mapping of the rela-tive attenuations in the thickness of the product over the whole line. This enables evaluationof the metallurgical uniformity of the finished product. The principle consists in measuringthe loss of UT intensity, in straight waves at 2 MHz, between the first and the second back-ground effect echo, on thickness in the region of 69 mm. This attenuation is 8 dB at the moston the whole pipe (including the elbow) and locally it goes to 14 dB on either side of thestub.

Moreover, the detection sensitivity was checked under the following conditions:

Straight waves

The artificial reflector, with 3 mm diameter flat bottomed hole corresponding to themarking threshold, is detected at 2 MHz through a thickness of 100 mm with a ratio 10 % ofthe amplitude of the signal from the reflector. The criteria applicable, given in Table 3, cantherefore be easily respected since their detectability is guaranteed.

Transverse waves at 45°The artificial reflector used is a 1.5 mm deep notch, less than I mm wide and approx.

50 mm long. It is detected at 70 mm deep with a signal-to-noise ratio of around 14 dB at 2MHz. That means diat the amplitude of the background noise is less than 20 % of that of thesignal from the reflector.

As the marking threshold of the readings is 50 % of the datum echo, and as the rejec-tion criterion is 100 % of the datum echo, their detectability is guaranteed.

The good permeability as well as the associated detection sensitivity are to be compa-red with the uniformity of the grain sizes measured.


MANUFACTURING OF THE FORGED AND BENT LOOP FOR CIVAUX 2 UNIT

As a sequel to the demonstration of the operational feasibility carried out for the cold li-ne of Civaux 1, EDF asked for this operation to be repeated and extended to the whole of aCivaux 2 loop. At this time the purpose was to optimise the manufacturing conditions to suitthe specific nature of each of the lines while proving the manufacturing organisation with thecomplex interweaving of subcontracting in a preseries production configuration.

The choice fell on loop No.1, each of the lines of which is illustrated in Figure 8.Cold line (BF1)

No change in the conditions for processing the steel was planned. CLl made a solid 87tons ingot.

The forging programme was optimised. Further, the tooling was improved by usingdie plates (profiled press tools).

The tube was bored by CLl and then Mannesmann did the bending of the 22° elbow inaccordance with the work done earlier.

Hot line (BCl)The processing of the steel in an electric furnace by CLl remains identical. The refining

of grade Z2CN 19-10 is also finished in a vacuum ladle with silicon killing. This phase endswith the casting of a solid 77 tons ingot.

The forging programme is based directly on that for the cold line with simplificationconnected with the absence of an incorporated stub pipe. Advantage was taken of using dieplates for the forging.

After drilling, also done by CLl, the induction bending of the 50° elbow was done byMannesmann in accordance with the qualification part made in 1988.

U-shaped-line (BUl)The processing of grade Z2CN 19-10 by CLl is identical : electric furnace + refining.

However, a major difference must be pointed out since the casting is done in an ingot mouldproducing a hollow 77 tons ingot.

The forging is broken down into two stages:- The first, done by CLl, consists of drawing on a mandrel, followed by parting the productinto three hollow billets.- The second consists of drawing each hollow billet to the forging dimensions on a press fit-ted with a manipulator (Tecphy). The lengths of blanks A, B and C at the forging dimen-sions are 5.2 metres, 4.6 metres and 3.9 metres respectively.

No drilling is necessary. Only the boring and machining operations are planned.Each of the blanks was bent and heat-treated by Mannesmann in accordance with the

same conditions as for the previous parts, to produce the following 3 elbow or elbow andstraight part assemblies to be assembled direct on site.- Blank A --> 90° elbow and straight part- Blank B --> 90° elbow and straight part- Blank C --> 40° elbow.




CONCLUSION

This work as demonstrated the feasability of the new design introducing large forgingcomponents in turn of castings:I - It is possible to forge large components in austenitic stainless steel. Several more or lesscomplex geometries have been tested for making the Primary Piping; they make it possibleto dispense with a large number of prefab welds (assembling).2 - Different ingot types have been experimented (solid ingots and hollow ingots). They alllead to having large sized forgings with excellent internal soudness which are metallurgicallyand mechanically uniform.3 - The good ultrasound permeability, and the associated UT testability, of large componentsin austenitic stainless steel have been proved at manufacturing level and will enable a follow-up in service which is better suited to most requirements.

REFERENCES

[1] Slama G., Petrequin P. and Mager T. "Effect of Aging on Mechanical Properties ofAustenitic Stainless Steel Castings and Welds", SMIRT Post Conference, seminar 6,

Assuring Structural Integrity of Steel Reactor Pressure Boundary Components,Monterey, August 1983.

[,2] Bonnet S., Bourgoin J., Champrerond J., Guttmann D., Guttmann M., "Evolution ofMechanical Properties of Various Cast Duplex Stainless Steels in Relation to

Metallurgical and Ageing Parameters: An outline of Current EdF Programmes" -Material Science and Technology, p. 221, March 1990.

Stephen Price1 and Graham A Honeyman2

THE OPTIMIZATION OF MECHANICAL PROPERTIES FOR NUCLEARTRANSPORTA TlON CASKS IN ASTM A350 LF5

REFERENCE: Price, S, Honeyman, G. A., "The Optimization of MechanicalProperties for Nuclear Tnlllspol"tation Casks in ASTM A350 LF5", Steel Forgings:Second volume, ASTM STP 1259, E.G. Nisbett and A.S. Melilli, Eds., American Societyfor Testing and Materials, 1997.

ABSTRACT: Transpol1 flasks are required for the movement of spent nuclear fuel. Dueto their nature of operation, it is necessary that these flasks are produced from forgedsteels with exceptional toughness properties. The material specification generally cited forflask manufacture is ASTM A350 Grade LF5 Class I, a carbon-manganese-nickel alloy.The range of chemical analysis permitted by this specification is very broad and it is theresponsibility of the material manufacturer to select a composition within this range whichwill satisfY all the mechanical propel1y requirements, and to ensure safe and reliableperformance.

Forgemasters Steel and Engineering Limited have experience in the manufacture oflargehigh integrity fuel elemen~ flask forgings which extend over several decades. Thisexperience and involvement in international standards in USA, Europe and Japan hasfacilitated the development of an optimised analysis with a low carbon content, nickellevels towards thc top end of the allowed range, a deliberate aluminium addition to controlgrain size and strictly controlled residual element levels The resultant steel has excellentlow temperature impact propcl1ics which greatly exceed the requirements of thespecification This analysis is now bcing adopted for the manufacture of all currenttransp0l1 flasks

KEYWORDS: nuclc"r, flak, manuf~1cture, ASTI\1 A350 LF5, improvement, properties,optimization

1Technical Dircctor (Designate), Forgeillasters Steel and Engineering, Sheftield, England.

"Managing Director Forged Rolls (UK) Ltd, Shefticld, England.

79


INTRODUCTION

In the last thirty years the demands imposed by increasingly stringent fuel specificationshave necessitated the development of improved designs of nuclear transport flasks. Theflasks used for the transport of Magno x fuel have a cuboid form of flask body fitted with atop lid (Figure I) The shipment of spent fuel from Light Water Reactor (L WR) powerstations in Japan and Europe had been conducted utilising Excellox 3, 3B and 4 flasksdesigned by British Nuclear Fuels PLC (BNFL) [I] These have a welded steel shell withcircumferential cooling tins and a separate steel-encased lead liner for gamma shielding.The fuel is carried in multi element bottle (MEB). A water annulus between the lead linerand the flask provides neutron shielding.

PRICE AND HONEYMAN ON THE OPTIMIZATION 81

During the 1990's higher burn up and mixed oxide fuels are increasingly being irradiated.The Excellox 6 and Excellox 7 flasks have been designed by Nuclear Transport Limited(NTL), under contract from BNFL in order to meet the demands of these fuelspecifications [2] The flasks are of forged steel monolithic construction and are cladinternally with stainless steel. The thick shell forming the body of the flask provides all thegamma shielding and thus dispenses with the need for a separate liner (Figure 2a and 2b).


The flasks are therefore easier to maintain and decontaminate. The lid can be bolteddirectly onto the body; this construction gives the flasks greater strength. Neutronshielding is provided by an exterior layer of carbonated rubber located between thecooling fins.


The flask has been designed in two Icngths to increase the range of fuels which can betransported. The Excellox 6 flask will accommodate 5.0 metre Icngth PWR fuel with a230 mm square cross section. The Excellox 7 flask takes fuel up to a length of 4.5 metreswith a cross section of 2l5mm x 215mm. The flasks are fitted with pairs of trunnions forhandling and support. The first Excellox 6 flasks have recently entered service, deliveringspent fuel to the Thermal Oxide Re-Processing Plant (THORP) at Sellafield from aGerman reactor.

MATERIAL REQUIREMENTS FOR LARGE NUCLEAR FLASKCOMPONENTS

The safe and reliable performance of nuclear transport flasks must be ensured. The flasksare designed to meet the requirements of the International Atomic Energy Agency (IAEA)Safety Series 6 "Regulations for the Safe transport of Radioactive Material", and thechoice of steel is limited by these requirements The specification selected by flaskdesigners is generally related to ASTM A350 Grade LF5 [3). This is essentially a carbonmangancse-nickel stecl: the chcmical requirements are given in Table I. It is immediatelyapparent that this permits a wide range of analysis. especially with respect to carbon,nickel, sulphur, and phosphorus contents. However constraints are imposed by the needto meet certain mechanical properties The required properties for the Magnox andExcellox 6 flasks are compared with the A350 LF5 specification in Table 2. The level ofthe strength needed is quite modest and it is the stipulated low temperature impact testvalues that are particularly onerous The differing test temperatures are a consequence ofthe operating conditions experienced by the transportation flasks

Table I--Chemical ReQuirements for ASTM A350 Grade LF5

ELEMENT COMPOSITION, WT %

Carbon 0.30 max

Manganese 0.60 to 1.35

Phosphorus 0.035 max

Sulpr1Ur 0.040 max

Silicon 0.20 to 0.35

Nickel 1.0to2.0

Chromium 0.30 max

Molybdenum 0.12 max

Copper 0.40 max

Niobium 0.02 max

Vanadium 0.03 max

The orientation of the test pieces required by A350 LF5 is in the direction of majorworking of the forging whereas the Magnox and Excellox components were evaluated inthe longitudinal, transverse and radial orientations. The test material must be extractedfrom the heaviest section of the forging usually at a location such that the central axis ofthe test specimen is at least a quarter of the maximum heat treated thickness from thenearest surface and at least a quarter of the maximum heat treated thickness from thenearest treated surface and at least IOOmm fi'om any second treated surface. Thesefactors combined'necessitate the careful consideration of chemical analysis; it is theresponsibility of the Forgemasters not the designer to impose a melting range which willyield a chemical composition that satisfies the pertinent specification

Although the ASTM A350 specification does not give any non-destructive testingrequirements, ultrasonic and magnetic pal1icle examination acceptance levels weredetailed by the flask designers. For the flask components supplied by Forgemasters Steel& Engineering Limited, ultrasonic indications of 5mm and above and magnetic particleindications of3mm and above would be cause for rejection of the forging, Asteelmaking and forging route was therefore selected to ensure compliance with theserequirements


MANUFACTURING PROCESS OPTIMIZA nON

Forgemasters Steel and Engineering Limited has been responsible for the production ofmachined body, base and lid forgings for the transportation flasks. The Excelloxcomponents were supplied rough)nachined to British Steel Engineering, the manufacturerof the complete flask assembly.

The melting of carefully selected scrap was carried out in a 90 tonne basic electric arcfurnace. rn addition to the melting of raw materials, the aim of this operation was toreduce the phosphorous to an acceptable level. This was accomplished by use of anoxidising slag. Reversions of phosphorous as a result of carry over of this slag into thevacuum arc degassing (V AD) unit can be a problem, care was therefore taken to ensureslag free tapping. For the manufacture of Excellox 6 flasks this was achieved using asliding gate on the electric arc furnace.

Secondary steelmaking was conducted in the VAD unit to reduce the sulphur content andto trim the analysis. The steel was deoxidised by silicon and aluminium additions, thelatter element being maintained at a high level for grain refinement purposes. Ladledegassing procedures were adopted to reduce the hydrogen content before final teemingof the molten steel into the ingot mould. For the majority of casts further degassing wascarried out by teeming under vacuum using the vacuum stream degassing method whichgave hydrogen levels of Ippm or less. For casts requiring more than 90 tonnes of steel,the temperature of the first heat of steel was maintained in a ladle furnace while the secondwas undergoing VAD treatment. The heats were then sequentially teemed through a ponyladle.

For all the flask components, forging was carried out using the 10,000 tonne press.Several forging heats were required to produce the forgings. During forging operationsthe temperature was carefully controlled to avoid the ductility trough associated withaluminium treated steels [4] The forging sequence for the Magnox flask bodies includedtwo upsetting and plating operations to ensure adequate consolidation of the structure.This was followed by cogging to a block and then punching using a tapered tool to form arecess in the block. The latter operation was conducted for two reasons, firstly to lessenthe amount of material requiring removal during subsequent machining and secondly toreduce the section size and thus facilitate the hydrogen diffusion operation. The Excellox6 flask bodies were produced as hollow forgings. The ingots were first upset to form adisc and a ho\e was then punched through with a ho\\ow punch ihis method of punchingremoves some of the central material as a core. As it is this area of the ingot which ismost likely to contain any segregation or porosity it is a particularly useful forgingprocedure for high integrity components such as nuclear transport flasks. After punching,the length of the forging was increased by working on a die with a mandrel in a cavity.

Forging is followed by a preliminary conditioning heat treatment. This included ahydrogen degassing treatment for which the forging was held at approximately 650°C toallow hydrogen to diffuse out of the steel. The Excellox 6 and Magnox lid forgings were


then taken up to austenitising temperature before quenching in water and tempering toachieve the appropriate mechanical properties. As ASTM A350 LF5 steel is low incarbon content and extremely tough as such it is highly resistant to quench cracking and istherefore generally given its quality heat treatment in the as forged condition. TheMagnox flask body forgings, however, were machined prior.to their quality heattreatment. The purpose of this being to decrease the mass of material and to bring theforging closer to its finished size, thus enhancing its mechanical properties. The removalof the forging and treatment scale also provides a satisfactory surface for an earlyultrasonic examination.

The ASTM A350 specification allows the forgings to be furnished in the normalised,normalised and tempered or quenched and tempered condition. The influence of coolingrate from austenitising has been studied using manganese-nickel forgings produced byForgemasters Steel & Engineering Limited [5]. Two casts, one with 0.04% carbon, theother with 007% carbon, were evaluated in the normalised and tempered and quenchedand tempered conditions For the 0.04% carbon material there was little difference inproperties for the two conditions. However for the 0.07% carbon steel the quenchingoperation increased the tensile strength by 40 Mpa and the impact energy at -30°C by over100 joules. The downward shift in fracture appearance transition temperature amountedto 25°C.

It is apparent that the magnitude of this change is composition dependent. The greatereffects of cooling rate observed in the 0.07% carbon material are thought to be related toits higher hardenability. This leads to the formation of a more completely bainitictransformation product after water quenching. It is now normal practice to water quenchASTM A350 LF5 forgings if the carbon content exceeds 0.04% to ensure that the impactpropel1ies are met.

On completion of the quality heat treatment cycle, test material was removed from theforgings to allow the properties to be evaluated in accordance with the relevantspecification The forgings were then machined to final dimensions and subjected to finalultrasonic and magnetic parlicle examination

RI-:'SULTS AND DISCUSSION

It has already been inferred that the selection of chemical composition to guaranteeexcellent low temperature impact properties is a crucial pal1 of flask manufacture. Thedeleterious influence of sulphur on impact strength has been well documented [6].Although the importance of reducing sulphur content to low levels during secondarysteelmaking is recognised, the etTect of sulphur on susceptibility to hydrogen crackingmust also be considered Manganese Nickel steels exhibit a high susceptibility toHydrogen Cracking This susceptibility of the steel combined with the large section sizeof many of the flask forgings could lead to unacceptable long post forge solid statehydrogen degassing treatments. It is known that hydrogen becomes trapped in the steel atinterior cavities such as those produced when sulphide content increases the possible trap


sites for hydrogen and so reduces the concentration at anyone site. For this reason aminimum sulphur content 01'0.005% is usually imposed for ASTM A350 Grade LFSsteels. Production constraints may lead to sulphur levels as high as 0.010% if theminimum of 0.005% is to be satisfied It is therefore essential to demonstrate that therequisite impact properties are achievable at this higher sulphur level

The impact energy values at -30DC have been collated for Magnox flask componentsmanufactured by Forgemasters Steel & Engineering Limited. These are plotted as afunction of sulphur content in Figure 3. Data are also included for other A3S0 LFSforgings with high sulphur contents for non-nuclear applications. The values presented aresolely from radially orientated test specimens as it was in this test direction that thepoorest results were obtained The plot illustrates that the required impact energy of 48joules is comfortably met with sulphur of 0010% and below. When the sulphur levelexceeds 0.013% the results are consistently poor.

Carbon levels are also known to influence toughness properties. This influence is thoughtto be less pronounced than that of sulphur and various relationships between the twoelements have been considered in order to quantifY their combined effect. The mosthelpful correlation developed is that between the impact energy and the sum of carboncontent and twenty times the sulphur content. This is demonstrated in Figure 4. This is aclearer relationship than that shown in Figure 3 and it suggests that it is insufficient to lookat sulphur conlent alone when designing a composition with good low temperature impactproperties Some of the out-lying data points at low impact energies on the earlier plotare explained by their relatively high carbon content combined with twenty times thesulphur content at a value of 0.30% or less if the impact test requirement is to be satisfied.Adopting this approach increases the options available when designing a range of chemicalcompositions to meet a specification This is par1icularly beneficial during manufacturewhere production constra~nts might be imposed.

Carbon Content + 20 x Sulphur Content, wI. %

FIG. 4--Effect of carbon and sulphur contents on impact toughness

In order to make optimum use of the relationship illustrated in Figure 4, the influence ofcarbon on strength must be quantified. The results of a review of data for LF5 flask body,base, and lid forgings are presented as lower bound line in Figure 5. It is considered that aminimum carbon level of 0.05% should be stipulated to achieve the requisite tensilestrength. It is clear that it is the specified tensile strength rather than the yield strengthwhich controls the permitted chemical analysis. As 0.03% is a workable range in thesteelmaking process, a melting aim of 0.05% to 0.08% carbon is usually requested. Byinference. a l11<lximumsulphur content of 0.0 II % can be tolerated. The preferred meltingrange already l11entioned of 0.005% to 0.0 I0% is thus acceptable.


It should be noted that these carbon and sulphur levels are significantly below themaximum values allowed by the ASTM A350 LF5 specification. The phosphorus level issimilarly restricted to assist with the production of good impact properties.

The addition of nickel to low carbon-manganese steels increases the steel's hardenabilityand produces a predominately bainitic microstructure. This improves both strength andtoughness. A minimum nickel content requirement of I .4% has been established to givegood through hardenability

After deoxidisation of the steel, the aluminium level is maintained at between 0.015% and0.025%. This results in a tine grained steel with enhanced impact properties.

The resulting aim melting range utilised by Forgemasters Steel and Engineering Limitedfor ASTM IU,;O Grade LF5 flask forgings is presented in Table 3.

TABLE 3--Aim melting range for A350 LF5 flask components

ELEMENT COMPOSITION, WT %

Carbon 0.05 to 0.08Manganese 1.20 to 1.35Phosphorus 0.010 max

Sulphur 0.005 to 0.010Silicon 0.20 to 0.25Nickel 1.40 to 1.70Chromium 0.30 max

Molybdenum 0.12 max

Aluminium 0.015 to 0.025Copper 0.40 max

Niobium 0.02 max

Vanadium 0.03 max

The improvement obtained in mechanical properties by water quenching has already beendescribed. As the carbon content of the steel used exceeds 0.04%, it is imperative thatforgings are produced in the water quenched condition if they are to realize their fullpotential with regard to optimizing toughness properties and giving adequate strength.


CONCLUSIONS

1. An optimized manufacturing route has been developed which makes it possible toproduce forgings for nuclear transport flasks in ASTM A3S0 Grade LFS withexcellent properties.

2. Chemical composition should be controlled within tight limits. In particularcontrol of carbon and sulphur contents have been shown to be are critical. Ideallyboth elements should be maintained at low levels; however it is possible to balancetheir levels and still achieve the requisite strength and toughness. This givesbenetits in the time required for hydrogen degassing.

3. Water quenching and tempering is the recommended quality heat treatment formanganese-nickel forgings to ensure a microstructure which will maximise lowtempcrature impact properties.

ACKNOWLEDGEMENTS

Forgclllaslcrs Steel and EngillL'cring Limited would like to thank British Nuclear FuelsPLC and Brit ish Steel Engineering l'or thcir technical assistance and for permission topublish this paper

REFERENCES

[I] Gowing, R. and Purcell, pc., "Further Experience and Developments in theTransp0l1 of Spent Fuel", Proc. PATRAM '92, Yokohama, Japan, 1992.

[2] Gowing, R. ct ai, "Introduction of New Flasks for High Burn-Up Spent Fuel",Proc. Third International Conference on Transpol1ation for the Nuclear Industry,Cumbria, England, 1994.

[3] ASTM A3'iO, "Standard Specification for Forgings, Carbon and Low-Alloy Steel,Requiring Notch Toughness Testing for Piping Componcnts", American Societyfor Testing and Materials, 1993

[4] Wilson, F.G. and Gladman, T, " Aluminium Nitride in Steel", InternationalMaterials Review, vo133, No), 1988

['i] RCYIH lids, P E and [)lIliclI, [), 8ritish Steel Technical Research Report, 1981.

[6] Pickering, F.B., "Towards Improved Toughness and Ductility", ClimaxMolybdenum Company Synopsiull1, Kyoto, Japan, 1971.

General Industrial Forgings

John E. Fielding!, Robert B. Focht', Kenneth F. Reppert', and EugeneL.Tihansky·

DEVELOPMENTS IN FORGING INGOT PRODUCTION AT BETHFORGE INC.

REFERENCE: Fielding, J.E., Focht, R.B., Reppert, K.F., Tihansky, E.L.,"Developments in Forging Ingot Production at BethForge Inc.," SteelForainas: Second Volume, ASTM STP 1259, E.G.Nisbett and A.S.Melilli,American Society for Testing and Materials, 1997.

ABSTRACT: BethForge manufactures precision custom forgings for diverseapplications including turbine and generator rotors, components fornuclear reactors and other pressure vessels, and rolls for steel andaluminum rolling mills. This paper will discuss the production offorging ingots for BethForge at the modernized steel making facility inSteelton, PA. The facility is operated as part of a joint effort withPennsylvania Steel Technologies, Inc. (also a wholly owned subsidiary ofBethlehem Steel) and consists of a 150 ton (136 metric ton) DC electricarc furnace, a ladle refining station including a ladle furnace andladle degasser, and teeming facilities which include bottom pouring andvacuum stream degassing. Steel is produced for BethForge ingots as wellas continuously cast for the Pennsylvania Steel Technologies productlines. Forging, heat treating and machining operations remain atBethlehem, PA and ingots as large as 130 inches diameter (3300 mm) and290 tons (263 metric tons) are hot transported the 90 mile (145 km)distance from Steelton to Bethlehem in specially constructed insulatedrailcars. In addition to describing the facilities and operations, thepresentation will focus on solutions to several of the uniqueengineering and technical challenges realized in successfully bringingthis operation on line.

KEYWORDS: ingot, ladle refining, bottom pouring, forgings,transportation, rotors, pressure vessels, rolls

INTRODUCTION

BethForge, Inc. is a fully integrated open die forging manufacturerlocated in Bethlehem, Pennsylvania, USA. BethForge operates modernmachining facilities, vertical and horizontal heat treatment furnaces,and 9,072 metric tons (10,000 short tons) and 2,268 metric tons (2,500short tons) open die forging presses. These operations are supported bya strong technical commitment including extensive efforts in product

I Manager, Technical Services, BethForge, Inc., Bethlehem, PA 18016

2 Manager, Primary Operations, BethForge, Bethlehem, PA

3 Senior Supervisor, Technical Services, BethForge, Inc., Bethlehem, PA

4 Chief Engineer, BethForge, Inc., Bethlehem, PA

93


and process engineering, dedicated chemical and mechanical testinglaboratories, and the support of the Bethlehem Steel Homer ResearchLaboratory. BethForge manufactures precision custom forgings includingturbine and generator rotors, components for pressure vessels (includingnuclear reactors) and work rolls and back-up rolls for steel andaluminum metal producers. The steel for these operations has been solelyproduced at the Pennsylvania Steel Technologies facility in Steelton,Pennsylvania since November 1995. Prior to this date, the steel wasproduced at the BethForge electric furnace operation in Bethlehem. Thisfacility, although capable of producing excellent quality steel, hadlimitations. Because ladle refining capabilities were not included, thequality demands of many of the ingots produced had to be achievedthrough the use of selective scrap practices combined with multiple slagprocess in the arc furnace (oxidizing followed by one or morereducing/refining slags). Therefore, arc furnace time was significantlylonger and scrap costs higher per ton than can be achieved with typicalladle refining operations. The cost to use the above practices as wellas process limitations of the multiple slag method had become a limitingfactor in the ability to produce the full range of forgings desired byour customers.The new facility in Steelton, Pennsylvania, is operated as part of ajoint effort with Pennsylvania Steel Technologies, Inc. (also a whollyowned subsidiary of Bethlehem Steel) and consists of a 150 ton (136metric ton) DC electric arc furnace, a ladle refining station includinga ladle furnace and ladle degasser, and teeming facilities which includebottom pouring and vacuum stream degassing. Steel is produced forBethForge ingots as well as continuously cast for the Pennsylvania SteelTechnologies product lines. This facility has the capability to producethe full range of ingot weights previously produced at the BethForgeelectric furnace operation in Bethlehem and meet the more restrictivequality demands of the current market. However, several engineering andtechnical obstacles required resolution prior to full implementation ofthis new melt facility as a supplier of BethForge ingots.

NEW MELT FACILITY CAPABILITIES

The flowline diagram on Figure 1 illustrates the production of BethForgeforging ingots at Pennsylvania Steel Technologies in Steelton, PA. Asummary of the operations sequence is presented in the precedingparagraphs along with a brief description of the respective facilityfeatures .

• Melting is conducted in the 150 ton (136 MT) DC electric arcfurnace. This furnace was manufactured by NKK-United and is a roofswing, top charge type and has a melting capacity of 150 tons (136MT) with a 16 ton (14.5 MT) heel. TOLal tap-to-tap time for thefurnace is approximately one hour. The furnace shell has an insidediameter of 23 feet (7010 mm) and a height of ca. 14 ft - 11 in.(4546 mm). Electrodes are either 28 in. (711 mm) or 30 in. (762

mm) diameter and the melting power input averages 74 MW.

The melting furnace charge consists of fragmented scrap with limeadded during melting. The ferro alloys, plus fluxes anddeoxidizers, are added to the furnace ladle during tapping of thefurnace, as well as incrementally during refining. The aluminum(deoxidizer) added at this point is made to lower the oxygencontent to facilitate the refining operation which will follow.

FIELDING ET AL. ON DEVELOPMENTS IN FORGING INGOT 95

The eccentric bottom tapping feature of this furnace providesefficient separation of the steel bath from the furnace slag,thereby reducing the carryover of slag and making the subsequentladle refining operation more efficient, resulting in cleanersteel.

• Following tapping from the DC electric arc furnace, the furnaceladle of steel is moved to the ladle furnace (LF) where steelrefining is accomplished by providing chemical adjustments,heating for temperature control, and interaction with refiningslag to promote desulfurization and enhance cleanliness (inclusionremoval). Direct readout of the liquid steel temperature andoxygen content is provided. A sketch of the LF is presented inFigure 2.

Argon stirring, through a porous plug in the bottom of the ladle,provides temperature uniformity in the steel bath as well aspromoting desulfurization through slag-metal interfacing andtransfer of sulfur to slag under low oxygen bath conditions andlow FeO-MnO slags. There is also a wire feeder system to allow forthe introduction of elements, such as calcium, to controlinclusion morphology.

• After achieving the prescribed temperature and chemistry for theheat of steel being processed, the ladle is moved to the tank typeladle degasser (refer to Figure 3). Argon stirring during ladledegassing (1) assists in the removal of dissolved gases (hydrogen,oxygen, and nitrogen) and (2) facilitates removal of aluminum thathad been added prior to ladle refining and thereby improve steelcleanliness. Micro alloying can be performed after degassing and,if necessary, the ladle of steel can be returned to the LF forreheating.

• After leaving the ladle degasser, the steel will either be bottompoured or vacuum stream degassed into the required ingots.BethForge has the capability to bottom pour ingots in the range of40 inch (1015 mm) to 92 inch (2335 mm) diameter, whereas thoseingots from 108 inch (2745 mm) to 130 inch (3300 mm) diameter mustbe vacuum stream degassed. However, the 92 inch (2335 mm) diameteringots may be vacuum stream degassed if required by specificationor if exceptionally low levels of dissolved gases are required.

• Following ingot teeming and solidification, the hot ingots aretransported from Pennsylvania Steel Technologies to BethForge inspecially designed insulated rail cars. These cars have beendesigned to maintain elevated ingot temperatures which willprotect the ingots from cracking and allow charging of the ingotsdirectly into the forge heating furnaces. Further discussionregarding these transport cars is presented in the followingsection.

OPERATIONAL AND TECHNICAL ISSUES

When BethForge made the commitment to having Pennsylvania SteelTechnologies provide ingots for our manufacture, there were a number ofexpectations and issues which confronted those responsible for thistransition. The major areas of concern were as follows:

• Chemical analysis control• Desu1furization and cleanliness


TABLE 2 -- Comoarison of Process Caoabilitv Data - Chemical Analvsis(0.7C-3.2Cr)

Melt C - Std. Mn - Std. Cr - Std. V - Std.Source Dev. Dev. Dev. Dev.

(%) (%) (%) (%)

EFM 0.02 0.04 0.09 0.01

PST 0.02 0.02 0.03 0.003


the pony ladle and pouring refractories up to a stable temperature.While the first heat is pouring, the second heat is being held and willlose approximately 40°F to the surroundings. When the first ladle isdrained, the second ladle is positioned to complete the pouring of thetwo ingots through the pony ladle. At the completion of pouring the twoingot combination, the second ladle will be repositioned to bottom pourthe third ingot located on a second bottom pour plate through its owndownfountain.

Development of the above practice, in addition to practices for bottompouring smaller ingots and vacuum stream degassing large ingots, haspermitted successful production of the full range of ingot sizespreviously produced at the electric furnace operation in Bethlehem. Asstated earlier, BethForge has the capability to bottom pour ingots inthe range of 40 inch to 92 inch diameter, whereas those ingots from 108inch to 130 inch diameter must be vacuum stream degassed. Although the92 inch diameter ingots may be vacuum stream degassed if required byspecification or if exceptionally low levels of dissolved gases arerequired, ingot sizes smaller than 92 inch diameter must be bottompoured.

Deaassina

All steel previously produced in the electric furnace operation atBethlehem was subsequently degassed either by ladle-to-ladle degassingor pouring in vacuum (ladle-to-mold). Ladle-to-ladle degassing wasutilized on 40 inch and 48 inch diameter ingots, whereas ingots of 54inch diameter through 130 inch diameter were poured in vacuum.

Although all steel produced for BethForge at Pennsylvania SteelTechnOlogies is subjected to at least one degassing operation, theladle-to-ladle degassing technique used in our previous melt facilityhas been replaced by degassing in the tank degasser prior to bottompouring of the ingots. As previously described, after achieving theprescribed temperature and chemistry for the heat of steel beingprocessed, the ladle is moved to the tank type ladle degasser where themolten steel is argon stirred during ladle degassing to assist in theremoval of dissolved gases (hydrogen, oxygen, and nitrogen). This typeof degassing operation can be extended to all ingot sizes through 92inch diameter, followed by bottom pouring.

Pouring ingots in vacuum is conducted on all ingot sizes of 108 inchthrough 130 inch diameter. The 92 inch diameter ingot may be produced byeither ladle-to-ladle degassing or ladle-to-mold degassing. Although thebasic ladle-to-mold degassing process is unchanged from the processpreviously used at Bethlehem. steel for ingots poured may be tankdegassed prior to ladle-to-mold degassing. This provides for doubledegassing should that be deemed necessary.

Although the degassing methods used at our previous Bethlehem facilitywere quite successful in producing products with low levels of dissolvedgases, initial chemical analysis data from material produced atPennsYlvania Steel Technologies are encouraging. In fact, hydrogenlevels after ladle degassing at Pennsylvania Steel Technologies appearquite similar to hydrogen levels in ladle-to-mold degassed steelsproduced in our previous electric furnace operation at Bethlehem (1.4ppm in molten steel after degassing and 0.7 ppm in fUlly heat treatedproduct) .


Deoxidation Process

During tapping of BethForge grades from the primary melt furnace atPennsylvania Steel Technologies, silicon as an alloy addition andaluminum as a deoxidant is added to the furnace ladle. The amount ofaluminum added varies from 100 pounds (45 kg) to 350 pounds (159 kg),depending on the aim silicon level of the particular grade beingproduced. Heats with higher aim silicon level (0.60 to 0.70%) requireless aluminum than heats with low silicon aims (0.10% maximum) because aportion of the silicon added takes part in deoxidizing the liquid steel.After tapping, the heat is sent to the ladle furnace to recover the heatloss from the addition of bulk alloys. At the time of the initialtemperature measurement at the ladle furnace, the oxygen level of thebath is also checked with the expectation that the oxygen will be lessthan 10 ppm. Any additional deoxidation necessary is achieved with theaddition of small amounts (30 to 40 pounds) of aluminum bars (90%aluminum). An oxygen level of less than 10 ppm is desired in order toenhance desulfurization and to promote recovery of additions ofoxidizable alloys. The final aluminum level specified to the melt shopis 0.010% maximum, and control of the aluminum level of the bath hasbeen excellent with actual final aluminum contents generally being0.005% to 0.008%.

Incrot Transportation

BethForge has a fleet of nine insulated railroad cars for transportinghot ingots from the ingot teeming facility at Steelton to its forgingshop in Bethlehem. Ingots are shipped hot (12500F) for a distance ofapproximately 100 miles. The ingots arrive at Bethlehem with atemperature loss of approximately 250°F after a track time of 16 hours.

Ingots can be shipped in a variety of methods ranging from multiplesmall ingots per car to a single ingot of 130 inch diameter weighing upto 290 tons (261 MT). Six railroad cars can handle ingots from 40 inchdiameter to 78 inch diameter with an ingot weight of 78 tons (70 MT).Two cars can handle ingots from 40 inch diameter to 108 inch diameterwith a combined ingot weight of 212 tons (191 MT). One railroad car canhandle ingots from 40 inch diameter to 130 inch diameter with a combinedingot weight of 320 tons (288 MT).

Since the various size ingots have different strip times, some ingotsmust be held while awaiting the strip of larger ingots. All of therailroad cars are equipped with heaters to hold ingots at temperaturewhile awaiting shipment. Just prior to shipment, the heaters aredisconnected and the insulated covers are sealed. No heat is addedduring transport.

Upon arrival of the insulated railroad cars at Bethlehem, the covers areremoved and the hot ingots are immediately placed in forge furnaces. Thecars are promptly returned to Steelton. All of the railroad cars averageone round trip per week.

SUMMARY

During the first year of ingot production for BethForge at thePennsylvania Steel Technologies melt facility, it has been demonstratedthat the full range of required ingot sizes can be successfully producedby either bottom pouring (ingots up to 92 inches diameter) or toppouring (92 inches to 130 inches diameter). It has been demonstratedthat these ingots can be stripped hot and transported to the BethForge


site with acceptably low heat loss (250°F) such that thermal cracking isavoided in even the most highly alloyed grades.

Although complete evaluation of all products from the variety ofavailable ingot sizes is not yet complete, it is apparent that the newmelt facility affords greater chemical analysis control resulting inmore consistent analyses. The ladle refining furnace has beenconsistently achieving the level of desulfurization expected for themore demanding products produced by BethForge. With the reheatingcapability and control available at the ladle refining station, heatsare being teemed within more narrow temperature ranges than experiencedat our prior melt facility and two consecutive heats may be poured toproduce ingots requiring greater heat weights.

All heats are subjected to at least one degassing operation, with thelarger ingot sizes being double degassed. BethForge is cOllecting gasanalysis data from virtually all bottom poured ingots and many of thelarger top poured ingots. Initial data indicate that acceptable hydrogenlevels are present in products from both bottom and top poured ingots.

While initial experience with the new melt facility at PennsylvaniaSteel Technologies has been favorable, BethForge plans to furtherevaluate all BethForge products until we are confident that the meltpractices have been optimized.

Gerald Stein', Wal ter Kirschner', Joachim Lueg J

APPLICATION OF NITROGEN-ALLOYED MARTENSITIC STAINLESS STEELS IN THEAVIATION INDUSTRY

ABSTRACT: Nitrogen in stainless martensitic steels has a beneficialinfluence on the mechanical as well as on the chemical properties.However the effect of nitrogen is limited due to the rather lowsolubility of this element. A special alloy development in combinationwith a pressurized melting technique lead to distinctly higher nitrogencontents. Stainless martensitic steels containing high nitrogen contentsare manufactured by VSG today on an industrial scale using the PESR-process (Pressurized Electroslag Remelting) . Depending on specialapplications these steels are available with different chemical analysisunder the trade mark CRONIDUR.

The basic composition of all CRONIDUR-alloys consists of about 15 %Chromium, 1 % Molybdenum, 0.15 to 0.35 % Carbon and 0.20 to 0.40 %Nitrogen. The combination of Cr + Mo + N leads to a superior corrosionresistance of these HNS-steels (HNS: High Nitrogen Steels) in comparisonto similar carbon based alloys. Focused on applications with a requiredminimum hardness of 58 HRC, like stainless bearings or screw shafts, theC+N-content is tuned between 0.60 to 0.80 % (Brand: CRONIDUR 30) .Additions of max. 0.3 % Vanadium and 0.1 % Niobium qualifies the brandCRONIDUR 20 for enhanced temperature applications like turbine disks orblades.

KEYWORDS: nitrogen-alloying, PESR-technique, corrosion resistance, highhardness, aerospace industry, aviation industry

1. NITROGEN METALLURGY

For the produceion of massively nitrogen alloyed steels, the nitrogensolubility in the molten condition is of major importance. In pure ironthe nitrogen solubility is rather poor, in particular regarding theliquid state - at a temperature of 1600 DC at atmospheric pressure only0.04 % N are soluble.

'Member of the Board (Technology) , VSG Energie- und Schmiedetechnik,Altendorferstr. 104, 45143 Essen, Germany

'Head of Research and Development Department, VSG Energie- undSchmiedetechnik GmbH, Westendstr. 15, 45143 Essen, Germany

'Head of Quality Department, VSG Energie- und Schmiedetechnik GmbH,Hattingen, Germany

104

STEIN ET AL. ON APPLICATION STAINLESS STEELS 105

In order to achieve higher nitrogen contents, melting and subsequentsolidification must take place at elevated N-partial pressure accordingto sievert's square root law and! or via adding alloying elements whichincrease the solubility of nitrogen. The sievert formula and additionallythe interaction parameters of the most important alloying elements aregiven in Fioure 1.

~ure 1: sievert's law and parameter for the calculation of thenitrogen solubility at 1600 C [1]

It is necessary to distinguish between the solubility-increasing elementsMO, Mn, Cr, Ti, Nb and V and the solubility-decreasing elements Ni, Siand C. Yet those elements which permit the highest nitrogen contents likeTi, Nb and V have a strong nitride forming tendency, promoting nitrideformation at rather low N-concentrations. Therefore only Cr, Mn and Moremain as solubility-increasing elements [2].

Due to the austenite stabilizing character of Mn, this element is notsuitable, if a martensitic structure is desired. So Cr and Mo are thebasic alloying elements for stainless martensitic nitrogen-steels, givingthem enhanced corrosion resistance and nitrogen solubility as well.

2. PRODUCTION OF HIGH NITROGEN STEELSThe most efficient and reliable way to produce HNS was and probably stillis the Pressurized-Electroslag-Remelting-Process (PESR).

The PESR process operates according to the same principle as the ESRprocess. Electrodes having the desired chemical composition are remeltedin a closed vessel under pressure (~).

Christian E. Fischer,' Jay S. Gunasekera,z James C. Malas3

PROCESS MODEL DEVELOPMENT FOR OPTIMIZATION OF FORGED DISKMANUFACTURING PROCESSES

REFERENCE: Fischer, C.E., Gunasekera, J.S., and Malas, J.C., "Process ModelDevelopment for Optimization of Forged Disk Manufacturing Processes," SteelForgings: Second Volume, ASTM STP 1259, E.G. Nisbett and A.S. Melilli, Eds.,American Society for Testing and Materials, 1997.

ABSTRACT: This paper addresses the development of a system which will enable theoptimization of an entire processing sequence for a forged part. Typically such asequence may involve several stages and alternative routes of manufacturing a given part.It is important that such a system be optimized globally, (rather than locally, as is thecurrent practice) in order to achieve improvements in affordability, producibility, andperformance. This paper demonstrates the development of a simplified forging model,discussion techniques for searching and reducing a very large design space, and anobjective function to evaluate the cost of a design sequence.

KEYWORDS: Forging, Optimization, Machining, Simplified Models

INTRODUCTION

Industrial forging processes generally involve a series of operations that transforma workpiece into a useful finished product characterized by specific shape, size, andproperties. A typical sequence of thermomechanical operations consists of multiple stagehot and cold forging processes interspersed with suitable heat treatment processes. Theshape of the product is achieved through the forging and machining, while the propertiesof the products are, in general, dependent upon the entire thermomechanical history. The

, ResearchAssociate, Center forAdvancedMaterialsProcessing,Ohio University,Athens, OH 45701.

2 Director, Center for Advanced Materials Processing, Ohio University, Athens, OH 4570l.

3 Wright Laboratories Materials Directorate,Wright PattersonAFB, OH.

116

FISCHER ET AL. ON PROCESS MODEL DEVELOPMENT 117

challenge is to optimize the entire thermomechanical processing route in order to achievethe best balance of manufacturing costs, delivery schedules, and properties in the fmalproduct.

Forging sequences have traditionally been designed through the use of variousempirical rules which have evolved as a result of many years of experience. Build andtest methods have commonly been used for die design and for the selection of processvariables such as forging temperatures and ram velocity. These methods result in hightooling and set-up costs and in long lead times before production which, in turn, adverselyaffect the manufacturing enterprise. Recently, many limitations of build and test methodshave been overcome through the use of process models based upon computer simulationtechniques. However, despite the success with modeling and optimization of individualmanufacturing processes, particularly machining [1,2,3] no consistent system capable ofmodeling and globally optimizing an entire multiple stage process has been developeduntil recentl y.

Researchers at Harvard University and Virginia Polytechnic Institute and StateUniversity (VPI) have recently proposed new methods for finding a near optimumsolution to problems such as this with enormous search spaces. Under both approaches,the problem is modeled as a Discrete Event Dynamic System. An initial workpiece anda candidate sequence of processing operations and parameters are selected based on somesearch rule, the results of the operation are estimated using some predictive model, anda cost value is assigned to these results. In the Harvard approach, the operations andparameters are chosen randomly, and ranked ordinally from best to worst. In the VPIapproach, candidate sequences are chosen by permuting one value in the incumbent(current best) sequence, and this new sequence is either accepted or rejected based onsome rule. Both methods rely on models which can provide quick estimates of resultsto trim down the solution space to a small sub-region where the application of moreaccurate, more time consuming analysis techniques is realistic.


interaction to set up the problem and evaluate the solution. Clearly, a solution which ismuch faster and better suited to automation is required even for these search algorithmsto be feasible.

Reducing the Search Space

The initial search space at the beginning of the optimization problem contains alldesigns resulting from all possible combinations of process parameters. A large subsetof these designs are infeasible because they will not produce a part with acceptablegeometry, microstructure, and mechanical properties. The remaining subset of designsare feasible n that is, they will produce acceptable parts. The optimal design is withinthe feasible solution space.

Traditional optimization seeks to identify a single best design which represents aglobal minimum of the objective function. By softening the requirement of optimum to"good enough," substantial reductions in computation times can be realized. Both of theaforementioned researchers (Ho and Jacobson) are implementing this goal softening,where the goal is to select a design within the top 5% instead of the. absolute optimum.

Simplified Models

In this approach to optimization, a large number of randomly generated designsis simulated using computer models of the manufacturing processes. The results of eachsimulation are evaluated using some objective function, and based on these resultsadditional designs are generated and evaluated. The process continues until the designis judged to be near enough to optimum.


It is desirable to move very quickly from the initial search space to the feasiblesubset, then to the near optimum subset (figure 1). Once the smaller near optimum subsethas been identified, a slower, more accurate search can be conducted to find a "goodenough" design. Readily available computer models of manufacturing processes such asFinite Element Analysis give very good detail at the expense of speed. In the early andmiddle stages of reducing the search space, there is no need for precise detail, but onlyfor generalized output parameters which can be used to identify a design which is clearlyinfeasible or not near optimal. Furthermore, all of the computer models should employconsistent means for representing geometry, mechanical properties, and design parametersof the process. As Yang [4] observes, despite the development of many successfulmodels for individual manufacturing processes, no consistent system that integrates all ofthose individual modules has yet been developed.

This paper describes the development of such a system of consistent forging,machining, extrusion, and heat treatment models which execute extremely quickly whilestill providing enough accuracy and detail to screen potential designs and arrive at a setof near optimum designs, which can be further analyzed using slower, more detailedmodels. The model relies on fundamental equations of the thermomechanical processesbeing modeled combined with simplifying assumptions about geometry and mechanicalproperty variations throughout the workpiece.

Simplified models which provide good estimates of average strain, strain rate, andtemperature are used to identify process designs which fall within the processing windowand eliminate those which include undesirable processing regions. A comparison of thesimplified forging model and a finite element model of the forging of an aircraft enginecompressor disk is given.

PROGRAM STRUCTURE

In order to implement a consistent system of models, the output from each modelmust be in the same format as the input to the next model. Furthermore, the order ofoperations is not defined a priori, so it should be possible to call the models in any order.The data required for each operation may be grouped into two classes: geometry andthermomechanical history: Geometry data includes only the shape of the part. Thecomponents currently being studied are axisymmetric, so the geometry can be stored asa 2 dimensional profile of the cross section. Thermomechanical history includes suchfield properties such as strain, strain rate, and temperature, as well as microstructure dataincluding grain size. All of these are continuum properties. However, since the objectivefunction evaluation is based on a few characteristic values, it is reasonable to divide theworkpiece into a small number of disks and rings, and calculate these characteristic valuesonly in these regions (see fig. 2).

The second input data set for each model is process parameters. Inputs to theforging model include die geometry, ram velocity profile, die temperature, and lubricationconditions. The model will produce process specific output parameters. For forging theseinclude total ram force and peak die pressure.


PART SPECIFIC PROCESS MODELSIn the manufacture of an aircraft engine compressor disk there are several possible

sequences of operations which will result in a component which meets all geometric,mechanical, and microstructural requirements. A schematic illustration of the possiblepaths is shown in figure 3.

Forging Model

All changes in continuum properties can be described as functions of geometry orthe change in geometry. Thus, it is necessary to have a reasonable model of the geometrychange during the forging process before strain, temperature, or microstructurecalculations can be made. Unfortunately, most methods such as FEM or UBET whichpredict geometry changes rely on the determination of a velocity field based on aniterative energy minimization. For a model which will be executed many thousands of


times for a single optimization, any iterative procedure is unacceptable if a direct closedform solution is available.

Based on study of several finite element simulations of the die fill process, distinctmaterial flow phases occur during die fill. (see figure 4). In each phase, the strain ineach ring is calculated from the height and diameter change'. Reasonable approximationsof these flow phases can be obtained by making certain assumptions about constantheights or volumes of the deforming regions. Lee [5] has shown that deformation energyis minimized if the height of ring 3 (figure 2) remains constant during free expansion.

The temperature change is calculated from heat generated due to plastic work lessthe heat lost to dies and the environment. Microstructure changes are a function of strain,strain rate, temperature and time, and can be found from semi-empirical models.

The die pressure can be determined from modified slab analysis, withmodifications for inclined surfaces, free surfaces, and flash extrusion. For arbitraryinclined surfaces, as illustrated in figure 5, the equilibrium equation is given by

Machining Model

The machining process is unique in that, unlike forging, it affects only geometryand not material properties of the workpiece. As such, the machining parameters may beoptimized independent of the other processes in the selected sequence. Once optimumtool feed rate, spindle speed, and depth of cut have been determined, an optimum materialvolume removal rate can be determined. Hence, machining cost can be estimated froma simple. cost/volume of material removed formula.

The machining model simply checks to insure that the specified machined shapecan be achieved from the specified workpiece, then calculates the volume of material tobe removed. The cost function calculates a dollar cost based on this volume, and returnsthe value to the main calling routine.

Optimization Techniques

The structure of the optimization programs is modular. The main optimizationroutine defines the process sequence. This includes the processes, the order in which theywill be performed, and process parameters such as times, temperatures, ram velocities,and die geometries. This information is passed to the cost (objective) function module.This model calls the process models in their appropriate sequence. The process models


compute the final geometry, strain, peak strain rate, and temperature of the part, as wellas process time if it is not specified, and return these values to the cost function. Fromthese results, the cost function module calculates the value of the objective function andreturns this value to the main calling function. The VPI code selects its next candidateprocess based on the value of the objective function. The Harvard code randomlychooses its next candidate process.

Validation of Model

The results from the simplified forging model were compared to the results offinite element analysis. There is good correlation between fill patterns shown in figure6, and between strain and temperature values. The FEM fill patterns shown in figure 7are 1) initial indentation, 2) workpiece contacts flange section of die, 3) outside cornerof flange fills, 4) inside comer of die backfills, 5) flash forms.

ILLUSTRATIVE EXAMPLE

Obiective Function

To address the affordability concerns of the Air Force, the total cost of the partis the logical function for minimization. The costs of materials and tooling, along withthe processing operations costs are well established. The probabilities of the introductionof latent defects during processing are determined by process operation statistics. Thecost of a failure to achieve final shape or of flaws introduced by processing the part ina region of process instability (outside the processing window) is the same as the cost ofa latent defect discovered at the end of the processing, scrapping the part. Microstructuralobjectives are more difficult to assess costs, but more important to the final cost sincethey determine the long term serviceability of the component. Here a Taguchi analysisis a appropriate, with a variation in a microstructural parameter that leads to prematurecomponent failure having the same cost as a critical latent defect, scrapping the part ifdetected, or the purchase price of an aircraft if undetected. As the total cost is dependenton the processing sequence, which represents discrete variables, evaluation is simple,while analytical minimization is not feasible.

The·total process cost is calculated as the raw material cost, plus the sum of theprocessing costs for each discrete operation, plus the sum of all penalty costs. Processingcosts include:

• Equipment costs. based on hourly depreciation over the life of themachine;

• Labor rates for the number of workers required to set up the process,monitor the operation;


• If the workpiece does not undergo sufficient strain to induce dynamicrecrystalization, an additional recrystallization heat treatment (RX heattreat) step is required. Hence, the penalty for low strain values is the costof this heat treat step.

Sample EvaluationsThree sample cases were evaluated using the objective function evaluation

techniques described here. The simplified process consisted of billet procurement, oneor two forging steps with preheat, heat treatment, and final machining. The part evaluatedwas a hypothetical turbine rotor made of TiAI 49-2 with a diameter of 60cm. Theprocesses evaluated were:

CASE 1: Isothermal Near net shape forging with a 1cm envelope around the final partshape. Mean strain rate was 0.001. Forging temperature was 2150K. High die stressesand forging temperatures lead to a relatively low die life and resultant high costs forforging. The minimum strain value was not reached in all sections of the workpiece, soa recrystallization heat treatment step was required.

CASE 2: Pancake forging of a rectangular cross section billet, then bulk machining tothe final shape. The die stresses were significantly lower, and the simpler die shape ledto significantly lower die costs, which offset the increased material and finish machiningcosts.

CASE 3: Pancake forging of preform, near net shape forging, and finish machining. Thenear net shape forging operation was conducted at a lower strain rate, which improveddie life, but moved the process out of the optimal processing window, and incurredsubstantial cost from the microstructure penalty function.

A graphical comparison of the cost function contributions for each sequence isshown in figure 8. More extensive runs of the VPI code have suggested that for smallerforgings, pancake forging followed by extensive machining is the most cost effectiveapproach. However, as the size of the forging increases, near net shape forgings becomeincreasingly economically feasible.

CONCLUSION .

This paper has presented a new approach to modeling the entire manufacturingprocess as a discrete event dynamic system optimizing the entire process based on thedesired properties of the end product. It is shown that the current method of optimizingindividual steps of manufacturing does not yield an overall optimum for the process. Theglobal cost function evaluation has been applied to a compressor disk and preliminaryresults show the influence of a) process selection and b) parameter selection on theindividual cost of discrete processing steps as well as the overall cost. In order to bestatistically significant, a huge number of designs have to be evaluated. If theoptimization is to be completed in a realistic time frame, the models must execute very


quickly. The models introduced here are several orders of magnitude faster than finiteelement analysis, but sacrifice detail and a bit of accuracy for that speed.

Figure 8 Comparison of cost contributions for 3 manufacturingexamples

Ordinal optimization is particularly well suited to work with less accurate models.The technique can quickly screen a very large search space and identify a small group ofpotential solution which will, with some statistical certainty, contain a "good enough"optimal solution.

REFERENCES

[1] P. Shvng, M. Srinivasan, 1995, "Multi-Objective Process Planning inEnvironmentally Conscious Manufacturing: A Feature-Based Approach", Annalsof CIRP, Vol 44/1/1995. p.433-437.

[2] K. Yamazaki, Y. Kawahara, J.e. Jeng, H. Aoyama, 1995, "Autonomous ProcessPlanning with Real-Time Machining for Productive Sculptured SurfaceManufacturing Based on Automatic Recognition of Geometric Features," Annalsof CIRP, Vol 44/1/1995, p.439-444.

[3] J.H. Zhang, S. Hinduja, "Determination of the Optimum Tool Set for a GivenBatch of Turned Components," Annals of CIRP, Vol 44/1/1995, p.445-450.


[4] M.S. Yang, 1995, "Optimization of the Production of Integral Blade and RotorComponents with a Hybrid Discrete Event Model based on Ordinal Optimizationand Genetic Algorithm," Unpublished.

[5] T.Y. Lee, "Improved Slab Method for Axisymmetric Forging", M.S. Thesis, OhioUniversity, Athens, OH 1996.

Daniel 1. Antos!, and Edward G. Nisbett2

MANUFACTURE AND PROPERTIES OF CONTINUOUS GRAIN FLOWCRANKSHAFTS FOR LOCOMOTIVE AND POWER GENERATION DIESELENGINES

REFERENCE: Antos, D. 1., and Nisbett, E.G., "Manufacture and Properties ofContinuous Grain Flow Crankshafts for Locomotive and Power Generation DieselEngines," Steel For~infis Second Volume ASTM STP 1259, E. G. Nisbett, and A. S.Melilli, Eds., American Society for Testing and Materials 1997.

ABSTRACT: The bulk of the large crankshaft production volume is associated with themedium speed diesel engine market These engines have seen intense development tosustain higher power outputs without change in the physical size of the crankshaft and atthe same time there has been continuing pressure to reduce costs. Fatigue and bearingjournal wear are the major technical hurdles that threaten the crankshaft life, andmeasures for dealing with these issues are described. Continuous grain flow (CGF)crankshafts are responsible for the continued integrity of these enhanced power outputengines and the production of these crankshafts is described. Comparisons are madewith the older slab forging crankshaft production method. The demand for the mediumspeed diesel engine and its natural gas derivative is strong and supports an aggressiveengine building industry serving locomotive, marine and power generation markets. Thisdemand in turn relies on practical national standards that serve the needs of the enginebuilder, material supplier and the end user

KEYWORDS: crankshaft, continuous grain flow, forgings, carbon steel, alloy steel,induction hardening, nitriding, locomotives, diesel engines, generators.

I Director of Quality Assurance, National Forge Company, 1 Front StreetIrvine Pennsylvania 16329 USA.

2Senior Staff Metallurgist, National Forge Company, 1 Front Street, IrvinePennsylvania 16329 USA.

129


Most domestic freight and many passenger train locomotives are diesel electricpowered, and in the midwestern states it is not uncommon to see long coal or mineraltrains being hauled by as many as five diesel electric locomotives. The medium speeddiesel engines used in locomotive service are currently rated up to about 3,275 kW (4,400hp), and new locomotive engines rated at 4,475 kW (6,000 hp) are expected to beginservice in 1997. Sixteen and eighteen cylinder Vee type engines are the commonly usedstyle and these generally operate in the 950/1 ,200 rpm speed range. This type of enginealso sees service as main propulsion units for trawlers and other small ships. Otherfrequent marine use is in power generation sets. Stationary power generation sets,commonly teamed with waste heat recovery equipment are used in central powergeneration equipment. In Europe where many of the rail systems are electrified, themain applications for the medium speed diesel engine lie in central power generation,frequently using natural gas conversions, and in marine applications that do include theoffshore oil producing platforms.

Universally the crankshafts for medium speed diesel engines are, by choice, of thecontinuous grain flow (CGF) type. This is because of the enhanced fatigue strengthassociated with this type of manufacture. The classification agencies in recognizing thesuperiority of the CGF crankshaft permit them to be designed to higher stress levels, thusreducing crankshaft size and engine weight. As an example one prominent classificationagency [1] permits the use of a factor of 1.15 for CGF crankshafts compared to 1.0 forslab forged shafts. This factor is increased to 1.25 if the crankshaft were to be nitrided.

MANUFACTURE: FORGING

Slab For~ing

Until the advent oflarge CGF crankshafts, the common method of forging largediesel and compressor crankshafts was by slab forging. This involved forging the shaft asa long rectangular shape with the output and auxiliary drive shafts offset forged at eachend, The central axis of the original forging ingot would then run the full length of theshaft close to the centerline of the rectangular or slab section. The slab was then notchedout at the crankpinjournallocations, to be followed in turn by local heating and twistingoperations to set the crankpins in the correct angular positions. Following this work thecrankshaft would be heat treated to develop the required mechanical properties. A sketchshowing these stages of manufacture is shown in Figure 1. The major disadvantage ofthis manufacturing method is that the ingot longitudinal axis lies close to the criticallyloaded parts of the crankshaft, so that adverse concentrations of nonmetallic inclusionsmay be present where they could cause maximum damage. Apart from reducing thefatigue strength capability of the crankshaft, there is a high manufacturing rejectionprobability that adversely affects costs and scheduling. However this manufacturingmethod is still used for low production volume crankshafts, especially for large slowspeed engines or compressors.

CGF Crankshafts

The prime objective of the CGF forging is to keep the ingot axis close to thecentral axis of the crankshaft throughout its length. This enables the metallurgicallycleaner steel to be located at the highly stressed, or critical zones of the crankshaft, andso enhance fatigue strength. There are several forging procedures that can accomplishthis task, and in all of them the starting material is in the form of a forged or hot rolledbar. This is to ensure that the ingot centerline corresponds to the bar centerline.

Closed Die Forgin~

After initial blocking or shaping of the bar the closed die crankshaft is forged toform the complete shaft with the crankpins in the required angular positions. Thephysical limitations of press size and power limits the size of this type of crankshaft,both in terms of crankpin and bearing journal diameter, and overall crankshaft length.The length limitation has been overcome in some engine designs by making the shaft intwo sections joined in the middle by means of bolted flanges. Crankshaft size can beincreased also by forging the shaft with the crankpins all in the same plane, followed byhot twisting as is done in slab forging to obtain the proper angularity.


TR (Tadeusz Rut) and RR For~ing Devices

These machines are used in conjunction with an open die forging press to formCGF crankshafts one crankpin throw at a time. They are fairly complicated and usesystems oflevers to give axial upsetting forces to form the webs well as the verticalmotion of the press ram to give the crankpin offset. Dies are'included to give the webshape although appreciable stock may have to be removed to arrive at the final shape.This is in contrast to the closed die forgings where the forging size is close enough to thefinished dimensions to permit the use of as forged webs. The TR device [2.] [3.] wasdeveloped in Poland and has been thoroughly described in the literature, but this devicewas predated by the French RR machine [1] from a concept originating in 1937. It isclaimed that the term CGF was coined for the RR process. However practical use of theprocess does not seem to have been reported until about 1960. Machines of both typesare in use in many countries [i], particularly for the production of the very large marinepropulsion engines. These crankshaft forging machines require the dedicated use of theopen die forging press during the crankshaft production period.

Work continues on crankshaft manufacturing development and two papers on thissubject [{i],[l] were presented at the Twelfth International Forgemasters Meeting held inChicago in 1994. One of these [7] reported reducing the time to make a 50 tonnemarine crankshaft forging from 15 weeks to 5 weeks, largely by reducing the number offorging reheats.

ANTOS AND NISBETT ON MANUFACTURE AND PROPERTIES 133

Installed over 35 years ago, the unique forging press at National Forge company,Irvine Pennsylvania was designed to make diesel engine crankshafts primarily for thedomestic railway locomotive industry. The press was intended for high volumeproduction based on the anticipated growth in diesel electric propulsion systems. Thepress, was described in a paper by Becker 1967, [8.]and is shown in Figure 2.

The press has a cruciform shape, with a tie rod press that has an upsetting ram of3,600 tonnes capacity at right angles to a cast steel C-frame press of2,000 tonnescapacity. The C-frame press has two opposed rams that provide the essential offsettingcapability. As with the RR and TR machines the crankshaft is built up one crankpinthrow at a time, so that the press is ideal for making several crankshafts of the same typeduring a single press run. A simplified forging sequence is shown in Figure 3.

Forge heating of the successive segments of the crankshaft bar is accomplished in apair of horizontal induction furnaces that enable one bar to be forged whilst the next inline is being heated. The plant layout is shown in Figure 4.

In operation a calculated length of bar is induction heated to the forgingtemperature before being transferred to the press by a special crane. The coupling flangeis forged first followed by the successive crankpin throws. Although the forgings madein a given press run must be of the same design, the number of crankpin throws can vary,for example eight throw crankshafts for a V-sixteen cylinder engine can be forged in thesame run as four throw crankshafts for a V-eight engine of the same design. A specialrotating die that is part of the crankpin die stack enables succeeding crankpins to beforged in the correct angular position.


Figure 4 - - Multidirectional press layout with theinduction furnaces at the bottom right side.

The grain flow of a typical CGF forging made in this press is shown in Figure 5

Figure 5 - - Macro etched longitudinal section through a CGF crankpinthrow showing how the grain flow follows the contour of the forging.


The process lends itself to high volume production with forging runs of 40 to 50crankshafts being not uncommon. Such capacity is needed to meet the demands of thedomestic railroads, as well as to supply the world wide export market. A typical COFcrankshaft forging is shown in Figure 6.

Manufacture: Heat Treatment

Heat treatment of COF diesel crankshafts follows forging directly unless weldedcounterweights are to be used, in which case the counterweights are attachedimmediately before heat treatment. The choice of heat treatment cycle is usually left tothe manufacturer and depends on the type of material used and the required mechanicalproperties. Normalizing and tempering is preferred from the manufacturing stand pointbecause the risk of distortion is reduced and the lower yield strength to tensile strengthratio makes for easier straightening. However the choice is usually dictated by therequired tensile strength and surface hardness, as well as the need for impact testing.Most classification societies require charpy impact testing so that when agencyinspection is a part of the order, then the crankshafts must be quenched and tempered.For some of the medium carbon - manganese steels water quenching is possible, but forthe alloy steels a water - polymer or oil quench is used.

Impact testing is rarely required for domestic railroad crankshafts and since failureinvariably.involves fatigue the test is considered to be of no particular value relative tothe life or performance of the crankshaft.

Stress relief after press straightening is generally considered to be an essential partof the crankshaft heat treatment process.

Manufacture: Material

The material used in diesel crankshaft manufacture is chosen partly on the basis ofthe minimum tensile strength and base hardness properties, and also on whether or notsurface hardening techniques are to be applied in finishing the crankshaft. Surface


hardening is intended to increase the wear resistance of the crankpin, main bearingjournals and thrust collars. If the hardening process is extended to the crankshaft mainbearing and crankpin fillets, then the fatigue strength of the crankshaft is increased verysignificantly by the introduction of residual compressive stresses. This improvementalso applies to the highly stressed areas at the bearing oil holes. The commonly usedcrankshaft materials are shown in Table I.


Material Cleanliness

With the trend towards higher tensile strengths for crankshafts, and coupled withthe improvements in steel quality that can accompany secondary refining and vacuumdegassing, minimum material cleanliness requirements are being asked for increasingly.This is usually done on a heat qualification basis, testing the product from the top andbottom of the first middle and last ingots in the heat. An equivalent plan is used forbottom poured ingots, taking the number of pouring plates used into account. The steelsare generally rated to ASTM E 45 Method D, Standard Practice for Determining theInclusion content of steel., DIN 50 6023

, microscopic examination of special steels usingstandard diagrams to assess the content of non-metallic inclusions, or 1IS G 0554,microscopic testing method for the non-metallic inclusions in steel. The samples areoften taken at the mid-radius position of crankshaft bars produced from the heat or at thenear surface and mid-radius positions. A maximum oxygen content in the steel also hasbeen used to help control cleanliness.

Tension Test

Because fatigue strength is the principal mechanical property in the design ofcrankshafts much emphasis is placed on the ultimate tensile strength of the heat treatedforging. The tensile strength of the crankshaft is often capped to require a workingrange, and a surface hardness range is sometimes imposed to get the same effect.However, this practice can lead to heat treatment difficulties if the chosen surfacehardness range is not compatible with the mid-radius tension test requirement. Thetension test is taken from an integral prolongation of the crankshaft forging so that thetest section will have the same diameter as either the crankpin or main bearing journals.

The yield strength requirement is usually specified as a minimum value, and isoften chosen depending on the type of heat treatment likely to be used. Occasionallyyield strength is required to be a minimum percentage of the actual tensile strength.Typical crankshaft tensile requirements are shown in Table 2 for medium speed dieselengine crankshafts currently in production.

Impact Test

As mentioned earlier, many of the classification agencies require charpy impacttesting of crankshafts. The need for this test is difficult to justify, since stressconcentration points that might give rise to brittle failure cannot be tolerated in acrankshaft because of the severe fatigue loading. Apart possibly for cold start-upsituations, the crankshaft runs at the lubricating oil temperature of about 100°e. Sufficeit to say that in-service brittle failure of medium speed diesel engines would be anextremely rare event. Be that as it may however, charpy testing of agency inspected

3 Verlag Stahleisen mbh, Postfach 8229, D-4000 Dusseldorf Germany4Japanese Standards Association 1-24 Aksaka 4, Minato-Ku Tokyo 107 Japan


crankshafts is virtually a universal requirement for crankshafts sold in Europe and someother offshore countries. This requirement has the effect of mandating a quench andtemper heat treatment cycle and can even force the use of an alloy steel instead of acarbon manganese steel that would be used if no impact testing were to be required.The minimum absorbed energy requirement for crankshafts is commonly in the range of22J to 30J at room temperature.


Hardness

As has been mentioned, hardness ranges are often used as a control measure forexcessively high tensile strengths but increased hardness also reduces bearing journalwear. This is an added advantage for the higher strength crankshafts. However fordomestic railroad application supplementary bearing journal hardening is generally usedto improve wear resistance.

MANUFACTURE: SURFACE HARDENING

The most widely used surface hardening techniques for the large CGF crankshaftsare induction hardening and nitriding. These systems can be expected to increase thetreated surface hardness in the commonly used crankshaft materials to between 48HRCand 56HRC. This significantly boosts wear resistance in the bearing areas. Anotheradvantage of these surface treatments is that they induce high residual compressivestresses in the treated surfaces, and these afford significant fatigue strength enhancementto the treated areas. In Table 3 the material grades, heat treatments and surfacetreatments are matched.

Induction Hardeninl:

Induction hardening of the bearing journal surfaces of medium speed dieselengine crankshafts has been a well established practice for very many years, and a casedepth of 6 mm to 9 mm is commonly applied to the main bearing and crankpinjournalsurfaces. great care is taken to ensure that the hardened zone does not encroach on thefillet areas to avoid the potentially harmful effects of residual tensile stresses at the edgesof the hardened track encroaching on the highly stressed fillet area. In Figure 7 a macro-etched section through an induction hardened bearing journal surface is shown. It can beseen that the hardening pattern ends clear of the edges of the fillets. Another part of theinduction hardened journal surface that has to be protected is the radius area around andinto the oil holes, since high torsional stresses are experienced in these areas.Unfortunately very short and hard to detect hardening cracks can occur at these oil holes,and great care qas to be taken during final magnetic particle examination to make certainthat none of these defects are present.

The hardening process involves rapidly austenitizing the surface layers of thejournal by means of induction coils and following this immediately with a liquid quenchin a water-polymer solution. In order to be effective the carbon content of thecrankshaft materials usually in the range of 0.45% to 0.50% Many of the inductionhardened crankshafts are made using carbon - manganese steels, and a common variant


is the 5046 grade. For the higher strength crankshafts, chromium - molybdenum lowalloy steels such as 4145 are often specified. It is possible to increase the fatiguestrength of an induction hardened crankshaft by extending the hardening to include thejournal fillets. This involves the use of special inductors such that the journal bearingsurfaces and the fillets are hardened at the same time, thus avoiding softened rings at theedges of the fillets. This type of procedure is widely used for small automotive andtruck crankshafts, however in these large diesel engine crankshafts the hardening processis drastic and distortion risks are large with consequent straightening difficulties. Afterhardening the crankshafts are tempered at about 260°C to avoid instability at the engineoperating temperature. Because ofthe good depth of hardening capability it is usual toleave appreciable grinding stock on the hardened surfaces. This assists in dealing withdistortion during hardening, but the producer must be on guard against grinding cracks.An example of a fully hardened large COF crankshaft journal section isshown in Figure 8.

An advantage of induction hardened bearing journals is that the depth ofhardening, of the order of 6 mm, will permit regrinding of a badly scored bearing and theuse of undersize bearing shells. However such damage incidents quite often involvebearing failure and surface heating. If this happens the induction hardened surface willbe softened, and worse, intergranular penetration by the copper or a copper alloy fromthe bearing might also be present.

TABLE 3 --Surface treatment

Grade Heat Treatment Surface Treatment

5046 (mod) Normalize, Quench, Temper Induction harden

S44 SY Normalize, Quench, Temper As heat treated

4145H Normalize, Quench, Temper Induction harden

4130 (mod) Normalize, Temper Nitride

4]40 (mod ]) Normalize, Temper Nitride

4]40 (mod 2) Normalize, Temper As heat treated

4]40 (mod 3) Normalize, Quench, Temper As heat treated, also Nitride

S42Crl Normalize, Quench, Temper As heat treated

34CrMo4 Normalize, Quench, Temper As heat treated

42CrMo4 Normalize, Quench, Temper As heat treated

BS 1504623 (mod) Normalize, Quench, Temper As heat treated

BS 970-817 M40 (mod) Normalize, Quench, Temper As heat treated


Figure 8 - Example of a fully induction hardened journal bearingsurface showing the continuous case from one bearing fillet to the other.

Nitriding .

Nitrided crankshafts are made from chromium - molybdenum steels such as 4130,although many of the higher strength variants shown in Table 1 can be used. Howeverbecause the critically loaded surfaces of the crankshaft are readily nitrided at the sam~time as the bearing journals, high strength base materials are not always required. Thismeans that a normalized and tempered heat treatment is often adequate, and the loweryield strength makes post heat treatment straightening easier to do.


Both gas nitriding and ion nitriding methods are applicable. The gas nitridingmethod is readily adaptable to batch processing several shafts in the furnace load. Thecycle is long, taking in all about five days from start to finish, and the nitridingtemperature is generally between 5100C to 540°C. Consequently the temperingtemperature for the crankshaft must be at least 570°C. Areas of the crankshaft that arenot to be nitrided are stopped off using special paints that usually contain tin. Thenitrided case depth relatively shallow, being about 0.6 mm, including a thin surfacecompound called white layer. When using the Floe nitriding procedure the white layerdepth is about 0.025 mm, and this must be removed from bearing surfaces because it isfriable in nature, and could cause bearing damage if left in place on the bearing journal.The white layer is removed by polishing the bearing journals, and this process also bringsthe journals and thrust collars to size, as well as providing the required surface finish.It is apparent from this that the crankshaft is close to final size immediately prior to

nitriding. The ion nitriding process does offer more control over the surface white layer.Metal shields are often used to prevent nitriding of specified crankshaft surfaces. becausealthough paint can be used, it can tend to foul the vacuum chamber of the furnace.Because of their complex shape ion nitriding of large crankshafts is often limited to oneshaft per furnace load.

Nitriding provides a most uniform case regardless of the geometry of the part, andthis is especially important in highly stressed areas such as oil holes. Interruption of thecase, for example by a surface stop off paint spot, as is shown in figure 9, will leave anon-nitrided area with a periphery of high residual tensile stress. This would be veryundesirable in a bearing journal fillet for example.


Effect ofNitridin~ on Fati~ue Stren~th

The beneficial effect of nitriding on fatigue strength is shown in Figure 10 fortests run on R.R. Moore rotating beam samples for 4130 normalized and temperedmaterial. Fatigue testing of full size crankshaft sections has confirmed the significantbenefits of nitriding in increasing fatigue strength, and improvements of 33% have beenreported [~]. As previously mentioned one of the classification societies [1] permits theuse of a 1.25 multiplication factor in calculations for fatigue strength. The general axiomthat one should use data from the fatigue testing of actual components, rather than relyingon results from very small specimens hold especially true for crankshafts.


CRANKSHAFT COUNTERWEIGHTS

In automotive practice the crankshaft is generally forged with integralcounterweights, but for large medium speed diesel engine crankshafts the counterweightsconstitute a separate component. Traditionally these weights were bolted onto thecrankshaft, and this practice is still followed in crankshaft design by most oftheEuropean builders. The counterweight fit onto the shaft web must be made to closetolerances, and with carefully produced locating devices such as dowel pins. High gradebolting must also be used to secure the counterweight. A typical counterweight locationarea is shown in Figure II.

Another manufacturing method that originated over 35 years ago for the highvolume domestic diesel locomotive market is to weld the counterweights into position asshown in Figure 12. This is done by the manual metallic arc process, or by semiautomatic welding using carbon dioxide (C02) as the shielding gas.

Following welding, the joints are ground for magnetic particle examination andthen after any repair!>the crankshafts are heat treated for mechanical properties.Although the strength of the weld is reduced by this procedure, the heat effected zone inthe crankshaft web is effectively removed, and post weld stress relief is also effective.The welds are examined again after machining, but while there is still time to stressrelieve any needed repairs. Dynamic balancing is done near the end of themanufacturing process, weight can be removed from the counterweights for this purpose.This method of attaching the counterweights has been very successful both in terms ofsafety in the engine, and in reducing manufacturing costs. The counterweightsthemselves are made from a low carbon - manganese steel such as 1020.


NONDESTRUCTIVE EXAMINATION

Ultrasonic Examination

Ultrasonic examination of crankshafts is required by many of the classificationagencies. This examination must be done after heat treatment for properties, but at astage in machining before certain design features, such as oil passages are included.However most CGF diesel crankshafts are not subjected to ultrasonic examinationbecause the critical material is located on the final crankshaft surfaces, rather than beingburied within the crankshaft cross section.

Ma~netic Particle Examination

The most significant nondestructive examination is by the magnetic particlemethod when the crankshaft is in the finished condition. Because of the difficulty inadequately demagnetizing the crankshaft when direct current (DC) magnetization is used,the alternating current (ac) method is preferred. The latest test method for this is toASTM A 966 Magnetic Particle Examination of Steel Forgings using AC Magnetization.The risk of attracting magnetic debris in the lubricating oil to the bearing journal surfacesis the reason for requiring the minimum level of residual magnetism. Aside fromdetecting crack-like indications that might be left after grinding an induction hardenedsurface, and surface nonmetallic inclusions, the wet continuous fluorescent method willreveal the presence of white layer and non nitrided spots on nitrided surfaces, and coldworked rings on polished surfaces.


CONCLUSIONS

The need for continuous grain flow crankshafts has been spurred by the dualneeds of improved quality to enhance fatigue life, and to economically supply a largemarket demand.

Improved and often ingenious production methods are continuing to evolve on aglobal basis, to improve crankshaft quality and reduce manufacturing costs.

The need to extract higher power outputs either from existing engine designs, orfrom new designs for a fiercely competitive market has required the use of more highlyalloyed steels, and more sophisticated machining methods for the crankshafts

Surface treatments to strategically enhance fatigue strength and wear resistancehave been shown to be highly effective, but many existing crankshaft designs have yet totake advantage of these systems.

Welded counterweights significantly reduce crankshaft manufacturing costs forthe medium speed diesel engine without sacrificing reliability or safety.


REFERENCES

[1] Lloyds Register Published Rules and Regulations Part 5, Main and AuxiliaryMachinery, Chapter 2, Oil Engines (Effective Date 29 January, 1994), Section 3Design, 3.8 Fatigue Strength.

[2] RUT, T. "New Forging Method of Semi Built-Up Crankshafts", Proceedinis of the8th International FOfllemasters Meetini October 23-28. 1977. Kyoto. Japan

[3] RUT, T. "Forging of Long Stroke Crankshafts by the TR Method", Steel For~in~sSymposium. Williamsbur~. Vir~inia ASTM STP 903 pp. 504-519.

[~] Ruget, G., "Development of the RR Continuous Grain Flow Process forCrankshafts", Proceedin~s of the 5th International Forllemasters Meetinll. May. 1970.Terni. Italy pp. 503-520

L5J Araki, S., Ochi, T. Fujita, H., Hirano, 1., Takahara, H, "New Forging Process forContinuous Grain Flow Crankshaft in Upsetting and Offsetting Operation",Proceedin~s of the 11th International Foriemasters Meeting. Terni. Italy. June 1991.

[2] Jilek, L., Dvorak, B., Molinek, B., "Big Crankshaft Forgings" Proceedin~s of the12th International Forgemasters Meetini. Chicago. Illinois. September 11-17. 1994.

[1] Bokelmann, D., Forch, K., "Forging of Heavy Continuous Grain Flow (CGF)Crankshafts without Intermediate Cooling," Proceedinlls of the 12th InternationalFoq~emasters Meeting. Chicago. Illinois. September 11-17. 1994.

L8.] Becker, 1.R., "Forging Crankshafts on a Horizontal Multi-directional Forging Press"Metal Formin~. January. 1968 pp.4-10.

Thomas E. Hebel *

REDUCING STRESS RELATED PROBLEMS IN STEEL PORGINGS USING SUB-HARMONICVIBRATIONAL ENERGY

REPERENCE: Hebel, T.E., "Reducing Stress Related Problems in SteelForgings using Sub-Harmonic Vibrational Energy," Steel Forqinqs: SecondVolume, ASTM STP 1259, E.G. Nisbett and A.S. Melilli, Eds. , AmericanSociety for Testing and Materials, 1997.

ABSTRACT: There are predominantly three residual stress relatedproblems plaguing steel forgings. They are 1) distortion followingmachining, 2) delayed distortion, and 3) premature cracking tendencies.The major source of these problems is the stress caused by a sharptemperature drop from hot forging or welding.

This paper will detail SUB-harmonic vibrational technology as aneffective stress relief solution to reducing these problems. Also, thispaper will outline the three step approach for using the sub-harmonicenergy process during welding. When used during welding additionalbenefits of reducing weld distortion and cracking will be achieved.

Data collected from seven Test Cases will be presented including reducedmachine distortion (43-76%) from forged cannon barrels and aerospacecomponents. Reducing premature cracking will be demonstrated usingexamples of longer fatigue life (200-400%) from components such asforged crankshafts and connecting rods used in racing engines. Datafrom the U.S. Department of Energy sponsored study on this technologywill be presented. This report concluded that there is evidence thatthe Sub-harmonic treated and thermal stressed relieved specimens are"comparable."

KEY WORDS: Sub-Harmonic, stress relief, vibration stress relief,mechanical stress relief

Steel forgings are desired because of several reasons inclUdinghigher quality, greater density and uniformity, higher strength in alldirections, good machinability, and nearer net size. Yet despite thesetypical advantages manufacturing and performance problems still existunless a stress relief process is employed.

In an effort to reduce these problems SUB-harmonic vibrations havebeen applied to forgings in two ways - as a stress relief process andduring welding as a weld conditioning process.

SOURCE OP THE PROBLEM

Steel forgings are subjected to a sharp temperature drop duringthe forging process. This induces thermal residual stress [1] . Otherexamples of thermal stress inducing processes are welding, casting,machining and grinding, hardening, and EDMing. Thermal stress causesthe vulnerability of metal to have certain types of problems in

* Vice President, Bonal Technologies, Inc. , 21178 Bridge St.,Southfield, Michigan, 48034

148

HEBEL ON REDUCING STRESS 149

manufacturing and in service. These problems are distortion immediatelyfollowing machining, grinding or heat treating, delayed distortion, andpremature cracking.

To reduce and eliminate these problems the workpiece can be stressrelieved by either thermal or mechanical means [2]. In the past thecommon method of stress relief has been by heat treating. But heattreat stress relief, although effective, has many drawbacks.Significant drawbacks to heat treat stress relief include time andexpense, changing mechanical properties, treatment distortion, surfaceoxidation, and limitations in size and weight.

Other stress relief methods have not been widely used on forgingsfor obvious reasons; natural ageing takes too long, cryogenics is tooexpensive and limiting in size, stretch and compression both requiresimple shaped parts. This leaves vibration. Vibration stress reliefdoes not have any of the above limitations.

When considering using vibration for stress relieving one mustrealize that, like heat, many different energy levels are possible butonly a small range will generate consistently effective stress reliefresults.SUB-HARMONIC VIBRATIONAL STRESS RELIEF

Bonal Technologies, Inc., Southfield, Michigan, developed avibrational stress relief process to accomplish the standard stressrelief benefits of:

1. reducing distortion immediately following machining andgrinding; (Fig. 1)

2. reducing delayed distortion; and3. reducing premature cracking.

Furthermore, Bonal discovered that this same process can be usedduring welding to reduce weld distortion and produce a very crackresistant weld.


The process, called META-LAX [TM], has two fundamental principlesto its operating procedure:

1. SUB-Harmonic energy must be used for the actual stress re-lieving frequency; and

2. the harmonic curve of a thermally stressed part will shiftand tabilize to a new frequency location as the workpiecebecomes relaxed of thermal stress.

PRINCIPLE #11 SUB-HARMONIC ZONEAll metal components will at some frequency exhibit a harmonic

reaction to induced energy (Fig. 2A). The harmonic curve occurs whenthe excited component cannot dissipate an~more energy from the forceinducer and responds with an out-of-proportion amplitude movement.

Bonal discovered through trial and error that at and near theleading edge of the harmonic curve is the optimum frequency for usingvibrational energy for stress relief. The harmonic curve clearly es-tablishes the sub-harmonic zone. The sub-harmonic zone is identifiedas the leading lowest 1/3 portion of the harmonic curve.

FIG. 2- Applying f -h frequency in sub-harmonic zone maximizesenergy absorption for stress relief as indicated by large areawithin hysteresis curve. At harmonic peak frequency fh, internalenergy dissipation falls to zero. P

TM Bonal Technologies, Inc. Southfield, MI


Caution: The sub-harmonic processing is in sharp contrast toother vibration techniques which attempt to stress relieve "AT" thepeak. Using the peak as the stress relief energy level may damage thestructure and almost always leads to inconsistent results [3].

It was later learned through vibration handbooks that the sub-harmonic energy level is the same energy level that corresponds to thematerial's highest damping energy dissipated [4]. When a metalcomponent is subjected to a sinusoidal (vibrational) force, a hysteresisresponse occurs. The resulting hysteresis loop is the stress-strainrelationship of a part from a dynamic energy source (Fig. 2B). It willbe different than the straight line generated in the elastic zone fromthe stress-strain diagram of a static loaded part. When the eccentricweight of the vibrator, for example, reaches the top and then begins tofall back there is a split second lag in response from the part. Thislag reaches a maximum delay at and near the foot of the harmonic curve.The area within the hysteresis loop is the damping energy dissipated andreaches a maximum at and near the foot of the harmonic curve - the sub-harmonic zone.

Therefore, the metal's highest energy dissipation level is alsoits greatest stress relief potential level (Fig. 2C).

In 1987, Richard Skinner, PE, from Lockheed Missiles andAerospace, mathematically proved that when using vibrations to achievestress relief a slightly lower frequency than the harmonic peakfrequency must be used [5]. Mr. Skinner further noted that this levelwould vary slightly depending on the strength of the metal treated.

In stress relieving the workpiece using vibrations the operatormust scan the workpiece to identify the harmonic peak frequency and thusdetermine where to set the sub-harmonic stress relieving frequency. Thetime for dwell at this frequency will generally vary between 15 to 30minutes depending on the weight and material of the workpiece.

PRINCIPLE #2: SHIFTINGAll metal components have a natural harmonic peak. If, however,

the part has been subjected to a thermal shock (causing residual stress)during manufacturing the harmonic peak will be in an unnatural frequencylocation. By applying sub-harmonic vibrations the component willneutralize the thermal stress. In doing so, the harmonic peak willshift and stabilize in a new frequency location. This would be itsnatural harmonic frequency (Fig. 3).

An analogy would be like having a musical instrument out of tune(thermally stressed). As it comes in tune the true natural note isheard (stress relieved).

Bonal discovered this phenomenon immediately after Principle #1was applied. Once the harmonic curve settles into a new frequencylocation then stress relief benefits were realized on a consistentbasis.

Professors T.E. Wong and G.C. Johnson at the University ofCalifornia - Berkeley issued a report in 1987 in which theymathematically demonstrated that the harmonic curve will be influencedby the residual stress level of the part. They concluded "The findingsregarding the shift in natural frequency as a result of residual stressmay provide a method for examining the effectiveness of the stressrelieving processes [6]."

As the harmonic curve shifts, the operator will take scansperiodically. After each shift the operator will reset the stress


FIG. 3- 1st scan displays reaction of workpiece toinduced vibrational energy. 2nd scan displays a shiftindicating a change in thermal stress. 3rd scan re-peats 2nd scan verifying stress relief is complete.

relieving frequency in the new sub-harmonic range and dwell at that newfrequency for 5 to 10 additional minutes. Eventually there will be nomore shifting. One scan will repeat the previous scan. When this occursstress relief is complete for that force inducer location. If theworkpiece is large (greater than l5-feet) one or more additional forceinducer placements will be needed to assure overall stress reliefthroughout the entire workpiece.

WHEN TO APPLY SUB-HARMONIC VIBRATIONS FOR STRESS RELIEVING

To reduce the problems normally associated with manufacturing itis recommended to apply the sub-harmonic process after forging and againin between rough machining and finish machining. If the forging must beheat treated for mechanical properties enhancement, including softness,it is recommended to use sub-harmonic energy immediately before heattreating to reduce the distortion that would normally occur through theheating process. However, if the heat treatment involved a sharp tem-perature drop, like hardening, then a sub-harmonic process should alsobe applied after the heating and quenching process. For long termperformance quality assurance it is also recommended to apply the sub-harmonic process after all manufacturing steps are completed and beforecomponents.are put into service. Many times this can be accomplishedwhen the forging is assembled into its final position as the case withmost forged engines parts.

Since there are no negative side affects associated with sub-harmonic processing, even on finished parts, its application should beconsidered prior to anticipated distortion and cracking problems.

Severely mechanically stressed parts do not represent good appli-cations for this process (e.g. cold rolled, cold formed, bent, stamped,etc.) .

Caution: Sub-harmonic stress relieving, like natural ageing, doesnot significantly change the mechanical properties of the metal includ-ing softness or temper (+4%). If changes are needed in softness, tough-ness, or grain size then a specialized heat treatment is required.


SUB-HARMONIC WELD CONDITIONING I

Frequently forgings need to be welded. Welding could be requiredduring manufacturing, final installation, or repairing.

When the sub-harmonic stress relief process is applied duringwelding (hence sub-harmonic weld conditioning) generally a weld grainrefinement results. This translates into several benefits for thewelded forging:

1. less weld cracking2. less weld distortion3. less porosity4. longer fatigue life

Other observations of sub-harmonic weld conditioning includedeeper weld penetration, less undercutting, smoother weld bead, easiercontrol of the weld puddle in all positions, and less spatter.

The U.S. Department of Energy sponsored a study in 1989 investi-gating sub-harmonic processing in comparison to heat treat stress reliefand control specimens. Sub-harmonic processing was applied in both ways- as a stress relief and during welding as a weld conditioning process.The report concluded "there is evidence that the META-LAX [sub-harmonic]system is performing comparable to the thermal stress relief process onA36 carbon steel [7]." See summary chart as published in DOE's own TechBrief about their findings [8, Table 1].


FIG. 5-- Sub-Harmonic Weld Conditioned Weld MetalPhotomicrograph. Note a finer weld grain.

2. Less distortion. The rate of solidification is slowed downallowing the weld metal to solidify more evenly from the rootto the face. This translates into less distortion and lessresidual stress in the HAZ. Therefore the HAZ will be crackresistant compared to a non-weld-conditioned HAZ.

3. Stress relief of base material. The sub-harmonic weldconditioning process stress relieves the base material whilethe welding is going on. This enables the base material to becrack resistant.

In applying sub-harmonic vibrational energy during welding threesteps are recommended:

1. Pre-weld Stress Relief for about 1/2-2/3 the time it wouldnormally take for a stress relief;

2. Weld Condition - reset frequency to foot of harmonic curveuntil welding is complete; and

3. Post Weld Cooldown until the last area welded is less than200-F.

FIELD STUDIESTEST CASE #1: CONTROLLING DISTORTION FOLLOWING MACHININGForged 413S cannon barrels, G-inch OD x 4-3/4-inch ID x GS-inch,tolerance is O.OOG-inch total indicator runout (TIR) with O.OOOG-inchconcentricity, 230 lbs. each (Fig. G).

Set #1 (30 barrels) processing:forged,heat treated,heat treat stress relieved,rough machined,heat treat stres3 relieved,finish machined,measured.

RESULTS: Fifteen barrels experienced at least O.OGO-inchdistortion during rough turning which split the cutter, jammed upand bent the barrel. This scrapped all lS barrels (SO% of 30).The other lS barrels were able to be processed as expected andmeasured after finish machining between O.OOS-.OlO-inch TIR.


FIG. 6-- Sub-harmonic stress relief was added before roughturning these cannon barrels. Final distortion averaged 50%less, a 50% scrap rate was eliminated, and the second heattreat stress relief was eliminated which saved $230,000/yr.

Set #2 (30 barrels) processing:

forged,heat treated,heat treat stress relieved,

ADD SUB-HARMONIC STRESS RELIEVED,rough machined,

DELETE heat treat stress relief,finish machined,measured.

RESULTS: Distortion level after finish machining were 0.002-.004-inch TIR on all 30 sub-harmonic treated barrels for an average 50%less distortion and NO scrap. Also time and money were savedsince the second heat treat stress relief was eliminated.

Set #3 (270 barrels) processing: Same as Set #2.

RESULTS: Same as Set #2. Further studies estimated that "fuelcost savings· associated with eliminating the second heat treatstress relief were $230,000 per year on the end product.

156 SiEEL FORGINGS: SECOND VOLUME

TEST CASE #2: REDUCING PREMATURE CRACKING

Forged Crank shafts for boat V-8 race engines (Fig. 7).

Set #1. Normal Performance: 10 race average to fatigue failure.

Set #2. Added sub-harmonic stress relief before finish machining andagain on finished crank.

RESULTS: 30 race average to fatigue failure for over a 200%fatigue life improvement.

FIG. 7-- High performance crankshafts see triple life(200% improvement) after being sub-harmonic treated.

TEST CASE #3: REDUCING DISTORTION POLLOWING HEAT TREATING

Forged Cams for V-6 race engines (Fig. 8).

Set #1. Normal distortion: 0.020-.030-inch after heat treating.

Set#2. Added sub-harmonic energy immediately before heat treating.

RESULTS: 0.002-inch distortion after heat treating for 90-93%less.


TEST CASE #4: REDUCXNG DXSTORTXON POLLOWXNG CARBURXZXNG AND HARDENXNGForged Gears, 8620, l-inch x 7-inch Dia. (Fig. 9).

Set #1. Normal Distortion: 0.015-inch after heat treating.

Set #2. Added sub-harmonic energy immediately before heat treating.

RESULTS: 0.003-inch distortion after heat treating or 80% less.

FIG.9-- Gears were sub-harmonic stress relievedbefore carburizing and hardening for 80% lessdistortion.

TEST CASE #5: REDUCXNG PREMATURE CRACKXNGForged Aluminum Connecting Rods for drag race engine (Fig. 10).

Set #1. Normal Performance: 70 passes. After 70 passes failure wasexpected at any time.

Set #2. Added sub-harmonic stress relieve (8) con-rods before initialinstallation.

RESULTS: Connecting rods were replaced after 354 passes without afailure for 405% fatigue life improvement.


TEST CASE #6: REDUCING DISTORTION POLLOWING MACHININGForged 7075 Aluminum Longeron, 10-inch x 18-inch x 18-feet, 1625 Ibs.(Fig. 11).

FIG. 11-- Sub-harmonic stress relief was appliedafter forging this longeron resulting in anaverage of 55% less distortion.

Set #1. Normal Distortion: 0.150-.375-inch TIR, 0.200-inch TIR ave.

Set #2. Added sub-harmonic stress relief before rough machining.

RESULTS: 0.090-inch TIR average for 43-76% less distortion,average 55% less.

TEST CASE #7: IMPROVE PATIGUE LIPEForged 4130 Mud Pump Fluid End, 18 x 18 x 36-inch

Set #1. Normal Performance (new pump): 1000-3400 running hours.

Set #2. Added sub-harmonic stress relief during welding of cracks.

RESULTS: 7000 running hours before replacement of ~ parts,106% minimum fatigue life improvement without any cracking.

CONCLUSIONSForgings are subjected to a thermal shock which in turn causes the

workpiece to be· susceptible to having distortion and cracking problems.By using sub-harmonic vibrations before machining, before heat treating,and as a quality assurance practice on finished parts, forgings havebeen shown to have up to 90% less distortion in manufacturing and lastup to 400% longer in service. Furthermore, by applying sub-harmonicenergy during welding the quality of the weld metal is improved and theproblems associated with welding are reduced.

REPERENCES[1] Van Vlack, Lawrence, Elements of Materials Science, 2nd Edition,

Addison-Wesley, 1967, p 407.

[£] Connor, Leonard P., ·Welding Technology," Weldinq Handbook, 8thEdition, AWS, 1989, p 219.


[3] Cheever, D. L. and Rowlands, E. W., "Vibrational Conditioning of- Castings and Weldments: An Exploratory Study," Control of

Distortion and Residual Stress in Weldments, ASM, 1977, p 22.

[4] Harris, Cyril and Crede, Charles, "Engineering Designs and- Environmental Conditions," Shock & Vibration Handbook," VOL 3,

McGraw-Hill, 1961, p 36-2.

[5] Skinner, Richard, PE, "An Investigation into the Theory Behind- Sub-Resonant Stress Relieve," 1987, p 10.

[6] Wong, T. E. and Johnson, G. C., "Ultrasonic Evaluation of the- Nonlinearity of Metals from a Design Perspective," 1987, p 15.

[7] Holdren, Richard, PE, "The Meta-Lax Method of Stress Reduction in- Welds," U.S. DOE DE-FGOl-89CE15412, 1990, p 33.

[~] Wilkinson, David, "Meta-Lax Stress Relief Process," Dept. ofEnerqy Tech Brief, SEMA, 1995.

Michael 1.Leapl, James C. Wingertl, and Charles A Mozdenl

DEVELOPMENT OF A PROCESS FOR TOUGHENING GRAIN-REFINED,HIGH-STRENGTH STEELS

REFERENCE: Leap, M. 1.,Wingert, 1.C., and Mozden, C. A, "Development of aProcess for Toughening Grain-Refined, High-Strength Steels," Steel Forgings:Second Volume, ASTM STP 1259. E. G. Nisbett and A S. Melilli, Eds., AmericanSociety for Testing and Materials, 1997.

ABSTRACT: There is a gradually increasing awareness of the deleterious effects ofgrain-refining precipitates on the Charpy V-notch impact toughness of high-strengthsteels, and recognition of these effects resulted in the development of a general processthat provides substantial improvements in the ductile fracture resistance of tempered mar-tensite. The results of this investigation indicate that the processing-induced refinement ofgrain-refining precipitates also provides significant improvements in the lower-shelf andtransition toughness of331O, 8219 and 4340 steels tempered at 180°C and 4340 steeltempered at 593°C. In contrast to the upper-shelf regime, where the application of theprocess produces a change in either the mode or character of unstable fracture at constantstrength and austenite grain size, the refinement of grain-refining precipitates is associatedwith an increased amount of ductile fracture and an increased resistance to the initiation ofunstable crack propagation at any given temperature in the lower-shelf regime.

KEYWORDS: High-strength steels, grain-refining precipitates, ductile fracture, brittlefracture, impact toughness, instrumented impact testing

The presence of various types of second-phase particles in tempered martensitemicrostructures is known to degrade the toughness of high-strength steels. Fracture atupper-shelf temperatures is generally dependent on a mechanism in which the initiationand growth of primary microvoids occurs at non-metallic inclusions [1] and other largesecond-phase particles [1-4]. Based on the application of the characteristic distance con-cept to the conditions associated with strain-controlled fracture [5,6], Garrison [7] has

IPrincipai Research Engineer and Research Analysts, respectively, Materials ScienceDepartment, The Timken Company, 1835 Dueber Avenue, S.w., Canton, OR 44706.

160

LEAP ET AL. ON A PROCESS 161

interpreted the Rice-Johnson model [8] for ductile fracture in terms of the interparticlespacing of void-initiating particles in the microstructure:

Grain-refining precipitates (e.g., aluminum nitride and microalloy carbonitrides) havebeen more recently identified as microstructural features that can significantly degrade theductile fracture resistance of grain-refined, high-strength steels [15,16], and a process hasbeen developed [17] to maximize the toughness of high-strength steels via the refinementof grain-refining precipitates. The process consists of four basic steps, Figure 1: (i) re-heating and hot deformation at high temperatures (i.e., preferably above the solution tem-perature of the least soluble nitride or carbonitride species present in the steel), followedby accelerated cooling to 500°C; (ii) subcritical annealing to promote the precipitation ofthe maximum volume fraction of aluminum nitride (AIN) and/or carbon-rich microalloycarbonitrides in ferrite; (iii) final austenitization at commercially practical hardening tem-peratures, followed by quenching, and; (iv) tempering to the desired strength level. Thisprocess has a great deal of applicability to the manufacture of forged high-strength steelcomponents since the solution pretreatment can be directly integrated into the reheatingand forging operations. A high-temperature pretreatment is attractive from the standpointof maintaining forgeability and die life, although the ultimate objective of reheating andhot deformation at high temperatures is to minimize the content of coarse precipitatesretained through the solidification and initial processing of a steel. Subsequent to forging,however, cooling must be conducted at an accelerated rate in order to limit the amount ofreprecipitation that can occur in austenite. The combination of high-temperature forgingfollowedby accelerated cooling has been most widely applied to vanadium-modified, med-ium-carbon steels [18], and as a result of the widespread adoption of these steels in a vari-ety oflow-to-intermediate strength applications, an increasing number offorgers have in-stalledconveyor lines with forced-air cooling capabilities. This general type of technologyin conjunction with subcritical annealing and austenitization at commercially practical tem-


FIG. l--Schematic illustration of the process for improving the toughness of grain-refined,high-strength steels.

peratures comprises the process for improving the toughness of grain-refined, high-strength steels.

The process has been shown to provide substantial improvements in the room-tem-perature impact toughness of high-strength steels [16] containing 0.1-0.4% C, 0.023-0.047% AI, and nitrogen in concentrations representative ofEAF steelmaking practices(66-105 ppm), Figure 2. In making this comparison between the conventionally-process-ed and pretreated/annealed conditions for each of a variety of alloy steels, the strength isconstant within about 40 MPa, the macroscopic mode of fracture is equivalent, and bothmaterial conditions exhibit fine-grained austenite microstructures. The process providesimprovements in impact toughness over a broad range of strength [15,16], although theeffectiveness ofthe process is expected to diminish with prolonged tempering at high tem-peratures. Preliminary-results [15,16] also spggest that the process improves both the lon-gitudinal and transverse impact toughness of high-strength steels over a relatively broadrange of sulfur content, but very little systematic work has been conducted to determinethe compositional limits associated with improved transverse toughness. Moreover, it isfairly well established that the brittle fracture resistance of ferritic materials is degraded bythe presence of comparatively large second-phase particles such as MnS [19], TiN [20-22], and grain-boundary cementite [23,24]; however, the extent to which smaller grain-refining precipitates will affect 10wer-shelf7transition toughness and the extent to which therefinement of these precipitates will improve low-temperature toughness is largely un-known [25].


ROOM-TEMPERATURE IMPACT TOUGHNESS (J)

(CONVENTIONALL Y-PROCESSED STEELS)

FIG. 2--The room-temperature impact toughness of pretreated/annealed steels is shown inrelation to the impact toughness of the corresponding conventionally-processed steels.Error bars represent estimates of the 95% confidence limits for the mean toughness associ-ated with each steel and material condition.

The purpose of this investigation is to evaluate the level of improvement in the low-temperature impact toughness of grain-refined, high-strength steels that results from theapplication of the process for refining grain-refining precipitates. The steels comprisingthis study include a 3310 steel (0.024% Al-85 ppm N-0.009% S), an 8219 steel (0.023%Al-105 ppm N-0.014% S), and a 4340 steel (0.006% Ti-0.026% Al-70 ppm N-0.002% S).The 3310 and 8219 steels are evaluated in the lightly-tempered (180°C) condition and the4340 steel is evaluated after tempering at 180°C and 593°C.

EXPE~ENTALPROCEDUREThe compositions of the steels utilized in this investigation are listed in Table 1. The

steelswere obtained in the form of 150 rom x 150 rom wrought billets. Sections of the


billets were reheated in the 1230-1260°C range for 3-4 hours, forged to cross-sectionaldimensions of70 mm x 140 mm, and air cooled to room temperature. After sectioningand milling to cross-sectional dimensions of60 mm x 130 mm, the forged sections werereheated at 1250°C for three hours and hot rolled to 16 mm plate in five passes. Thereduction per pass ranged from 17% to 23%, and the last pass was completed at temper-atures in the vicinity of lOOO°C. Hot-rolled plates were either air cooled to room temper-ature or oil quenched directly off the rolling mill. The direct-quenched plates were subse-quently annealed at 700°C for 1.5-2.0 hours and air cooled to room temperature. Theprocessing associated with the air cooled plates is hereafter referred to as conventionalprocessing, whereas the processing of the direct-quenched and subcritically annealed steelswill be referred to as the pretreated/annealed condition.

TABLE I-Steel Compositions (weight percentages)

Steel C Mn Si Cr Ni Mo S P AI Ti N (ppm)

3310 0.11 0.55 0.30 1.55 3.39 0.04 0.009 O.OlO 0.024 0.003 85

8219 0.20 1.34 0.24 0.58 0.31 0.18 0.014 0.009 0.023 0:002 105

4340 OAO 0.74 0.30 0.79 1.70 0.27 0.002 0.006 0.026 0.006 70

Oversized test specimen blanks were extracted from the mid-plane of the hot-rolledplates in both the 10ngitudinaI/LT and transverse/TL orientations. The specimen blankswere austenitized at 900°C for one hour, quenched, and tempered at 180°C for one hour.Specimens of the hardened 4340 steel were also tempered at 593°C for one hour.

Mechanical Testinl!

The room-temperature tensile properties of the steels were determined from speci-mens with a 9 mm diameter and 36 mm gage length in accordance with ASTM StandardTest Methods for Tension Testing of Metallic Materials (E 8). Standard Charpy V-notchimpact tests were conducted at temperatures between -100°C and 120°C in accordancewith ASTM Standard Test Methods for Notched Bar Impact Testing of Metallic Materials(E 23).

Instrumented impact testing was conducted on longitudinally-oriented, Charpy V-notch specimens at -100°C and -196°C. Charpy specimens were fractured in a Tinius-Olsen Low-Blow Impact Machine containing an instrumented striker. The electrical signal


from the striker was amplified with a Dynatup Model 500 Dynamic Response Module andstored on a digital oscilloscope. The tests were conducted in accordance with the generalguidelines proposed by various investigators [26-29], and following the recommendationsof Hoover [30], signal filtering was not employed in the acquisition of the load-time (P-t)data. A pendulum rotation angle of32.5°, corresponding to an initial impact energy, Eo,of37.4 J and an initial impact velocity, Vo, of 1.65 mis, was utilized in the testing of speci-mens at -196°C. Depending on material toughness, pendulum angles between 32.5° and70° (Eo = 157 J and Vo = 3.38 mls) were required to fracture specimens at -lOO°C. Fin-ally, analysis of the instrumented impact test data was conducted in the manner describedby Leap [31].

Metallol!raohv

The austenite grain structures of the steels were qualitatively evaluated via lightmicroscopy. Sections of Charpy specimens were tempered at -480°C for 24 hours, pre-pared for examination, and etched in a saturated picric acid solution containing sodiumtridecylbenzene sulfonate as a wetting agent.

Fractography was conducted on Charpy V-notch specimens using a Cambridge In-struments Stereoscan 250 scanning electron microscope operating at 20 kV. A represen-tative indication of the micromechanisms controlling fracture was obtained by viewing theCharpy specimen with an impact toughness closest to the mean toughness for each steeVprocessing condition.

RESULTSThe tensile properties of the steels are summarized in Table 2. The conventionally-

processed and pretreated/annealed conditions for each steel exhibit similar levels of yieldstrength, tensile strength, and total elongation. The mean values of tensile reduction inarea for the pretreated/annealed steels tend to be slightly greater than the correspondingvalues for the conventionally-processed steels, but in an overall sense, the tensile ductilityof the two material conditions is similar in magnitude for each steel. The steels all exhibitaustenite grain structures that can be generally classified as fine grained, although the con-ventionally-processed steels tend to exhibit more frequent occurrences of isolated, slightlylarger grains in the microstructure [15], Figure 3. The minor differences in grain structureare stochastically consistent with the processing-induced refinement of grain-refining pre-cipitates in the pretreated/annealed steels and the relative contents of grain-refining precip-itates predicted [32] to be present in the three steels during final austenitization. Never-theless, comparisons of differences in the toughness of the conventionally-processed andpretreated/annealed material conditions are, for all practical purposes, made at similarlevelsof strength and tensile ductility in steels possessing fine-grained austenite micro-structures.


TABLE 2-Summary of Room-Temperature Tensile Testand Impact Toughness Datal

SteeV SpecimenCondition2 Orientation cr.(MPa) crT(MPa) A (%) 'P (%) E (J)3 He

3310/CP LILT 1050 1290 16.3 63.7 89.9 ± 9.5 39T/TL 1060 1280 13.0 47.2 32.0 ± 2.8 39

3310/AP LILT 1060 1290 16.3 66.9 121.0 ± 2.0 39TITL 1060 1300 13.4 49.9 39.9 ± 1.8 39

8219/CP LILT 1250 1520 14.0 55.5 57.5 ± 3.9 45TITL 1260 1570 11.1 35.7 18.2 ± 1.2 45

8219/AP LILT 1250 1520 15.4 58.1 80.0 ± 5.8 45TITL 1250 1520 11.1 34.9 19.5 ± 1.5 45

4340/CP LILT 1540 2070 12.3 39.8 15.7 ± 1.5 54TITL 1570 2070 10.3 32.0 14.6 ± 0.8 53

4340/AP LILT 1560 2060 12.5 45.0 28.2 ± 0.9 54TITL 1550 2070 11.5 37.9 26.0 ± 1.5 54

4340/CP LILT 1050 1140 17.0 57.5 44.9 ± 4.9 37(TT= 593°C) TITL 1060 1140 15.1 48.9 46.2 ± 1.6 37

4340/ AP LILT 1010 1110 19.1 60.3 76.6 ± 1.4 37(TT= 593°C) T/TL 1030 1120 16.1 52.3 70.8 ± 2.2 36

IAll specimens were austenitized at 900°C for one hour, quenched, and tempered at 180°Cfor one hour unless indicated otherwise.

2CP = conventionally-processed steel; AP = pretreated/annealed steel.3Mean values and standard deviations.

Nomenclature

cro= 0.2% offset yield strength crT= tensile strengthA = elongation in 36 mm 'I' = reduction in areaE = room-temperature impact toughness TT= tempering temperature


FIG. 3--Austenite grain structures exhibited by the (a) conventionally-processed and(b) pretreated/annealed 8219 steel after final austenitization at 900°C for one hour.

Impact transition-temperature curves for the steels are shown in Figures 4-7. Theapplication of a solution pretreatment and subcritical anneal prior to final austenitizationproduces significant improvements in longitudinal impact toughness over the entire rangeoftest temperature examined in the 3310,8219 and 4340 steels tempered at 180°C andthe 4340 steel tempered at 593°C. The application of the process also improves the trans-verse impact toughness of the low-sulfur 4340 steel at both strength levels, Figures 6b andTh. However, the 3310 steel (0.009% S) exhibits much more moderate increases in trans-verse impact toughness (i.e., the relative improvement in toughness varies from ~14% atupper-shelf temperatures to ~28% as the test temperature approaches -100°C), Figure 4b,and the extentoofprocessing-induced improvements in the transverse toughness of the8219 steel (0.014% S) is rather limited over the entire range oftest temperature, Figure5b.

The results of instrumented impact tests at -196°C are summarized in Table 3 forlongitudinally-oriented, Charpy V-notch specimens. Impact testing at -196°C revealedtwo types ofload-time (P-t) behavior in the steels, Figure 8. A type I P-t response is char-acteristic of linear-elastic behavior in which macroscopic fracture occurs in a predominant-lybrittle manner. This type of behavior is exhibited by all the conventionally-processedsteels and the pretreated/annealed 4340 steel after tempering at 180°C. Crack initiation inthese specimens occurs via the formation of fine-scale microvoids in the region immediate-


FIG. 4--Impact transition-temperature curves for conventionally-processed and pre-treated/annealed specimens of the 3310 steel: (a) LT specimen orientation and (b) TLspecimen orientation. All specimens were austenitized at 900°C for one hour, waterquenched, and tempered at 180°C for one hour.


FIG. 5--Impact transition-temperature curves for conventionally-processed and pre-treated/annealed specimens of the 8219 steel: (a) LT specimen orientation and (b) TLspecimen orientation. All specimens were austenitized at 900°C for one hour, waterquenched, and tempered at 180°C for one hour.


FIG. 6--Impact transition-temperature curves for conventionally-processed and pre-treated/annealed specimens of the 4340 steel: (a) LT specimen orientation and (b) TLspecimen orientation. All specimens were austenitized at 900°C for one hour, oilquenched, and tempered at 180°C for one hour.


FIG. 7--Impact transition-temperature curves for conventionally-processed and pre-treated/annealed specimens of the 4340 steel: (a) LT specimen orientation and (b) TLspecimen orientation. All specimens were austenitized at 900°C for one hour, oilquenched, and tempered at 593°C for one hour.



ly below the notch root, Figure 9a. This ductile shear zone extends along the entire lengthof the notch root, although the amount of ductile crack extension below the notch root isextremely limited (50-150 J.lIll). Subsequent unstable fracture through the ligament occursin a brittle manner, Figure 9c.

The pretreated/annealed 3310 and 8219 steels as well as the 4340 steel tempered at593°C exhibit both type I and type IT P-t behavior at -196°C. In comparison to the linear-elastic behavior, the occurrence of a type IT P-t response is associated with the formationof a larger ductile shear zone below the notch root in the 3310 steel and small shear lipsthat emanate a short distance along the ligament in both the 3310 and 8219 steels, Figure9b. However, fracture in the pretreated/annealed steels primarily occurs by the more brit-tle type of behavior, and the values ofEl associated with the type IT P-t curve appear to besimilar in magnitude to the upper end of the range ofEl for specimens exhibiting type I

FIG. 9--The morphology of fracture at -196°C is shown for the (a,c) conventionally-processed and (b,d) pretreated/annealed 8219 steel: (a,b) notch-root region includingthe ductile shear zone and (c,d) central section of the specimens. SEM fractographs oflongitudinally-oriented, Charpy V-notch specimens.

FIG. lO--Normal distribution representations of the nominal fracture strength, obtainedfrom instrumented impact tests at -196°C, are shown for conventionally-processed andpretreated/annealed specimens of the (a) 3310, (b) 8219, and (c) 4340 steels. The fracturestrength values are normalized by the mean fracture strength for the conventionally-pro-cessed steel. All specimens were austenitized at 900°C for one hour, quenched, and tem-pered at 180°C for one hour.


behavior, Table 3. These data suggest that the energy to initiate a ductile tear in the notchroot (i.e., an energy comprising a portion ofEl) is relatively constant, and any microstruc-tural differences that allow additional ductile separation are manifested as the expenditureof energy E2• Fracture in the conventionally-processed 4340 steel occurs much in thesame manner as in the pretreated/annealed 3310 and 8219 steels, but fracture in the pre-treated/annealed 4340 steel is characterized by the formation of a somewhat larger shearzone and shear lips that extend over a significant portion of the ligament. These observa-tions are consistent with the relative magnitudes ofE:z/El in specimens exhibiting type IIbehavior; that is, the specimens with smaller ductile shear zones and limited shear lip for-mation exhibit comparatively low values ofE:z/El (0.20-0.39), whereas larger amounts ofmacroscopic plasticity in the pretreated/annealed 4340 specimens are associated withhigher ratios ofE2 to El (0.50 ~ E:z/El~ 1.08). Regardless of the extent of ductile crackextension and the relative size of shear lips on the specimens, unstable fracture ultimatelyoccurs in a predominantly brittle manner at -196°C, Figures 9c and 9d.

Notwithstanding specific details regarding the dependence of fracture strength onmicrostructure (i.e., the evaluation of fracture strength in terms of weak-link and charac-teristic distance concepts [23,24] when brittle fracture occurs below the root of a bluntnotch) or the presence of a small ductile shear zone below the notch root, a nominal frac-ture strength can be calculated from the maximum load for specimens exhibiting type I P-tbehavior. Normal distribution representations of the nominal fracture strength are shownin Figure 10 for the 3310,8219 and 4340 steels tempered at 180°C, where the values ofnominal fracture strength are normalized to the mean value of the fracture strength for theconventionally-processed steel. The processing-induced increase in the mean value ofnominal fracture strength is statistically significant at the 99% confidence level for eachsteel.

The results of instrumented impact tests at -100°C are summarized in Table 4 forlongitudinally-oriented, Charpy V-notch specimens of the steels. The application of asolution pretreatment and subcritical anneal prior to final austenitization produces a transi-tion from type II to type III P-t behavior in the 3310 steel, Figure 8. The value ofEl issimilar in magnitude for the two material conditions, although the magnitude ofE2 for thepretreated/annealed steel is significantly greater than the corresponding value for the con-ventionally-processed steel. This large difference in E2 is related to the differences in theamount of ductile tearing below the notch root and the size of the shear lips on the Charpyspecimens. In particular, both the conventionally-processed and pretreated/annealed steelsexhibit shear zone formation below the notch root, Figures Ila and lIb, but after the for-mation of the shear zone, unstable crack propagation in the conventionally-processed steeloccurs by cleavage intermixed with small amounts of ductile rupture, Figures IIc and lIe.Shear zone formation in the pretreated/annealed steel is followed by tensile separation in aductile manner, Figure lId, and subsequent unstable fracture occurs by a mixed-modemechanism of cleavage and ductile rupture, Figure 11f The ductile component of un-stable fracture in the pretreated/annealed specimens is comprised oflarger (primary)microvoids located at large non-metallic inclusions and smaller (secondary) microvoidsthat form at smaller inclusions and other second-phase particles in the microstructure.


FIG. 11--The morphology of fracture at -100°C is shown for the (a,c,e) conventionally-processed and (b,d,t) pretreated/annealed 3310 steel: (a,b) notch-root region includingthe ductile shear zone, (c,d) notch-root region below the ductile shear zone, and (e,t)central section of the specimens. SEM fractographs of longitudinally-oriented, CharpyV-notch specimens.


Finally, the P-t curves for the pretreated/annealed steel exhibit a post-unstable fractureenergy, E3, that reflects the formation of a ductile compression face on the Charpy spec-imens. This energy accounts for between 8% and 23% of the total energy to fracture thepretreated/annealed 3310 specimens at -100°C.

The fracture behavior of the 8219 steel at -100°C is similar to that of the 3310 steel,although the P-t curves for pretreated/annealed specimens of the 8219 steel do not exhibitan energy following the unstable (maximum) fracture load. Once again, the values ofEIare similar in magnitude for the two material conditions, the primary difference in tough-ness resulting from the different amounts of energy expended in the propagation of theductile shear zone below the notch root and the formation of shear lips on the Charpyspecimens. Unlike the 3310 steel, unstable fracture at -100°C primarily occurs by cleav-age in both the conventionally-processed and pretreated/annealed 8219 specimens.

Fracture in the lightly-tempered 4340 steel initiates by microvoid coalescence in theshear zone below the notch root. Unstable crack propagation from the end of the shearzone occurs by cleavage/quasicleavage in the conventionally-processed steel and by a mix-ture of cleavage, quasicleavage and fine-scale ductile rupture in the pretreated/annealedsteel. Similar to the transition from type I to type IT P-t behavior at -196°C, the value ofEI associated with type II behavior at -100°C appears to coincide with the upper end ofthe range ofEI values for the conventionally-processed specimens exhibiting type I P-tbehavior at -100°C. The energy E2 reflects the additional ductile tearing that occurs be-low the notch root, the formation of shear lips on the Charpy specimens, and the morpho-logicalchanges in the unstable fracture region of the specimens.

The application of the process is associated with a transition from type II to type IVP-t behavior in 4340 specimens tempered at 593°C. Crack initiation at -100°C occurs in aductilemanner for both material conditions, although the pretreated/annealed steel alsoexhibitsmore ductile crack extension, significantly larger shear lips, and the formation of alarge compression face on the Charpy specimens. Unstable fracture primarily occurs bycleavage in both material conditions, but there is more microplasticity in the form of tearridges on the pretreated/annealed specimens. All of these factors combine to yield thelargedifference in the impact toughness ofthe conventionally-processed and pretreated/annealed4340 steel at -100°C.

Fracture in the upper-shelf regime primarily occurs through changes in the mode ofunstable fracture in the 8219 steel, consistent with previous observations [15-17] of otherlightly-tempered, low-alloy steels exhibiting upper-shelf behavior at room temperature.The application of a solution pretreatment and subcritical anneal prior to final austenitiza-tion specifically changes the unstable fracture morphology from a mixture of quasicleavageand ductile rupture with isolated occurrences of transgranular cleavage to a predominantlyductilemode of separation, Figure 12. The fracture surfaces of the pretreated/annealedspecimensare comprised of three types of microvoids: (i) primary microvoids that nuc-leate at large MnS inclusions, (ii) a comparatively low density of intermediate-sized micro-voids that nucleate at smaller MnS inclusions and any larger, residual second-phase par-


FIG. 12--The morphology of unstable fracture at 24°C is shown for the (a) convention-ally-processed and (b) pretreated/annealed 8219 steel. SEM fractographs oflongitudin-ally-oriented, Charpy V-notch specimens.

ticles in the microstructure (e.g., TiN and iron/alloy carbides), and (iii) extremely finemicrovoids that form at smaller second-phase particles in the microstructure during thefinal stages of separation. The replacement of quasicleavage facets with fine-scale micro-voids is primarily responsible for the 17 J increase in the longitudinal, room-temperatureimpact toughness of the 8219 steel.

The lightly-tempered 3310 and 4340 steels fracture in a completely ductile mannerat upper-shelf temperatures, irrespective of prior processing history. However, there areprocessing-induced changes in the character of the ductile fracture in both the notch rootand unstable fraCture regions of the Charpy V-notch specimens. The conventionally-pro-cessed steels exhibit primary microvoids at non-metallic inclusions and a moderate densityof small-to-intermediate sized (secondary) microvoids that are connected by smaller voidsin the intervening material, Figures 13a and 13b. Fracture in the pretreated/annealed steelsalso occurs through the formation of primary microvoids at non-metallic inclusions andintermediate-sized microvoids at larger second-phase particles in the microstructure, butthe final separation of the intervening material between the secondary microvoids occursby the formation of networks of extremely fine microvoids, Figures 13c and 13d. Fracturein these regions presumably reflects the presence of grain-refining precipitates and residualalloy carbides that are not large enough to initiate quasicleavage or allow a significant


FIG. 13--The morphology of unstable fracture at upper-shelf temperatures is shown forthe lightly-tempered (a,c) 3310 (24°C) and (b,d) 4340 (l00°C) steels: (a,b) convention-ally-processed and (c,d) pretreated/annealed specimens. SEM fractographs oflongitudin-ally-oriented, Charpy V-notch specimens.

amount of growth in the high density of microvoids that nucleate immediately prior to finalseparation at extremely high levels of strain-energy density. Finally, the changes in theductile fracture morphology of 4340 specimens tested at lOO°C,which are basically equi-valent in the notch root and unstable fracture regions of the Charpy specimens, are some-what different from the near upper-shelf behavior exhibited by 4340 specimens tested atroom temperature [15,16]. In this latter case, the refinement of grain-refining precipitatesproduces a change in the mode of unstable propagation from predominantly quasicleav-age to ductile rupture in the central section of the Charpy specimens.

Upper-shelf fracture in 4340 specimens tempered at 593°C is also associated withfullyductile behavior for both material conditions. The appearance of the fracture belowthe notch root of the specimens is similar in both the conventionally-processed and pre-


treated/annealed steels, although the character of the fracture in the central section of thespecimens is different for the two material conditions. The conventionally-processed steelexhibits primary microvoids at non-metallic inclusions and fine-scale ductile rupture in theintervening matrix material, Figure 14a. The pretreated/annealed steel also exhibits pri-mary microvoids, but separation of the material between these features occurs by the for-mation of intermediate-sized microvoids interconnected by extremely small microvoids,Figure 14b. The smallest microvoids are primarily present in the form of void sheets, butthere are regions in which the microvoids are widely spaced (large void spacing relative tothe size of the microvoids) and connected by ductile tears that have the appearance of ex-tremely shallow voids (Le., the fracture surface is rumpled and exhibits many small cusps).

DISCUSSION

The most commonly applied model for brittle fracture [23,24,33,34] is based on thepremise that unstable fracture occurs when the maximum stress below the notch root ex-ceeds the fracture strength of the material over a characteristic microstructural distance.For the case oflow-strength ferritic steels, it is implicitly assumed that crack initiationoccurs at grain-boundary carbides while the critical propagation event is defined by theextension of the microcrack across a grain and through the adjacent grain boundary. In

FIG. 14--The morphology of unstable fracture at 24°C is shown for the (a) convention-ally-processed and (b) pretreated/annealed 4340 steel tempered at 593°C. SEM fracto-graphs of longitudinally-oriented, Charpy V-notch specimens.


accord with the general deformation and fracture behavior of notched bend specimens[35-37], Yan et al. [38] suggested that unstable fracture requires the attainment of both acritical stress and critical plastic strain. Considering the distributions of notch stress [23,39,40] and notch strain [40-42] in three-point bend specimens, this model for brittle frac-ture defines a range of position below the notch root in which unstable fracture is likely toinitiate by a weak-link mechanism from a single defect in the material. However, thesemodels do not account for the effects of the size and density of the relevant microstruc-tural feature(s) on the resistance to brittle fracture.

The stress-controlled initiation of unstable fracture over a range of position below anotch root or crack tip can also be described solely in terms of the statistical occurrence ofa sufficiently large microstructural feature located at an appropriate position in the notch-root/crack-tip plastic zone [43]. Lin, Evans, and Ritchie [44,45] modeled the tempera-ture-dependent roles of grain size and particle size on the brittle fracture behavior oflow-strength ferritic steels with a function of the form:


microstructural feature1 in a single test specimen. Unstable fracture can also initiate frommultiple sites as ",No increases beyond INp, but the initiation of microcracks at an increas-ing density of sites corresponds, in the limit, to the transition from a weak-link to a criticaldamage mechanism for brittle fracture with decreases in test temperature [47,48].

The second method of toughening corresponds to a critical damage mechanism forbrittle fracture, and this method provides the basis for the processing-induced improve-ments in the low-temperature toughness of the steels evaluated in the present investiga-tion. Specifically, crack initiation at -196°C occurs by the formation of a ductile shearzone, and subsequent unstable fracture initiates uniformly along the leading edge of theshear zone, Figure 9. Both the ductile and brittle modes of separation are consistent witha critical damage type of fracture. The presence ofa ductile shear zone in the Charpyspecimens also indicates that the initiation of transgranular cleavage requires a substantialincrease in stress intensification over that provided by the notch, and this increase is real-ized through the effective increases in geometric acuity and crack length associated withthe formation of the ductile shear zone. Although this mode of separation is characteristicof ductile crack initiation [49,50] along the deformation slip lines that emanate from thenotch root ofa Charpy V-notch specimen [51-53], the presence ofa ductile shear zone isnotable from the standpoint that the yield strength in shear for each material must be lowerthan the corresponding fracture strength for ductile crack initiation to be possible at-196°C .

.The presence of a ductile shear zone and the initiation of unstable fracture in terms ofa critical damage mechanism provide a unique opportunity to examine the effects of pro-cessing on the resistance to both ductile and brittle crack propagation at low test tempera-tures. For specimens exhibiting linear-elastic (type I) P-t behavior at -196°C, Table 3, theimpact toughness provides a rough measure of the energy required to initiate and propa-gate the shear crack below the notch root. However, the processing-induced differencesin impact energy correlate with differences in the size of the ductile shear zone, Figures 9aand 9b, such that the effects of grain-refining precipitates on the ductile propagation resis-tance of the steels cannot be isolated2. Considering the potential for blunting at second-phase particles that are larger than a critical size [21,22,45], larger particles in the micro-structure (e.g., MnS inclusions, oxide inclusions, and TiN precipitates) may provide pref-erential initiation sites for the microvoids that constitute the ductile shear zone. If thisspeculation is correct, then normalizing the impact energy values for each steel by the sizeof the shear zone would probably yield values of ductile propagation resistance that aresimilar in magnitude for the conventionally-processed and pretreated/annealed conditions.

lIn this context, the definition for a single defect could very well encompass anagglomeration of stable microcracks that combine to effectively form a single microcrack.

2QuantifYingdifferences in the resistance to ductile crack propagation would requirereliable measurements of crack initiation load.


Moreover, since these larger particles are basically unaffected by processing, the resistanceto ductile shear propagation at low temperatures should be somewhat independent of theextent of AlN refinement in the steels.

In comparing the two material conditions, the differences in E1 represent differencesin the amount of ductile crack extension below the notch root, whereas the fracture loadprovides a quantitative indication of the stress required for unstable propagation once theshear crack has grown to a critical size, Figure 10. However, there are two potentialdrawbacks to representing this stress as the fracture strength of the material: (i) the initia-tion of unstable fracture follows subcritical ductile crack propagation and (ii) the actualcrack length (i.e., depth of the notch and ductile shear zone) is not utilized in the calcula-tion of the fracture strength. The most common methods of determining fracture strengthrelyon the constrained plasticity associated with the low-temperature deformation ofnotched specimens. At extremely low temperatures, where the fracture strength is lessthan the macroscopic yield strength [37], unstable fracture is nucleation-limited [45] andensues from the notch root [54] after the imposition of small amounts of microplasticity inthe form of twinning and/or slip [55,56]. Since the elastic/rigid-plastic and elastic-plasticstress distributions have been thoroughly characterized for various notched, three-pointbend specimens [39,40,53], estimates of fracture strength are more typically obtained inthe range of temperature where the unstable fracture load is less than the general yieldload [23,33,35,37,39,40]. Various investigators [57-59] have also calculated a "cleavagecharacteristic stress" at a temperature corresponding to the upper bound of the fracture-initiationtransition1. In contrast to the results of the present investigation, the aforemen-tioned methods of evaluating fracture strength all represent test conditions in which un-stablefracture occurs prior to ductile tearing in the notch root.

Unstable fracture following ductile crack propagation has been studied for the case offracture toughness testing in the high-temperature end of the transition regime. Rosenfieldand Shetty [60,61], after evaluating the fracture toughness data for an ASTM A508 steelby a variety of methods, found that Seidl's available-energy method [62] provided lower-bound values of transition toughness with the least amount of variability. The available-energy method is based on the premise that unstable (cleavage) fracture is driven by thestored elastic energy in a specimen at the moment of instability, regardless of the extent ofsubcritical crack extension preceding unstable fracture. This method of analysis partitionsa non-linear, load-displacement response into components that represent the plastic energyassociated with stable crack growth and the elastic energy available to propagate the crackin an unstable manner. Although subcritical crack growth prior to unstable fracture is notmanifested as non-linearity in the (type I) P-t curves, the analysis of the fracture load datainthis investigation is also based on the implicit assumption that the stored elastic energy

1Thefracture-initiation transition [52] is generally defined as the range of temperaturebounded by the temperature at genern:lyield and the temperature corresponding to thetransition from type I to type II P-t behavior [31] in low-strength steels.


in a specimen is equivalent to the unstable fracture resistance of the material at the mo-ment of instability. In conjunction with this assumption, the effects of differences in thedepth of the ductile shear zone on the calculation of the fracture strength are ignored,which will effectively provide conservative estimates of differences in the brittle fractureresistance of the two material conditions. The results of this investigation indicate that themean values offracture strength for the conventionally-processed and pretreated/annealedmaterial conditions of each steel are significantly different at the 99% confidence level,Figure 10. These differences exemplifYthe beneficial effects of refining grain-refiningprecipitates on increasing the brittle fracture resistance of tempered martensite under con-ditions in which a critical damage mechanism is operative.

An increase in test temperature from -196°C to -100°C produces a variety of changesin the fracture behavior of the steels, Tables 3 and 4. The values ofPMAX for type I behav-ior and the values ofPoy for type nbehavior increase over this range of temperature in thesteels. These increases in load reflect the increasing amounts of constrained plasticity inthe notch root that accompany increases in temperature [33-37,39-41,56-59,63]. Unlikethe lower-strength steels evaluated in previous investigations, where brittle fracture occursin the range of temperature associated with constrained plasticity, ductile crack growthprecedes unstable fracture in the lightly-tempered martensite microstructures at tempera-tures much below what would be considered the fracture-initiation transition in ferritic ma-terials (i.e., the range encompassing the temperature at which general yielding occurs inthe specimen and the temperature where ductile crack initiation is observed). With theexception of the conventionally-processed 4340 steel, increases in PMAX are accompaniedby a transition from type I to type nP-t behavior as test temperature is increased from-196°C to -100°C. Considering the localized and well-defined nature of ductile crackinitiation in these steels, it would be reasonable to speculate that ductile shear zone forma-tion at low temperatures occurs prior to general yield in the Charpy specimens as a resultof the high material strength in combination with lower absolute levels of material tough-ness. However, the occurrence of type III and type IV behavior at -100°C is character-istic of deformation and fracture at temperatures above the fracture-initiation transition[31,37,52] in the pretreated/annealed 3310 and 4340 (593°C temper) steels. The manifes-tation of these latter P-t responses is consistent with the comparatively low strength andhigh (processing-induced) toughness of the two steels.

Differences in the maximum load and the morphology of unstable fracture indicatethat the application of a solution pretreatment and subcritical anneal prior to final austen-itization increases the resistance to both the initiation and propagation of unstable fractureat -100°C in the lightly-tempered 3310 and 4340 steels. Subtle differences in the unstablefracture morphology of 4340 specimens tempered at 593°C also suggest that the processprovides improvements in the resistance to unstable propagation, but the changes assoc-iated with the transition from type n to type IV P-t behavior [63-65] make it difficult todistinguish any relative differences in the resistance to the initiation of unstable fracture. Itis also not possible to directly quantifYany processing-induced differences in the initiationor propagation resistance of the 8219 steel at -100°C since the values ofPMAX are virtuallyidentical and the unstable fracture morphologies are comprised of transgranular cleavage


for both material conditions. However, based on the significant differences that exist inthe size of the ductile shear zone, it can be inferred that the resistance to the initiation ofunstable fracture is greater in the pretreated/annealed material condition.

The application of a solution pretreatment and subcritical anneal prior to final austeni-tization changes the morphology of fracture in the resultant, tempered martensite micro-structures at upper-shelf temperatures, Figures 12-14. Both the conventionally-processedandpretreated/annealed steels exhibit similar densities of the primary voids that nucleate atnon-metallic inclusions and the intermediate-sized microvoids that nucleate at MnS andother large second-phase particles retained through heat treatment. However, the growthofthese features is interrupted well before impingement by the activation of comparativelysmallerparticles as critical damage fracture-initiation sites in the microstructure. The acti-vationof coarse AlN precipitates appears to be largely responsible for the initiation ofquasicleavage in the lightly-tempered martensite microstructures of conventionally-pro-cessed, low-alloy steels tested at upper-shelf temperatures and the conventionally-process-ed4340 steel tested at near upper-shelf temperatures [15,16,66]. The refinement-inducedincreasesin particle strength relative to the matrix strength produce a transition to a moreductilemode of separation [45] in the pretreated/annealed steels, thereby resulting in theincreasedlevels of upper-shelf impact toughness. For the 4340 steel, which contains sub-stantialamounts of both TiN and AlN, the transition in fracture morphology from quasi-cleavageto ductile rupture also appears to be associated with a change in the species ofactivatedparticle from AlN in the conventionally-processed steel to smaller TiN precipi-tates in the pretreated/annealed steel [66].

Non-metallic inclusions and coarse grain-refining precipitates more than likely pro-videsites for the initiation of the secondary microvoids observed on fracture surfaces ofthe conventionally-processed, high-nickel steels, Figures 13a and 13b. The relative refine-mentof microvoids in the intervening material, however, is stochastically consistent withtheprocessing-induced refinement of grain-refining precipitates in the pretreated/annealedsteels(i.e., decreases in precipitate size and increases in particle density at similar levels ofvolumefraction). The increases in void initiation resistance that accompany the refine-mentof grain-refining precipitates effectively postpones the fracture of the interveningmaterialbetween the larger microvoids, such that when final separation occurs at higherlevelsof strain-energy density, a higher density of smaller particles is activated as sites formicrovoidinitiation"

The large content of elongated MnS inclusions appears to limit the extent of improve-mentsin the transverse toughness of the 8219 steel (0.014% S), Figure 5b. At lower-shelftemperatures, where the toughness of this low-alloy steel is governed by the depth of duc-tileshear zone extension below the notch root, elongated MnS inclusions oriented alongthe direction of crack propagation appear to provide numerous sites for the prematureinitiationof unstable fracture, Figure 15. The preferential growth of the crack along thesulfideinclusions effectively produces increased levels of stress intensification along thecrackfront, thereby limiting the extent of ductile shear zone extension required to propa-gate the macroscopic crack in an unstable manner (i.e., crack propagation along the MnS


FIG. 15--The fracture appearance of the notch-root region is shown for the (a) conven-tionally-processed and (b) pretreated/annealed 8219 steel after testing at -90°C. SEMfractographs of transverse-oriented, Charpy V-notch specimens.

inclusions provides a more energetically feasible means of inducing a change in the modeof fracture than the complete extension of the ductile shear zone). Thus, the similar levelsoflow-temperature transverse toughness exhibited by the conventionally-processed andpretreated/annealed 8219 steel result from the inability to completely extend the ductileshear zone prior to unstable fracture, and this basically limits the potential benefits that canbe derived from the refinement of grain-refining precipitates. The presence of elongatedMnS inclusions affects ductile shear zone formation in an equivalent manner at upper-shelftemperatures, although the processing-induced differences in the mode of unstable fracturein the matrix material between the inclusions, Figure 12, translate into moderate increasesin transverse toughness!, Figure 5b. It appears that the process for improving toughnessmay become ineffective at low temperatures once the volume density ofMnS inclusionsexceeds a critical value, whereas the level of improvement in transverse upper-shelf tough-ness may vary in a more continuous fashion with the relative proportions of matrix mater-ial and MnS that comprise the fracture surface.

The results of this investigation indicate that the refinement of grain-refining precipi-tates provides Hnprovements in the impact toughness of grain-refined, high-strength steels

IThe improvements in transverse upper-shelf toughness are somewhat inconsistentwith respect to sulfur content. For example, past work [15,16] has shown that the processprovides improvements of 102% (6.1 Jto 12.3 J) and 27% (9.4 Jto 11.9 J) in the trans-verse room-temperature toughness ofan 8620 steel (Rc 44) containing 0.80% Mn-0.036%S-4 ppm Ca and a 4340 steel (Rc 54) containing 0.75% Mn-0.016% S, respectively. Theimprovement in the mean toughness of the 4340 steel is not significant at the 95% con-fidencelevel.


over the range of temperature encompassing the lower-shelf, transition, and upper-shelfregimes of fracture behavior. However, the lack ofa processing-induced improvement inthe low-temperature transverse toughness of the 8219 steel (0.014% S) emphasizes theneed to gain a fundamental understanding of the interrelationships between the content ofMoS, the dispersion ofMnS, base steel composition, and the potential to improve trans-verse toughness through the refinement of grain-refining precipitates.

CONCLUSIONS

1. The application ofa solution pretreatment and subcritical anneal prior to final austen-itization provides substantially improved levels oflongitudinal impact toughness inhigh-strength steels over the range of test temperature encompassing the lower-shelf,transition, and upper-shelf regimes of fracture behavior.

2. The process for improving matrix toughness provides increased resistance to the initia-tion of unstable fracture via transgranular cleavage in grain-refined, high-strengthsteels.

3. The process for improving matrix toughness increases the resistance to unstable crackpropagation at high temperatures in the lower-shelf regime for steels containing sub-stantial amounts of nickel (e.g., the 3310 and 4340 steels).

4. Although the process provides increased levels of transverse impact toughness in the4340 (0.002% S) and 3310 (0.009% S) steels, improvements in the transverse tough-ness of an 8219 steel containing 0.014% S are limited in extent. In view of pastresults, these data suggest that processing-induced improvements in transverse tough-ness are dependent on both the content and dispersion ofMnS in a steel.

ACKNO~LEDGMENTS

The authors would like to thank The Timken Company for permission to publish theresults of t~s investigation. We would also like to express our gratitude to Ken Casnerand the staff at Quality Circle Machine for machining the test specimens for this research.

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16.Leap, M. 1. and Wingert, 1. C., "Application of the AdvanTec Process for Improvingthe Toughness of Grain-Refined, High-Strength Steels," paper to be published in the38th Mechanical Working and Steel Processing Conference Proceedings, 1996.

17. Leap, M. J., "Prevention of Particle Embrittlement in Grain-Refined, High-StrengthSteels," United States Patent 5,409,554, April 25, 1995.

18.Fundamentals of Micro alloying Forging Steels. Krauss, G. and Banetji, S. K., Eds.(Warrendale, PA: The Metallurgical Society, 1987).

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20.DeLoach, Jr., 1. 1. and Franke, G. L., "Microstructural Features Controlling TransitionBehavior in Martensitic, High Strength Steel Weld Metals," Welding 90 (Gournay-Sur-Marne, France: Institute for Industrial Technology Transfer, 1990), pp. 99-107.

21.Linaza, M. A., Romero, 1. L., Rodriguez-Ibabe, J. M., and Urcola, 1. J., "Influence ofthe Microstructure on the Ductile-Brittle Behavior of Engineering Steels ContainingBrittle Particles," 36th Mechanical Working and Steel Processing Conference Proceed-W (Warrendale, PA: Iron and Steel Society, Inc., 1995), pp. 483-494.

22.Linaza, M. A., Romero, 1. L., Rodriguez-Ibabe, 1. M., and Urcola, 1. 1., "CleavageFracture of Micro alloyed Forging Steels," Scripta Metallurgica et Materialia, Vol. 32,1995, pp. 395-400.

23.Ritchie, R. 0., 'Knott, 1. F., and Rice, 1. R., "On the Relationship Between CriticalTensile Stress and Fracture Toughness in Mild Steel," Journal of the Mechanics andPhysics of Solids. Vol. 21,1973, pp. 395-410.

24.Ritchie, R. 0., Francis, B., and Server, W. L., "Evaluation of Toughness in AISI 4340Alloy Steel Austenitized at Low and High Temperatures," Metallurgical Transactions,Vol. 7A, 1976, pp. 831-838.

25. Chang, E., Chang, C. Y., and Liu, C. D., "The Effects of Double Austenitization onthe Mechanical Properties ofa 0.34C Containing Low-Alloy Ni-Cr-Mo-V Steel,"Metallurgical Transactions. Vol. 25A, 1994, pp. 545-555.


26. Ireland, D. R., "Procedures and Problems Associated with Reliable Control of theInstrumented Impact Test," ASTM STP 563: Instrumented Impact Testing (phila-delphia, PA: American Society for Testing and Materials, 1974), pp. 3-29.

27. Saxton, H. 1., Ireland, D. R., and Server, W. L., "Analysis and Control ofInertialEffects During Instrumented Impact Testing," ASTM STP 563: Instrumented ImpactTesting (philadelphia, PA: American Society for Testing and Materials, 1974),pp. 50-73.

28. Ireland, D. R., "Critical Review ofInstrumented Impact Testing," Dynamic FractureToughness. An International Conference (Cambridge, England: The Welding Insti-tute, 1977), pp. 47-62.

29. Server, W. L., Wullaert, R. A., and Scheckherd, 1. W., "Evaluation Procedures forDynamic Fracture-Toughness Testing," ASTM STP 631: Flaw Growth and Fracture(philadelphia, PA: American Society for Testing and Materials, 1977), pp. 446-461.

30. Hoover, W. R., "Effect of Test System Response Time on Instrumented CharpyImpact Data," ASTM STP 563: Instrumented Impact Testing (philadelphia, PA:American Society for Testing and Materials, 1974), pp. 203-214.

31. Leap, M. 1., The Effects of Forging on the Microstructural Development. Strength.and Dynamic Fracture Behavior of Microalloyed Ferrite-Pearlite Steels (Golden, CO:Colorado School of Mines, M. S. Thesis, 1987).

32. Hillert, M. and Staffansson, L. I., "The Regular Solution Model for StoichiometricPhases and Ionic Melts," Acta Chemica Scandinavicll, Vol. 24, 1970, pp. 3618-3626.

33. Tetelman, A. S., Wilshaw, T. R., and Rau, Jr., C. A., "The Critical Tensile StressCriterion for Cleavage," The International Journal of Fracture Mechanics, Vol. 4,1968, pp. 147-157.

34. Malkin, 1. and Tetelman, A. S., "Relation Between Kro and Microscopic Strength forLow Alloy Steels," Engineering Fracture Mechanics, Vol. 3, 1971, pp. 151-167.

35. Wilshaw, T. R. and Pratt, P. L., "On the Plastic Deformation ofCharpy SpecimensPrior to General Yield," Journal ofthe Mechanics and Physics of Solids. Vol. 14,1966, pp. 7-19.

36. Turner, C. E., "Measurement of Fracture Toughness by Instrumented Impact Test,"ASTM STP 466: Impact Testing of Metals (philadelphia, PA: American Society forTesting and Materials, 1970), pp. 93-114.


37. Wullaert, R. A., "Applications of the Instrumented Charpy Impact Test," ASTM STP466: Impact Testing of Metals (philadelphia, PA: American Society for Testing andMaterials, 1970), pp. 148-164.

38. Yan, C., Chen, 1. H., Sun, 1., and Wang, Z., "Critical Assessment of the Local Cleav-age Stress orin Notch Specimens ofC-Mn Steel," Metallurgical Transactions, Vol.24A, 1993, pp. 1381-1389.

39. Ewing, D. 1. F. and Griffiths, 1. R., "The Applicability of Slip-Line Field Theory toContained Elastic-Plastic Flow Around a Notch," Journal of the Mechanics andPhysics of Solids, Vol. 19, 1971, pp. 389-394.

40. Griffiths, 1. R. and Owen, D. R. 1., "An Elastic-Plastic Stress Analysis for a NotchedBar in Plane Strain Bending," Journal of the Mechanics and Physics of Solids, Vol. 19,1971, pp. 419-431.

41. Wilshaw, T. R., "Deformation and Fracture of Mild Steel Charpy Specimens," Journalof the Iron and Steel Institute, Vol. 204, 1966, pp. 936-942.

42. Lequear, H. A. and Lubahn, 1. D., "Root Conditions in a V-Notch Charpy ImpactSpecimen," Welding Research Supplement, December, 1954, pp. 585s-588s.

43. Evans, A. G., "Statistical Aspects of Cleavage Fracture in Steel," Metallurgical Trans-actions. Vol. 14A, 1983, pp. 1349-1355.

44. Lin, T., Evans, A. G., and Ritchie, R. 0., "Statistical Analysis of Cleavage FractureAhead of Sharp Cracks and Rounded Notches," Acta Metallurgic!l, Vol. 34, 1986, pp.2205-2216.

45. Un, T., Evans, A. G., and Ritchie, R. 0., "Stochastic Modeling of the IndependentRoles ofParticIe Size and Grain Size in Transgranular Cleavage Fracture," Metallur-gical Transactions, Vol. 18A, 1987, pp. 641-651.

46. Weibull, w., "A Statistical Distribution Function of Wide Applicability," Transactionsof the American Society of Mechanical Engineers (Journal of Applied Mechanics),September, 1951, pp. 293-297.

47. Landes, 1. D., The Effect of Size. Thickness and Geometry on Fracture Toughness inthe Transition (Geesthacht, Germany: GKSS Research Center, Report 92/E/43,1992).

48. Landes, 1. D., "A Two Criteria Statistical Model for Transition Fracture Toughness,"Fatigue & Fracture of Engineering Materials & Structures. Vol. 16, 1993, pp. 1161-1174.


49. Lereim, 1. and Embury, 1. D., "Some Aspects of the Process Zone Associated with theFracture of Notched Bars," What Does the Charpy Test Really Tell Us? (Metals Park,OH: American Society for Metals, 1978), pp. 33-53.

50. Zia-Ebrahimi, F., Matlock, D. K., and Krauss, G., "On Ductile Crack Initiation inNotched Bend Specimens," Scripta Metallurgica. Vol. 16, 1982, pp. 987-992.

51. Green, A. P., "The Plastic Yielding of Shallow Notched Bars Due to Bending," Jour-nal of the Mechanics and Physics of Solids. Vol. 4, 1956, pp. 259-268.

52. Green, A. P. and Hundy, B. B., "Initial Plastic Yielding in Notch Bend Tests," Jour-nal of the Mechanics and Physics of Solids, Vol. 4, 1956, pp. 128-144.

53. Alexander, J. M. and Komoly, T. 1., "On the Yielding of a RigidIPlastic Bar with anIzod Notch," Journal of the Mechanics and Physics of Solids, Vol. 10, 1962, pp. 265-275.

54. Hendrickson, 1. A., Wood, D. S., and Clark, D. S., "The Initiation of Brittle Fracturein Mild Steel," Transactions of the American Society for Metals, Vol. 50, 1958, pp.656-676.

55. Griffiths, 1. R. and Cottrell, A. H., "Elastic Failure at Notches in Silicon-Steel,"Journal of the Mechanics and Physics of Solids. Vol. 13, 1965, pp. 135-140.

56. Knott, 1. F. and Cottrell, A. H., ''Notch Brittleness in Mild Steel," Journal of the Ironand Steel Institute, Vol. 201, 1963, pp. 249-260.

57. Li, D. M. and Yao, M., "Effect of Notch Depth on Low-Temperature FractureBehavior of Mild Steel," Scripta Metallurgica. Vol. 21, 1987, pp. 593-596.

58. Huang, Z. and Yao, M., "An Approach to Cleavage Microcrack Propagation," ScriptaMetallurgic!l, Vol. 23, 1989, pp. 1335-1340.

59. Yan, X. and Lei, W., "Research into Fracture Behavior of Mild Steel in Crack-LikeNotch Impact Test," Scripta Metallurgica et Materialia. Vol. 29, 1993, pp. 797-800.

60. Rosenfield, A. R. and Shetty, D. K., "Lower-Bound Fracture Toughness of a Reactor-Pressure-Vessel Steel," Engineering Fracture Mechanics. Vol. 14, 1981, pp. 833-842.

61. Rosenfield, A. R. and Shetty, D. K., "Cleavage Fracture of Steel in the Upper Ductile-Brittle Transition Region," Enlrineering Fracture Mechanics, Vol. 17, 1983, pp. 461-470.


62. Seidl, w., "Specimen Size Effects on the Detennination ofKIc-Values in the Range ofElastic-Plastic Material Behavior," Engineering Fracture Mechanics, Vol. 12, 1979,pp. 581-597.

63. Feamehough, G. D. and Hoy, C. 1., "Mechanism of Deformation and Fracture in theCharpy Test as Revealed by Dynamic Recording of Impact Loads," Journal of the Ironand Steel Institute. Vol. 202, 1964, pp. 912-920.

64. Server, W. L., Norris, Jr., D. M., and Prado, M. E., "Ductile Crack Initiation in theCharpy V-Notch Test," What Does the Chaq>yTest Really Tell Us? (Metals Park,OH: American Society for Metals, 1978), pp. 187-200.

65.Norris, Jr., D. M., "Computer Simulation of the Charpy V-Notch Toughness Test,"Engineering Fracture Mechanics. Vol. 11, 1979, pp. 261-274.

66. Leap, M. 1., unpublished research, The Timken Company, 1996.

Martin A. Walsh' and Stephen Price2

THE CHEMISTRY MODIFICATIONS TO ASTM A 707 FOR OFFSHORESTRUCTURAL INTEGRITY

REFERENCE: Walsh, M.A., Price, S., "The Chemis~ry Modifications to ASTMA 707 for Offshore Stmcturalintegrity", Steel Forgings: Second Volume, ASTMSTP1259, E.G. Nisbett and A.S. Melilli, Eds, American Society for Testing andMaterial s, 1997.

ABSTRACT: The development of low alloy steels for offshore applications hasgenerally led to the selection of increasingly "leaner" chemistries which are tailored tomeet specific design requirements. Low carbon, low alloy, copper bearing steels basedon the ASTM A 707 grade system have attracted considerable interest for applicationscombining high strength and toughness with good weldability. Forgemasters Steel andEngineering has been intimately involved in the development of A 707 variants whichhas led to contracts for the production offorged components for the Auger and Marsplatforms. A thorough review of structure property relationships with regard to therole of copper precipitation during ageing, has been undertaken. Comparisons havebeen drawn with conventional low carbon, low alloy steels where similar properties areachieved with tempered bainitic microstructures.

KEYWORDS: copper bearing steels, offshore, tension leg platform, weldability, lowalloy steel

INTRODlJCTION

Forgcmasters Steel and Engineering Limited (FSEL) is involved in the production oflow alloy steels for critical applications in offshore structures One of the mostdemanding markets is that of forged components for tension leg platforms (TLP). TheTLP is a relatively modern type of oil production platform in which a floating structureis tethered by vertical buoyant tubular sections to a template on the sea bed.

]Product Metallurgist. Forgel1lastcrs Stcel and Engineering, Sheftield, England

'Technical Director (Dcsign;lte). Forgclllilstcrs Steel and Enginecring, Sheffield, England

196

WALSH AND PRICE ON THE CHEMISTRY MODIFICATIONS 197

Currently there are six TLP's in service with further platforms under construction seeTable ]. In deep sea fields their design offers significant cost reductions and improvedreliability over the more conventional "jacket" or "semi-submersible" type rigs, seeFigure ].

TABLE l--Materials of deep sea offshore platforms

TLP LOCATION DESIGNER RISER MATERIAL

Hutton North Sea CONOCO 3.5 NCMV

,Joliet Gulf of Mexico CONOCO CMV

Snorre North Sea SAGA HY80

Auger Gulf of Mexico SHELL A707 MOD

Heidrun North Sea CONOCO HY80

Mars Gulf of Mexico SHELL A70? MOD


The mooring and riser systems of a TLP require exceptional heavy wall strength andtoughness due to the high dynamic and static stresses endured during service.

Development of suitable low alloy steels has become increasingly important as TLPdesigns have evolved and consideration of the weldability of the components has beennecessary. Material development has focused on improving the weldability of3% NiCr Mo steels by reducing the alloy and carbon levels and therefore reducing the carbonequivalent. For the Auger and later, the Mars platforms a more sophisticated approachhas been adopted, with the development of a low carbon, copper precipitationhardening steel which offers significant advantages over conventional low alloy steels.

The relative weldability of various alloy chemistries can'be represented graphically onthe Graville Diagram, Figure 2 [I] The diagram describes the material susceptibly toHAZ cracking as a function of carbon content and carbon equivalent.

FIG. 2--Weldability of steel as a function of carbon and carbon equivalent from Graville

This paper reviews the development of compositions which lie in region of 1 of theGraville diagram, resulting in excellent weldability whilst maintaining high strength andtoughness.


HISTORICAL REVIEW

The development of high strength steels during the twentieth century gave an optimumcombination of low cost, strength, toughness and weldability for fabricated structuralcomponents

These steels were principally carbon manganese or "mild steel" grades, but it was notuntil the 1950's that the benefits of micro-alloying were realised Micro-alloying withaluminium, vanadium, niobium or titanium either singularly or in combination produceda fine dispersion of carbonitrides which pinned grain boundaries and dislocationsresulting in fine ferrite grain size A reduction in ferrite,grain size had the advantage ofa positive effect on both strength and toughness in accordance with the Hall Petchrelationship and in contrast to the addition of alloying elements which increase strengthat the expense of toughness. The Hall Petch relationship can be described by thefollowing equation:-

More recently the performance of micro-alloyed steels has been further enhanced by theintroduction of thermo-mechanical processing techniques which generate ultra finegrain size, resulting in superior strength and toughness,

In the 1970's and 80's thick wall forgings for the offshore industry were generallymanufactured from low alloy steels, typically Ni Cr Mo V or Cr Mo types. Thesesteels had been developed from the power generation and pressure vessel industries andwere generally not welded The Hutton TLP mooring system utilised 3.5% Ni Cr MoV type steels as detailed in Tables 2 and 3. This design required yield stress in exc~ssof 850 Mpa with excellent cryogenic toughness

The desire for forging grade steels to match the weldability of pipeline steels hasbrought about a gradual reduction in carbon and alloy contents and consequentlyreduced carbon "equivalents and hardenability. The etfect of carbon equivalent on yieldstrength is shown schematically in Figure 3 [2). The Snorre TLP was the first to usewelded tethers and risers instead of conventional mechanical fasteners. These weremanufactured from modified grades of HY80 and 2.25% Cr Mo to achieve theminimum yield strength of 550 Mpa and retain good weldability. Typical chemistriesand properties are shown in Tables 2 and 3. The potential for increasingly "leaner"steels of this type is limited by their hardenability and the necessity to maintain a fullybainitic microstructure.

An alternative route has been the development of a very low carbon, precipitationhardening steel for the Auger platform, subsequently also used in the Mars platform

This skel is hdsed on ASTI\1 Forged Cdrhon dnd Alloy Steel Flanges for Low-Temperature Service (A 707 grade L'i) and relics on copper precipitation to offset thereduction in c;llhon content

Although both genel"ic types of steel ale readily weldable, the copper hearing steelshave a low calhan content which allows fabrication withoul preheat or post weld heatlr(~alll1ent

FIG.3--Relationship between yield strength, carbon equivalent and heat treatment

ALLOY DEVELOPMENT

FSEL has developed and optimised various alloy grades for offshore applications.These steels are primarily based on two generic systems; ASTM A707, used as a basisto develop A707 modified and HY80 based types. They have been developed throughlaboratory trials and full size prototypes prior to the manufacture of actual productionforgings.

The alloy systems are fundamentally different in the mechanism by which they achievetheir strength. The HY80 alloy relies on the formation of a bainitic transformationprodu'ct whilst the A707 modified alloy relies on copper precipitation hardening.

In general the development of both alloy systems is similar, with improved weldabilitybeing achieved by a decrease in carbide forming elements such as chromium and carbonand an increase in solid solution strengthening elements such as manganese.

The control of carbon content is extremely important as a slight variation can haveconsiderable effect on the strength and toughness. In general, the effect of increasingcarbon is to raise the strength in a proportional manner. Conversely, very low carboncontents promote increased toughness and weldability.


Molybdenum and chromium additions retard copper precipitation by influencing thetransformation kinetics and consequently prevent auto-ageing. Furthermore theyincrease alloy hardenability and add resistance to softening during ageing. Carbideforming elements provide an incremental increase in strength.

The addition of nickel increases the degree of crystallisation during hot working andconsequently has a grain refining influence Nickel also has a profound effect onhardenability and promotes higher toughness.

Manganese is the major solid solution strengthening agent and partly offsets thestrength reduction common with very low carbon conte~ts.

Alloying additions specific to the A707 modified grade are principally precipitateformers such as copper and niobium The primary effect of copper is as a secondaryhardening precipitate whilst niobium facilities austenite grain size control. Table 4summarises the modifications made to the hasic ASTM A707 grade L5 specificationduring the development of A707 modified

TABLE 4--ASTM A707 Grade LS and typical production analysis of A707 modified

SPECIFICATION CHEMICAL COMPOSITION· HEAT ANALYSIS (WT%) MAXIMUM UNLESS STATED

C Si Mn P S Cr Mo Ni V Cu Nb AI

ASTM 0.40 0.60 0.15 0.70 003 1.00A707L5 0.07 0.35 0.70 0.025 0.025

0.90 0.25 1.00 . 1.30 0.03

A707MODIFIED 0.03 020 1.30 0.006 0002 070 0.45 2 1 1.2 045 0.02(TYPICAL)

Experience at FSEL in the production of tendon extensions for the Auger TLP andtaper stress joints for the Mars TLP has shown that the properties of A 707 modifiedmeet the requirements of the most demanding otfshore specifications. Typicalproduction test results are shown in Table S Experience from production was backedup with work to characterise the material and a Continuous Cooling Transformationdiagram (CCT) was constructed. This is shown in Figure 4. The expected phases arepresent which have been widely reported in the literature previously for A 707 typematerial. Laboratory trials were also carried out to further demonstrate the propertiesof A 707 moditied


TABLE 5-- Typical production test results in A707 Modified

Component R 0.2 Rm A Z Kee Pelluii FATTMPa MPa % % KsiJin P3 Drop Weight °c

@ 4.4 °c Test @ -23.3 °c

Tendon Extension 540 633 27 79 344/300/369 NO BREAK -45

Taper Stress JOint 586 660 27 78 331/362/368 NO BREAK -60

FIG. 4--Continuous cooling transformation diagram of A707 Modified

LABORATORY TRIALS

An induction melted 50kg split cast was manufactured to characterise the effects ofcopper on the base chemistry of the A707 modified. The product check analysis for thecast is given in Table 6.


TAB LE 6--Analyses of trial split cast

STEEL C 5, Mn P S Cr Mo Ni Cu Nb AI CE

A707 MODIFIED 0.05 0.31 1.56 0.010 0.005 0.71 0.45 0.96 1.06 0.043 0.03 0.68+Cu

A707 MODIFIED-Cu 0.045 0.29 1.55 0.010 0.005 0.70 0.45 0.96 0.01 0.047 0.03 0.49

The ingots were forged in the range I220ae to 900ae to a 20mm diameter bar giving aforging reduction of approximately'): I. Thc bars were allowed to air cool andtransform.

A simulation quality heat treatment was given to each bar to replicate the cooling ratesand hold periods of a water quenched 300mm section. Various tempering/ageingtemperatures were applied to characterise further each grade.

Further tests were conducted to examine thorough thickness hardenability of each alloyby Jominy end quench tests.

RESULTS

The tensile test data is shown graphically in Figures') to 7. The tempering/ageingresponse of each alloy has been investigated in terms of tensile strength, 0.2% yieldstrength, and hardness.

The optimum ageing response for the copper bearing steel was approximately 470aegiving a peak tensile strength of900MPa with a corresponding yield strength of7')OMPa.

The softening response of each stccl increased significantly at temperatures abovesooae, and although they generally followed a similar trend, the copper bearing steelwas higher in st.rength in each heat treatment condition.

The hardness response is shown in Figure 7 and generally follows the same trendsdescribed above for tensile propel1ies The hardenability of both steels wasinvestigated using the Jominy end quench test, the results of which are showngraphically in Figure 8. The copper bearing steel demonstrates higher "as quenched"hardness and greater depth of hardness




The impact toughness of both alloys was also investigated and showed the non coppercontaining material to have a lower transition temperature and greater upper shelfenergy as detailed in Figure 9.

FIG. 9--lmpact energy transition curve

DISCUSSION

The capability of several alloys to achieve yield strengths in excess of SSO Mpa incombination with excellent ~cIdability has been reviewed Their weldabiltiy as afunction of carbon equivalent and carbon content are superimposed on the Gravillediagram in Figure 2.

The 2.25% Cr Mo and HY80 grades are used for similar applications and, althoughthey lie in zone III, are readily weldable when suitable preheat and postweld heattreatments are applied. The 2.25% Cr Mo steels are preferred for sour gasenvironments where a restriction in nickel content prohibits the use ofHY80 grades.Nickel levels of the order of3.0% increase the tendency for stress corrosion crackingdue to the influence of retained austenite Early A707 variants with nickel contents lessthan I% were developed specifically lor sour service components such as export gasrisers. Unfortunately no comparative corrosion data is readily available.A focal point in the development of "leaner" hybrids has been compliance with the zoneI requirements of the Graville diagram combined with yield strength in excess ofSSOMPa. This is of special interest as very low carbon contents, below 0.07% virtuallyprohibit hard, crack sensitive heat affected zones on welding, making preheat prior towelding unnecessary


Modifications made to the HYSO analysis has reduced the carbon and carbonequivalent values to the extent that these alloys are tending towards the boundarybetween zone I and zone III. It is considered that beyond this position the limitingfactor is attaining yield strength values in excess of 550 Mpa.

This limitation has been overcome by the development of a low carbon steel based onASTM A707 grade L5. Its chemistry satisfies the Zone I requirements and its strengthreadily exceeds 550 Mpa A typical production analysis is shown in Table 4. The CCTcharacteristics have been investigated and the expected phases are present as have beenpreviously repol1ed for A707 type grades [3].

The role and influence of copper on the mechanical properties and microstructures havebeen investigated utilising a laboratory trial melt [4]. The base alloy containing only aresidual copper level is similar to the HYSO hybrid grades, albeit with considerablylower carbon.

Clearly the ageingltempering curves show lhatthe copper containing material offerssignificantly higher strengths and hardcnability but does suffer from reduced toughness.The addition of copper influences the strength by two mechanisms, firstly as a solidsolution strengthening element it enhances both the "as quenched" and overagedstrength by approximately IOOMPa and secondly as a precipitation hardening element itpromotes a peak hardening condition during ageing where yield strengths up to 750Mpa are achieved.

The reduction in toughness shown by the copper bearing alloy is in part due to theeffect of copper precipitation and higher strength which depresses the upper shelfenergy and increases the transition temperature. However, the toughness issignificantly greater than the most severc current design requirement.

A feature of copper bearing steels has been occasional variations in impact energywhich has been related to microstructure It would appear that copper precipitationpromotes a non uniform carbide microstructure by influencing carbon segregation,which in turn, results in localised high carbon transformation products and coarsecarbide preci'pitation The combined effect of these features is to produce low energyfracture initiation sites resulting in variable Charpy impact behaviour. Previous studies[5] have shown that this behaviour in copper steel can either be eliminated orsignificantly reduced by introducing double water quenching during the quality heattreatment or by increasing the nickel content to greater than 1.6%

Several components have been successfully manufactured in A707 modified for theAuger and Mars platforms, ie. Tendon extensions and Taper Stress joints and a typicalproduction analysis is shown in Table 4. Typical mechanical properties are included inTable 5 and compare favourably with the typical acceptance criteria als9 detailed.


CONCLUSIONS

1. The weldability of conventional low alloy steels can be enhanced by a reductionin carbon and alloy content consistent with achieving the minimum throughthickness, strength and toughness requirements.

2. For designs where yield strengths in excess of 550 MPa are required only A707alloys can currently be welded without preheat and be safe from HAZ cracking.

3. Levels of approximately I% copper increase solid solution strengthening andpromote precipitation hardening.

4. Copper increases carbon segregation and results in a non homogenous carbideprecipitation and localised high carbon transformation products. However, thishas been shown to be eliminated by thermal treatments or by increasing thenickel content to above 1.6%.

5. FSEL has been involved with the development of A707 modified based onASTM A707 grade L5 and production experience using the modified grade hasshown that it is capable of meeting the most demanding offshore specifications.

REFERENCES

[1] Graville, B.A., "Cold Cracking, Cracking in Welds in HSLA Steels", Welding ofHSLA Structural Steels ASM, Rome, Italy Nov 9-12, 1976.

[2] "High Strength Low Alloy Steels", International Iron and Steel IndustryCommittee on Technology, Brussels, 1987.

[3] Krauss, G, and Thompson, SW, "Ferrite Microstructure in ContinuouslyCooled Low and Ultra Low Carbon Steels", ISIJ International, VoL35, N08,pp.937-945, 1995

[4] Willard, Louisa, University of Sheffield, PhD Thesis unpublished work.

[5] Smith, JD.,Kearney, MG., and Dedmon, S.L, "Production and Properties ofForged Copper Precipitation Strengthened, High Toughness Low Alloy Steelsfor Tension Leg Platform Components", 11th International Forgemasters meeting11114 June, Terni, Italy.

Test Methods

John R. M. Viertl!

ULTRASONIC SIGNAL PROCESSING USING INDICATION SETS FORDETECTION AND CHARACTERIZATION

REFERENCE: Viertl, John R. M.; "ULTRASONIC SIGNAL PROCESSING USINGINDICATION SETS FOR DETECTION AND CHARACTERIZATION"; SteelForgings: Second Volume, ASTM STP 1259, E.G. Nisbett, and A.S. Melilli, Eds.,American Society for Testing and Materials, 1997.

Abstract: The potential for discontinuities in turbine and generator rotors must be dealtwith in the manufacture of these components. In recent years it has become practical toacquire and analyze large quantities of ultrasonic data obtained from testing solid rotorforgings.

This paper describes an ultrasonic system for testing solid rotors, and discusses thedetection, characterization, and imaging capability developed for evaluating rotorindications.

For detection of rotor indications, the ultrasonic signal path is divided into rangewindows. The signals within each window are compressed and stored for detection andcharacterization algorithms. Detection is performed by associating echoes from oneultrasonic path with ti}.edata from neighboring paths. These associations are constructedwith respect to the cylindrical coordinate system used during the process.

Characterization is performed using the indication data sets constructed during thedetection process. The final results determined by these automatic data processingalgorithms are presented to the user as potential indication lists which can be correlated toimages constructed from the original data.

Keywords: ultrasonic rotor testing, ultrasonics, solid rotors, indication sets,characterization, detection, nondestructive testing, low alloy steel

]Senior Engineer, GE Co. Power Generation Engineering, I River Road, Schenectady, NY 12345

213


Introduction

The ultrasonic testing offorgings for the production of turbine and generator rotors hasbeen practiced for may years by various forging manufactures and by their usersI1-ll.Anumber of authors have described the manual testing of forgings and efforts to automatethem[±]. Ideally, the ultrasonic examinations are applied as early as possible so that anattempt can be made to eliminate the indications by additional forging. However, thecoarse grain structure limits the early application of an ultrasonic test to frequencies lessthan 1 MHz. After forging and heat treatment the grain size is reduced. At this stage, thehigher frequencies of 2 to 5 MHz are used for ultrasonic testing on prepared rotorsurfaces. The ultrasonic inspection directions and tests used are divided into radial andaxial.

The ultrasonic testing of rotor forgings typically uses two orthogonal testing directions.These are the radial and the axial tests. The most significant test direction is radial. Theforging process tends to align indication structures along the axial or longitudinaldirection of the rotor. The ultrasonic beam propagating in the radial direction will tendto be perpendicular to these indication structures of some particular radii. Thus scanningthe surface of a rotor with a radially propagating ultrasonic wave uses a beam directionthat is likely to produce the highest ultrasonic echo. This is the direction that is mostlikely to yield the most sensitive test.

In axial ultrasonic testing structures are usually not favorably oriented for receiving astrong ultrasonic echo. The propagation paths are usually so long that only the largestindication structures are detected. Sometimes false indications are produced by echoesfrom the larger grain structure near the center of the rotor.

This paper describes the use of ultrasonic indication sets for the analysis of the largequantities of data from an automated ultrasonic rotor forging test system used to test solidrotors and operating at 2.25 MHz. The motivation to automate the radial ultrasonic rotortest was to improve detection of indication structures and reduce operator variability oftest data. Manual testing relied upon an operator to observe the CRT and record alloccurrences of an ultrasonic echo exceeding a predetermined threshold. The automatedsystem removes the operator variability, but requires signal processing and analysis todistinguish between real signals and noise.

The rotor test configuration for automatic testing is shown in Figure 1. The rotor issupported upon a set of hard faced power rolls. During testing the rotor is rotated whilethe transducer is held stationary, following a circular path on the rotor surface. After afull revolution of data is obtained, the transducer is stepped axially to the next positionfor a circumferential scan. Thus the transducer moves in a circumferential-axial pathover the rotor surface. This revolution of data forms a data slice from the rotor. Aftercompletion of a data slice, an axial step is performed and the circumferential scan isrepeated. This operation continues until the pre-selected axial scan length has been

VIERTL ON ULTRASONIC SIGNAL PROCESSING 215

tested. While this axial step is occurring, the data is compressed and stored on an opticaldisk. It is also displayed upon the system monitor.

Fig. I Schematic of ultrasonic system for rotor testing.

In order to understand the construction of ultrasonic indications sets, the ultrasonicsystemand the raw data capture are described. A spike pulser excites an oil coupled PZTtransduceroperating at 2.25 MHz as shown in Figure 1. The ultrasonic pulse propagatesI alonga rotor diameter and is detected by the same transducer (pulse-echo testing). TheI signalis amplified with a logarithmic amplifier with approximately 78 dB dynamic range.Thisamplifier is adjusted so that 20 dB signal change produces a one volt change inoutput. The amplitude record is full wave rectified and digitized into 8 bits with a 25MHz sampling rate.

Thevoltage range of the analog to digital converter spans 4 volts and provides 64 outputlevelsfor each volt. The analog output on the logarithmic amplifier has been adjusted so

at each decade of signal level corresponds to one volt.


Sampling the ultrasonic return at 25 MHz corresponds to a one-way metal path of 0.12mm (0.0047 inches). The range of the ultrasonic test pulse has a length of approximatelythree to five oscillations at 2.25 MHz. This corresponds to a range extent ofapproximately 0.0079 m (0.311 inches) to 0.0132 m (0.0.520 inches). From this point ofview, any ultrasonic pulse is represented by 33 to 55 data points. Each completeoscillation of the ultrasonic pulse is represented by approximately 11 data points.

Although manual ultrasonic testing from the rotor surface follows a spiral path, acylindrical surface scan was selected for the automatic testing system. The axial step sizeis 0.0127 m (0.5 inches), a distance which is approximately 1/8 th of the transverse oraxial beam spread at the center of a 925 mm (36.4 inch) diameter rotor. After thetransducer has completely an axial move of one step, it is pulsed every 0.5 degrees until afull revolution of data has been acquired. The angular reference both for the zero degreesposition and the one-half degree step is based upon data from the rotor angle encodershown in Figure 1.

The digitized ultrasonic return is divided into 512 range bins which span the time fromultrasonic pulse generation to approximately one microsecond beyond the time to receivethe back wall echo. For instance, this time is approximately 315 microseconds for thepulse to travel one round trip in this 925 mm (36.4 inch) diameter rotor. This assures thatthe ultrasonic system will always capture the back wall echo. Monitoring of the backwall echo is performed to insure that the ultrasonic test amplitude for that record isadequate for evaluation.

The raw data set is quite large. For instance, for the rotor diameter being discussed, asingle pulse-echo record consists of approximately 7875 data points. One 0.0127m (0.5inches) cylindrical test slice has 6.2 Mbytes of data. The data produced by a 6.09m ( 20ft.) long section is approximately 3.0 Gbytes.

To reduce the data storage requirements, the data is compressed by selecting andretaining the maximum amplitude in each range bin. Compressed in this way, the rawultrasonic record retains all the significant amplitude information. The error of the rangeinformation is increased slightly, but this is not significant in this application. Aftercompression, the storage space required for the data is only approximately 275 Mbytes.The ultrasonic beam geometry is shown is Figure 2. This geometry assumes that therotor is approximately 0.925 m (36.4 inches) diameter, and the ultrasonic transducer sizeis 6 mm x 25 mm (1/4" x I"). The ultrasonic beam spread is greater in the transverseplane than in the axial direction of the rotor. This is desirable because it insures that therotor interior is well covered by the ultrasonic beam. The fact that the beam spread ismore restricted in the axial direction is not a hindrance for good coverage due to the smallaxial step size.


Fig. 2 Ultrasonic beam geometry.

Thus an ultrasonic test system for inspection of un-bored rotor forgings was developedwhich captures the ultrasonic return signals, provides detection and characterizationcapability of those signals.

Detection of Ultrasonic Signals

The ultrasonic indication sets are built from the compressed ultrasonic records bysearching each range bin for a maximum signal value. Each bin which originallycontained 15 data points is replaced by the maximum value in the bin. Comparisonsbetween the original record and the compressed record show that all the instantaneouspeak values are retained. There is a small time shift in that the range is now determinedby the bin number, i.e. a number between 0 and 512. For instance, for the 0.925 m (36.4inches) diameter rotor, the range resolution implied by an error of one bin isapproximately +/- 0.0036 m (0.14 inches). Since the primary purpose of this system is todetect ultrasonic indications, some of the range accuracy has been relaxed so that the datastorage requirements are reduced. This data compression approach guarantees that noindication signal will be missed while allowing the compression of the data by a factor ofapproximately 11 to 1. This is the form which is saved for archival purposes and for


subsequent data analysis. The data analysis function is performed on the same field testsystems that recorded the original raw rotor data. The compressed ultrasonic data istypically recorded on an optical disk.

The amplitude of the data is recorded as the logarithm of the original signal. In additionto the ultrasonic amplitude data, both axial and angular position are recorded in thecompressed data records.

Data Analysis and Ultrasonic Indication Set Building

This section describes the sequence of operations used to build indication sets. Theprocess begins with the selection of one of the ultrasonic records obtained from pulsingthe transducer on the rotor surface at one of the 0.5 degree angular locations asdetermined by the encoder attached to the rotor. A range dependent threshold for dataanalysis is selected by the operator. It is applied during set building. The operator alsoenters the number of adjacent range windows that are expected to be associated with atarget. The number of range windows selected typically corresponds to the ultrasonicpulse length. Usually several range windows are required to span the range of theultrasonic pulse. Any real target must satisfy this minimum number of adjacent rangewindows to be considered a target for set building.

The first requirement the ultrasonic echo should satisfy is that at least one value in eachof these range windows exceeds the selected threshold. If it does, an indication set isstarted for the range coordinate. Any adjacent range window that has an ultrasonicamplitude that exceeds threshold is also added to this range set. The end of a range setoccurs when an adjacent range window does not have an ultrasonic amplitude thatexceeds threshold. If another ultrasonic amplitude farther along a range record exceedsthreshold, a new set is started. In this way, several sets may be been started and ended forthis ultrasonic record. The ultrasonic record is replaced with these range sets. At thispoint, several pieces of data are retained with each set. These data are the maximumamplitude for each range set along with the starting and ending range window numbers,and the angular and axial coordinate values for the transducer. This process is repeatedfor each 0.5 degree ultrasonic record until 360 degrees of data have been analyzed.

At this stage in the set building process, range sets have been formed for each ultrasonicrecord. The next step is to associate the range sets according to their angular coordinate.This starts with anyone of the range sets and compares its range window values to thoserange sets of an adjacent ultrasonic record. If the adjacent range sets overlap with theselected set, a new set is formed. This new set, a range theta set, contains the maximumand minimum range window values, the starting and ending angular coordinates of therange sets of the adjacent ultrasonic records, and the larger of the range set maximumamplitudes. Also the coordinate information for this maximum amplitude is recorded inthis new set, namely, the range, angular, and axial position of the transducer.


The final set formation process builds range theta 'z' sets from the range theta sets bysearching the data of the range theta sets for sets that have coordinate values that areadjacent to those of other range theta sets and were formed from ultrasonic records takenat different axial locations. New extreme values for the coordinates are determined fromthe range theta sets. In addition, the maximum amplitude and its coordinates are recordedin the new range theta 'z' set.

Since these sets are formed from targets that move through the ultrasonic beam both onthe near and far side of the rotor centerline, duplicate indication sets may be formed.These duplicates are combined by comparing the completed indication sets with thosethat are separated from each other by 180 degrees. If the sets overlap sufficiently, the onewith the larger amplitude is selected. Its coordinate extremes are adjusted to reflect thiscombining of sets. An asterisk is used by the system to mark data that has been reducedby this folding operation. The peak amplitude and coordinate values are compared andused to determine the effective flat bottom hole diameter. Data for an operator tomanually confirm the final target is also provided. With this conservative approach,target detection and evaluation

Ultrasonic Calibration Rotor Test Results

A rotor test piece of 0.925 m (36.42 inches) diameter was constructed with various flatbottom test holes located at six separate transverse planes. The planes are separatedsufficiently to insure that the ultrasonic echoes only come from the targets in the testplane. Each plane has three radially directed, flat bottom holes drilled 2 inches deep andspaced approximately 120 degrees apart. Four planes of test holes were constructed withholes of 1.19, 1.59, 3.175 and 6.25 mm (3/64, 1/16, 1/8 and 1/4 inch) diameter. Theseholes are used to verify system performance in detection and evaluation.

Table 1a and 1b present the data obtained from this test rotor for the 3.175 mm (1/8 inch)diameter flat bottom holes. There are eight ultrasonic indication sets showing differentvalues for the 'equivalent flat bottom hole' (EFBH) size. The column headings labeledpeak refer to the maximum measured ultrasonic amplitude and its location in angularcoordinates. The angular coordinates are referenced to a stamped symbol on the rotor testpiece which is defined as zero degrees. The EFBH calculation is based upon formula thatcan be derived from reference (5).

EFBH = ((4/rc)*Ref. Area*(T. Amp./Ref. Amp.)*(T. Range/Ref. Range)"2)1!2

whereEFBF = effective flat bottom hole diameter, inches,Ref. Area = area of reference target, inches2T. Amp. = target amplitude, volts,Ref. Amp. = reference amplitude, volts,T. Range = distance to indication, inches,


Ref. Range = distance to reference area, inches.

Notice that all three holes were detected, and that the characterization method providesuseful sizing data.

Table la. Example ofindication Set Analysis(3.175 mm (118inch) flat bottom hole targets)

Indication Peak Back Wall Peak Peak Peak EFBHNo. Amplitude Amplitude Theta Axial Radial Size

(%) (%) (') Range Range (mm)(cm) (mm)

1 5.2 777.4 352.0 59.7 408 3.052 6.7 805.8 356.5 61.0 410 3.403 6.5 777.4 4.5 59.7 416 3.404 6.0 805.8 119.5 58.4 414 3.235 6.3 805.8 131.0 61.0 416 3.306 5.6 835.4 13 \.5 62.2 412 3.057 7.8 805.8 240.5 61.0 416 3.668 6.3 835.4 245.5 62.2 416 3.23

Table lb. Example ofindication Set Analysis(Continuation from Table la)

Indication Range Range Theta Theta Z ZNo. Min. Max. Min. Max. Min. Max.

(mm) (mm) (') (') (cm) (cm)1 407 414 349.5 356.0 59.7 59.72 410 416 356.0 358.5 61.0 61.03 412 418 0.0 19.0 59.7 62.24 408 425 109.5 122.5 58.4 62.25 412 420 124.5 133.0 61.0 61.06 408 416 13\.5 140.0 62.2 62.27 410 418 234.0 249.5 58.4 63.58 412 418 237.0 252.5 62.2 62.2

Figure 3. shows'one of the raw data slices that was used to build the indication sets.

Discussion

Ideally, the operator would select analysis parameters so that each target is represented byone range theta 'z' set. However, with small, low level targets with noise, there will besome portion of an indication set that may be treated as another set.

Examination of this data confirms the very wide beam spread in the transverse plane.Each of these holes is detected for approximately an eighth of a revolution of the rotor.This implies that there are of the order of 90 ultrasonic records for each target, or a total


of 180 records with indication data in this data slice. The 'Z' values of Table 1b showthat as much as 5.1 cm ( 2.0 inches) of axial data contributed to building these sets. And,since the axial scan step size is 1.27 cm, on the order of 5 data slices also contributed tothese sets. Thus, approximately, 900 ultrasonic records with indication data were used tobuild the sets shown in Tables la & lb. This is a data reduction of the order of 112 times.

These indications sets are used to confirm the presence of ultrasonic targets in rotorforgings. For this purpose, two additional data parameters are calculated and includedwith the indication set data. These parameters are the location for a transducer that maybe used to confirm a target and the range to the target in microseconds. This has provedto be very helpful for target confirmation by an operator.

Although the raw ultrasonic data is recorded without a threshold value, one is applied tothe ultrasonic records for set building. The threshold algorithm calculates a value basedupon the ultrasonic range being processed. It is based upon the statistics of the entire dataslice. Typically, for regions beyond the rotor centerline, the threshold is a constant basedon statistics and bounded by operator selected values.

In the case of the data in Tables la & lb, the threshold range was selected to be between3 and 5 %. A sampling procedure is used to select the threshold if the ultrasonic recordsshow that it should be between these values. The minimum value of 3% was selectedbased upon the noise seen in the background of the raw data records. Since the raw datais available, other threshold values can be selected for analysis.

The mean value of the effective flat bottom hole calculations is 3.29 +/- 0.20 mm. Thenominal calibration hole size is 3.175 mm. This amounts to a difference ofapproximately 3.6 %.

The image of the raw compressed data shown in Figure 3 illustrates the nature of theultrasonic returns from a flat bottom hole.

Conclusions

Indication sets of ultrasonic data are useful for the detection and characterization ofultrasonic data. The reduction of the large quantities of data into lists of a few targetsprovides a significant improvement over manual methods. Comparisons of EFBH andnominal hole sizes agree within 4 %. Indication confirmation by manual methods iseasily achieved. Storage of the raw, ultrasonic data provides the operator with the abilityto examine the quality of the data before performing data analysis. Automation of thissolid rotor test has improved the detection of indication structures and reduced operatorvariability. After analysis, the indication lists can also be compared to the raw data.


Fig. 3 Compressed, raw ultrasonic data taken at the axial location of 61.0 cm (24.0inches). Under the data display, the upper horizontal axis is travel time in microseconds.The lower horizontal display is the color palette in percent. The vertical axis is degrees

rotation. At the adjustable, movable electronic cross hairs the range (microseconds) andamplitude (percent) are measured and displayed along the right-hand side of the image.

Acknowledgments

The author wished to thank GE Materials and Processes Engineering, the NondestructiveEvaluation Services Operation of GE Power Generation Services Department, and SteamTurbine and Generator for their support this project, and specifically, Ms. E. Dixon forcomputer programming, Mr. M. Auger for electronic design, and Mr. M. Martincich fortransducer fixture design.

REFERENCES

[I] Krautkramer, J. and Krautkramer, H., "Heavy Forgings," Ultrasonic Testing of Materials,2nd Edition, Springer-Verlag, Heidelberg, New York 1977, pp. 357-371.


[2] Goldman R.G. and Auger, M.E., "Improved ultrasonic instrumentation for the qualitycontrol of turbine-generator components," 3rd Internat. Conf. Nondestr. Test., Tokyo 1960.

[3] Cook, G. G. and Walker, W. D., "Ultrasonic testing of gas turbine discs and heavy rotorforgings," Proc. 4th Internat. Conf. Nondestr. Test., London, 1963. London: ButterworthsSci. 1964, pp. 215-217.

[4] Goldman R.G. and Auger, M.E., "Automatic ultrasonic examination of large rotorforgings," IRE Convention Record, Vol. 9, No.6, 1961, pp. 316-326.

[5] Krautkramer, J. and Krautkramer, H., "Heavy Forgings," Ultrasonic Testing of Materials,2nd Edition, Springer-Verlag, Heidelberg, New York 1977, pp. 357-371.

William Reese Ill, David C. Ronemus2, Keyang Huang3, Emory W. Zimmers Jr.4,

and Thiru Veerabadrans

APPLICATION OF FEM-BASED MODELING OF OPEN-DIE FORGING TOPRODUCT AND PROCESS DEVELOPMENT AT BETHFORGE, INC.

REFERENCE: Reese II,W., Ronemus, D.C., Huang, K., Zimmers, E.W., andVeerabadran, T., "Application ofFEM-Based Modeling of Open Die-Forging toProduct and Process Development at BethForge, Inc.," Steel Forgings: SecondVolume. ASTM STP 1259, E.G. Nisbett and A.S. Melilli, Eds., American Society forTesting and Materials, 1997.

ABSTRACT: This paper describes a FEM-based computer system implemented by BethForge tomodel the open-die forging process and its application to product and process development. A briefoverview of BethForge, the markets it serves, and the typical product development cycle is given.Alternative methods for modeling open-die forging are discussed and the advantages offered by FEMtechniques due to recent advances in computer hardware and software are presented. The FEM-basedmodeling system that BethForge has developed in partnership with the Lehigh University ComputerIntegrated Manufacturing Laboratory (LU-CIM) and UES, Inc. is described along with the associatedmodeling methodologies. Several case studies are presented including 2D-axisymrnetric simulation ofhead-fonning, punch-fonning, upsetting, and back extrusion, and the 3D simulation of ingot cogging.

KEYWORDS: finite element method, open die forging, computer modeling, product andprocess development, FEM modeling, CAD modeling

'Forging Engineer, BethForge Inc., 1275 Daly Ave., Bethlehem, PA 18016

2Process Modeling Engineer, BethForge Inc., 1275 Daly Ave., Bethlehem, PA 18016

3Professor, Lehigh University, Bethlehem, PA 18015

4Professor, Lehigh University, Bethlehem, PA 18015

sSenior Engineer, UES Inc., 4401 Dayton-Xenia Rd., Dayton, OH 45432

224

REESE II ET AL. ON FEM-BASED MODELING 225

INTRODUCTION

BethForge, Inc. is a wholly-owned subsidiary of Bethlehem Steel Corporation.BethForge is the largest forgemaster in the United States and has been manufacturing forover 100 years. BethForge best serves those segments of the forging market requiringlarger-sized forgings. These market segments include:

1) The metals market with both hot and cold mill work rolls and backup rolls;2) The general industrial market with shafting, die blocks, etc.;3) The pressure vessel market with formed heads (with and without integral

nozzles), and special shell forgings with integral nozzles;4) The electric power generation market with rotors and shafts.Maximum ingot weight at BethForge is 261 tons (290 US tons) and 3302 mm

(130 inches) in diameter. Capabilities include a 9,000 (10,000 US) ton press, andcomplete machining, heat treatment, and testing facilities.

Historical Product/Practice Development--For many years BethForge developednew products and practices through large-scale experiments and trials on productionforgings. Large scale, pre-production trials on products of the size produced at BethForgecan cost tens of thousands of dollars. If production forgings are used for developmentwork, both dollars and delivery are at risk. To minimize these risks, changes in practicesmust be done in a very conservative manner. Both of these alternatives for developmentare very time consuming. Today's competitive market does not favor those who are slowto respond to sales inquiries, changes in the marketplace, or the emergence of newtechnology.

Modeling of Ooen-die Press Forging--In an effort to reduce developmental leadtimes and cost, over the last 15 years BethForge engineers have worked with a variety ofresources to develop and apply modeling techniques to forge shop processes, including theheating, forging, and cooling of steel. These modeling tools have afforded BethForge thefreedom to try many new ideas easily, cost-effectively, and with no risk to product quality.They have been used to develop new products and methods, as well as refine existingones. The major focus of these efforts has been modeling ofthe open-die forging processusing both physical and computer modeling techniques.

PHYSICAL MODELING AT BETHFORGE

Bethlehem Steel's Research Department supports BethForge and has been amongthe leaders of metal-processing modeling efforts in the U.S. The Homer ResearchLaboratory (HRL) features a world-class physical modeling facility [1]. In work carriedout at HRL, both plasticine and lead have proven to be excellent tools in simulating theworking of hot steel [1,2,3]. When processed under strictly controlled conditions, thesephysical modeling materials can provide, at room temperature, information and insight intothe material flow, stress/strain behavior and forming loads for steel at hot-workingtemperatures [3-6]. BethForge has developed a close working relationship with HRLengineers and has made this relationship the cornerstone of numerous modeling efforts inboth the plasticine and lead modeling media [7-11].


Both of these modeling media provide good similitude with only a few exceptions.For example, in plasticine a significant difference can be seen compared to hot steel in thethe amount of elastic strain experienced during a deformation. The resulting large amountof elastic recovery after the deformation can make plasticine modeling results difficult tointerpret [3].

Lead shows a low elastic strain, similar to that of hot steel, which makes leadsimulations dimensionally more accurate than plasticine [2,3]. However, lead exhibits selflubricating qualities which can lead to erroneous results in cases where the effects offriction on material flow are severe.

Neither of the above media can accurately model the effects oflarge changes inforging temperature, alloy composition, or phase transformation dynamics, so considerableincentive exists to use other modeling technologies to analyze these situations.

COMPUTER MODELING AT BETHFORGE

For many years BethForge has used various computer-based heat transfer modelsfor the mainframe and PC to model heating and cooling processes. These models weredeveloped both in-house and in partnership with others [12,13]. They have proven to beaccurate and helpful in predicting temperature profiles, and have helped optimize theseprocesses.

Several years ago BethForge made a strategic commitment to expand thedevelopment and use of computer modeling technologies as an aid to improving itsoperations. As part of this effort, FEM-based modeling of the open-die forging processwas explored as a possible complement to existing physical and computer modelingcapabilities. During initial feasibility studies, it became apparent that certain trends hademerged making FEM modeling offorging attractive as a cost-effective, stand-alonedevelopment tool. For example, improved software compilers and convergence algorithmshad led to modeling codes that were fast, cost-effective, and user friendly. Better pre-/post-processing software tools had become available to facilitate the visualization andanimation of time-dependent data generated from computer models. Work had beenconducted on the characterization of the deformation and heat transfer properties ofcommon forging materials. Lastly, and most importantly for BethForge, tremendousprogress had been made in the price-to-performance ratio of the computer hardwarerequired to run computation-intensive forging codes in a timely manner.

In view of the above trends, a decision was made to partner with the CIM lab atLehigh University and apply for a state-funded grant under the Ben Franklin Partnership(BFP) program to specifYand implement a FEM-based modeling system at BethForge.Lehigh University would provide expertise in CAD, FEM, and computer science in theselection of hardware and software packages to make up the proposed system. A grantrequest was submitted to, and approved by, the Commonwealth of Pennsylvania. Workwas started on evaluating potential hardware and software components for the system.


THE BETHFORGE FEM MODELING SYSTEM (BFMS)

Initial design guidelines established by BethForge for the proposed BFMS were:* Lowest possible initial and recurring costs.* Must be usable by personnel who are not FEM/CAD/Unix experts and whose

time is not dedicated 100% to CAD/FEM efforts.* Hardware to be adaptable/usable for other applications.* Maximum leverage of existing skills wherever possible.* Maintain flexibility for future additions/changes.* Capability to do multi-pass 3D simulation of cogging/blocking, roughing, etc.

To help meet these guidelines, the following design philosophies were adopted:* Maximum use of low-cost PC-based hardware and software.* Leasing of hardware and software as opposed to outright purchase.* Modular approach to both hardware and software functionality.* Use of multi-purpose network-based peripherals.

After considerable study a 2D/3D FEM bulk deformation process modelingpackage called ANTARESTM(described in a subsequent section) was selected as thedeformation modeler. The other elements of the system were designed around itscharacteristics. An overview of the hardware configuration that BethForge and LU-CIMengineers selected for the BFMS is shown in Figure 1. Installation of the system atBethForge was completed during 2Q96.

Modeling Svstem Hardware

Unix Workstation--The primary requirements for running multiple-stroke 3D-forging simulations are a fast processor with superior floating point performance, largeamounts of RAM and ample disk swap space. A 64-bit DEC AlphaStation 600 5/266MHzwith 256MB RAM, 4MB cache, three 4.3GB hard disks, CD-ROM, internal4mm tapedrive, 21" color CAD monitor and externa11 GB removable-platter hard drive wasspecified. Its primary function is to run the FEM code. Floating-point performance israted at SP~Cfp92 > 400. This workstation is located in the Modeling Engineer's office.

PC's--Two P5-133MHz PC's with 48MB RAM, 2.1GB SCSI-II hard drive and21" monitors were specified. One is located in the Modeling Engineer's office and one inthe Forging Engineer's office. They are used to run CAD, meshing, graphics, and pre-/post-processing software. The Modeling and Forging Engineers can also logon to theBFMS workstation from their PC's using X-terminal software and run the FEM program.These PC's are also used for word processing, spreadsheets, etc.

Network Server--A high-speed, lOBase-T Ethernet network connects the PC'sand the workstation and is controlled by a PC-based server running networking software.The server handles print spooling and data transfer between the various computers.


FIG. l--Overview of the BethForge FEM Modeling System hardware.

Peripherals--SCSI-II removable-platter hard drives (1GB platter) on each of thePC's and the workstation provide warm and hot storage of model runs and facilitatetimely exchange oflarge data files with BethForge's partners. Several4mm DAT tapedrives serve as back-up for files and cold storage of old model runs. An inkjet colorprinter/copier provides contour plots and copies ofFEM simulation results while aD-sizeprinter/plotter is used for CAD hard copy of drawings.

Modeling System Software

The following sections and Table 1 descnbe the software related to the BFMS.

TABLE I--Sumrnary of software related to the BethForge FEM Modeling SystemEnvironment Software Function

Unix-workstation~ ANTARESTM Primitive 2D CAD modelingDigital Unix™ Version 3.2 2D mesh generator

Simulation database inputFEM solver (2D and 3D)Results viewing/AnimationIGES, PATRAWM, I-DEASTMoutput

XVTM Graohics conversionPentium PC's: AutoCADTMR13 2D/3D CAD modelingWindows 95 IGES Translator CAD data file conversion

FEMApTM/AMM Meshing, Pre-/Post -processingHiJaak™ PRO Graphics conversion (Unix to PC)WS_FTP95LE File Transfer Protocol (ftp)X-server Remote logon to workstation (telnet)

Network Server (PC) NetWareTM3.12 File transfer, print spooling, communicationNetWare™ 3.12


CAD ModelinQSoftware--AutoCAD R13-c4 from Autodesk, Inc. was selected asthe CAD package to create 2D tool and workpiece geometries for 2D simulations and the3D solid models required for 3D simulations. It is a full-featured program which supportscreation of2D drawings with spline contours and three types of 3D models: wire-frame,surface, and solid. In order to get high-quality meshed geometries, a solid model with ajoining operation is used. AutoCAD supports both the IOES and ACIS output formats.The IOES 2D format can be imported directly into ANTARES. The 3D IOES and ACISformated solid models are imported via a program called FEMAPTM.

Pre-ProcessinQ and Meshing--FEMAP with Advanced Meshing Module (AMM)from ESP, Inc., was selected for pre-processing and meshing of the CAD models prior toa FEM run. FEMAP is a software package designed specifically to serve as a complementto FEM solvers to enhance their ease of use. In the BFMS its main function is to generatehigh-quality triangular surface meshes for input to ANTARES, but it offers other keyfeatures. It can directly import geometries from many different CAD formats, and haspowerful "CAD-like" tools to modify geometries or generate new ones. FEM-compatiblemeshes can be generated in FEMAP by several methods ranging from manual creation tomapped meshing between key points. Appropriate materials and section properties can becreated or assigned from libraries.

The Advanced Meshing Module (AMM) is an optional extension to FEMAP. Itprovides totally automatic 2D surface meshing and 3D solid-mesh generation. AMM candirectly import and mesh ACIS solid models created by AutoCAD. It offers the capabilityto customize the mesh as it is constructed. This feature facilitates the creation of thehighest-quality FEM mesh models, an important prerequisite for good FEM simulations,especially for 3D models. Figure 2 shows a sample of a high-quality 3D mesh for a headforming blank generated in FEMAP using AMM.

FEMAP also provides a comprehensive set of tools to evaluate the mesh modeland identify errors that are often not visually obvious. The proper meshing of the die andbillet objects into discrete elements can be verified and any errors found can be correctedin FEMAP. Items checked for include: occurrences of coincident geometries, nodes, andelements; incorrect direction of the normal vector of surface elements; and impropersequencing of the element and node numbers.

FIG. 2--Segment of a forging blank with 3D mesh generated using AMM.


FEMAP provides input and output interfaces to most popular analysis programssuch as PATRANTM,IDEASTM,ANSYSTM,etc., through an appropriate translator. ThePATRAN neutral file format and the IDEAS universal file format output can be importedinto ANTARES. The balance of the pre-processing setup is then done in the inputinterface provided by ANTARES.

FEM Solver __ANTARES™ is a special purpose software package developed andmarketed by UES, Inc., for modeling 2D and 3D bulk deformation processes [14,15]which fits well with BethForge's forging operations. It is a FEM code based on the rigid-thermoviscoplastic model and can be used for coupled thermal and stress analyses on boththe workpiece and dies. ANTARES possesses material workability modeling capability[16,17,18] and is a powerful tool to model the forming process and optimize processparameters [15,19]. The software is designed around six different modules in an integratedinterface as shown in Figure 3.

FIG. 3--ANT ARES Integrated System Interface.

Salient features of the ANTARES™ package include the following:

* CAD geometry import via IGES, PATRAN, IDEAS and ARIES interfaces.* User-friendly X-windows-based pre-processor and and post-processor.* Flow stress database of common engineering materials.* Automated simulation with automatic 2D meshing and 2D/3D remeshing.* Animated display of material flow and evolution of stress, strain, temperature, etc.* Automated thermomechanical history transfer for multistage forming.* Material workability data generation.* Material point tracking for improved preform design.


During the initial set up of a simulation, the characteristics of the formingequipment used must be identified. ANTARES can handle a wide range of single andnmltiple-step processes including hydraulic and mechanical presses, hammers and rolls.ANTARES' modeling of the complex die/workpiece interaction takes into accountfrictional heating, plastic heating, die chilling and heat loss to the environment. Thecoupled thermal and stress analysis is based on the rigid-viscoplastic behavior of theworkpiece material and the elastic behavior of the tool materials and the lubricantmaterial's frictional and heat transfer properties. In conjunction with the workability mapof the material deformed, one can select the press speed and billet temperature to avoid,for example, wedge cracking in the cogging process [20].

Post-ProcessinQ--After an FEM run is completed, the results are displayed either inthe Results Interface module of ANTARES or in FEMAP. The output capabilities ofANTARES include color contour plots of many different field quantities which can bedisplayed at user-specified steps along the path of the simulation and others which can beplotted as a function of die travel. Variables available include: metal flow, stresses, strains,strain rates, temperature distributions, heat generated, etc., in the workpiece; and position,velocity, load, stresses, deformations, etc., for the dies. Polygon and circle "pointtracking" visualizations are used to follow the moment of a location within the workpiece.Postscript™ files of the contour plots can be generated and exported to any type ofdocumentation software. Screen captures are done with a graphics utility programprogram called XV™.

Results can also be output as PATRAN files which can be imported into FEMAP.FEMAP provides visualization tools to display nodal and elemental results in a contour orX-Y plot format. FEMAP also provides direct, high-quality printing and plotting of bothtext and graphics data for documentation of the modeling results. It can export both textand graphics to non-engineering programs such as spreadsheets, databases, wordprocessors, publishing, or presentation software.

Communications/File Transfer--Communication and file transfer betweenBethForge and its partners is accomplished in a variety of ways. The PC's and theworkstation have 28.8kbps modems, telecommunication software, and X-server software.PC users can telnet to the workstation, run an ANTARES session simultaneously withother users, and display the results in X-windows environment on the Pc. WS_FTP95LEsoftware, a file transfer protocol, is installed on the PC's to transfer files between the PC'sand the workstation. Exchange of large files with remote partners is done using 1GBremovable-platter HDs and 4GB DAT tape drives, all of which are available to the PC'sand the workstation.

Determination/Selection of Material Properties Used in Simulations

Knowledge of the flow strengths and heat-trans fer-related properties of materialsto be forged is an important prerequisite for good FEM-modeling simulation results. Overits history BethForge has processed a number of different alloy systems and has developed


much of this information internally. BethForge also utilizes the capabilities of its partnersfor this information. Relationships with HRL and the National Center for Excellence inMetalworking Technology (NCEMT) have proven valuable in this effort. HRL has runseveral material characterization studies for BethForge [21]. NCEMT regularly publishesAtlas of Formabilitv bulletins [22] containing stress/strain data for different materials manyof which are processed at BethForge. NCEMT and UES, Inc., also offer materialcharacterization services to the public and we expect to use these services in the future.

Tvnical Modeling Methodologv

The specific methodology for conducting an FEM simulation depends on thecomplexity of the die and workpiece geometries. For simple 2D geometries such ascylindrical upset specimens, all work is generally done totally within ANTARES. Formore complex geometries the typical modeling methodology is as described below andshown in Figure 4.

FIG. 4--Diagram of typical modeling methodologies

For 2D simulation the initial model geometry development is done in AutoCAD onone of the engineer's PC's. The model is saved in IGES file format and converted intox-y-r format by the Utilities function in ANTARES. The x-y-r geometry file is then readinto ANTARES through the Input Interface.

For 3D simulation the initial solid model construction is also done in AutoCAD,but saved as ACIS file format. The ACIS solid model file is imported into FEMAP andmeshed with AMM. The meshed geometry is then translated to either PATRAN orIDEAS file format and imported into ANTARES. Input of the simulation parameters,running of the simulation, and initial visualization of the results are all carried out inANTARES. Results are transferred back to FEMAP if additional post-processing isrequired.


APPLICA nONS AND CASE STUDIES

Some specific examples of use of the BethForge FEM Modeling System forproblem solving and/or new product development within BethForge are as follows:

2D Axisymmetric Simulations

Head forrning--This example illustrates how the BFMS was used to help developan improved forging blank design for a formed-head closure. Figure 5 shows section viewsof the original and improved workpiece configurations for a forging blank.

ORIGINAL BLANK DESIGN IMPROVED BLANK DESIGN

FIG. 5--Section view of axisymmetric forging blank and tooling.a) original blank design. b) improved blank design.

The modeling methodology used was as follows: CAD models of the tool set (topdie and bottom die) and the first iteration of the forging blank, Figure 5a, were generatedusing AutoCAD. The blank and tool geometries were imported via the IGES translatorinto ANTARES. A 2D-axisymmetric simulation was possible due to the symmetry of thedies and workpiece. Hydraulic press parameters, materials for the objects, the lubricant tobe used and the process parameters such as the initial temperature conditions as well ascontact boundary conditions were assigned to the blank and dies. The 2D-axisymmetricgeometries were meshed in ANTARES as shown in Figure 6a and submitted to theANTARES 2D-axisymmetric FEM solver. The results of the simulation, Figure 6c, werecompared to the desired final contour shown in Figure 7. An area ofunderfill waspredicted by the simulation results shown in Figure 6c. A second iteration, an improvedblank design as shown Figures 5b and 6b, was designed with AutoCAD. Additional blankthickness was added in the underfilled area and the overall diameter on the blank wasreduced. The die design was not modified. The simulation was then repeated with theimproved blank geometry. The results are shown in Figure 6d.


c) d)FIG. 6--Mesh contours of simulation results.

a) at initial step for original blank design. b) at initial step for improved blank design.c) at last step for original blank design. d) at last step for improved blank design.

FIG. 7--Results comparison of predicted contours and desired final contour.


Figure 7 illustrates the trend predicted by the FEM model for the new blank vs. theoriginal design. This trend agreed with the real-world results when field measurementswere taken on the actual forgings, i.e., the improved blank design resulted in an as-forgedworkpiece more dimensionally suitable for heat treatment and final machining. The area ofunderfill was eliminated and the cut-off waste on the top of the circumference was greatlyreduced.

Numerous other 2D axisymmetric simulations have been conducted as part ofBethForge's process/product development efforts. A brief overview of some of this workis as follows:

Punch forming with flattening--Figure 8 shows a case study of a two-stage,punch-then-flatten operation. In this study, FEM analysis was conducted to help predictthe severity of an edge curl-up condition (Figure 8b) that was expected to occur in thepunch stage and the ability of a subsequent flattening operation (Figure 8c) to correct thecurl-up. The trends predicted by the model agreed well with those seen in real-worldproduction runs of similar products. A deformation strategy was planned and successfullyexecuted in production.

BEFORE PUNCH AFTER PUNCH FLATTENING

FIG. 8--Mesh contours of punch forming with flattening.a) before punch. b) after punch. c) flattening.

Ingot upset die design--Figure 9 illustrates the FEM modeling of an upsetoperation for a large ingot. Modeling runs were carried out to assess the effect of varyingthe ingot "chuck" design relative to the ring die design on void consolidation in the "deadzone" of the ingot. The study suggested areas for potential improvement so fieldevaluations of a modified design have been initiated. Preliminary results are encouraging,although more data are required. Additional field evaluations are planned.


a) b)FIG. 9--Simulation results of upset. a) before upset. b) after upset.

Back Extrusion-- Figure 10 illustrates simulation of a back extrusion operation.This work is part of an ongoing new product development study.

a) b)FIG. lO--Simulation results of back-extrusion. a) before extrusion. b) after extrusion.


3D Non-Axisymmetric Simulations---In¥ot Co¥¥in¥

Figure 11 shows an area that we are actively pursuing, i.e., 3D simulationstrategies for ingot cogging processes. This type of process is used to forge an ingot into arotor or a roll. This application is difficult to model because it consists of multipledeformations along the body of the ingot with the ingot being rotated and/or translatedafter each stroke of the press.

In a simulation of this nature the workpiece deformation and stresses are ofprimary interest, and not the dies. Using CAD, a solid model for the workpiece and skinmodels for the dies are built. FEMAP/AMM is used to discretize the boundary surfacesusing triangles. The resulting initial mesh is shown in Figure lla. The workpiece and diesurface meshes are imported into the pre-processor of ANTARES™. The boundaryconditions, mesh specifications, and simulation control parameters are assigned. Specialkeyword files are written to disk to specify the details of the 3D simulation.

In this simulation there are four deformation strokes along the length of the ingotin each pass. After finishing one pass down the body of the ingot, the ingot is rotated 90degrees or 45 degrees around its longitudinal axis and translated back to the originalposition. This sequence is repeated again and again until the necessary passes have beencompleted. During simulation of this type offorging operation the FEM results of eachpress stroke are automatically set up to be transferred as input to the next run. Figure 11bshows the simulation results of the first fifteen deformation strokes of the press.

FIG. Il--Simulation results onD rotor forging. a) before rotor forging. b) third strokeof the fourth pass of rotor forging (15th deformation stroke).

We have made good progress in simulating this type of operation, and hope tosoonbe able to assess the effect of various drafting strategies on void consolidation duringingot cogging.


SUMMARY/CONCLUSION

BethForge and its partners have developed and implemented a low-cost, FEM-based computer system for modeling open-die forging, and have succesfully applied it toreal-world problems. Based on the results achieved to date, it is concluded that FEMmodeling is a cost-effective and worthwhile tool for forging product and processdevelopment.

FUTURE DIRECTION FOR FEM MODELING EFFORTS

Based on the initial successes described above, we plan to continue work on 3Dcogging applications and expand the range offorging-related issues that can be addressedusing ANTARES. We also plan to extend the use of2D/3D FEM-based modeling to otherareas, including ingot teeming. We are evaluating ProCAST™ [23,24], a 2D/3D FEMpackage for the modeling of casting processes.

We are confident that in the not-too-distant future all new BethForge processesand products will be modeled in some way on the computer prior to their implementation.

ACKNOWLEDGMENTS

The authors are grateful to Bethlehem Steel Corporation and BethForgemanagement for their support, and for permission to publish this paper. In addition,thanks are due to the Commonwealth of Pennsylvania for their financial support of thiswork through the Ben Franklin Partnership Program.

REFERENCES

[1] R. L. Bodnar, D. C. Ronemus, B. L. Bramfitt, and D. C. Shah, "Physical Modeling of Hot-Deformation Processes - Using Plasticine," Iron and Steelmaker, Vol. 13, No.8, August1986, pp. 35-46.

[2] R.C. VanKuren, et ai, "Lead and Plasticine Modeling of Blooming Mill Operation," internalreport, Bethlehem Steel Research Department, May 18,1987.

[3] H. Ortiz and M.M. Vyas, "Physical Modeling Technology to Study Hot DeformationProblems," 2nd IntI. Conf. on Appl. of Mathematical and Physical Models in the Iron andSteellndustI)', April 1992.

[4] S. Shida, et aI, "Simulation of the Hot Rolling of Steel using Lead," 1. Janan Society forTechnology ofPlasticitv. Vol 9, 1968.

[5] H. Tsukamoto, et aI, "Simulative Model Test on Metal Forming using Plasticine as a ModelMaterial," SAE Technical Paner MF 74-179, 1974.


[6] E. Erman and D. C. Ronemus, "Visits to Japanese Steel Companies and 3rd ICTP," InternalBethlehem Steel Research Department Report, March 8, 1991

[7] E. Erman, N. M. Medei, A. R. Roesch, and D. C. Shah, "Physical Modeling of the BlockingProcess in Open-Die Press Forging," J. Mechanical Working Technology, Vol. 19,1989,pp. 165-194.

[8] E. Erman, N. M. Medei, A. R. Roesch, and D. C. Shah, "Physical Modeling of the UpsettingProcess in Open-Die Press Forging," 1. Mechanical Working Technology, Vol. 19, 1989,pp. 195-210.

[9] E. Erman and D. C. Shah, "Development of Forging Practices for Economical Production ofSound Heavy Forgings," 30th Mechanical Working and Steel Processing ConferenceProceedings, ISS-AIME, Vol. XXVI, 1989, pp. 239-253.

[10] N. M. Medei, K. A. Taylor, J. E. Fielding, and E. Erman, "The Manufacture of LargeAlloy 600 Contour-Forged Discs," 29th Mechanical Working and Steel ProcessingConference Proceedings, ISS-AIME, Vol. XXV, 1988, pp. 225-233.

[11] E. Erman, N. M. Medei, S. J. Plesko, and D. C. Shah, "Development ofInnovative Open-DiePress Forging Techniques for Pressure Containment Components," 29th MechanicalWorking and Steel Processing Conference Proceedings, ISS-AIME, Vol. XXV, 1988,pp.205-214.

[12] A. M. Basirco, "DFLUX n A General Purpose One Dimensional Heat Transfer Program"Internal Bethlehem Steel Corporation Report, March 2, 1979.

[13] C.J. Van Tyne, R.B. Focht, T.D. Nelson, and W. Reese, "Modeling of Heating Cycles ForLarge Forgings Using A PC-Based Program," 35th Mechanical Working and SteelProcessing Conference Proceedings, ISS-AIME, Vol. XXXI, 1994, pp. 217-222.

[14] "ANTARES Primer Manual," D.E.S. Inc. 4401 Dayton Xenia Road, Dayton, Ohio 45432.

[15] T. Veerabadran, N. Hattangady, and S. Manna, "Simulation Based Design of MetalForming--a New Way to Profit," Proceedings of Forge India'95, April 1995, pp B2/l-8.

[16] Y. Prasad, H. Gegel, et aI, "Modeling of Dynamic Material Behavior in Hot Deformationand Forging ofTi-6242," Metall. Trans, 15A, 1983.

[17] H. Gegel, J. Malas, et aI, "Materials Modeling and Intrinsic Workability," Int'l Conf. onTech. of Plasticity, Stuttgart, 1987.

[18] K. P. Rao, M. Hua, and w.L. Xu, "Integrated Computer Modeling and Simulations of MetalForming Processes," Research Report #70012, City University of Hong Kong, March, 1995.

[19] W.L. Xu, and K. P. Rao, "Role of Computer Simulation Tool in the Design of Forging,"Proceedings of Forge India'95, April, 1995, pp B3/l-8.

[20] H. Gegel, and T. Veerabadran, unpublished research conducted by UES, Inc. for Wright-Patterson Air Force Base, June 1996.


[21] R. L. Bodnar, and R.E. Steigerwalt, "High-Temperature Stress-Strain Data for HeavyForging Materials," Internal Bethlehem Steel Corporation Research Report, May 30, 1986.

[22] W. L. Moore and P.K. Chaudhury, "Atlas of Formability Data for Hot DeformationSimulation," presentation at Rapid Materials Forming Technology Conference, Worcester,MA, October, 1994.

[23] J.N. Pennington, "The Die Casting Engineers' Virtual New World," Modem Metals. May1995, pp 58B.

[24] T. Veerabadran, S. Manna and H. Gegel, "Simulation Based Process Design ofCasting-ltPays Off," IIF Transactions 1995. Proceedings of the 43rd Indian Foundry Congress.3-5February 1995, Jamahedpur, The Institute ofIndian Foundrymen.

John E. steiner1 and Edward L. Murphy1

A RARE TYPE OF FLAKE-LIKE FORGING BURST IN HEAVY FORGINGS

REFERENCES: steiner, J.E., and Murphy, E.L. "A Rare Type of Flake-LikeForging Burst in Heavy Forgings", steel Forqinqs: second volume. ASTM STP1259, E.G. Nisbett and A.H. Melilli, Eds., American society for Testingand Materials, 1997.

ABSTRACT: Forging bursts are infrequently reported in the publishedliterature. One type of forging burst, which has been observed only onfour occasions in the combined experience of the authors, is described.The similarity of this defect with hydrogen-caused thermal flakes isremarkable. It is characterized by a fracture appearance similar to thatof flaking when broken open, by orientation and distribution similar toflaking, but within a specific section of a large forging, generally inultrasonic test response and by its good response to healing duringreforging, as with flaking.

Case histories of two 3.5Ni-Cr-Mo-V forgings, one 1Cr-1Mo-O.25Vforging and one martensitic stainless steel forging are presented, andprocessing factors related to these bursts and means for avoidance of theproblem are discussed. A fifth case of possible misidentification ofthis phenomenon as hydrogen flakes in a small gear component forging isdiscussed. Approaches to differentiation of this defect from flaking arediscussed.

KEYWORDS: forge bursts, flake-like defect, hydrogen flakes, ultrasonictesting, large forgings, forging too cold.

Flaking has a characteristic pattern, Figure 1. It seeks thecentral zones of a heavy section of a forged or rolled piece. Manyexceptions to this rule have been observed. They may be related to suchfactors as non-uniform cooling, especially of long pieces in safety heattreatments, and to favoring product from the top end of the ingot wherepositive segregation of hydrogen occurs. Beyond such explainableexceptions, a small family of flake-like defects in heavy forgings hasbeen observed. These are attributed to a rather uncommon if not raretype of forge bursting. Initially, these bursts are discussed on a casehistory basis.

1 consultant, Heavy Forgings, Engineering Materials & Processes, Inc.,121 Edgewood Avenue, pittsburgh, Pennsylvania 15218

241


FIG. 1--Distribution of ultrasonic Indications in aHeavily Flaked Forging.(Metals Progress, July, 1959

reprinted with permission)

CASE I A469 cl6 - Ni-Cr-Mo-V generator rotor.

This generator rotor, produced in the late 1970's, is unique inthat it represents the authors' first encounter with this phenomenon.That is, the first instance where the flake-like defect was judged to bebursts and not flakes. Unfortunately, no documentation of the case isavailable. This history is largely anecdotal.

The forging was a large 2 pole generator rotor about 1100 mm (44inches) in diameter as forged. The turbine end flange represented the topof the original ingot and the collector end (no flange), the bottom ofthe ingot.

After forging, preliminary heat treatment and machining, ultrasonicinspection revealed a large number of traveling flake-like indications inone diameter-length of the collector end, Figure 2.

FIG. 2--Zone of ultrasonic indications in Case I rotor.

The entire rotor was otherwise free of reportable ultrasonicindications.

STEINER AND MURPHY ON HEAVY FORGINGS 243

There were no irregularities in the melting or degassing of theingot or safety treatment of the forging. The steel was presumed tocontain 1.0 ppm hydrogen. Furthermore, the indications were in steelrepresenting the bottom portion of the ingot where hydrogen, if present,would be the lowest.

Examination of cores drilled into the affected section revealedsmall flake-like cracks approximately 6 to 12 rom (0.25 to 0.5 inch) indiameter. These cracks exhibited the transgranular fracture appearancecharacteristic of hydrogen flakes.

Examination of the forging records of this rotor suggested that theaffected set-down diameter had been subjected to forging work in thelatter stages of the final forging heat. The usual practice was tofinish the journal ends and make a final pass over the entire body. Atthat time, if final measurements showed that a bit more overall lengthwas needed, it is conceivable that the affected section was forged downfurther to achieve such length. Such work would have been done afterconsiderable time had expired at the press and consequently, when thesmall diameter journal would have been at a relatively low temperature.This was a distinct possibility, because it was the practice not toreheat after finishing the body of the rotor forging.

Although the defects responded like flakes to ultrasonic probing,had a distribution within the affected journal section characteristic offlaking and exhibited all the characteristics of flakes uponmetallographic examination, it was concluded that they were not flakes.This conclusion was based upon (1) the fact that the steel had beendegassed to a level of about 1 ppm hydrogen, and would have beenimpossible to flake even upon air cooling from the press in this smallsection, (2) the fact that the pattern of distribution within the overallbulk of the forging was not characteristic of flaking, (3) the fact thatthe location of the small diameter affected area was between two otherrelatively small diameter areas each of which was unaffected, and (4) theobservation that the forging combinations described above stronglysuggest that some degree of forge work was locally employed attemperatures below the acceptable forging temperature range.

CASE II - A470 cIS - Cr-Mo-V turbine rotor (lCr-Mo-V)

Processing records on this rotor represent the most completedocumentation on the phenomenon of flake-like forge bursts. It isparticularly interesting because the lCr-Mo-V steel does not flake.

The forging was produced as a part of the Electric Power ResearchInstitute (EPRI) project on Advanced Rotor Forgings for High-TemperatureSteam Turbines [2]. The forge-crack problem had no relation to thetechnical objectives or results of the EPRI project and was reported byEPRI solely to complete the record of the processing of the forging.


The forging was an HP turbine rotor with a main body of 1260 mmdiameter as forged, Figure 3.

FIG. 3--zone of ultrasonic indications in Case II rotor.(Courtesy of EPRI, reprinted with permission)

The rotor was forged from a 2300 mm diameter ESR ingot. Hydrogencontent was 1.3 ppm in the top portion of the ingot and 1.9 ppm in thebottom. The integrally forged test ends were required for the EPRIproject. The affected journal between the body and the bottom test endwas forged to 815 mm diameter. ultrasonic testing after preliminary heattreatment, because of a strong noise level, revealed no indications inthis area. After machining to 539 mm diameter and quality heattreatment, ultrasonic testing revealed a number of individual sonicindications around the axis of the bottom journal and in the adjacentregion of the adjoining extra test material. The remaining part of therotor showed no recordable sonic indications. Due to these sonicindications, a reforging of the bottom journal became necessary.

The extra test materials were removed from the rotor beforereforging of the bottom journal. specimens for metallographicexamination were taken from the adjacent region of the bottom testmaterial containing sonic indications, Figure 3. Examination of thesespecimens under an optical microscope and under a scanning electronmicroscope showed that the sonic indications were due to small cracks.The surface of the cracks was very smooth and free from any non-metallicinclusions or segregation. From the location and from appearance ofthese small cracks it was concluded that they were formed due to veryunfavorable forging conditions of the bottom journal and of the adjoiningarea in the extra test material during the finish forging operation.

Although the defect, as in CASE I above, responded like flakes toultrasonic probing, had a distribution within the affected journalsection characteristic of flaking and exhibited all the characteristicsof flakes upon metallographic examination, it was concluded that theywere not flakes. This conclusion was based upon (1) experience that the1Cr-Mo-V rotor steel does not flake, (2) the fact that the pattern ofdistribution within the overall bulk of the forging was notcharacteristic of flaking, (3) the fact that adjacent larger diametersections were unaffected, with the exception of a short length of the test


end section and, (4) that the low forging temperatures, for the finishingof this last-to-be-forged journal suggest that considerable forge workwas done below the optimum forging range, Figure 4.

FIG. 4--Forging sequence of the ESR Ingot.(courtesy of EPRI, reprinted with permission.)


CASE III A 470 cl7 Ni-cr-Mo-V flywheel forging.

This forging was a 70 ton piece that, because of its large sizeand unwieldy shape, Figure 5, was difficult to handle at the forgepress.

FIG. 5--Flywheel Forging and affected zone.

ultrasonic testing after final heat treatment revealed numerousindications in the bottom journal. There were indications in the topjournal, but considerably fewer. The main body of the wheel and theinboard, adjacent-to-the-wheel portions of each journal were free ofindications. The hydrogen content of the forging as determined byclosely monitored deep seated sampling and testing from both journals wasbelow 1 ppm. Metallographic examination revealed close flake-likecracks.

Although, as in cases I and II above, the defects appeared to beflakes, for similar reasons, they were judged, in this case, not to be.The very low hydrogen content, the fact that the main body was clear ofall ultrasonic indications and the probability of extremely lowtemperature forging of the bottom journal strongly suggest forge bursts.The journal was hot worked at estimated 650C surface and 1050c centertemperatures and was at 520c surface temperature at finish.

AS to cause of the defects, there was agreement among theinvestigator~ that they were not flakes. However, one party held thatthey were the result of incipient faceting or overheating. The authorsconcluded that they were forge cracks.

CASE IV Type 420 stainless steel shaft

This shaft was forged straight down to a long, approximately 330 romdiameter piece. Toward the center of the shaft, a set-down to about 260rom over a length of about 230 rom was forged. Figure 6. ultrasonictesting of this set down revealed numerous small traveling indicationsmarkedly similar in response to hydrogen flakes.


FIG. 6--Shaft Forging and Affected Set-down Zone

There was an identical set-down some distance down the shaft thatexhibited a long axial single defect ultrasonic pattern. This is typicalof the initial stage of a Mannesmann-effect type forge burst. There wereno other ultrasonic indications over the full length of the shaft.

Again, for much the same reasons cited in the preceding cases, thedefects observed were not flakes. The ultrasonic indications of theaffected area in Figure 6 were the flake-like forge bursts that are thesubject of this paper.

Although the forging sequence and drafting practice of the twoset-down sections are not known, it is interesting to speculate on theorigin of the defects. The piece was finished with round dies, top andbottom. It is probable that the Mannesmann-type defect section wasforged first with rolling action under the dies. This, because theMannesmann-type defect is not temperature dependent. It can be producedin properly heated steel. It is probable that the piece was turned endfor end under the dies and the affected set-down forged at a lowertemperature, possibly much lower.

This steel does not flake, and even if it did, the short set-downsection could not have flaked in preference to the larger diametersections on either side.

DISCUSSION

The phenomenon of flake-like forge bursts has probably been with usfor a long time. It is rare. The cases presented chronologically inthis paper occurred over a period of twenty years, a period during whichthe authors have been active in flaking and hydrogen related problems andhave to their knowledge seen no others. No others, that is, until thecase of a "flake~" gear component forging described by Gensure, Melilliand Brueggemann [3] recently came to our attention. Although we have nobackground processing data on the forging from the immersion ultrasonicinspection results shown in Figure 7 (Gensure et aI, Figure 17), it seemsobvious that the defects at the journal position 9 are not flakes, butforge bursts. The defects at 3 through 6, where one might expect flakes,are identified as non-metallics. This example fits the pattern of thecited cases.


FIG. 7--Locations of ultrasonic indications in gear component forging.(Copyright ASTM, reprinted with permission)

Forging die configuration does not affect this phenomenon. Theabove cases included flat, Vee and round dies.

To forge a large rotor, for example, a large heavy press isrequired. with the accepted practice of not reheating after the body isfinished, a situation is created where a small journal section, which hascooled faster, may be forged when almost black only because the heavypress is available to do the stiff work. This situation is exacerbatedin shops that traditionally forge on the cold side or where transfers andhandling at the press are slow due to facility limitations orunforeseen emergencies.

In summary, the results suggest that care must be taken to forge asmall diameter section of a large forging at a reasonably high hotworking temperature. As noted, the occurrence of small flake-like forgebursts is probably quite rare. This paper, although rather heavilyanecdotal, hopefully will contribute to a better interpretation of theorigin of unusually "flaked" forgings and stimulate publication of otherexamples or even refutation of this premise.

REFERENCES

[1] Steiner, J.E., "Hydrogen in Heavy Forgings", Metals Progress,(ASM) Metals Progress, Vol 76, No.1, July, 1959.

[2] Swaminathan, V.P., steiner, J.E. and Mitchell, A.,"AdvancedRotor-Forgings for High-Temperature Steam Turbines"EPRI CS-4516 Volume 1, project 1343-1 Final Report, May 1986

[3] Gensure, J.G., Melilli, A.S., and Brueggemann, S.M.,"High Sensitivity, Immersion, ultrasonic Testing of steel Forgings"ASTM STP903, E.G. Nisbett and A.S. Melilli, American Society forTesting and Materials, philadelphia, 1986, pp 553-572

Yunrui Lil, Keyi Yang2, and Zhenlin Ne

A NEW MODEL FOR CALCULATING MAXIMUM BLOWFORCE OF DIE-FORGING HAMMER

REFERENCE: Li, Y.,Yang, K., and Ni, Z., "A New Model for Calculating Maximum BlowForce of Die-Forging Hammer", Steel Forgings: Second Volume, ASTM STP 1259, E.G.Nisbettand A.S.Meli11i,Eds., American Society for Testing and Materials, 1997.

ABSTRACT The blow procedure of a die-forging hammer is analysed theoretically takingaccount of the elastic deformation of the cushion under the anvil block, and the new formulae forcalculating the maximum blow force are deduced from a new model proposed in the paper. Theresults calculated by the formulae agree well with experimental data. Thus, the formulae and theresults provide a basis for designing and using die-forging hammers properly.

KEY WORDS Die-forging hammer Blow force Cushion

INTRODUCTION

The blow force of a die-forging hammer is an important technology parameter in die-forgingproduction. Exactly calculating its value is significant for determining plastic deformation process,improving the quality of the workpiece and protecting the main components of the hammer.Therefore, it has been a problem which people have been deeply interested in, but has not beensolved satisfactorily up to now.

The conventional formulae [1-3] are just used to evaluate the average blow force of a die-forging hammer, which can be used for upsetting or preforging a normal hot workpiece. However,they are not suitable to fmish forging and precise trim.The reason is that the elastic deformation ofthe cushion under the anvil, which will affect the blow procedure, has not been considered in theseformulae. When upsetting and preforging, the deformation value of workpiece is one order ofmagnitude greater than that of the cushion. So the cushion deformation may be ignored. But to fmish

1 Professor, 2 Doctor student, DepartmentofMaterials Science and Engineering, NorthwesternPolytechnical University, Xian, Shaanxi, 710072, P.R.China.

249


forging and precise trim, the deformation of the workpiece is quite little, and the value is of the sameorder of magnitude as the cushion deformation, or even more less. In the case, the deformation ofthe cushion is a important forging parameter affecting the blow procedure and can not be neglected[4.5] • Therefore, researching the effect of the cushion deformation on the blow procedure is necessaryfor determining the maximum blow force exactly.

In this paper, the blow procedure is analysed theoretically in consideration of the elasticdeformation of the cushion, and the formula of the maximum blow force is deduced from a newmodel presented in this paper. The results from calculating examples based on the formulae are ingood agreement with experimental data. Thus, the formulae and the results provide a basis fordesigning hammers and making forging processes.

THE DERIVATION OF THE MAXIMUM BLOW FORCE OF DIE-FORGING HAMMER


TEST OF THE MAXIMUM BLOW FORCE

The maximum blow force of typical hammers are calculated from Eqs (19), (20) and (21), andthe results are shown in Tab.3. The measured values are from Ref.[7 ] and from the experimentalresults tested on the I-ton hammer by the authors. The errors between calculated results andexperimental ones are about ten percent. Thus, the calculated data can meet the engineering need.The reason that the measured vaiues of the blow forces are low is that the blow velocity of hammerstested is lower than normal. For examples, the tested velocity of I-ton hammer is 5.2-5.7m/s, andfor 10-ton hammer, 6.S-7.Im/s. According to these data, the errors are 4.6-12.6 percent and 6.1-2percent respectively to the two kinds of hammer.

LI ET AL. ON A NEW MODEL 255

CONCLUSIONS

(1) Taking account of the elastic deformation of the cushion, the formula of maximum blowforce of die-forging hammer is deduced.

Turbine and Generator Forgings

Karl H. Schonfeld,1 Ralf Levacher,1 Michael P. Manning, 2 Paul F. Murley2

MARTENSITIC 11% CrMoNiNb STEEl FOR TURBINE ROTORS IN GEOTHERMALPOWER STATIONS

REFERENCE: SchOnfeld, K. H.• Levacher, R., Manning, M. P., Murley, P. F.,'Manensitic 11 % CrMoNiNb Steel for Turbine Rotors in Geothermal Power Stations,'Steel Forqinqs: Second Volume, ASTM STP 1259, E.G. Nisbett and A.S. MeUIIi,Eds.,American Society for Testing and Materials, 1997

ABSTRACT: Turbine rotors in a geothermal power station were required in high alloyed12% Cr steel. After some preliminary investigations it was found that a modified11% CrMoNiNb steel will fulfil the required mechanical properties as well as asufficient resistance to corrosion.Two LP rotors with appro 1 295 mm (51 in.) pre machined diameter and ungashed weightof appro 31 mt were manufactured. The steel was melted in a 125 mt electric arc fumaceand subsequently remelted into a 113 mt ESR ingot with 2 300 mm (90 in.) diameter.After forging and preliminary heat treatment the rotors were quality heat treated to a yieldstrength (0.2% and 0.02%) of appro600 MPa (87 ksi) and appro 550 'MPa (80 ksQrespectively and tensile strength of appro 780 MPa (113 ksi). This resulted in a FAIT ofappro 16°C (60°F). Low hardness is important to susceptibility to stress corrosion. It wasachieved to appro 20 HRC.The results have met the assumption with respect to this steel for application as rotormaterial for geothermal power stations.

KEYWORDS: martensitic stainless CrMoNiNb steel, turbine rotor, geothermal powerstations, low temperature application

INTRODUCTION

Two, five stage, double flow, bored LP rotors (1 295 mm/51 in. dia.)were ordered to replace existing Ni-Mo-V steel rotors in a geothermal power plant.The replacement rotors were requested to be fabricated from 12% chromium stainlesssteel for resistance to corrosion in the geothermal steam environment. The rotormaterial was required to have the following chemical composition and mechanicalproperties described in the tables below (Tables 1 + 2). These figures show the originalmaterial as well as the selected replacement material.

1 Manager and engineer, respectively, Quality and Materials Engineering,Saarschmiede GmbH Freiformschmiede, 0-66330 Volklingen, Germany

2 Lead Engineers, Structural Materials Engineering, General Electric Company,One River Road, Schenectady, NY 12345

259


element original requested selectedreplacement replacementl

material !

0.14-0.200.15-0.300.50-0.80 i0.015 max.0.012 max.10.5-11.2510.90-1.100.60 max.0.05-0.120.15-0.250.04-0.08

TABLE 1--0riginal, requested and selected material.

Tensile strength 621

26Impact, CVN 34 (25) (30)

FATT

TABLE 2- -Mechanical properties of original, requested and selected material.

SELECTION OF STEEL TYPE AND SPECIFIED PROPERTIES

There was significant experience using straight 12% Cr steel rotors for highpressure rotors in mechanical drive turbines. However, the 1 295 mm (51 in.) barrel diameterof the subject replacement rotor was significantly beyond the body diameters of rotorsmade previously in straight a 12% Cr material. For this reason, a material cooling ratesimulation test was performed to determine the expected bore properties with thelarger barrel diameters.

The heat treat study was run to simulate the cooling rate achieved at thecenter of 1 270 mm (50 in.) diameter forging during water quench from the austenitizingtemperature. The results of the heat treat simulation study showed the straight 12% Cr steelwould not have sufficient through hardening in a 1 270 mm (50 in.) diameter rotor forging.The failure to achieve full martensitic transformation resulted in a bore yield strength(0.02% offset) of approximately 345 MPa (50 ksi). Therefore, this material was notconsidered acceptable for the Geysers application.

Alternate 12% Cr steels such as standard soft martensitic steels,X5CrNiMo 16-5 and X5CrNi 13-4 (F6NM) were considered. After a thorough search offorging experience with these steels, it was decided to produce a lower strength versionof standard 10% chromium HP rotor material. There was significant experience with amodified 12% Cr steel for rotors in the 1015-1270 mm (40-50 in.) diameter range.

SCHONFELD ET AL. ON CrMoNiNb STEEL 261

However, it is typically processed to a 895 MPa (130 ksi) strength level [1].GE felt very confident that this material would be suitable for the

geothermal application. The main advantage of this material over straight 12% Cr steelswas its ability to achieve good deep seated microstructure and associated mechanicalproperties at the centerline of the rotor.

The chemical composition was modified. Molybdenum (0.35% max. to0.90-1.10%) and vanadium are added to increase hardenability. To compensate for thehigher levels of hardening elements, the chromium level has been adjusted downward(upper limit 11.25%) to minimize the formation of ferrite. Nickel levels also had to bealtered slightly (0.60% max.) to balance the composition.

The minimum Charpy V notch energy was set at 20 J (15 ft-Ibs) versusthe requested 41 J (30 ft-Ibs). This was based on experience with this material insmaller sizes at the 862 MPa (125 ksi) strength level. In these smaller rotors the CVNenergy ranged from 28 to 58 J (21 -43 ft-Ibs) with an average of 41 J (30 ft-Ibs). Due tothe large size of these rotors and the requirement for bore testing an extrapolatedminimum of 20 J (15 ft-Ibs) was set.

The rotor FATT was specified as +65°C (+ 150°F) maximum. This is incontrast to no requirements defined by the customer and the expected value of less than30°C (100°F) for the production rotor. For information, normal HP 1O%Cr steel rotors havea maximum FATT of +65°C (+150°F) after oil quenching as well as CrMoV steel rotors+150°C (+250°F) after air hardening. These are the limiting parameters for startup. Eventhough the maximum specified FATT for the geothermal rotor is defined as +65°C (+150°F),it was expected that the actual FATT would be substantially lower.

The requirements for hardness was set at 26 HRC maximum. Limits of18-24 HRC were requested to minimize susceptability to stress corrosion. The heattreat studies showed a variation of hardness from 21-29 HRc, with the higher valuescoming from samples with tensile strengths above 827 MPa (120 ksQ level.

PRELIMINARY INVESTIGATIONS

For verification of tensile properties, the cooling rate study was repeated andshowed that bore yield strengths of 552-620 MPa (80-90 ksi) could be achieved dependingto the tempering cycle. Various tempering cycles were studied to evaluate balance of tensilestrength/hardness and impact.

Chemical composition of testing material is shown in Table 3.

C ~n P S Cr,

0.19 0.68 0.0111 0.001 10.96I

TABLE3- -Chemical composition of testing material.

Samples were heat treated under conditions simulating the core of a 1 500 mm(59 in.) diameter rotor. Hardening temperature was choosen to 1 050°C (1 920°F) withcooling simulating oil cooling. Tempering was performed between 650 and 750°C (1 200 -1 380°F).The mechanical properties of tensile and impact tests are concluded in Fig. 1aand 1b. It shows that needed minimum yield strength, maximum tensile strength andhardness will be achieved after tempering in the range of 725°C (1 340°F).



MANUFACTURE OF GEYSER ROTOR

Saarschmiede has experience for many years in manufacturing very largeforgings also in high alloyed steels. Lots of rotors and discs in 9-12% Chromium steels aremanufactured, most of them for high temperature service [2] but also for low temperatureservice, e. g. geothermal application [3]. These large forgings are made in ESR quality withingot sizes up to 2 300 mm (90 in.) diameter and weight over 100 mt.

For the new LP rotors the steel was melted in a 120 mt electric arc fumace,cast into electrodes and subsequently remelted into a 113 mt ESR ingot with 2 300 mm(90 in.) diameter. Remelting of the ingot was performed in a 160 mt ESR plant with ingotretracting device.

The chemical composition of ESR ingot (heat analysis) is shown in Table 4.

C 5i Mn P S Cr Mo Ni Nb V N

0.18 0.19 0.68 0.010 0.001 10.97 0.92 0.59 0.061 0.19 0.061

TABLE 4--Chemical composition (heat analysis) of ESR melt.

After heating up to forging temperature the soaked ESR ingot was forged on a 60 MNforging press. The forging sequence is outlined in Fig. 2: stretching to 1 600 mm (63 in.)octogonal cross section, upsetting to 2200 mm (87 in.) and further stretching to finalforging shape with a total stretch reduction of 5.3.


After forging the rotor was transformed in the pearlite phase at 760·C (1 400·F). Before qualityheat treatment the forging was premachined. The heat treatment contour is shown in Fig. 3.

FIG. 3--Heat treatment contour of LP rotor.

The rotor was vertically austenitized at 1 OSO·C(1 920·F) and hardened in oil.Tempering was performed at 72S·C (1 340·F) with air cooling.

TEST RESULTS

Test specimens were taken from the body (radial) and prolongation of journal(tangential) as well as from axial core to determine product analysis and the mechanicalproperties. Fig. 4 shows the test locations of the rotor.

FIG. 4- - Test locations for LP rotor.

The product analysis is shown in Table S.


Test I C I Si I Mn P S I Cr Mo Ni : Nb V Nloc. I

I

heat 0.18 0.19 0.68 0.010 0.001110.97 0.92 0.59 10.061 0.19 0.061I

GEC 0.18 0.20 0.68 0.009 0.001110.90 0.911 0.5910.058 0.19 0.06BC 0.18 0.18 0.68 0.009 0.001 10.90 0.91 0.57 10.058 0.19 0.06

'I TE , 0.18 I 0.19 I 0.68 0.009 0.001110.9010.9210.5810.0610.19 0.06

TABLE 5--Chemical composition (product analysis) of LP rotor.

The mechanical properties are summarized in Table 6. All expected, neededand specified properties are achieved. Significant is that the mechanical properties are veryhomogenous over cross section. Radial body values as well as axial core values are within asmall scatterband. Radial tensile strength was determined to 762-783 MPa (11 0.5-113.5 ks~while core tensile strength was 769-800 MPa (111.5-116 ks~. Yield strength (0.02%) wasthen 521-541 MPa (75.5-78.5 ksi) respectively 538-572 MPa (78-83 ksi). At this levelreduction of area was appro 55% and impact Charpy V appro 70 J (50 ft-Ibs).

The specified FATT as +65°C (+150°F) was really substantially lower asexpected. It was determined to +16°C (+60°F).

Hardness HRC was with 19.7-21.3 on the lower range of specification and soon the safe side to stress corrosion properties.

Test ~~st IV.S IV'S' IT.S. EL ICVN FATTIHRCloc.

Idlr. 10.02% 0.2% I I

, :ksi I ksi ksi %A I radial! 76.5 186 110.5 24 119.7

I

B radial 178.5 189.5 113.5 21.5 20.7 I

C I radial ,75.5 '88.5 110.5 25 20.0AC transv J 78 86 111.5 20.5 20.2

IBC i transv~ 79.5: 87.5 113 20 20.8 I

I CC ,transv J 83 91 1116 118.5 20.0 III GEC ItranSVJ80.5 ,90.5 21.2

I1114.5 ,20

I TE ! transv J 82 90 1113.5 121 '21.0

Test i Test IV.S IV'S' IEL. ,R.A. ICVN !FATT HRC I

loc. I dir. ',Q.02% 0.2%I 'MPa IMPa J ·C I

A radial 1527 593 96 16 19.7

B \ radial' 541 '617 63 19 20.7i C I radial1521 610 70 19 20.0,

I AC i transv J 538 593 74 10 20.2

IBC 'I transv J 548 603 74 15 20.8CC I transv J 572 1627 56 16 20.0 i

1

GEC :transv J 555 :624 60 19 21.2 I

TE I transv, 565 '621 66 8 21.0

TABLE 6- - Mechanical properties of LP rotor.


SUMMARY

Two LP rotors for geothermal power stations were manufactured of modified11% CrMoNiNb steel. The rotors made of ESR ingot of 2 300 mm (90 in.) diameter and113 mt weight were forged on a 60 MN press and heat treated to geothermic specificationproperties. The results can be summed up with the following conclusions:

* Application of high temperature 11% CrMoNiNb steel for geothermal serviceis possible. Necessary is modified chemical composition and heat treatment.

* Through hardening up to diameter of 1 500 mm (60 in.) is given.* Mechanical properties over cross section are homogeneous.* Heat treated to a tensile strength of appro 780 MPa (113 ksij and yield

strength (0.02%) of appro540 MPa (78 ksi) results in a FATT of appro 16°C(60°F)

REFERENCES

[1] Newhouse, D. L., Boyle, C. J., Curran, R. M., 'A Modified 12-PercentChromium Steel for Large High- Temperature Steam Turbine Rotors,'ASTM Annual Meeting 1965

[2] Berger, C., Potthast, E., Bauer, R., Honeyman, G.A. 'Development of HighStrength 9-12% CrMoV Steels for High Temperature Rotor Forgings,'Proceedings IX, 5 of 11. Int. Forgemasters Meeting, Terni, 1991

[3] SchOnfeld, K., H., Potthast, E., 'Soft Martensitic Stainless CrNiMo Steel forTurbine Rotors in Geothermic Power Stations,' ASTM Symposium on SteelForgings, 1984

Yoichi Tsuda,l Masayuki Yamada, 1 Ryuichi Ishii,l Yasuhiko Tanaka/ Tsukasa Azuma,2

and Yasumi Ikeda3

DEVELOPMENT OF HIGH STRENGTH 12% CR FERRITIC STEEL FORTURBINE ROTOR OPERATING ABOVE 600°C

REFERENCE: Tsuda, Y., Yamada, M., Ishii, R., Tanaka, Y., Azuma, T., Ikeda, Y.,"DEVELOPMENT OF HIGH STRENGTH 12% CR FERRITIC STEEL FORTURBINE ROTOR OPERATING ABOVE 600°C,", Steel Forgings: Second Volume,ASTM STP 1259, E.G. Nisbett and A.S. Melilli, Eds., American Society for Testing andMaterials, 1997.

ABSTRACT: An advanced 12% chromium ferritic steel has been developed for thehigh-temperature rotors in improved fossil-fired power plants. The development of thissteel is based on the experimental studies of various 12% chromium heat resistant steels.Effects of alloy elements on the creep rupture strength have been investigated and the10Cr-1.8W-0.7Mo-V-3Co-Nb-B steel was selected to be a hopeful candidate. Thisrevealed the creep rupture strength was higher than that of 1OCr-1Mo-1 W-V-Nb rotorsteel being used at 593°C.The producibility and properties of the new steel have been verified with the medium sizetrial forging manufactured from a 20 ton ingot. Although a few technical subjects areremaining, these are expected to be solved with the modification of the manufacturingprocess.This newly developed high strength 12 % chromium ferritic steel is expected to beapplicable to high-temperature rotors operating at 630°C or above, and the significantimprovement o~thermal efficiencies would be obtained in fossil-fired power plants.

KEYWORDS: heat resistant steel, 12% Cr ferritic steel, high-temperature rotor, fossil-fired power plant, creep rupture strength

1 Specialist, manager, and engineer, respectively, Heavy Apparatus EngineeringLaboratory, Toshiba Corp., Yokohama, 230 Japan.2 Senior research engineer, and research engineer, respectively, Muroran Laboratory, TheJapan Steel Works, LTD., Muroran, 051 Japan.3 General manager, Muroran Plant, The Japan Steel Works, LTD., Muroran, 051 Japan.

267


Recently, the higher efficiency of thermal power plants has been extremelyimportant to protect the earth environment. In fossil-fired power plants, the increase ofsteam conditions (temperature and pressure) has been conducted to gain higher efficiency.In order to increase the steam temperature, high strength materials are necessary for hotturbine components especially such as high-temperature rotors. These components aredesirable to consist offerritic steels such as 1% Cr-Mo-V and 12% Cr, because theaustenitic materials have disadvantage of thermal expansion and thermal conductivity.

The" 12% Cr rotor" and the "modified 12% Cr rotor" have been used for turbinesoperating up to 593°C because of their good elevated temperature strength [1], [2].Current tendency of fossil-fired power plants revealed that more stronger rotor steels arerequired for advanced power plants with the steam temperatures of 600°C or above.Consequently, various investigations are being conducted to attain such steels [J.], [1].

Research and development (R&D) have been carried out on a new rotor steelwhich has excellent creep rupture strength suitable for high-temperature rotors ofadvanced power plants. This paper covers the optimization of chemical composition,the verification study of material properties and producibility of a trial rotor forging, andthe future plan.

OPTIMIZATION OF CHEMICAL COMPOSITION

TABLE 1 shows the nominal chemical composition of the new rotor steel and1OCr-l Mo-I W-V-Nb steel which is one of "modified 12% Cr rotor" steel. Comparedwith 1OCr-l Mo-l W-V-Nb steel, the feature of chemistry of the new rotor steel is asfollows:

(1) Lower carbon, manganese and nickel contents(2) Higher tungsten and lower molybdenum contents(3) Addition of boron and cobalt

To determine the above mentioned chemical composition, the evaluation of 50 kglaboratory heats was carried out.

TABLE I--Nominal chemical composition of rotor steels (mass%)

C Si Mn Ni Cr Mo V Nb Co W N B FeNew 12% Cr

0.11 0.03 0.1 0.2 10.0 0.7 0.2 0.05 3.0 1.8 0.02 0.01 Bal.Rotor SteelModified 12% Cr

0.14 0.03 0.6 0.7 10.0 1.0 0.2 0.05 1.0 0.04 Bal.- -Rotor Steel

Test Materials

Nineteen kinds of IOCr-0.2V-0.05Nb steels were manufactured using the high-frequency vacuum induction melting process to investigate the influence of alloy

TSUDA ET AL. ON DEVELOPMENT OF HIGH STRENGTH 269

elements. These steels had varied chemical composition of carbon, nickel, boron,tungsten, molybdenum and cobalt as shown in TABLE 2. The manganese content ofevery steel was reduced to be about 0.1% to improve creep rupture strength [ll InTABLE 2, the Series 1 is respected to evaluate mechanical properties and the Series 2 isrespected to evaluate weldability for the journal over-lay.

After forging and preliminary heat treatments, those were austenitized at 1070°Cfor 3 hours and cooled at a rate of 100°C/hour. Afterwards, a double tempering wasapplied to reduce residual austenite structure. The first and the second tempering was570°C for 10 hours and 680°C for 20hours, respectively

TABLE 2--Chemical composition of test materials

Series No. C Si Mn Ni Cr Mo V Nb Co W N B Fe

1 0.13 0.04 0.07 0.17 9.96 0.68 0.18 0.04 - 1.76 0.0342 - Bal.2 0.13 0.04 0.06 0.18 9.95 0.66 0.18 0.04 3.07 1.76 0.0324 - Bal.3 0.13 0.03 0.08 0.19 9.96 0.66 0.18 0.04 5.03 1.72 0.0353 - Bal.4 0.16 0.04 0.06 0.18 9.97 0.66 0.18 0.04 3.11 1.77 0.0323 - Bal.5 0.11 0.04 0.07 0.19 9.97 0.68 0.18 0.04 5.16 1.78 0.0319 - Bal.6 0.09 0.03 0.08 0.24 10.17 0.67 0.19 0.05 2.98 1.81 0.0217 0.005 Bal.

1 7 0.11 0.03 0.07 0.24 10.35 0.71 0.20 0.07 3.00 1.81 0.0225 0.011 Bal.8 0.11 0.03 0.08 0.23 10.20 0.67 0.19 0.05 5.07 1.80 0.0210 0.012 Bal.9 0.09 0.03 0.08 0.62 10.17 0.66 0.19 0.05 3.10 1.77 0.0245 0.012 Bal.10 0.110.030.08 0.22 9.97 0.66 0.20 0.05 2.01 1.850.03140.011 Bal.11 0.12 0.04 0.09 0.21 10.16 0.65 0.21 0.06 1.03 1.81 0.0314 0.009 Bal.12 0.11 0.04 0.08 0.21 10.02 1.02 0.21 0.05 3.03 1.01 0.0306 0.011 Bal.13 0.11 0.03 0.06 0.23 9.90 0.22 0.20 0.05 3.03 2.57 0.0268 0.011 Bal.14 0.10 0.02 0.08 0.22 9.90 0041 0.20 0.05 3.10 2.22 0.0258 0.010 Bal.15 0.09 0.06 0.11 0.21 10.16 0.64 0.21 0.05 3.08 1.79 0.0188 - Bal.16 0.11 0.06 0.09 0.21 10.27 0.66 0.21 0.05 3.10 1.80 0.0200 0.006 Bal.

2 17 0.11 0.05 0.09 0.21 10.21 0.66 0.21 0.05 3.08 1.81 0.0197 0.009 Bal.18 0.11 0:05 0.09 0.21 10.27 0.66 0.21 0.05 3.10 1.800.0193 0.012 Bal.19 0.10 0.05 0.08 0.20 10.20 0.65 0.21 0.05 3.08 1.84 0.0204 0.016 Bal.

Influence of Alloying Elements on Mechanical Properties

Tensile test, Charpy impact test and creep rupture test were conducted with testmaterials of Series 1. In every materials, tensile strength and elongation were almostfixed to be 900 MPa and 20%, which are normal values of 12% Cr rotor steel. On theother hand, toughness and creep rupture strength were extensively changed along with thechemical compositions. FIG. 1 and FIG. 2 show the influence of each alloying element


on toughness and creep rupture strength.

FIG. l--Effect of carbon, nickel, boron on toughness and creep rupture strength

~-- The increasing of carbon content from 0.11 % to 0.16% has favorableeffect on toughness. However, the higher carbon decreases creep rupture strength. Thedesirable carbon content was determined to be 0.1 % level because the creep rupturestrength is more important factor for the high temperature turbine rotors.

Carbon forms various types of carbides such as M23C6 and NbC containing theother several elements. It is expected that the lower carbon prevents the growth of suchcarbides during creep.

Nickel--Lower nickel is known to have a beneficial effect on creep rupturestrength. This knowledge was obtained also in the test results shown in FIG. I.Moreover, since FA TT rises slightly even though nickel content decreases from 0.6% to0.2%, the lower nickel content is considered to be desirable for the new rotor steel.


Boron--It is apparent that a little quantity addition of boron extensively improvesthe creep rupture strength and reduces the toughness. The higher boron content isconsidered to be desirable concerning the creep rupture strength. However, since theforgeability extensively decreases by addition of 0.015% boron, about 0.01% boron wasselected as the desirable composition.

FIG. 2--Effect of tungsten, molybdenum, cobalt on toughness and creep rupture strength

Tungsten. Molybdenurn--When the molybdenum equivalent is fixed to be 1.5,higher tungsten is known to increase creep rupture strength [~]. In this paper, themolybdenum equivalent is thus:

Mo equivalent = Mo + 1/2W (mass%)

The-greatest part of the strengthening effect of tungsten and molybdenum isseemed to be achieved by the time the tungsten level is reached to 1.8% in our


investigation. On the other hand, toughness extensively decreases by the addition oftungsten exceeding 1.8%. Thus, the 1.8% tungsten and 0.7% molybdenum contentswere selected as the desirable composition.

Tungsten and molybdenum are known to be solid solution strengthening elements.Moreover, the precipitation strengthening effect by intermetallic compounds containingtungsten recently becomes to be noticed [1], [8l It is considered that the highertungsten content maintains the sufficient solute tungsten and increases the quantity ofintermetallic compound such as Laves or mu (Jl) phase contributing to precipitationstrengthening.

Cobalt--Because of the possible tendency for increasing chromium equivalent dueto above mentioned chemistry modifications, the formation of delta ferrite was of concern.To offset this potential problem, an addition of cobalt was considered. Moreover, thebeneficial effect on creep rupture strength by increasing cobalt up to 3% is apparent.Cobalt is expected also to act as a solid solution strengthening element.

Influence of Alloying Elements on Weldability for Journal Over-lay

12% Cr rotors often cause a bearing problem called "shaft galling". The journalover-lay welding with low alloy steels is considered to be an effective means to preventthis problem. There is a possibility that the new rotor steel has an undesirableweldability because of the boron addition. Thus, the verification study on theweldability was conducted using the five kinds oftest materials with varied compositionof boron.

TABLE 3 summarizes the test results. The cracking sensitivity is affected alsoby the boron content and the heat input. The addition of 0.016% boron causes thecracking in the case of gas tungsten arc welding (GTAW) with a small heat input of 9360J/cm. On the other hand, in the case of sub-merged arc welding (SAW) with a large heatinput of 30900 J/cm, the cracking occurred by the addition of 0.009% boron. Since thenew rotor steel contains about 0.01% boron, these results indicates that at least the initialwelding layers of over-lay should be made by welding method such as GTAW with smallheat input.

TABLE 3--Weldability test results

MethodB (%)

0 0.006 0.009 0.012 0.016GTAW no crack no crack no crack no crack crackSAW no crack no crack crack crack crack

GTAW : Gas Tungsten Arc WeldingSAW: Sub-merged Arc Welding


VERIFICATION STUDY ON MATERIAL PROPERTIES ANDPRODUCIBILITY OF TRIAL ROTOR FORGING

A medium size trial rotor forging was produced on the basis of knowledgeobtained so far. The purpose is confirming of the material properties, particularly creeprupture strength, as well as verifYing producibility.

Production of Trial Rotor Forging

The drum diameter of the trial rotor forging was 1060 mm as shown in FIG. 3 andis large enough for covering the high-pressure or intermediate-pressure rotors for turbinesin 1000 MW plants.


The outline of the manufacturing process of the trial rotor forging is shown in FIG.4. To reduce impurities as little as possible, the ladle refining was carried out after basicelectric furnace melting. The electro-slag remelting process was applied to get an ingotof 20 ton with minimal segregation. The quality heat treatment consists of quenching anddouble tempering. The rotor forging was austenitized at 1070°C and afterwards oilquenched. The first tempering was carried out at 570°C, and the second tempering at680°C was applied.

Evaluation of Trial Rotor Forging

Chemical CompositionnResults of detailed chemical analysis in various regionsare shown in FIG. 5 and FIG. 6 as the original ingot shape.


The chemical compositions were almost uniform except boron. In both centerand surface regions of the ingot, boron tends to be gradually more diluted as approachingthe top end from the bottom end. The boron content near the top end was about 0.005%,which is lower value than desirable one. However, near the bottom end, the higherboron content of about 0.02% was observed. This boron content was not predictedvalue and undesirable for also forgeability and weldability.

In ESR process, various chemical reactions occur in the slag. In relation withboron, the oxidization and the deoxidization are conceivable. If these reactionsequilibrate during ESR process, a flat distribution of boron content could be attained.


However, the oxidization of boron progresses in actual ESR process because the oxygencomes into the slag from the electrode and the atmosphere. As a result of thisoxidization, the distribution of boron content in ingot would not become uniform.

Metallography-- The macrostructure revealed no apparent segregation. Themicrostructures in various regions are shown in FIG. 7. The microstructure is temperedmartensite which is typical of 12% Cr rotor steel and no delta ferrite was found.

FIG. 7--Microstructure of trial rotor forging

Tensile and Charpy Impact Properties--Tensile and impact properties in variousregions of the fully heat-treated trial rotor forging are shown in FIG. 8. The tensileproperties are uniform throughout the forging. The FATT of the body was about 11DOC


which is inferior to that of "modified 12% Cr rotor". However, this value is expected tobe allowable practically because it is almost equivalent to that of 1% Cr-Mo- V high-pressure rotor steel.

FIG. 8--Tensile and impact properties in trial rotor forging

Creep Rupture Properties-- The smooth and notch rupture data from the trial rotorforging are summarized in FIG. 9, which include the mean line for 10Cr-IMo-l W-V-Nb"modified 12% Cr rotor" and results of the laboratory heat. The improved rupturestrength of the new rotor steel is observed. Moreover, since the results of the trial rotor


forging are almost consistent with those of the laboratory heat, the mass effect isconsidered to be negligible. The notch rupture data show notch strengthening.

FIG. 9--Creep rupture strength in trial rotor forging

Other Properties--In addition, such tests as low cycle fatigue, creep rate and longtime aging have been conducted. The results obtained so far show the favorableproperties as high-temperature rotors.

FUTURE PLAN

Since a technical subject on the boron distribution is remaining, thecountermeasures for getting the uniform boron distribution are required. For the present,it is expected that a better result will be attained by the modifications of ESR parameterssuch as the slag composition and the melting speed. And the manufacturing andevaluation of a product-size rotor forging will be planned for the final verification.


SUMMARYLower carbon, manganese, nickel, molybdenum contents, higher tungsten content

and addition of boron, cobalt significantly improve the creep rupture strength of 1OCr-IMo-l W-V-Nb "modified 12% Cr rotor" steel. From this knowledge, the lOCr-1.8W-O.7Mo-V-3Co-Nb-B steel with excellent creep rupture strength has been developed forhigh-temperature rotor forgings.

Results from the production and evaluation of the trial rotor forging indicate thatit is possible to produce new 12% Cr rotor forgings with desirable properties andmicrostructure. It has also been confirmed that superior creep rupture strength can bereadily obtained. Although a question on the boron distribution remains, this isexpected to be solved with the modification ofESR process.

The use of the high-temperature rotor forgings made from this steel would resultin significantly improved efficiency due to higher operating temperature of 630°C orabove. It can be considered the newly developed steel may be widely applied to high-temperature rotors of improved fossil-fired power plants in future.

REFERENCESill Tsuda, Y., Miyazaki, M. and Kaplan, A., "Advanced 12% Cr Steel for High-

temperature Rotors," Proceedings ofEPR! Third International Conference onImproved Coal-Fired Power Plants, San Francisco, California, April 1991

ill Ito, F., Kuwabara, K., Miyazaki, M., Fukui, Y. and Takeda, Y., First InternationalConference on Improved Coal-Fired Power Plants, 1988, pp 6-127

ill Fujita, T., 3rd International Charles Parsons Turbine Conference, Newcastle, UK,April 1995, pp 493

ill Metcalfe, E. and Baker, W. T., The EPR! / National Power Conference, London, UK,May 1995, pp 1

ill Kitagawa, I., Morinaka, K. and Fujita, A., 11th International Forgemasters Meeting,Terni~Spoleto, Italy, June 1991, pp IX.6

[Ql Fujita, T., The Thermal and Nuclear Power, Vol. 42, No. 11, November 1991, pp 21

ill Kimura, K., Ishii, R., Matsuo, T. and Kikuchi, M., 123rd Committee on HeatResisting Metals and Alloys Report, Vol. 34, No.7, 1993, pp 127

ill Igarashi, M. and Sawaragi, Y., 123rd Committee on Heat Resisting Metals andAlloys Report, Vol. 35, No. 11, 1994, pp 285

Ramaswamy Viswanathan1

HISTORICALOVERVIEWOFIMPROVINGCLEANLINESSOFROTORSTEELS FORELECTRICUTILITYApPLlCAnONS

REFERENCE Viswanathan, R., "Clean/Superclean Steel Rotors for Electric UtilityApplications," Steel For~in~s: Second Volume, ASTM STP 1259, E.G. Nisbett andA.S. Melilli, Ed., American Society for Testing and Materials, 1997.

ABSTRACT Improved cleanliness is perhaps the only approach that results insimultaneous improvement in strength and ductility at elevated temperatures as wellas toughness at low temperatures of steels. In addition, superclean steels in whichmanganese and silicon have also been reduced, provide greater resistance to stresscorrosion cracking than conventional steels. Major projects are underway worldwideto promote the use of clean/superclean steel rotor and disk forgings, both for lowtemperature and for high temperature applications in steam and combustion turbines.An international workshop sponsored by EPRI was held in 1995 in London, at whichturbine manufacturers and steelmakers discussed ongoing activities with respect toclean steels. This paper summarises EPRI sponsored research on rotor steelchemistries and provide an overview of developments over the last 15 years.

KEYWORDS steel, impurities, clean, rotors, embrittlement, creep, stress corrosion.

INTRODUCTION

The adverse effects of impurity elements phosphorus (P), antimony (Sb), tin (Sn),arsenic (As), sulfur (S), oxygen (0), and deoxidants aluminum (AI) and silicon (Si)on the mechanical properties of steels have been known for many decades. P, Sb, Sn

IManager, Materials Application Technology, Electric Power Research Institute, 3412Hillview Avenue, Palo Alto, CA 94304.

280

VISWANATHAN ON HISTORICAL OVERVIEW 281

and As interactively with Si and Mn cause temper embrittlement and lead toreduction in the fracture toughness (Kd and increase in the ductile-brittle transitiontemperature (FATT). Presence of sulfide inclusions, and nonmetallic inclusionscontaining Al and Si can facilitate cavity nucleation at the grain boundaries and inthe grains, thus facilitating creep fracture at high temperatures and ductile fractures inthe upper shelf region. These changes result in reduced creep ductility at hightemperatures and reduced fracture toughness at lower temperatures. Advancementsin the steel making technology during the last two decades have enabled reduction ofsome of these impurities and deoxidants to levels as low as 20 ppm leading to whatmight be called "clean" steels. Further realization that in view of the low sulfurlevels achievable, even Mn is no longer necessary to "fix" the sulfur and can,therefore, be reduced to levels as low as 0.01 to 0.02% has resulted in "superclean"steels.

The development that has made clean and superclean steels possible wassecondary steel refining via ladle furnaces in conjunction with vacuum degassing inthe ladle and during casting. Ladle treatment of molten steel for the purpose ofdesulfurization and vacuum carbon deoxidation were developed only in 1975. Thesetechniques provided a production means of manufacturing high purity steels thatpreviously could only be made in the laboratory as a control for temperembrittlement research studies. Thus, the practical solution to the problem in alloysteels was to refine the steels from Mn, Si, P, Sn, As, and Sb during steel makingoperations. Manganese, silicon, and phosphorus are easily removed during theoxidizing stage of steel making, because they oxidize preferentially to iron, and enterthe oxidizing slag. This generally is done in the electric arc furnace. Tin, arsenic,and antimony are controlled by scrap selection, with basic oxygen furnace steel scrapused as the starting material for electric arc furnace melting if possible. Afterseparation of the oxidized steel from the oxidizing slag, it is transferred to a reducingslag in a ladle refining furnace, which removes the sulfur. Vacuum treatment of thedesulfurized steel in the ladle furnace accompanied by argon bubbling provides ameans for deoxidation via the vacuum carbon deoxidation (VCD) process. Thisleaves fewer oxide particles dispersed in the steel as would occur if deoxidation weredone with silicon. The alloying elements are added at this stage. Casting into moldscontained in a vacuum chamber, called vacuum stream degassing (VSD) completesthe double degassing treatment. The extensive degassing combined with argonbubbling removes nitrogen and hydrogen as well as oxygen. The result is asuperclean steel free to the maximum extent possible of grain boundary segregatesand of non-metallic oxides and sulfides. Alternate routes which may not be aseffective to produce clean/superclean steels also include use of Al deoxidation in lieuof VCD and conventional pouring in air with argon shrouding.

The primary driver in cleaning up steels has been the desire to achieve improvedfracture toughness, although benefits in other properties such as resistance to creep,


stress corrosion, etc. have accrued incidentally and serendipitously, makingcleanliness even more desirable. In general, any modifications to alloy content orheat treatment to improve toughness of the steels have an opposite effect on the creepstrength. Reduction of the impurities (tweaking) provides one of the few viable waysof achieving the best results with respect to all the mechanical properties.

Good fracture toughness and freedom from impurity induced temperembrittlement are crucial requirements for rotors and discs. Regardless of theoperating temperature, the final failure scenario for all catastrophic rotor/disc burstsinvolves either failure due to peak thermal stresses during a start-up/shut-downtransient, during over-speed testing or during steady operation at low temperatures.Hence FATT and fracture toughness considerations at low/intermediate temperaturesare important [1,2].

Reduced fracture toughness of rotors and discs arising either inherently due tosulfides and inclusions or due to impurity-induced embrittlement results in a decreaseof the component life under base load and cyclic operation due to increased risk ofbrittle fracture, and is one of the major causes of early component retirement.Embrittlement can also have a major constraining effect on the operating procedure.To keep transient stresses low, stringent controls have to be exercised with respect tothe start-stop cycles. For instance, during each cold start, the rotors may need to bepre-warmed to a temperature above the ductile-to-brittle fracture appearancetransition temperature (FATT) over a period of several hours prior to imposition offull load. These require-ments lead to increased operational costs for the plant anddecreased flexibility and avail-ability. Reduced toughness also increases inspectioncosts due to the more frequent inspections required. In the case of low-pressure (LP)rotors, the LP cross-over temperature in fossil plant turbines is kept below about370°C (700°F) to avoid embrittlement, even at the risk of reduced efficiency.Increase of the LP cross-over temperature to 400 to 42rC (750-800°F) would resultin major efficiency gains. Improving the toughness of LP rotor steels also means thatthe increasing of the strength such as the tensile strength and the proof strength canbe obtained along with maintaining the conventional toughness level by themodification of the tempering condition. Accordingly, the length of the last stageblading can be increased to improve the turbine thermal efficiency. Toughness ofsteels thus has a significant effect on the reliability, availability, cycling ability,efficiency and operating and maintenance costs (O&M) of turbines.

Table I shows the principal 'driver' requiring advancements in the differenttypes of turbine rotors. For a more complete summary of rotor developments thereader is referred to [~]. In the case of HP/IP rotors operating at elevatedtemperatures allowing greater efficiency and improved toughness providing greaterreliability, longevity and cycling ability have been the drivers. For LP rotors,increasing the LP cross-over temperature from 370°C (700°F) to 400°C (750°F)


without the risk of temper embrittlement has been the principal driver since such achange would increase the turbine efficiency substantially. Increased longevity understress corrosion conditions has also been an incidental benefit. For UP/LP combinedrotors, reduced installed cost/kwh by housing the entire rotor inside a single casinghas been the driver. This becomes possible if the creep strength requirements in thehigh temperature end and toughness requirements in the low temperature end canboth be satisfied in a single forging. In all cases, improved reliability of the rotorleads to decreased inspection costs since the inspection intervals can be extended.

EPRI has been active in sponsoring and catalyzing research pertaining toimproved UP, LP and UP/LP rotor steels primarily through improvements incleanliness since the mid 70s. Although the deleterious effects of impurities wereknown prior to that, there were no systematic studies of the synergistic effects ofimpurities. It was known that Mn contributed to embrittlement, but high Mn levelswere routinely maintained although no longer required in view of the steady decreasein the S content of steels [4]. In order to make clean steels a reality, the steelmakers,especially in the U.S. needed to be convinced of three things:

1. That impurity elements do in fact adversely affect the toughness, creep ductility,stress corrosion and other key properties of steels.

2. That clean steel components in the sizes needed can be successfully made withavailable technologies, without driving up the costs to unrealistic levels.

3. That superclean steels with a very low Mn could be made and forged into productshape.

All of these concerns were systematically addressed through a series of research anddevelopment and demonstration projects. The objective of this paper is to review thevarious EPRI-sponsored developments relating to clean/superclean steel technology.

UP AND IP ROTORS

These rotors typically operate at a maximum temperature of either 540°C(lOOO°F) or 565°C (l050°F) at the steam inlet end with temperature falling to about345 °C (650 OF)at the steam exit end of the high pressure (UP) orthe IntermediatePressure (IP)' turbines. For temperatures up to 540°C (lOOO°F), a lCrlMol/4V steelis generally used and for temperatures up to 565°C (l050°F) variations of a 12%Crmartensitic steel are used.

Figure 1 shows the reduction in the concentration of impurity elements inlCrlMol/4V steels over the last few decades, due to advancements in steel makingand refining technologies.


Fig. 1 Trends in impurity levels inCr-MoV rotor steels [1]

Numerous improvements have been made since the late 1940s in terms of rotorcompositions, fabrication practices and cleanliness, as illustrated in Figure 2 for hightemperature forgings, based on the extrapolation of results cited in [~, Q]. Thetoughness and rupture ductility have seen corresponding improvements as shown inFigure 2 [~, Q]. The key events in this evolution include:

1. Switching to higher creep strength CrMoV rotors from the NiMoV rotors in1948.

2. Introduction of rotors austentized at 954°C (1750°F) in place of the earlier rotorsaustenitized at 10100C (1850°F) to eliminate notch sensitivity and low ductilityproblems around 1953.

3. Use of electric furnace melting and vacuum pouring introduced around 1956 toreduce bore segregation of nonmetallic inclusions and sulfide streaks.

4. Introduction of 12 percent Cr rotors for 1050°C service in 1960.5. Application of vacuum arc remelting process in 1970.6. Demonstrafion of clean steel CrMo V rotors about 1984.7. Demonstration of higher creep strength in modified 12 percent Cr rotor

compositions about 1987.

The integrity of rotors has also been improved by advances in welding procedures forthe initial fabrication or repair incorporating smaller forgings of higher quality.


Fig. 2 Evolution of HP rotor properties in relation tochanges in steel making technology (based on[~, .6])

The first EPRI project dealing with clean steel rotors was started in the wake ofthe catastrophic failure of the Gallatin HP/IP rotor, wherein presence of sulfidestringers at the bore had been implicated in a major way [1]. To reduce the sulfurlevel in HP/IP lCrlMol/4V type (ASTM A470 - Class 8) rotors, Swaminathan andJaffee demonstrated the effectiveness of three advanced steel makingtechnologies [~, 2]. Three full-size 30 ton Cr-Mo-V rotor forgings were producedfrom 100 ton ingots each by vacuum carbon deoxidation (VCD), electroslagremelting and low sulfur vacuum silicon deoxidation (low S) processes. Ladlerefining was applied to the liquid metal in the low Sand VCD processes, whereas inthe ESR process refining of the steel was achieved during electroslag remelting. Allthe processes reduced sulfur to very low levels in the range 10-20 ppm resulting inexcellent bore quality and cleanliness. In addition, the VCD process also resulted inlow P (30 ppm) and low Si levels. The centerline cleanliness of the rotors wereappreciably better than conventional rotors.

Compared with conventional forgings, the fracture toughness was found to besuperior, as shown in Figure 3. In addition, creep strength and ductility and lowcycle fatigue properties were also improved. Considering the fact that, in general,improvements in rupture strength and ductility and in rupture strength and fracture


toughness are mutually opposed, the simultaneous improvements achieved withrespect to all of these properties must be considered a major milestone in rotortechnology. The three trial rotors made in this project were installed at threeoperating plants in the U.S. rated at 520 MW each. Rotors with low levels of Siproduced by the VCD process have been used in several commercial plantsrecently [lQ]. In Germany, on the other hand, fracture toughnesses comparable to thelow sulfur steels have been achieved since 1973 even in commercial cleanlinessCrMoV steels by oil quenching, in contrast to the practice of air quenching in the U.S[il]. However, the extreme reduction of Mn content required in "superclean" steelshas not been performed in 1CrlMol/4V rotor steels. The application of supercleansteel technology to this steel was tried with a view of increasing the toughness, butthe results were not promising [12], The lCrlMoV steel lost harden ability as a resultof the absence of Mn and compensation to restore harden ability was necessary inorder to achieve a fully bainitic structure at the slow cooling rates characteristic ofthe center of rotors. The alloy modifications made for compensation eventually ledto the 2.5% NiCrMoV steel discussed later as a candidate for HP/LP combined shaft.

With respett to 12%Cr steels for applications at and above 565°C (l050°F), allthe available information has been compiled and summarized LU]. There is limiteddata on what appear to be clean steel versions of 12%Cr steels in the literature Ll1-11], Several modified versions of 12%Cr steels for use above 593 °C(llOO°F) havebeen investigated. The most recent compilation of these compositions may be foundelsewhere Ll5.-1aJ. Many of these compositions investigated tend naturally to be lowin Si, Sand P since many of the rotors are produced by VCD, ESR or other refinedprocesses. For all practical purposes these may be called "clean" steels. Sincealloying element modifications and not specifically improved cleanliness is the focus


of these studies, they are not reviewed here. There are, however, no reported studiespertaining to superclean versions of 12% steel for steam turbine applications. Tworecent studies have reported results of evaluation of superclean trial discs of 12%Crsteels for combustion turbine disc applications. The superclean steels were shown tohave significantly improved toughness and freedom from long term embrittlementcompared to conventional12%Cr steels [12,2Q]. This development would clearlypermit use of superclean steel discs at higher temperatures with attendant gains inefficiency.

LP ROTORS

LP rotors are typically made of a 3.5%Ni, 1.5%Cr, 0.5%Mo, 0.1%V steel. Themaximum permissible temperature for these rotors has been limited to about 370°C(700°F) due to the possibility of in-service temper embrittlement. Cleanlinesstherefore has been a focus primarily to eliminate temper embrittlement so thatincreased efficiencies could be realized by operating at higher temperatures.

Figure 4 illustrates the decrease in FATT as a function of the year of manufacturefor LP rotor materials. This is the core FATT in the as-fabricated condition. Theplot shows that the FATTs have been quite low and that toughness has not beenmuch of a concern since the early 60s. The concern has been mainly in-servicetemper embrittlement and post exposure FATT. With the introduction of VCD andreduction of Si in the early 60s, the embrittlement problem was reduced. With theadvent of the superclean steels, where both Mn and Si are reduced using a number ofalternative techniques, the embrittlement problem has been virtually eliminated.

The application of clean steel and superclean steel technologies with respect to3.5Ni 1.5Cr 0.5Mo 0.1V steels typically used for rotor steels was first investigated inEPRI project RP2060 [12]. The superclean steels contained approximately 0.001 %S,0.002%P, 0.02%Si and 0.02%Mn, 25 ppm oxygen and 50 ppm N. Laboratory ingotsweighing 100-200 Kg were made, forged and evaluated. These evaluations showedthe superclean version of NiCrMo V LP rotor steels had superior ductility andtoughness properties and high temperature creep strength compared to theconventional NiCrMoV rotor steel. The next step was to make and evaluate trialrotors of the 3.5NiCrMo V steel used for low pressure steam turbine rotors. Two trialrotors (34 tons 1186 mm dia.) were cast by Vereinigte Edelstahlwerke A.G. (nowBoehler A.G.) in Kapfenberg, Austria in 1985 [21-23J. Japan Steel Works andToshiba sponsored a third trial rotor (106 tons 1800 mm dia.) made by JSW inMuroran, Japan. Sections of these trial rotors were widely distributed throughout the


world. Some of the results from these studies conducted have been reported byWatanabe, Yoshioka, Schwant and others [~-25.J. On the basis of the successfultrial rotor demonstrations, superclean 3.5NiCrMoV was selected for the ChubuElectric Company's 700 MW units for the Kawagoe and Hikenen plants. Other trialand production 3.5NiCrMoV rotors were manufactured by Kobe Steel and JapanCasting and Forging. Two recent retrofits for Duke Power Company's River BendStation were manufactured by Kobe Steel. One Japanese steel manufacturer reportedproduction and delivery of 28 superclean LP rotors in the 700 MW class forimproved efficiency and SCC resistance.

In the case of the Kawagoe USC (Ultra Super Critical) plant of Chubu ElectricCompany in Japan, the increase in the inlet steam temperature of LP turbine wasapplied without the cooling by the extracted steam in order to improve the thermalefficiency. In this USC plant, the superclean LP rotors were used for the first time inthe world. The LP inlet temperature of this plant was about 30°C (54 OF)higher thanthat of the conventional plants and therefore the thermal efficiency was improved byabout 0.1% with no change in the turbine operation and the inspection interval [lQ].

The superclean 3.5NiCrMoV steels have lower FATT, higher fracture toughnessand higher upper shelf energy compared to conventional grade as shown in Figures 5and 6. The FATT ranges from -20°C (_4°F) to -70°C (-94°F) at the core and mostly


below - 110°C (-166 OF)at the rim. These values are somewhat below thoseachievable for conventional grade. Most importantly, the superclean steels areimmune to temper embrittlement (see Figure 7). The higher fracture toughness andlower FATTs translate into larger critical crack sizes and hence increased longevityunder steady and cycling conditions.


The increased creep strength of superclean 3.SNiCrMoV steels compared toconventional grade steel is shown in Figure 8. The increased strength was derivedwithout compromising the creep ductility. This is thought to derive from the absenceof MnS particles at the prior austenite grain boundaries which might cavitate andinitiate grain boundary creep fractures. The service temperature ceiling forsuperclean gratle 3.SNiCrMoV is limited by creep strength to a maximum of SOODCinstead of to a maximum of 3S0DClimited by susceptibility to temper embrittlement.This makes it a good candidate for use in single shaft HP/LP configurations for steamtemperatures up to SOODC(930DF). Superclean steels have also been usedextensively for Combustion Turbine disk applications.


An incidental but nevertheless key benefit of superclean 3.5NiCrMoV steels istheir greater resistance to stress corrosion cracking (SCC) initiation. The results areclear that crack propagation rates are the same for superclean and conventional3.5NiCrMoV. However, data from many sources indicate that crack initiation ismuch more difficult in superclean grade steel, probably because of its greatercleanliness. In applications where rotors and discs must operate in the presence ofcondensed steam, this characteristic could spell the difference between success andfailure, particularly if the steam would be contaminated by salt impurities or oxygen.Figure 9 presents initiation data on 3.5NiCrMoV heat treated to 860 Mpa yield fromScarlin and Denk [2..Q].

EPRI has prepared a superclean steel guide for utility use [21]. This guidesumma-rizes a very large body of available literature and provides a samplecompositional sp.ecification as shown in Table 2.

Superclean 3.5NiCrMoV has become widely produced and accepted for LP aswell as single shaft HP-LP rotors with steam temperatures up to 500°C (930°F), thepresently accepted temperature limit for adequate creep strength. Superclean3.5NiCrMo V steels have also found a "niche" market for combustion turbine parts.There seems to be considerable interest in this application especially in Europe.Numerous papers describing production and properties of superclean steels arecontained in the symposium volumes listed under [1:1:. 11].




HPILP SINGLE SHAFT ROTORS

In many combined cycle plants, it is cost effective to have a single shaft HP/LProtor configuration for the steam turbine. The need for multiple casings iseliminated, the unit is more compact and simpler to install. The rotor used in thisfashion should combine in a single forging high creep strength at one end with hightoughness at the other end. EPRI has sponsored two major development efforts, oneinvolving a superclean version of a 2.5%NiCrMo V steel jointly with the GeneralElectric Company, Toshiba Corporation and Japan Steel Works under projectRP1403-55 and a second project involving KWU-Siemens in the lead with Saarstahl,and several European turbine manufacturers (Asea Brown Boveri, GEC-Alsthom,MAN Energie and Parsons) as partners under RP1403-21 to evaluate a "clean steel"version of a 2%CrMoNiWV steel. A superclean version of the European steel wasnot attempted due to hardenability reasons.

In the GECffoshibalJSW project, based on the successful results of laboratorywork by Bodnar and Hansen of BethForge and Jaffee of EPRI on an improvedCrMoV rotor steel [lEl, a superclean 2.5%Ni-CrMoV trial high pressure-lowpressure rotor shaft was produced by the Japan Steel Works Ltd. (JSW) [2,2l. Theshaft consisted of a 1270 mm (50.0 in) diameter high pressure (HP) section and a1750 mm (68.9 in) diameter low pressure (LP) section. The chemistry anddimensions of the ingot were optimized based on a study at JSW with specialattention being directed to the reduction of chemical segregation in the ingot. Thesteel was refined in a basic electric arc furnace and a ladle refining furnace, and thencast into a 90 metric ton (99 ton) ingot using the vacuum carbon deoxidation (VCD)process. The final chemistry was 0.23%C, 0.03%Si, 0.03%Mn, 0.003%P, 0.001 %S,2.5%Ni, 1.6%Cr, 1.2%Mo and 0.25% V. After the preliminary heat treatment bydouble normalizing, the HP and LP portions were quenched differentially. The LPportion was water spray quenched from 935°C (1715°F) and the HP section wasforced air cooled from 950°C (1742 OF). The LP portion of the rotor shaft wastempered to two tensile strength levels of 860 Mpa (125 ksi) and 790 MPa (115 ksi)by two stage tempering for evaluation. The FATT at the center of the LP and HPsections were approximately 22°C (72°F) and 3°C (3rf), respectively, for astrength level of 790 MPa (115 ksi). Increasing the tensile strength of the LP sectionto 860 MPa (125 ksi) increased the FATT to 49°C (120°F). Creep rupture testsrevealed that the creep rupture strength of the HP section was equivalent to thestrength of a conventional CrMo V rotor steel. Generally, other characteristics suchas fatigue properties and fracture toughness were equal to or in excess of those of thestandard CrMoV turbine rotor steel levels. However, notch bar creep rupture testsshowed a tendency for notch weakening and the notch rupture ductility wassomewhat marginal and could interfere with the use of this alloy in steam turbineapplications. Figure 10 illustrates the crossover of the notch bar stress rupture curvewith respect to the smooth bar curve. The apparent cause of this difficulty was the


removal of niobium from the original proposed chemistry because of the concernabout the potential adverse effects of chemical segregation. The proposed solution tothis problem would involve the restoration of the niobium in conjunction withelectroslag remelting (ESR) to minimize the segregation tendency.

The EPRI-Europe rotor project involves evaluation of a 2%CrMoNiWV steel-which was proposed by Saarschmiede in 1980 [lQ]. The alloy has been meanwhilesubjected to extensive testing in collaboration with turbine manufacturers. In thepast 6 years, 25 ro(ors for HP/LP and 20 rotors for HP/IP combination turbines havebeen manufactured. The purpose of the EPRI project was to optimize a clean steelversion of this alloy using lab heats, followed by production and evaluation of a trialrotor using the optimized composition. A combination rotor of 1250 mm (50")diameter in the HP section and 1750 mm (69") diameter in the LP section has beenproduced from a 95 ton ESR ingot of 2300 mm (91") diameter [31.]. The rotor wasfirst quenched and tempered to an 0.2% yield strength of approx. 703 MPa (102 ksi)by differentially cooling the HP section (air/water quenching simulating oilquenching) and the LP (water quenching) section from 950°C (1745 OF) austenitizing


temperature. This resulted in a FATT of approx. -15°C (_10°F) in tangential rimlocations and of +55 °C (131 OF)in the core of the HP section as well as of +57°C(135°F) in the core of the LP section. Following this testing, a through-section radialcore of 395 mm (15.5") diameter was taken from the HP section and from the LPsection for further testing by the turbine manufacturers. Subse-quently, the rotor wastempered to a lower 0.2% yield strength of approx. 628 MPa (91 ksi) and tested inthe same way as the higher-strength material. The FATT values in this case werefrom -28 to -40°C (-18 to -40°F) in tangential rim locations, + 14°C (5rF) in thecore of the HP section and +19°C (66°F) in the core of the LP section. The resultsavailable so far have met the assumptions with respect to this steel. They need,however, to be substantiated by the further investigations, particularly the long-termtesting, to be conducted by turbine manufacturers.

In the absence of long term creep data on the clean steel version of the2%CrMoNiWV trial rotor, we could still use the data on the various productionrotors compiled previously for purposes of comparison [.32] with the assumption thatthe clean steel version will exhibit higher rupture strength. Figure 11 shows acomparison of the Larson-Miller rupture data for the 2.5%Ni superclean CrMoV andthe production version of the 2%CrMoNiWV alloys. The data for both fall withinthe General Electric Company's scatterband of conventionallCrlMo 0.25V HP/IProtor steels. Slightly higher long term rupture strength is indicated for the2.5%NiCrMoV alloy compared to the 2%CrMoNiWV alloy. Figure 12 compares therupture ductilities of the two alloys. The tendency for the decrease in ruptureductility with time is evident for the 2.5%NiCrMoV alloy consistent with the notchweakening observed to this alloy.

Several parallel developments have been taking place in the meantime withrespect to HP/LP single shaft development and have been reviewed byViswanathan [31]. With the exception of the EPRI-European 2%CrMoNiWV alloy,the experience base and long term performance data is insufficient at this time. Allof the alloys have HP bore FATTs that are much lower than conventional HP rotoralloy lCrlMol/4V (80 to 100°C).

INDUSTRY SURVEY RESULTS

A worldwide survey of steelmakers, equipment manufacturers and utilities wasconducted by Nutting in behalf of the Electric Power Research Institute [.3.4], Resultsof his survey show that the capability for producing clean/superclean steels existsquite extensively. There would appear to be no great difficulty in obtaining suitableforgings, the initial cost premium is not high and there appears to be no difficulty infabrication. Out of a total tonnage of 200,000 tons during the four year period 1990-1993 supplied by the responding steelmakers to the power industry, roughly 7,000 to10,000 tons (about 5%) were produced to EPRI superclean steel guidelines.



The steelmakers, whether they supplied steel or not, reported values of theresidual element contents of their conventional steels which were not widelydifferent from those required by the superclean guidelines. The European producershad greatest divergence in relation to tin and arsenic whilst the Japanese producershad greatest divergence in relation to arsenic with only one producer havingproblems with tin. Antimony levels were all very close to the guidelines. All thesteelmakers produced power generation steels to meet their own internal standards.

There was some demand from the steelmakers to see a relaxation of the EPRIguidelines for superclean steels. One company wished to have the Mn specificationraised from 0.02-0.05% to 0.08-0.12%, whilst another company expressed a similarview with a Mn relaxation to 0.1%.

The European steelmakers favored a relaxation in the residual element limitsparticularly for Sn, As and Sb. Such a request can clearly be related to the problemsof obtaining low residual scrap in Europe, but the request in relation to antimony isstrange in view of the fact that their normal production steels meet the required limitas mentioned above.

Another interesting view was that chemistry control was the concern of the plantbuilder, or plant user, and that the present guidelines could be met without too muchdifficulty. It must be remembered, however, that compositional requirements doinfluence the price. There was surprising unanimity amongst the steelmakers thathad supplied superclean steels that a premium of 15-20% would be charged inrelation to conventional grades for ingot weights up to 100-150 tonnes - somewhathigher premium of 40% was quoted for ingot weights in the 300-600 tonnes range.Companies that had not supplied superclean steels anticipated a premium of 30%- avalue widely quoted in the literature.

Views were expressed about the influence of a relaxed specification on theexpected premium. The same company that suggested a relaxation in the manganesecontent to 0.08 to 0.12% believed that this would reduce the premium to about 5%.A European view was that if the residual content could be relaxed but keeping theexisting siljcon, sulfur and manganese superclean levels, the premium would bereduced to 5%. The steelmakers were unanimous in their views that any increase indemand for superclean steels coming either from the power generation or otherindustries, would not influence the premium.

There was a surprising divergence of views as to whether a demand forsuperclean steels would develop from other than the power generation industries.About half anticipated some demand - about half did not. This divergence could bedue to differences in the customer base of the different companies and from the


details given in the replies to the questionnaire it was impossible to assess thisaccurately. But where positive answers were given, the increased demand was seenas likely to come from the chemical industry, and for pressure vessels and foroffshore applications.

In reply to a question about the accuracy of analytical procedures, there weresome remarkable differences in views. In some cases the answers were givenspecifically in relation to 3.5 Ni-Cr-Mo- V steel and in others they covered the wholesteel product range. The steelmakers whose answers covered the whole range oftheir steel products in general gave replies which showed lower tolerances than thosewho gave specific tolerances. The analytical tolerances were in general lower (that isto say more accurate) for the Japanese steelmakers than the others. In replying to thequestion about the need for improved analytical methods, all the steelmakers felt thatthese were accurate enough for dealing with the specification limits. There was onerequest for standardized methods to be established and another for improvedanalytical methods for hydrogen.

In commenting on the need for future work, all the steelmakers supported theviews put forward in the questionnaire about the possible study of the synergismbetween the normal alloying elements and the impurities such as S, P, As, Sb and Sn,and an assessment of the level of impurity levels below which no furtherimprovement of in service performance occurs.

It was also thought that further studies should be carried out on rotors coming outof service to give realistic data on in service embrittlement. Two steelmakersrequested further work on establishing the "safe" upper limit for manganese. This islinked to the view that a relaxation of the manganese specification would lower thesuperclean premium.

Further views were linked with the development of a superclean specification forCr-Mo- V steel and for a standard for HP-LP combined rotor. Whilst papers havealready been written on the development of superclean 12% Cr steels, there were anumber of requests for further work in this area with a view to establishing asuperclean standard for this class of material.

It was generally agreed that the need for clean/superclean steels must be placed inthe context of design and operation requirements. It is not our objective to makesteels as clean as they can get, but as clean as we need. For instance, what FATT,K1C' etc. are needed (rather than the lowest we can achieve at all costs) for differentapplications and what cleanliness is adequate to meet these requirements will have tobe decided by the turbine manufacturers based on customer needs on a case-by-casebasis. Some broad classifications of various applications and the correspondingcleanliness required would however be very useful in helping the utility cUstomer


understand his needs better. Further, more detailed cost/benefit analysis todemonstrate the value of superclean steels would also go a long way in promotingindustry acceptance of these steels.

SUMMARY

The general benefits of cleaning up steels have been amply demonstrated usinglaboratory heats as well as full-sized rotors. Significant improvements in fracturetoughness, FATT and freedom from temper embrittlement have been achieved whilemaintaining or improving creep strength and ductility of steels. In the case ofsuperclean LP rotor steels, improved SCC resistance has also been achieved. It hasbeen well recognized that these improved properties can lead to extended life,avoidance of forced outages, increasing inspection interval, more rapid start-up andcycling ability, avoidance of need to pre-warm before loading and increasedefficiency due to higher LP rotor cross over temperature. The capability forproducing and fabricating superclean steel forgings exist worldwide. During theyears 1990-93 alone nearly 10,000 tons of steel forgings had been supplied tosuperclean standards representing about 5% of total tonnage sold to power industry.Nearly 30 superclean steel LP rotors in the 700 -1000 MW class have beenmanufactured in Japan alone to combat stress corrosion. Superclean steel forgingshave found wide-spread acceptance in combustion turbine components in Europe. Inthe U.S., superclean steel standards have nearly been met in several recent forgingsordered even though not required by the purchase specification. The premium forsuperclean steels in terms of price seems to be in the 20 to 30% range, although manysteelmakers believe that this could come down to as low as 5% with slight relaxationin the specification for Mn or for the other tramp elements. Demonstration ofbenefits based on quantitative cost/benefit analysis, establishment of analyticalstandards for tramp elements, definition of realistic threshold levels for impuritiesbased on study of impurity-alloying elements synergisms in specific classes of steels,and extension of technology to 9-12% Cr steels seem to be some of the steps towardgaining more widespread acceptance of clean/super-clean steels.

ACKNOWLEDGMENTS

The author is grateful to Dr. E. Potthast, Mr. W. Wiemann, Professor J. Nuttingand Mr. Asa Kaplan for reviewing this manuscript and suggesting numerous usefulrevisions.


REFERENCES1. Viswanathan, R., and Jaffee, R. I., "Toughness of Cr-Mo- V Steels for Steam

Turbine Rotors," Journal of En~ineerin~ Materials and Technolo~y. Vol. 105,December 1983, p. 286.

2. Viswanathan, R., and Jaffee, R. I., "Metallurgical Factors Affecting theReliability of Fossil Steam Turbine Rotors", Journal of En~ineerin~ for GasTurbines and Power, Vol. 107, July 1987, p. 642.

3. Viswanathan, R., Dama~e Mechanisms and Life Assessment of Hi~hTemperature Components, ASM International, Metals Park, 1989, pp. 303-306,400-410.

4. Nutting, J., "Overheating Characteristics of Clean 3.5NiCrMoV", Final ReportRP2741-01, Electric Power Research Institute, Palo Alto, 1987.

5. Timo, D. P., Curran, R. M., and Plazec, R. 1., "Rotor Forgings for SteamTurbine Generators", R. I. Jaffee, ed., Electric Power Research Institute ReportWS79-235, 1979.

6. Curran, R. M., Newhouse, D. L., and Newman, J. c., "The Development ofImproved Rotor Forgings for Modern Large Steam Turbines", ASME Paper82JPGCPWR-Vol. 25, Joint Power Generation Conference, Denver, 1982.

7. Kramer, L. D., Randolph, D. D., and Weisz, D. A., "Analysis of The TennesseeValley Authority, Gallatin Unit No.2, Turbine Rotor Burst", Winter Meeting ofASME, New York, December 5-10,1976.

8. Swaminathan, V. P., Steiner, J. E., and Mitchell, A., "Advanced Rotor Forgingsfor High Temperature Steam Turbines", Report CS-4516, Electric PowerResearch Institute, Palo Alto, May 1986.

9. Swaminathan, V. P., and Jaffee, R. I., "Significant Improvements in Propertiesof CrMo V HP Rotors by Advanced Steel Making", Life Assessment andImprovement of Turbogenerator Rotors for Fossil Power Plants, R.Viswanathan, ed., Pergamon Press, New York, 1985, p. 657.

10. Yamada, M., Toshiba Corporation Personal Communication to J. Nutting,September 2, 1994.

11. Ewald, J., Berger, c., Keinburg, K. H., and Wiemann, W., "Present QualityLevel of Heat Resistant Forgings Made of 1%CrMoV Steels", Steel Research51, 1986, pp. 83-92, 172-177.

12. Ohhashi, 1., Bodnar, R. L. and others, "High Purity Steels for UtilityComponents", Report NP-5399, RP2060-1, -2, Electric Power ResearchInstitute, Palo Alto, September 1987.

13. Newhouse, D. L., "Guide to 12Cr Steels for High and Intermediate PressureTurbine Rotors for the Advanced Coal-Fired Steam Plant", Report CS-5277,RPI403-7, Electric Power Research Institute, Palo Alto, July 1987.

14. Coulon, P. A., "Improved Superclean 11%Cr Steel for High TemperatureTurbine Rotors", Superclean Rotor Steels. Proc. of a Workshop at Sapporo,Japan (EPRI GS-6921), R. I. Jaffee, ed., Pergamon Press, New York, 1989.


15. Sato, K., et aI., "High Purity 12%Cr Martensitic Steel", Superclean RotorSteels. Proc. of a Workshop at Sapporo. Japan (EPRI GS-6921), R. I. Jaffee,ed., Pergamon Press, New York, 1989.

16. Coulon, P. A., "Influence of Impurities in Improved 12% Chromium Steel forHigh Temperature Steam Turbine Rotor", Symposium on Residual andUnspecified Elements in Steel, ASTM Al Proceedings, Bal Harbor, Florida,November 11, 1987.

17. Sato, K., et aI., "Manufacturing of High Purity 12% Cr Rotor Forgings", Ck.anSteels Technology, R. Viswanathan, ed., ASM International, Metals Park,1992.

18. Berger, C., Mayer, K. H., and Beech, S. M., "High Temperature rotor Forgingsof High Strength 10% CrMoV Steels", 12th International ForgemastersMeeting, Forging Industry Education and Research Foundation, Cleveland,1994.

19. Taira, 1., et aI., "Development and Production of High Purity 12%Cr Steels forGas Turbine Disc Material", Proceedings of the 12th InternationalForgemasters Meeting, Chicago, published by Forging Industry Education andResearch Foundation, Cleveland, 1994.

20. Shiga, M., et aI., "Development of 12%Cr Steel for Gas Turbine Disc", 12thInternational Forgemasters Meeting, Forging Industry Education and ResearchFoundation, Cleveland, 1994.

21. Jaffee, R. I., Machner, P., Meyer, W., and Steiner, 1. E., "Production andProperties of a Superclean 3.5NiCrMoV LP Rotor Forging", EPRI ProjectRP1403-8, lronmaking Steelmaking, 13, 1986,322-326.

22. "Production and Properties of a Low Pressure Model Rotor with RelaxedSpecification", YEW Final Report to EPRI on RP2471-3, 1986.

23. Richman, R. H., and McNaughton, W. P., "Superclean Steel Development:Status Report", GS-661O, RP1403-32, December 1989.

24. Watanabe, 0., Yoshioka, Y. and Schwant, R. c., "Evaluation of a SupercleanNiCrMoV Low Pressure Turbine Rotor Forging for Advanced SteamConditions", Report of RP1403-15, Electric Power Research Institute,December 1991.

25. Yoshioka, Y., Watanabe, 0., and Miyazake, M., "Production and Evaluation ofSuperclean LP Rotor Forgings", Proc. of the Sapporo Workshop, ElectricPower Research Institute, GS-6921, 1989.

26. Scarlin, R. B., and Denk, J., "Stress Corrosion Cracking Behavior of a Clean3.5% NiCrMoV Steel", Report GS-6907, RP1403-18, Electric Power ResearchInstitute, Palo Alto, July 1987.

27. Richman, R. H., and McNaughton, W. P., "Superclean Steel Development: AGuide to Utility Use", Report GS-6612, RP1403-32, Electric Power ResearchInstitute, December 1989.


28. Bodnar, R. L., Hansen, S. S., and Jaffee, R. I., "Improved Superclean NiCrMoVRotor Steel", Report ER-6887, RP2426-4, Electric Power Research Institute,Palo Alto, July 1990.

29. Tanaka, Y., Watanabe, 0., Kaplan, A. and others, "Production and Propertiesof a Superclean 2.5%NiCrMoV HP/LP Rotor Shaft", Report TR-103689,RPI403-55, Electric Power Research Institute, Palo Alto, February 1994.

30. Finkler, H., and Potthast, E., "Presentation of a New Steel for High PressureShafts", Rotor Forgings for Turbines and Generators, R. I. Jaffee, ed.,Pergamon Press, New York, 1980.

31. Potthast, E., Viswanathan, R., and Wiemann, W., "Advanced 2%CrMoNiWVSteel for Combination Rotors", Proceedings of the 12th InternationalForgemasters Meetin~, Chicago. Published by Forging Industry Education andResearch Foundation, Cleveland, 1994.

32. Wiemann, W., "The Development of an Improved 2%CrMoNiWV HP/IPRotor", COST Project Report EVR 13452 EN, J. B. Marriott, ed.; 1991.

33. Viswanathan, R., "Application of Clean Steel/Superclean Steel Technology Inthe Electric Power Industry - Overview of EPRI Research and Products",Clean Steel' Superclean Steel, J. W. Nutting and R. Viswanathan, ed.; TheInstitute of Materials, London, 1995, pp 1-33.

34. Nutting, J. W., "The EPRI Survey on Superclean Steels", Clean Steel'Superclean Steel, 1. W. Nutting and R. Viswanathan, ed.; The Institute ofMaterials, London, 1995, pp 34-53.

Akihiro Itoh,' Hitohisa Yamada,' and Tomoo Takenouchi2

PREDICTION AND CONTROL OF SEGREGATIONS IN CrMoV STEEL INGOT FOR MONOBLOCKHLP ROTOR FORGINGS USING EXPERIMENTAL RESULTS OBTAINED FROM 8 TON SANDMOLD INGOTS

REFERENCE: Akihiro, I., Hitohisa, Y., and Tomoo, T., "Prediction and Controlof Segregations in CrMoV Steel Ingot for Monoblock HLP Rotor Forgings UsingExperimental Results Obtained from 8 ton Sand Mold Ingots, " SteelForoinos:Second Volume. ASTM STP 1259. E.G. Nisbett andA.S. Melilli, Eds.,American Society for Testing and Materials, 1997.

ABSTRACT: Remarkable segregation was observed in the modified super cleanCrMoV steel forgings for electric power generation applications. Therefore,to make the mechanism clear, effect of such elements as Mn, Ni, Cr and Moon segregation was studied, using 8 ton sand mold ingots, the solidificationtime of which corresponds to that of 100 ton ingot.

As a result, we found that factors controlling the segregation are 6-solidification ratio, mean partition coefficient, density difference ofmolten steels between bulk and segregated liquid at the solidification frontand so on. These factors can be calculated from chemical composition ofsteels.

Then, based on prediction model obtained from the experimental results,chemical composition and shape of ingot were tried to be changed. As forchemical composition, such elements as Mo which is heavy and 6 -former werekept low in the specification range. And, as for ingot shape, height todiameter ratio HID was kept high to shorten the solidification time. Thecarbon segregation along the axis of ingot was kept relatively low by thisingot design.

The eutectic Nb(C,N) inclusions which give bad effect on the toughnesswere also investigated. The conditions for the formation of such inclusionswere made clear and then predicted and controlled by the calculation fromchemical composition.

By this technical development, quality of ingot for HLP rotor forgingswas extremely ~mproved.

KEYWORDS: clean steel, solid fraction at peritectic temperature, meanpartition coefficient,6 and y solidification ratio, density difference,height to diameter ratio, prediction of segregation, eutectic Nb(C,N)

Introduction

CrMoV steel grades [1], used as rotorshaft in the electric powergeneration plant as a result of high creep rupture strength, also required

'Research engineer and senior research engineer, respectively, The JapanSteel Works, Co. Ltd., 4 Chatsu-machi, Muroran, Hokkaido, Japan.

'A vice-head of management and planning, The Japan Steel Works, Co. Ltd.,1-1 Nikko-cho, Fuchu, Tokyo, Japan.

305


high toughness [2] because of their new application to monoblock HLProtorshaft in the combined cycle power generation systems. To meet thisdemand, chemical compositions such as Ni, Cr and Mo were modified [1] , however,remarkable segregations were observed [.1] in the commercial ingot and becausethe expected mechanical properties such as toughness and ductility couldnot be obtained. Therefore, effect of alloying elements on the macro andchemistry segregations were studied using 8 ton sand mold ingots. Basedon the prediction model obtained from experimental results, chemicalcomposition and shape of ingots were optimized to reduce segregations.Furthermore, formation of the eutectic Nb(C,N) inclusions which give badeffect on the toughness was also studied.

By these technical developments, the quality of ingots for HLP rotorforgings was extremely improved.

Experimentals

1.Production and investigation of 8 ton sand mold ingotsTable 1 shows the chemical compositions of 8 ton sand mold ingots, in

which Mn, Ni, Cr, Mo and Nb contents were mainly changed in CrMoV systemtogether wi th carbon steel for comparison. Hereafter, steels with di fferentMn and Nb content in 1.7Ni-l.7Mo system are called as 1.7Ni-l.7Mo-Mn and1.7Ni-1. 7Mo-Nb.

Steels melted in the EAF and refined in the LRF were poured into sand moldwith dimensions of 840mm ¢ X 1015mmh for body and 1030mm ¢ X 600mmh for hottop. The solidification time of 8 ton sand mold ingot was assumed to beabout 30h[2J which corresponds to that of 100 ton ingot poured into ironmold. These ingots were then cut horizontally and vertically at the centerand macrostructures and chemical compositions were investigated.

2.Condition of formation of eutectic Nb(C,N)Electric resistance furnace shown in Fig. 1 was used in this experiment.

The sample, whose chemical compositions are based on 1.7Ni-1.7Mo-Nb steelwith different Nb and C contents, was melted at the temperature of 1823K

in Alz03 crucible. The sample weight is about BOg. After mel ting, the samplewas cooled at cooling rate of IBOK/h to match that of the center of largeingot. When the sample reached the desired temperature, it was quenched.The sample was cut at 15mm from the bottom, and area ratio of Nb(C,N) wasmeasured for the investigation of formation condition.

Experimental results

I.Segregation behavior of sand mold ingotsFig.2 shows representative macrosegregations observed in the ingots.

As shown, though relatively thin segregation streaks usually observed inthe conventional ingots were seen in O.2Ni-O.lMo steel, remarkably heavystreaks having inverse inclination angle were observed at the location fromthe near surface to the center of the ingot. However, no macrosegregationswere seen in 3.6Ni-O.4Mo steel. Therefore, formation of macrosegregationswas found to be much influenced by the chemistry of steels.

Next, segregation indexes of carbon Is along the longitudinal axis ofingot are shown in Fig.3, which indicates that though O.4Ni-l.3Mo and3.6Ni-O.4Mo steels show relatively small increase in Is toward the top ofingots, 1.7Ni-l. 7Mo base steels and 2. 6Ni -1. 4Mo steel show remarkableincrease in Is. In this case, analysis was done on the locations excludingmacrosegregations. It was assumed that carbon segregations were restrainedby increasing Mn and enhanced by increasing Mo and Nb.


2.Formation Condition of eutectic Nb(C,N) inclusionsFig.4 shows the effect of quenching temperature on the area ratio of

eutectic Nb(C,N) inclusions, showing that eutectic Nb(C,N) does not formover the temp~rature of l748K. The area ratio rapidly increases by decreaseof the temperature from l748K to l698K, then decreases toward l573K, belowwhich the area ratio is kept constant.

Fig.5 shows the effect of carbon and niobium contents on the formationof Nb(C,N). It can be seen from the figure that the critical condition forthe formation of eutectic Nb(C,N) is NbXC=O.012, under which Nb(C,N) doesnot form, so that Nb and C contents should be kept low in the range themechanical properties required are satisfied.

ITOH ET AL. ON PREDICTION AND CONTROL 309


Fig.6 shows the effect of difference between critical formationtemperature of 174 8K obtained from Fig. 4 and solidus temperature Tson the area ratio of Nb (C,N), showing that the larger temperaturedifference increases Nb(C,N) formation. We found from the mechanicaltest that area ratio of Nb(C,N) less than 0.02% does not give badeffect on the mechanical properties. It can be found from Fig.6 that, whenthe temperature difference is less than about 8SK, the area ratio ofNb(C,N) can be kept less than 0.02%. Therefore, the chemical compositionsof steel should be controlled to make sure Ts of that is over 1663K.

ITOH ET AL. ON PREDICTION AND CONTROL 311

Discussions

1.Factors controlling segregationFig.7 schematically shows the relationship between carbon segregation

index and distance from bottom. Distance from bottom withoutsegregation (Is=l) ~x and gradient of carbon segregation dIs/ dx are affectedby chemical composition of steels, then, factors controlling ~x and dIs/dxare investigated.


References[l]H.Usuda and I.Tsuji:Tetsu-to-Haqane, 57(1971), P.1213.[2]Y.Kadoya and T.Goto:Tetsu-to-Haqane, 76(1990), P.131.[3] H. Yoshida, Y. Ikeda and Y.Tanaka: EPRI Super Clean Steel Review, Sapporo,- Japan, Aug. 1989.[1] H. Yamada, T. Sakurai and T. Takenouchi: Tetsu-to-Haqane, 75 (1989), P. 97.[2] H. Yamada, T. Sakurai, T. Takenouchi and Y. Iwanami: 1mQ.nQ, 59 (1987), P. 85.[Q]T.Takahasi,M.Kudoh:Report of the 19th Committee. Japan Soc. for the

Promotion of Sci.,Rep.No.10227(1980).[1] H. Yamada, T. Sakurai and T. Takenouchi: Tetsu-to-Haqane 76 (1990), P. 438.[~JW.G.Pfann:Journal of Metals, 4(1952), P.747.[2] H. Yamada, T. Sakurai and T. Takenouchi: Tetsu-to-Haqane 76 (1990), P. 132.

Ryuichi IshiV Yoichi Tsuda,1 and Masayuki Yamada I

HIGH STRENGTH 12% CR HEAT RESISTING STEEL FOR HIGHTEMPERATURE STEAM TURBINE BLADE

REFERENCE: Ishii, R., Tsuda, Y., and Yamada, M., "HIGH STRENGTH12% CR HEAT RESISTING STEEL FOR HIGH TEMPERATURE STEAMTURBINE BLADE, " Steel Forgings: Second Volume, ASTM STP 1259, E. G. Nisbettand A. S. Melilli, Eds., American Society for Testing and Materials, 1997.

ABSTRACT: The factors to attain the high creep rupture strength for 12% chromiumferritic steels were studied. Large amount of Laves phase and high content of soluteelements during creep were important to maintain the creep resistance, and thedegradation correspond to the recovery of matrix. Based on these studies, the chemicalcomposition of 1O.5Cr-2.5W-lCo-0.2Re-VMoNbNB steel with large amount of Lavesphase and high solute tungsten content was determined for the new blade material. Theproducibility of this steel were verified with 2 ton ESR ingots which were successfullymanufactured by various industrial processes. The mechanical properties of the trial barswere sufficient to be applicable to the blade, especially the creep rupture strength was veryhigher than that of llCr-lMo-l W-VNbN steel being used. In near future, this new steelwill be applied to commercial power plants with steam condition of 593°C or higher as theblade material.

KEYWORDS: steam turbine, blade, high chromium ferritic steel, creep rupturestrength, Laves phase, tungsten, rhenium

An increase in operating temperature improves the thermal efficiency of fossil-frred

IEngineer, Specialist, and Manager, respectively, Toshiba Corp., Heavy ApparatusEngineering Laboratory, 2-4 Suehiro-cho, Tsurumi-ku, Yokohama 230, Japan

317


power plants. The enhanced operating temperature, however, requires the advancedmaterial with high creep rupture strength for high temperature rotating components suchas blades. For these components, ferritic steels are desirable because of having suitablethermal expansion and thermal conductivity.

The recent development of "modified 12% chromium steel" made it possible toincrease the steam temperature to 593°C [1]-[~].Current tendency of fossil-frred powerplants, however, reveal that stronger ferritic steels than "modified 12% chromium steel"are required for advanced power plants with the steam temperature of 600°C or higherL4],[~].

The authors have developed a 1O.5Cr-2.5W-ICo-0.2Re-VMoNbNB steel for hightemperature blade material. This steel has superior creep rupture strength to "modified12% chromium steel", due to solid solution strengthening by tungsten and rhenium, andprecipitation strengthening by Laves phases and carbides.

This paper describes the strengthening and degradation mechanism of 12%chromium steel with tungsten content, the concept for the development of new steel andthe evaluation results of trial bars for the blade.

FUNDAMENTAL INVESTIGATION

Stren&thening mechanism

Strengthening mechanism of 12% chromium steel with tungsten content werefrrstly studied by using TOS202 and TAF650. Typical chemical compositions and creeprupture strength of both steels are shown in Table 1 and Fig. 1.

TOS202 wa~ developed by Toshiba in 1980's for high temperature blade [~]. It hasbeen used as blades for high and intermediate stages in several commercial power plants inJapan. TAF650, one of the commercial steels for blades, was developed by Hitachi metalsLTD. and professor Toshio Fujita [Q].

Fig. 2 shows the amount of solute tungsten and extracted residue of TOS202 afterbeing aged and crept. All specimen show the depletion of the solute tungsten down toabout 0.5 mass % for more than 3,000h and the extracted residue gradually increased. Fig.3 shows the analysis results for TAF650. In this steel, almost the same results wereobtained for more than 1,000h except that the amount of residue was larger than that ofTOS202. They seem to become almost stable in both steels for more than 3,000h.


Fig. 4 shows the carbon extracted replicas and the transmission electronmicrographs for TOS202 and TAF650 crept at 650°C-176.4 MPa. There was littledifference in precipitation morphology between both steels except that the amount ofblocky Laves phase which were consumed to cover grain boundaries and martensitic lathinterfaces was very larger in TAF650 than that of in TOS202. The shape of lath orsubgrain was clearly identified in TAF650 because of continuous formation of Lavesphases and Mz3C6 carbides at these sites. In TOS202, however, it was obscure because theamount of the precipitates was not enough to cover these sites.

From these results, increasing the initial tungsten content increase not the solutetungsten but the amount of Laves phase precipitated during creep/aging mainly at grainboundaries and martensitic lath interfaces. It was considered that the annihilation ofdislocation, which caused the creep resistance to decrease, was suppressed by continuousformation of Laves phases and Mz3C6 carbides. This phenomena seems to be similar to"intergranular precipitation strengthening" due to secondary phase, which was confirmedquantitatively in nickel-base alloy [1],[~].Therefore, the precipitation of large amount ofLaves phase is one of the important factors to maintain high creep strength.

Degradation mechanismnDegradation mechanism of 12% chromium steel withtungsten content was investigated by using TAF650. Fig. 5 shows the change in Vickershardness of several heats of TAF650 which were subjected to different heat treatments.

ISHII ET AL. ON TURBINE BLADE 321

Up to 3,OOOhat 600°C, the hardness of each specimen was variously changed with theresult of annihilation of excess dislocation and precipitation of Laves phases. For morethan 3,OOOh,however, the hardness decreased and became almost the same valueregardless of the difference of initial conditions.

FIG. 5--Change in Vickers hardness of several heats ofTAF650 with aging time at 600°C.

Fig. 6 shows the relation of minimum creep rate between as tempered specimenand aged specimen crept at the same testing condition. Though the amount of theprecipitates became about constant for more than 1,OOOh,the increase in the minimum

FIG. 6--Relation of minimum creep rate between as tempered specimen and aged specimen.


creep rate was recognized in the specimen aged for more than 3,OOOh.The transmission electron microscopy observation revealed that the change in the

morphology offme precipitates in lath was scarcely identified for more than I,OOOh[2].These results suggest the degradation of this steel mainly correspond not to theprecipitation behavior before or during aging but to the softening of matrix. Fig. 7 showsthe schematic drawings of this consideration.

FIG. 7--Schematicdrawingsof the relationbetweenthe changeinmicrostructureand the creepresistance.

Concept of Ne~ Steel--Based on the above mentioned results, the concept for thedevelopment of new steel was considered. Laves phase has been regarded as harmfulphase which has detrimental effect on toughness, rupture strength and so on. However,current researches for 9-12% chromium steels with high tungsten content reveal that theprecipitation of this intermetallic phase may be effective on creep properties under somecircumstance [2]-[12]. In addition to this, detailed studies of precipitation behavior of hightungsten steel have been performed [U]. On the other hand, the "sigmoidal behavior"seems to be unavoidable as same as in other ferritic steels [14]-[16]. Then, in order tomaintain the creep resistance for long-term, it is important to take account of thefollowing:- Maintaining large amount of Laves phase;- Continuous formation of Laves phases and M23C6 carbides at grain boundaries and


martensitic lath interfaces;- Increasing the content of solute elements contributing to solid solution strengthening;- Precipitation of the fme particles (MX carbonitride and Laves phase) in martensitic lath.Fig. 8 shows the schematic drawing of this concept.


Fig. 9 shows one of the important results of these investigation. The amount ofsolute tungsten in rhenium doped steel was about twice as much as rhenium free steel atthe same creep time. Based on the test results, 1O.5Cr-2.5W steel with 0.2% rheniumcontent was selected to be the proper composition for the trial 2 ton bars.

MANUFACTURE AND EVALUATION OF 2 TON TRIAL BARS

Products

The chemical compositions, melting processes and heat treatment of trial bars areshown in Table 3 and Fig. 10. These steels were melted by vacuum induction furnace andcast into about 2 ton ingots. The forging or rolling was carried out as hot working process.The [mishing size was 135x80xL mm. One bar was manufactured by electroslag remeltingprocess (ESR). The other bars were manufactured by vacuum electroslag remeltingprocess (YSR = ESR under vacuum or low pressure condition). The products werenormalized at 1,120°C for 3h followed by oil quenching, and then tempered at 680°C for5h followed by air cooling. In addition, all bars were reformed and stress relieved at 650°Cfor 5h followed by air cooling.

Producibility--Castability and forgeability of this new steel were equal to theconventional 12% chromium steel. A target shape was successfully obtained. Some ofthem, however, were of low cleanliness in microstructure. In case of manufacturing large


size ingots made of low silicon, low aluminum, high tungsten, boron and nitrogen dopedsteel, it is very difficult to control boron content before and after electroslag remelting, andit is also difficult to prevent the formation of large inclusions which deteriorate itsmechanical properties under some circumstances. By choosing proper slag composition,however, it is easy for the steel containing above mentioned elements and the size like thisto control its composition and to gain the high cleanliness in microstructure. Distributionof each elements was homogeneous in both locations of all 2 ton bars.

Microstructure--Fig. 11 shows the optical micrographs of one of the 2 ton bars. It,revealed the tempered martensitic microstructure without o-ferrite formation. The prioraustenitic grain size number (ASTM) of the bars manufactured by different methods wereabout between No.3 and No. 4.5.

Fig .12 shows the as-polished microstructure of2 ton bars. In Fig. 12 (a) and Fig.12 (b), many large inclusions, such as BN and WB, were recognized. In Fig. 12 (c),however, these types of inclusions were scarcely identified. The reason of this was due


to the change in the slag composition at the remelting. However, it is hard for large ingots,for example rotor forging, to avoid the formation of such inclusions. The alteration of thechemical composition may be taken into account.

Tensile--Fig. 13 shows the actual strength of several 2 ton bars and averagestrength of TOS202 in the curved line. Tensile strength of 2 ton bars were lower than thatof TOS202, and 0.02% yield stress were almost equivalent to that of TOS202. These

values, however, are enough to design the blades. Fig. 14 shows that 2 ton bars haveenough ductility at elevated temperature.

FIG. 14--Tensileductilityof 2 ton trial bars.

Tou~hness--Fig. 15 illustrates the Charpy impact test results for several 2 ton barscompared with that of TOS202 in the curved line. The initial impact value for 2 ton bars

FIG. 16--Creeprupture strengthof newlydevelopedsteelwith 0.2 mass% rheniumcontent.

The rupture strength of 0.2% rhenium containing steel revealed higher than that ofTOS202 at high parameter value. The results from 2 ton bars were nearly equal to that of50 kg test melt. The influence of a little change in the chemical composition and the size ofingot on creep rupture strength were not recognized. From these results, the strength for0.2% rhenium containing steel is expected over 150 MPa as lOsh at GOOoe by using the


Larson-Miller parametric analysis (C=25). The "sigmoidal behavior" was not confIrmedyet.

APPLICATION

The chemical composition of the new steel was modifIed a little, and was namedTOS203. TOS203 is planned to apply to 700 MW steam turbine for commercial powerplants with high temperature steam condition (24.1 MPa, 593°C/593°C) as high andintermediate pressure first stage blades [17]. This turbine is a 3-casings tandem compoundtype with the rotational speed of 3,600 rpm. In the first stage blade of this machine, asingle flow type design will be realized instead of the double flow type design. Accordingly,the flow rate and the output with respect to the first stage blade have been doubled ascompared with the conventional double flow type. Though the blade length should also bedoubled and the centrifugal force results in increase, this machine is able to realize withoutchange in basic designs by using TOS203. Moreover, this steel will be applicable to over600°C class steam turbine blades.

CONCLUSION

In order to obtain a material having improved creep rupture strength for hightemperature steam turbine blade, strengthening and degradation mechanism wereinvestigated on 12% chromium steel with tungsten content. From this investigation, itshowed that the 1O.5Cr-2.5W-ICo-0.2Re-VMoNbNB steel with large amount of Lavesphase and solute tungsten provides the superior creep rupture strength. Based on the testresults, 2 ton trial bars were manufactured and the material properties were evaluated. Theresults are summarized as follows.(1) Continuous formation of precipitates (Laves phases and Mz3C6 carbides) at grainboundaries and martensitic lath interfaces was one of the important factors to maintain thecreep resistance.(2) Keeping the high content of solute tungsten during creep was one of the importantfactors to increase the long-term creep rupture strength, and this was attained by notincreasing initial tungsten content but doping rhenium.(3) Trial bars with the dimension of 135x80xL mm and the weight of 2 ton weresuccessfully manufactured by various industrial processes.(4) The homogeneous distribution of chemical composition, the microstructure with highcleanliness, and the sufficient properties for blade material were confirmed.(5) The 10.5Cr-2.5W-ICo-0.2Re-VMoNbNB steel with high creep rupture strength issuitable to the high and intermediate first stage blades for the advanced steam turbines.

REFERENCES

[1] Ito, F., Kuwabara, K., Miyazaki, M., Fukui, Y., and Takeda, Y., EPRI 1st Inter. Conf.


on Improved Coal-Fired Power Plants, (1988), p.6-127[2] Tsuda, Y., Miyazaki, M., and Kaplan, A., "Advanced 12% Cr Steel for High-Temperature Rotors", Proc. of EPRI 3rd Inter. Conf. on Improved Coal-Fired PowerPlants, (1991)[3] Yamada, M., Watanabe, 0., Yoshioka, Y., and Miyazaki, M., Tetsu to Hagane, vol.76(1990), p.1084 (written in Japanese)[4] Fujita, T., Proc. of 3rd inter. Charles Parsons Turbine Conf .. "Materials Engineering inTurbines and Compressors", Newcastle, UK, April 1995, p.493[5] Metcalfe, E. and Bakker, W. T., The EPRI / National Power Conf .. "New Steels forAdvanced Plant up to 620°C", London, UK, May 1995, p.1[6] Uehara, T., Ohno, T., and Fujita, T., Proc. of3rd inter. Charles Parsons Turbine Conf .."Materials Engineering in Turbines and Compressors", Newcastle, UK, April 1995, p.391[7] Matsuo, T., Kikuchi, M., and Takeyama, M., Heat-Resistant Materials. Proc. of the1st Inter. Conf., Wisconsin, USA, September 1991, p.601[8] Matsuo, T., Nakajima, K., Terada, Y., and Kikuchi, M., Mater. Sci. and Eng., A146(1991), p.261[9] Ishii, R., Tsuda, Y., and Yamada, M., CAMP-ISIJ, vol.8 (1995), p.673 (written inJapanese)[10] Kimura, K., Ishii, R., Matsuo, T., and Kikuchi, M., 123rd committee on HeatResisting Metals and Alloys Report., vol.34 (1993), p.127 (written in Japanese)[11] Ishii, R., Tsuda, Y., Yamada, M., Watanabe, 0., Nakamura, S., and Endoh, H.,CAMP-ISIJ, vol.6 (1993), p.1630 (written in Japanese)[12] Igarashi, M. and Sawaragi, Y. 123rd committee on Heat Resisting Metals and AlloysReport., vol.35 (1994), p.285 (written in Japanese)[13] Mimura, H., Ohgami, M., Naoi, H., and Fujita, T., Materials for AdvancedEngineering. Part I, Liege, Belgium, October 1994, p.361[14] Kimura, K., Kushima, H., Yagi, K., and Tanaka, C., "Creep and Fracture ofEngineering Materials and Structures". The Institute of Materials, Swansea, UK, 1993,p.555[15] Chikwanda, H., Strang, A., and Mclean, M., Materials for Advanced Engineering.Part I, Liege, Belgium, October 1994, p.291[16] Foldyna, V., Jakobova, A., and Kubon, Z., Proc. Inter. Sympo. on Materials Ageingand Component Life Extension, Milan, Italy, October 1995, p.15[17] Shinozaki, S., Kuroki, Y., and Yamaguchi, K., Proc. vol.2 CSPE-JSME-ASME-Inter.Coni. on Po~er Engineering, Shanghai, China, May 1995, p.671

Tsukasa Azuma, 1 Yasuhiko Tanaka, 1 Tohru Ishiguro, 2Hajime Yoshida, 2and YasumiIkeda2

MANUFACTURING AND PROPERTIES OF NEWLY DEVELOPED 9%CrMoVNiNbN HIGH-PRESSURELOW-PRESSURE ROTOR SHAFT FORGING

REFERENCE: Azuma, T., Tanaka, Y., Ishiguro, T., Yoshida, H., and Ikeda, Y.,"Manufact.uring and Propert.ies of Newly Developed 9%CrMoVNiNbN High-pressureLow-pressure Rot.or Shaft. Forging", steel Forqinqs: Second Volume, ASTM STP1259, E.G. Nisbett and A.S. Melilli, Eds., American Society for Testing andMaterials, 1997.ABSTRACT: In order to obtain the improved strength and toughness forhigh-pressure low-pressure rotor shaft forging, fundamental studies usinglaboratory heats were performed on the 9CrMoV base materials, and effectsof chemistry on toughness and creep rupture strength were investigated. Fromthe investigation, it is showed that the superclean 9CrMoVNiNbN steel withreduced si and Mn contents and Ni addition provides a superior strength versustoughness balance. Based on these fundamental studies, a trial high-pressure low-pressure rotor shaft forging with diameter of low-pressuresection of 1750 mm and diameter of high-pressure section of 1200 mm wassuccessfully manufactured from the diameter of 1800 mm, and the weight of65 ton ESR ingot. From the evaluation test results of this trial rotor forging,homogeneous distribution of chemistry was confirmed and low impuri ty contentswas observed in the whole forging. The superior strength and toughness wereconfirmed with good creep rupture strength. The FATT at the center oflow-pressure section was -3°C with the tensile strength level of 870 MPa.From the results of fracture toughness test, low cycle fatigue test, andisothermal aging test, superior mechanical properties were demonstrated.Thus, the superclean 9CrMoVNiNbN steel with reduced Si and Mn contents andNi addition, is particularly suitable to the high-pressure low-pressure rotormaterial for advanced combined cycle power plants.KEYWORDS: high Cr steel, Ni, steam turbine, rotor, forging, superclean,tensile strength, toughness, creep rupture strength, aging, low cycle fatigue,fracture toughness

In order not only to conserve the resources but also to reduce CO2emission, improving thermal efficiency of power plant have come to be required[1], and then combined cycle (C/C) power plant are highlighted. In the caseof those C/C power plants, high-pressure (HP) low-pressure (LP) rotor shaftis usually used in the steam turbine side. Recently the enlargement of thesize of HP-LP rotor shaft is required with increasing plant size, and severalnew modified CrMoV steels [2]-[6] are beginning to apply to HP-LP rotormaterial for the advanced type C/C power plant.

In these newly developed low alloy steels for the HP-LP rotor material,the chemistry is designed to have high tensile strength and good toughnessat the LP portion, and superior creep rupture strength at the HP portion,simultaneously. Therefore hardenability of these materials are increasedby adding alloying elements, such as Ni, Cr, Mo and W. These modificationof chemistry leads to improvement of toughness. In addition, in order to

lResearch engineer and senior research engineer, respectively, MuroranResearch Laboratory, The Japan Steel Works, 4 Chatsu-machi Muroran Japan.

2General manager, manager and general manager, respectively, MaterialsEngineering Department, The Japan Steel Works, 4 Chatsu-machi Muroran Japan.

330

AZUMA ET AL. ON ROTOR SHAFT FORGING 331

preserve the creep rupture strength which may decrease by alloy additionmentioned above, Si, Mn, and other residual elements were reduced as lowas possible.

Though, these alloy steels have been developed successfully, furtherimprovement have come to be requested since the larger of rotor diametersare to be expected.

In the present study, fundamental research by using laboratory heatswas performed concerning the effects of chemistry on the toughness and creeprupture strength in the 9CrMoV base mater ials, which may have some advantagein tensile strength, toughness and creep rupture strength. From theinvestigation, it is showed that the superclean 9CrMoVNiNbN steel with Niaddition provides the superior strength versus toughness balance.Furthermore, a trial rotor has been successfully manufactured and superiormechanical properties are confirmed. This paper introduces these evaluationtest results of this trial HP-LP rotor forging.AIM OF MECHANICAL PROPERTIES

In the HP-LP rotor, the material is required to have high creep rupturestrength at the HP portion and the good center toughness at the LP portionsimultaneously.

According to the increase of diameter and blade length at the LP portion,the material toughness and strength must be increased. As mentionedpreviously, with recent increase of C/C power plant outlet, the size of HP-LProtor, have been increasing and this requests high strength and toughnessfor turbine material. On the other hand, good creep strength equivalent tothat of CrMoV steel is also required in HP portion.

Fig. I shows a relationship between FATT at center portion of rotorand tensile strength for conventional CrMoV steel [7] and modified CrMoVsteels [2]-[6]. These data shows that the typical FATT is around 50°C withtensile strength of 900 Mpa in low alloy steels. However, in case of futureHP-LP rotor, the required toughness for LP section may reach below 20°C atthe center portion with tensile strength of 880 MPa. Considering thissituation, the aim of mechanical properties in this fundamental researchwas assumed to be tensile strength of more than 880 MPa, FATT of less than20°C, and creep rupture strength of CrMoV steel level.


FUNDAMENTAL INVESTIGATIONOptimization of Chemical Composition

Thirty seven heats of 9CrMoV steels with various alloy contents weretested as shown in table 1. Effects of chemistry on tensile property,toughness, and creep rupture strength were investigated. For example, fig.2 shows the effects of Si, Mn, and Ni contents on toughness and creep rupturestrength. From this figure, addition of Ni contents increases the toughness,but inversely decreases the creep rupture strength. Reduction of Mn contentshas no effect on the toughness. Reduction of Si contents increases toughness.Reduction of Mn, and si contents are effective for improvement of creeprupture strength. Similar investigations were performed for steels withvarious C, Cr, Mo, V, Nb, and N contents, and then the effects of chemistryon mechanical properties as show in table 2 were obtained. Based on thesetest results, the superclean 9CrMoVNiNbN steel with increased Ni contentwas selected to be the best composition for the trial HP-LP rotor materialas showed in table 3.



the weight of 65 ton was made (Fig. 5). Forging was performed by 8000 and10000 ton capacity presses in order to give homogeneous and soundmicrostructure (Fig. 6). After the forging, preliminary heat treatment andquality heat treatment was performed as shown in fig. 4. At quality heattreatment, oil quenching was performed from austenitizing temperature of10750C followed by the first tempering at 580°C for 36 hours and the secondtempering at 650°C for 60 hours. After this quality heat treatment,ultrasonic inspection was performed from the outer surface and no indicationswere detected. Fig. 7 shows the appearance of the trial HP-LP rotor shaftforging after rough machined.


EVALUATION OF TRIAL HP-LP ROTOR FORGINGFig. 8 shows the locations of the evaluation for the distribution of

chemical composition and mechanical properties in surface and center portionof the trial rotor forging.

Sulfur Print, Macrostructure, and MicrostructureFig. 9 shows the sulfur print, and macrostructure of the body cross

section at LP portion of the trial rotor forging. In the sulfur print, nosulfur indications appeared. The macrostructure was fine and homogeneous.Fig. 10 shows the microstructure of the surface and the center location ofLP portion of the trial rotor forging. Homogeneous tempered martensiticstructure with the prior austenite grain size of around 3.5 is attained inboth locations.

HardenabilitvFig. 11 shows the continuous cooling transformation (CCT) diagram of

the trial rotor material. Fig. 12 shows the comparison of hardenability ofthe superclean 9CrMoVNiNbN steel with the conventional 12CrMoVNb steel.Though transformed ferrite is observed at cooling rate less than 20 - 40oC/hin conventional 12CrMoVNb steel [9], transformed ferrite is only observedat cooling rate of 7°C/h in superclean 9CrMoVNiNbN steel. Thus, it isconfirmed that superclean 9CrMoVNiNbN steel has enough hardenability forthe material of HP-LP rotor with large diameter.


Tensile and Impact PropertiesFig. 13 shows tensile and impact properties at various locations in

the trial rotor forging. Uniform mechanical properties were attained in thewhole forging from the surface to the center portion. Especially, at thecenter of LP portion, FATT of -3°C with tensile strength of 870 MPa isachieved.

The obtained balance between tensile strength and toughness of the trialrotor forging is plotted in fig. 14. In this figure, the data of tensilestrength and toughness, which is obtained by simulated quality heat treatmentfor the excess materials at the bottom and top side of the trial rotor forging,are plotted. This balance between tensile strength and toughness falls onand near the lower bound of the result of fundamental study [8]. It shouldbe noted that the similar balance is appeared in HP portion and LP portion.All these data shows that the present 9CrMoVNiNbN steel was a significantadvantage in the balance between strength and toughness compared with thelow alloy steers [2]-[7].



Aqinq test resultsFig. 20 shows the results of iso-thermal aging test up to 10000 hours

by using the materials of the trial rotor forging. The aging test wereperformed between 343°C and 566°C. After iso-thermal aging treatment, thetoughness slightly decrease with increasing temperature. But each FATTwhich aged after 10000 hours at various temperature, is less than roomtemperature.

CONCLUSIONIn order to obtain a material having improved strength and toughness

for advanced HP-LP rotor forging, effects of chemistry on the toughness andcreep strength were investigated on 9CrMoV base materials. From thisinvestigation, it is showed that the superclean 9CrMoVNiNbN steel providesthe superior strength versus toughness balance. Based on the test results,the trial HP-LP rotor shaft forging was manufactured and properties wereevaluated. The results are summarized as follows.(1) Laboratory investigation indicates that the superclean 9CrMoVNiNbN steelwith Ni content of 1.3% is the most promising material which has the superiorstrength versus toughness balance to those of advanced CrMoV steels and withgood creep strength.(2) A trial rotor shaft forging with diameter of LP section of 1750 mm anddiameter of H~ section of 1200 mm was successfully manufactured from 1800mm diameter ESR ingot.(3) The homogeneous distribution of chemical composition, high tensilestrength, and superior toughness were confirmed.(4) Creep rupture strength, fracture toughness, hot tensile strength, lowcycle fatigue properties, and aging properties of the superclean 9CrMoVNiNbNsteel are equivalent and/or superior to those of conventional CrMoV steel.(5) The superclelan 9CrMoVNiNbN steel which has excellent mechanical andimpact properties, is particularly suitable to the HP-LP rotor material foradvanced C/C power plants.

REFERENCES[1] Kougami, K., and Isaka, H., "The Development of Ultra-Super-Critical

Thermal Power Plant," Testu to Haqane, vol.76 (1990), pp. 1043-1052[I] Yamada, M., Tsuda, Y., Watanabe, 0., Miyazaki, M., Tanaka, Y.,


Takenouchi, T., and Ikeda, Y., "HLP Single Cylinder Steam Turbine RotorForgings for Combined Cycle Power Plants," Proceedinqs of CleanMaterials Technoloqy, Chicago, USA, November 1992, pp. 161-168

[1] Fukui, Y., Shiga, M., Hidaka, K., Kaneko, R., and Tan, T., "Developmentof Superclean 0.2Mn-l.8Ni-Cr-Mo-V Steel for HP-LP Steam Turbine,"Proceedinqs of Clean Materials Technoloqy, Chicago, USA, November 1992,pp. 249-257

L1] Kadoya, Y., Kitai, T., Tsuji, 1., Matsuo, A., Tanaka, Y., Azuma, T.,and Ikeda, Y., "Effects of Cr, Mo, W, Mn, and Ni on Toughness of2.25Cr-Mo-V Rotor Steel," Testu to Haqane, vol.79(1993), pp. 980-987

L~] Potthast, E., Poppenhager, J., Wiemann, W., and Mayer, K. H., "Advanced2%CrMoNiWV Steel for Combination Rotors," Proceedinqs of 11thForqemasters Meetinq, Terni, Italy, June 1991, No.IX-8

[.§.] Tanaka, Y., Ikeda, Y., Ohnishi, K., watanabe, 0., Kaplan, A., Schwant,R. C., Jaffee, R. 1., andPoe, G., "Development of a Superclean 2.5NiCrMoVRotor Steel for HP and HP-LP Application," proceedinqs of 11thForqemasters Meetinq, Terni, Italy, June 1991, No.IX-7

LI] Swaminathan, V. P., Steiner, J. E., and Mitchell, A., "Advanced RotorForgings for High-Temperature Steam Turbines," EPRI Report CS-4516,Vol.1,2(1986)

[~] Kimura, K., Azuma, T., Tanaka, Y., and Ishiguro, T., "Dependence oftoughness and Strength of Alloying Element in High Purity 9Cr-1Mo-VSteel," CAMP-The Iron and Steel Institute of Japan, Vol. 4(1991) , p. 2032

[~] Amano, K., Fujida, T., Itoh, M., Nakamura, S., Yoshida. H., Report ofthe 123rd Committee on Heat Resistinq Metals and Alloys, Japan Societyfor the Promotion of Science, Vol.14(1973), No.1, pp. 1-6

John E. Steiner, 1 Edward L. Murphy, 1 Robert D. Williams, 1

HYDROGEN AND FLAKING AFTER 40 YEARS OF VACUUM DEGASSING

REFERENCE: Steiner, lE., Murphy, E.L., and Williams, R.D. "Hydrogen andFlaking After 40 Years of Vacuum Degassing", Steel Forgings: Second Volume.ASTM STP 1259, E.G. Nisbett and A. S. Melilli, Eds., American Society for Testing andMaterials, 1997.

ABSTRACT: Background on the early urgent push for hydrogen control, primarily inheavy rotor forgings, as a result of the rotor failures of the mid 50's, is presented in a briefretrospective. Recent experience with the several degassing processes, their efficiency andcosts in terms of safety treatment, delayed deliveries and claim losses is covered. Industryconcern with hydrogen analysis of vacuum degassed steel solid product is discussed withregard to end customer need, justification and reliability of analysis. with the increasinglywide-spread use of the ladle refining furnace, some producers are reporting hydrogenproblems with low sulfur vacuum degassed steels. Interaction among several factorspossibly related to this phenomenon is discussed.

KEYWORDS: hydrogen, steel forgings, flaking, vacuum degassing, hydrogen analysis,ladle refining, low sulfur-flaking.

The mid 1950's was the era of the dramatic large steam turbine and generator rotorfailures. Two of.these, the Ridgeland and Arizona rotors [1,2] were flaked. Hydrogenembrittlement, although discussed in the investigation of a third failure, the Cromby rotor[2], was not considered a verifiable cause of the failure. These failures acceleratedtechnical work in hydrogen embrittlement and hydrogen control that had been in progressor under consideration for several years at the heavy forge shops and large steam turbinebuilders. The introduction of mechanically pumped vacuum stream degassing atBochumer-Verein [3] in Germany in 1954, application of steam ejector pumping to thistechnology in 1956 and the technological and commercial impetus generated by these

I Consultant, Heavy Forgings, Engineering, Materials & Processes, Inc., 121 Edgewood AvenuePittsburgh, Pennsylvania 15218

344

STEINER ET AL. ON HYDROGEN AND FLAKING 345

failures thrust the industry into the age of vacuum degassing in essentially one year's time.Fortunately, the tools were at hand to do the job, and the system worked even beyondexpectations.

In addition to the 1 ppm hydrogen achieved in vacuum stream degassing,subsurface transverse tensile ductility was markedly and consistently improved because ofthe essential exclusion of stream oxidation in the vacuum tank. Reproducibility of tensileductility and elimination of any problem with hydrogen embrittlement of tension tests wereachieved virtually across the board.

Within months of mid 1956, a1lrotor producers had vacuum stream degassing forlarge steam turbine and generator rotor forgings. If you didn't vacuum degas it, youdidn,t se1lit. By 1960 the use ofladle-to-Iadle degassing coupled with argon shielding ofsubsequent direct pouring yielded 1.5 ppm hydrogen when all went we1l,but often muchhigher. This method became common practice for smaller ingots.

Recognizing the limitations in effectiveness of this latter process, US Steel evenevaluated a "lazy susan" turntable within a large vacuum tank to achieve pure vacuumstream degassing of small ingots [4]. Although a technological success, it was notadopted for production. The Gero process [5] employs the mold ofa small ingot as thevacuum chamber in vacuum stream degassing and is in successful production use today.

Concurrently, and we1linto the 1960,s, the lift processes, Dortmund-Hoerder (D-H) and RheinstaW-Heraeus (R-H) became popular for small ingot shops. As pointed outby Murphy and Steiner at the Williamsburg symposium [6], these processes are susceptibleto the vagaries of day-to-day operating control. At best, they yield 1.5 ppm hydrogenwith occasional excursions into the 2.5 to 3.5 range. As with ladle-to-Iadle degassing,they rely on shielding of the stream in subsequent pouring.

By the 1970's, with the introduction of sliding gate pouring, ladle refining furnace(LRF) processing was made feasible. The application of efficient porous plug technologyand magnetic stirring made use of a vacuum hood over LRF units attractive to severalproducers. The Finkl VAD furnace came on stream in the mid-60's. Insofar as the LRF-vacuum method is not as temperature dependent as the lift processes, it yields hydrogencontents closer to 1.5 ppm and often somewhat lower.

By the 1980's and to date, a1lseems to have been done. Vacuum treatment at the1 mm Hg (1 Tor) level ofliquid steel seems to have become a1lpervasive. "Vacuumtreated" is in virtua1ly a1lspecifications. Vacuum degassing seems to have become moreof a religious credo than a simple metallurgical process.

To assess how much something so popular is used when it might not be needed, areview of the effects of vacuum treatment on the low a1loyforging steels is in order. Keepin mind that VIM and VAR vacuums are about 3 microns (0.003 mm Hg), modern largetonnage steam ejector degassing systems discussed above are about 800 microns (0.8 mm


Hg) and the original Bochumer-verein mechanical pump system was about 3,000 microns(3.0 mm Hg.). These are quite different levels of vacuum. Herein, we will be discussingthe steam ejector 0.8 mm Hg technology, although, in common parlance it sometimesseems that some characteristics of the 0.003 mm Hg are loosely attributed to anythingwith the "Vacuum" label. To keep things in perspective, throughout this paper 0.8 mm Hgwill be used with steam ejector pumping, 0.003 mm Hg will be used with the complexmechanical diffusion pumping of VIM and VAR and 3.0 mm Hg will be used with the highvolume mechanical pumping of the Bochumer-Verein system.

Steam ejector (0.8 mm Hg) stream degassing was adopted to control hydrogenembrittlement oflarge rotor forgings as manifest in embrittlement of unaged tension tests.Incidentally, the concern was almost exclusively with the Ni-Mo- V steels. A major sidebenefit was the complete elimination of flaking at 1.0 ppm hydrogen, and greatly reducedflake sensitivity with the processes yielding 1.5 ppm hydrogen. When the process wasintroduced in USSteel, few dreamed that it would have this dramatic effect. In fact, in theshops, the only concern with flaking was the extent to which costly safety treatmentsmight be shortened. A third less heralded benefit is a marked reduction in quench cracksensitivity. For example, the ladle-to-Iadle degassing of the standard 52100 modified rollsteel completely eliminated the hazard and occurrence of spontaneous bursting or spallingof rolls in process. These benefits of vacuum processing are directly related to reducedhydrogen content and are, to our knowledge, irrefutable.

Steam ejector (0.8 mm Hg) degassing also reduces nitrogen. Few have everexpressed concern with nitrogen in the low alloy forging steels. Experience with twosteels, 4340 and 1.35 Mn-Mo, both basic electric furnace melted and aluminum killed isrecalled. These were much more flake sensitive and, for the 4340, more forge-crack-at-the-press sensitive than when produced in the basic open hearth, where the nitrogencontent is one half that of the BEF.

Steam ejector (0.8 mm Hg) degassing also reduces oxygen. This is a differentstory from nitrogen because of oxygen's pervasive metallurgical importance. Thepercentage reduction is about the same as nitrogen when the steel has been fully killedprior to degassing. But with no prior deoxidation, magical things happen. Although thefinal oxygen is about the same, almost everything about the final steel product isimproved. This is deoxidation with carbon rather than with a metallic precipitationdeoxidant. It is commonly called vacuum-carbon deoxidation (VCD).

All 0.8 mm Hg degassed steel undergoes some carbon deoxidation at this relativelylow pressure. The extent depends upon prior deoxidation. To the extent that aluminum,silicon and manganese serve no real purpose in the Ni-Cr-Mo- V family of steels, theelimination of these metallic deoxidizers in favor of carbon deoxidation is highly desirable.Great benefits in terms of base toughness and virtual elimination of temper embrittlementhave been achieved. In vacuum stream pouring, no prior deoxidation (open steel) resultsin wild foaming, which is especially undesirable in the hot top at the end of the pour.Similar foaming is encountered in the ladle processes. Because aluminum seems to be the


most innocuous while the most powerful of the deoxidizers named above, it would seemprudent to control the degree of foaming or wildness of carbon deoxidized steel by someprior deoxidation with aluminum. This deoxidation would be applied at an intermediatestage well before teeming and hence would be a type of aluminum intermediatedeoxidation system that should easily permit compliance with a restrictive maximum limitof, say 0.010 percent, on aluminum content.

The oxygen contents achievable from 0.003 mm Hg degassing are far below thosefrom 0.8 mm Hg degassing. In the VIM and VAR steels one sees the true effects oflowoxygen. In the 0.8 mm Hg range we enjoy primarily the effect that the pressure is lowenough to make carbon the strongest deoxidizer in the bath. This is what permits us toeliminate metallic deoxidizers and achieve 0.02 Mn and 0.02 Si to produce a tough, nonembrittling "super clean" steel.

Without question, the large rotor and nuclear forgings have benefitedimmeasurably from 0.8 mm Hg vacuum stream degassing. Other steels have too.

Listing the rotor steels involved in order of benefit from vacuum degassing wouldinclude:

1. A469 and A470, Ni-Cr-Mo-V - dramatic response to carbon deoxidationwith concomitant marked improvements in toughness and resistance totemper embrittlement, elimination of flaking, elimination of hydrogenembrittlement, improved resistance to quench cracking.

2. A469 and A470, Ni-Mo-V - elimination of flaking, elimination of hydrogenembrittlement, marked improvement of surface and subsurface transversetensile ductility (shielded pouring atmosphere effect).

3. A470 lCr-Mo-V improved resistance to quench cracking.

Although the Ni-Mo- V rotor steel arguably enjoyed the greatest benefits from0.8 mm stream degassing, it has been effectively phased out in favor of Ni-Cr-Mo- V, firstfor LP turpine rotors and finally for generator rotors over a period of 10 to 20 years afterdegassing was introduced. Presently, it is seldom used. However, the processingsimplicity of Ni-Mo- V steel may be of interest here. It is an extremely forgiving steel.

In the 1954-1957 transition from air cast to 0.8mm Hg vacuum cast rotor ingots,only the Ni-Mo-V and ICr-Mo-V steels were produced for rotors in the US. An extensiveprogram [7] was conducted on the Ni-Mo- V steel to measure hydrogen content and flakesensitivity on many heats produced in the period before and after 0.8 mm Hg vacuumstream degassing was introduced into production. The bulk of these heats were Ni-Mo-V,of interest because of the known flake sensitivity of the grade. only a few were lCr-Mo-V,which were quickly dropped from the program because of an absence of flaking of this


grade and, more dramatically, because of the explosive, "dynamite" steel [8] crackingbehavior encountered upon fast air cooling or water quenching from the press.

As a result of the considerable data on the 0.8 Mm Hg vacuum treated Ni-Mo-Vsteel and its absolute freedom from flaking or cracking in 400 mm (16 inch) square blocks,water quenched directly from the forge, and concurrent observation oflarge heavy sectioningot discards, a dramatic experiment was conducted. A Ni-Mo- V steel turbine rotorshaft, about 1260mm (50 inches) diameter was air cooled to ambient temperature directlyfrom the press. It was machined to final pre-quality treatment dimensions, normalized andtempered and tested. All test results were excellent. Total treatment time was about 7days. When retreated to the standard approximately 50 day schedule, the properties wereessentially identical, with toughness poorer, but only a bit.

For obvious reasons, no such heroic trials were conducted on turbine rotors of theICr-Mo- V dynamite steel. It is interesting to note that while cracking was furious, flakeswere never observed in test blocks from the few heats of air melted ICr-Mo-V steel ofthisearly program, or in the later lazy susan program [4] with 300 mm [12 inch] squareblocks. The authors have never found flakes in ICr-Mo-V rotor steel products.

Also, this ICr-Mo-V steel exhibited no improved toughness nor embrittlementresponse to carbon deoxidation versus Si-Mn deoxidation. This may be due to the needfor these elements in the composition for strength and elevated temperature properties.Probably, because it was so brittle to begin with, it had no toughness to lose by temperembrittlement. Also, its service life was to be spent well above the embrittlement range, infact, in the de-embrittlement range.

Thus, it might be that, if all rotors and generators had been exclusively of 1Cr-Mo- Vsteel, there may never have been any vacuum degassing.

Thus, the real impetus for hydrogen removal and its side benefits of carbondeoxidation came primarily from Ni-Mo-V rotor steel in the USA and from the Ni-Cr-Mo-V rotor steel in Europe. Ofthese two, the surviving Ni-Cr-Mo-V steel enjoyed by far themost striking benefits.

Largely because of tradition and necessarily ultra conservative specifications, thepotential benefits of drastically shortened heat treatments have not been fully realized.The behavior ofNi-Mo-V steel at I ppm hydrogen was especially well documented. Itpermits the most extreme shortening of safety treatments. The presently dominant Ni-Cr-Mo-V steel, however, came in at a time when no air-melted production heats wereavailable for comparison. The lazy susan work [4] stands as about the only quantitativemeasure of the flake sensitivity of this steel. This study shows that it is much less sensitivethan the Ni-Mo-V steel. Safety heat treatments for both these steels have been modifiedthrough the years with a wide variation in treatment times and philosophy evident amongheavy forge shops.


But what of other steels, especial1ythose used in forgings from small ingots where1.5 ppm or higher is the rule? How much could be saved in degassing treatment costs andin slow cooling safety treatment costs? The tonnages involved here are many times thetonnages of rotor steels. For the carbon steels [4], 1080 begins to flake at about 3.5 ppmhydrogen in 300 x 300 mm (12 x 12 inch) air cooled sections. The 1045 steel air cooled inthe same section size, does not flake even at 5.1 ppm hydrogen. The authors have neverseen flakes in 0.30 plain carbon steel, air melted and air cast, in any section size when aircooled directly from the press. An adequate safety treatment for plain carbon steelforgings under 0.30C in any size is to air cool from the forge press. In this regard, theManganese content becomes important. For example, at the high end of A266 C14, the1.35Mn content renders these steels higWyflake sensitive, even though the steel isspecified as a carbon steel. Plain carbon steel here refers to steels with about 0.65 Mn,certainly below 0.90Mn.

Thus, it seems economical1y questionable to both degas to 1.5 ppm and then inaddition, slow cool or safety treat any carbon steel product.

For the many low al10ysteels, similar if not so graphic analyses can be made. If1.0 ppm can be achieved, probably no subsequent safety treatment is necessary for thegreatest percentage of low alloy steels. Only when quench cracking upon air cooling isencountered should a simple tempering treatment be necessary.

Presently, many shops are degassing to 1.5 ppm and safety treating the product ofsmal1ingots. A typical producer, with about 40 percent of its product in flake sensitivegrades may experience annual costs on the order of $100,000 in claims and $200,000 insafety treatment in addition to the cost of degassing, probably another $300,000. In fact,many shops are compel1ed by custom, commercial pressures or guilt into vacuumdegassing everything, even low carbon steels.

For these steels at 1.5 ppm, which assumes, because of process variability a rangeof 1.5 to 3.0 ppm, a more conservative approach is suggested. Again, probably nothingflakes below 2.0 ppm [6]. Even the extremely "hot" Ni-Mo-V rotor steel can squeeze outonly a few flakes at this level. Thus, if the range can reliably be compressed into the 1.5 to2.0 ppOlhydrogen that most claim, there appears to be little justification for subsequentsafety treatment.

Hence, the ultimate solid state hydrogen analysis method is proposed. Air cool itcold after forging. If it flakes, the hydrogen is too high, indicating that the degassingprocessing had to have been inadequate. The steel has not been degassed to less than2,0 ppm. Because most forgings are routinely ultrasonically tested, control costs shouldbe low. Because failure is detected early in the process, delivery time is only moderatelyaffected.

If the melt shop is not confident of2.0 ppm maximum, perhaps they shouldabandon degassing and apply some of the degassing costs into a bit more safety heat


treatment after forge. If forced to degas for commercial reasons, as virtually all shops aretoday, they must then suffer the costs of extended safety treatments to the extent that theirdegassing is ineffective. Remember, at room temperature after forging, whateverhydrogen that remains in the piece is innocuous. Furthermore, at each stage ofprocessing, a bit more is lost as the cross section becomes smaller. In 300 x 300 mm (12 x12 inches) square blocks [4] on average 1.0 ppm hydrogen was lost from air-cast steels,upon slag cooling.

This is characteristic, especially with subsequent forge or heat treating processing.This might be at least part of the origin of the old rule, "Ifit doesn,t flake from the firstforge, it will not flake later" .

Hydrogen embrittlement, the bete noire of the early 1950's in heavy rotortechnology, has proved to be oflittle consequence. Hydrogen residuals in steel forgingsare not the source of the surface induced hydrogen embrittlement encountered in theexposure of steels to chemical processes, plating, and oil country problems.

Another phenomenon has been related to the 1.5 ppm steels. It has been widelyreported that contemporary ladle refined low sulfur steels have insufficient voids toaccommodate molecular hydrogen and therefore are more prone to flaking.

Most references to this sulfur effect are generalizations, but we are aware offorgings producers who have specifically reported flaking problems related to low sulfurcontent. In one case, four flaked heats had lower sulphur than four unflaked heats. In asecond case, tiny cracks observed at mid-radius were eliminated by avoiding "double zero"sulfur heats. In these instances, consideration of other pertinent factors such as theinfluence of winter/summer melting, reliability of hydrogen sampling and analysis, crewattention to melting parameters, comparative degree of safety treatment and identificationand quantification of flaking would be of interest. Thus, we believe that the sulfur theorymay be an over simplification of a complex mix of circumstances. Steels of the samegrade and hydrogen content do not necessarily flake at low sulfur levels, nor are theynecessarily flake free at relatively high sulfur contents. This suggests that there are othervariables at work.

It seems, however, that resulfurizing is a "too easy" solution. What else waschanged in processing when flaking was noted in the low sulfur steels? Surely, closerattention was undoubtedly paid to shop compliance with often neglected safety treatmentpractices.

Unfortunately, the following case does not involve forgings with both high and lowsulfur. Both were reasonably low, but it is mentioned here as it relates to compliance.Reasonably large vacuum treated forgings of 4140 steel were rejected for flakes. Theyreportedly had l.3ppm hydrogen. At this level, 4140 cannot flake. Could we question thehydrogen analysis? Furthermore, these forgings were slag cooled. Slag cooling, properlydone precludes flaking regardless of hydrogen content in any grade. This gets down to a


definition of slag cooling and shop compliance in adequate coverage of the pieceand total time buried. All this in the face of commercial pressures for moving product andmeeting delivery.

Remember, major producers of rotor forgings world wide produce extremely lowsulfur steels, but at the 1 ppm hydrogen content of stream degassed steel, they have noflaking, low or high sulfur and voids or no voids.

Certainly the concept of accommodating voids is appealing, especially for shopsfinding it necessary to accommodate up to 2.5 to 3.5 ppm hydrogen, but the unknowns aretoo numerous to support the generalizations that are rampant within the industry regardingthe effect of sulfur levels. A controlled investigation on this subject would be a valuablecontribution to our knowledge of the flaking mechanism.

In summary, what is vacuum degassed steel? If only it were as simple as to classifyit according to final hydrogen level as we have done in this paper. Certainly, specificationsfor solid steel analysis would open a can of worms [6] that would plague the industry forthe next quarter century. Thus, it appears that a process dependent specification, withproducer monitoring ofliquid hydrogen content at the latest possible stage, largely aspracticed now, is the only practicable approach, if not the only approach.

Perhaps tightening control of the dilution processes might permit large processingcost reductions. These processes can and do produce 1.5 ppm steel, plus or even minus abit. Reproducibility here is the key. This is largely up to the producer.

Over these 40 years of steam ejector vacuum degassing, hydrogen in forging steelshas been essentially no problem. There has been a very low incidence of flaking, justenough to keep us on our toes. There has been enough complacency to sort ofinstitutionalize the process and perhaps over apply it, or at least not to take full advantageof its benefits. The authors hope that these comments will stimulate enough thought andaction that a great technology might continue to improve over its next 40 years.

REFERENCES

[1] Emmert,H.D., "Investigation of Large Steam-turbine Spindle Failure" ASME paperNo 55-A-I72 presented at the ASME Diamond Jubilee Annual meeting, Chicago,IL, November 13-18, 1955

[2] Schabtach, c., Rankin, A.w., Fogelman, E.L. and Winnie, D.H., "Report of theInvestigation of Two Generator Rotor Fractures" ASME paper No 55-A-208presented at the ASME Diamond Jubilee Annual Meeting, Chicago, IL, November13-18, 1955.

[3] Tix, A., Stahl u, Eisen, 76, 61-68 (1956)


[4] Hodge, lM., Orehoski, M.A, and Steiner, lE., "Effect of Hydrogen Content onsusceptibility to Flaking", volume 230, August 1964, Transactions of theMetallurgical society of AIME.

[5] Warner, lH. and Hay, G.l, "Low Investment Degassing of Alloy Steel", AlSEYearly Proceedings, 1970, 144-150.

[6] Murphy, E.L. and Steiner, lE., "Hydrogen - Its Occurrence, Determination andControl in Steel Forgings" Steel Forgings, ASTM STP 903, E.G. Nisbett and AS.Melilli, Eds., American society for Testing and Materials, Philadelphia, 1986, pp573-582.

[7] Steiner, lE. "A Test for Flake Sensitivity", presented at AIME St. Louis Meeting1958.

[8] Danner, G.E., private communication on nomenclature coined at ErieForge & Steel, Co. ca.1950.

Michael Gold!

INTERNATIONAL BUSINESS, CODES, AND MA TEIDAL SPECIFICATIONS

REFERENCE: Gold, M., "International Business, Codes, and MaterialSpecifications", Steel Forgings: Second Volume. ASTM STP 1259, E. G. Nisbett andA. S. Melilli, Eds., American Society for Testing and Materials, 1997.

ABSTRACT: The nature of the business environment oflarge forging manufacturers andusers has changed significantly in the twelve years since the first ASTM forgingsymposium. The largely domestic nuclear market has become a global fossil andpetrochemical market. Further, the primary market is on the other side of the world fromNorth America. To participate in this market, original equipment manufacturers arerequired to form joint ventures, with manufacturers in customer's countries, just to obtainmarket access. This has a significant effect on how their businesses operate, on theconstruction codes to which they design and build equipment, and on material sources andmaterials specifications they are required to use to be competitive. There will also besignificant changes required in ASME construction codes and ASTM materialsspecifications, if North American manufacturers are to continue to use these and stillcompete effectively.

KEYWORDS: ASME Code, ASTM material specifications, global market place,installed capacity, joint ventures

INTRODUCTION

I am pleased and honored to have been asked to present the keynote address at theSecond ASTM Forging Symposium. My purpose is to indicate how changes in the natureof the business have affected how original equipment manufacturers use AmericanSociety of Mechanical Engineers (ASME) construction codes and ASTM materialsspecifications in this new environment. I will do this from the perspective of a materials

! Manager, Materials Technology and Standards, Utility, Environmental & IndustrialDivision, Babcock &Wilcox, 20 S. Van Buren Ave, P.O. Box 351, Barberton,OH 44203-0351.

353


engineer in a large boiler manufacturing company, and as an active participant in thedevelopment of ASME codes and ASTM materials specifications.

All of us at this symposium are interested in forgings, either as a manufacturer, or as auser, for their application in boilers, turbines, pressure vessels, and other large equipment.For most of us, there have been dramatic, even catastrophic, changes since the first ASTMForging Symposium in 1984. Then, our businesses were mostly domestic, and werelargely nuclear oriented. Now, our businesses are almost totally international, and, exceptfor the replacement nuclear steam generators segment, largely fossil and petrochemical. Inbrief, we have a global market, and, particularly from my perspective as a representative ofa boiler manufacturer, it is on the other side of the world (whether we slice the globe East-West or North-South). This is reflected in the increased participation in this symposiumfrom representatives outside of North America.

IT'S A GLOBAL MARKET PLACE

The domestic OEM boiler business has vanished, at least for the present, essentiallybecause it is saturated. Even the repair-replace-upgrade market has stagnated, due to theuncertainties of utility deregulation and consolidation. The Babcock &Wilcox PowerGeneration Group still has major engineering and manufacturing facilities in NorthAmerica, at: Melville, Saskatchewan; Cambridge, Ontario; Barberton, Ohio; West Point,Mississippi; and Paris, Texas. Babcock & Wilcox's corporate parent, McDermott, hastheir corporate headquarters here in New Orleans, Louisiana, where the symposium isbeing held. On the other hand, Babcock &Wilcox has closed or sold significantmanufacturing and engineering facilities in: Little Rock, Arkansas; Wilmington, NorthCarolina; Canton, Ohio; and Lynchburg, Virginia. Instead, major joint venture facilitieshave been opened in Mexico, Egypt, Turkey, India, China, and Indonesia.

THE MARKET IS ON THE OTHER SIDE OF THE WORLD

This shift has occurred for very fundamental reasons. As a simple comparison of installedelectric gener~ting capacity versus population reveals, the new boiler market is largely inthe Far East, Latin America, Middle East, and in Eastern Europe (Table 1).

GOLD ON INTERNATIONAL BUSINESS 355

TABLE 1 -- Comoarative installed capacities

The conclusion to be drawn from these data is that there will be essentially no new coalfired boiler construction in North America, Western Europe, or Japan. This is reflected inFigure 1, which shows that 80% of the coal fired boilers to be installed in the period of1994-2004 will occur in India, China, the Pacific Rim, and other locations (largely EasternEurope and Latin America). The effect on Babcock &Wilcox is that in 1995, bookingswhere 81% international and only 19% in North America (Figure 2). This situation hasbeen developing for some time, as seen in Figure 3, which identifies the location ofBabcock &Wilcox OEM projects (boilers, replacement nuclear steam generators, andscrubbers) over the last ten years.



JOINT VENTURES ARE REQUIRED

Participation in joint venture companies in the customer's country is necessary, just to beconsidered a qualified vendor. Joint ventures with foreign partners can have benefits ofsignificant cost reductions, equity contributions, and influence with local customers.Selecting the right partner provides local market knowledge and an understanding of thephilosophy oflocal business. Contracts for the boilers at Suralaya and Tanjung Jati, inIndonesia, would not have been possible without the joint venture with PT Babcock &Wilcox, on Batam Island (Figure 4). Babcock & Wilcox's participation in the 2S boilersin China would not have been possible without the joint venture with the Babcock &Wilcox Beijing Company. This plant, ten miles west of Tain'an men Square has hundredof acres of land, 1 200000 sq. ft. (111 000 m2) of production buildings, 3 900 Chineseemployees, and 8 Babcock &Wilcox ex-patriot employees. Participation in the Egyptianmarket (Figure S) would not be possible without the joint venture, Babcock &WilcoxEgypt, in Cairo. In some of these markets there is still a significant growth potential asshown by Figure 6.



The increased manufacturing and engineering capacity, represented in these joint venturecompanies, will eventually be brought to bear in other ventures, and in our domesticmarket when it recovers. This will have a significant impact on Babcock & Wilcox'sNorth America manufacturing and engineering resources. But the alternative is not to bein this business at all.

ASME CODE ACCEPTED BUT NOT ALWAYS COMPETITIVE

In these new international markets, the ASME Boiler Code is accepted, but notexclusively, as it is in the United States and Canada. Other codes may be acceptedequally, or may even be preferred. Some countries permit design to the ASME Code, aslong as local codes are also met (e.g. India and the Czech Republic). Other internationallyaccepted codes are occasionally less conservative than the ASME Code, which can putNorth America companies, used to using the ASME code only, at some disadvantage.This is illustrated in Figure 7, which shows the criteria used by the ASME Boiler andPressure Vessel Committee, to establish the allowable stresses for the new grade 91(9Cr-1Mo- V) alloy product forms, which have recently been introduced in boilerapplications. The allowable design stress at any given temperature is the lowest value ofany of these curves. Thus, the maximum allowable design stress for fossil boilerconstruction, for A336 F91, at lOOO°F(540°C), under the ASME Code is 14.3 ksi(98.6 MPa). This value is controlled by the tensile strength criterion in the ASME Code.Some major European codes do not use a tensile strength criterion and their allowablestress value, at lOOO°F(540°C), is controlled by the rupture criterion. This leads to avalue closer to 18 ksi (124 MPa), which is about 25% higher than the ASME value. Thisdifference leads to a significant difference in component thickness, weight, and cost. The


conservatism of the European design is much less than that using the ASME Code. This iscounter balanced, to some extent, by tighter controls on materials manufacturers,amounting to virtual certification ofmateriaIs manufacturers. Even with these additionalcontrols, the finished European component is considerably less expensive, and is perceivedto be equally safe.

WORLD SOURCING OF MATERIALS

Even for components manufactured in Babcock & Wilcox's North American facilities, themajor sources of pressure part raw materials supply (primarily tubing, pipe, plate, castings,and forgings) has shifted from almost totally domestic to largely international. In order oftonnage, major countries of supply are Japan, the United States, France, and Germany.Babcock & Wilcox's purchasing department anticipates that China will also become amajor source, within a few years. Likewise, for new boiler facilities in the world market,fabricated components wiII be manufactured in many of the joint ventures locations, forassembly at the plant site. Figure 8 shows that, for the current facilities being built at EIKureimat, in Egypt, major components are being supplied by B&W facilities in Canada,China, and India, as well as those in Egypt.


LOCAL MATERIALS AN OPTION

Some markets permit design and manufacture to the ASME Code, with a significantamount oflocal material content. So far, this has meant that local materials have beenused for non-pressure components. But this situation will not last long. Most non-domestic materials producers will readily make their products to ASME and ASTMmaterials specifications. However, in some locations, it would be possible to use materialsto ns, DIN, or CEN specifications, and in a few cases these maybe preferred.

CHANGES TO CODES AND SPECIFICATIONS

To facilitate use of materials made to non-domestic specifications, the ASME Boiler &Pressur~ Vessel Committee has removed the structural impediments to the use of materialsmade to international standards other than ASTM. Steps have been initiated toincorporate two such materials, a structural steel made to Canadian Standards Associationspecification CSA G40.21-92, for a material similar to A36 but with better toughness atlow temperatures; and a European Committee for Standardization (CEN) Specification,EN 10028-2:1992, for pressure vessel steel similar to ASI6. Further, ASME has takensteps to review factors of conservatism that appear to make the products designed to itscodes less competitive in the world market. For example, an active study is underway to


consider reducing the criterion for tensile strength, from 1.1 -;-4, to 1.1 -;-3.5. In my viewthis is too small a step: I think ASME ought to consider eliminating the tensile strengthcriterion, as some European codes have done. This will require other steps, to deal withstress concentrations, strain concentrations, and the increased use of high yield strength-to-tensile strength ratio materials that will result. Elimination of the tensile strengthcriterion is still likely to lead to components that are much more competitive with thosedesigned to European codes.

Similar avenues ought to be pursued by ASTM specification writing committees. Someitems that ought to be considered are: increased uniformity of requirements across productforms for the same alloy; tolerances that recognize those of the same product produced bymanufacturers in other major industrialized countries (particularly those in Japan andWestern Europe); and more freedom to use test methods permitted by other internationalcodes.

My final message is that it is a small and rapidly changing world, and if we, in the majormanufacturing industries, and on code and specification writing committees, do notrespond to the changes quickly, we will not survive the competition, much less have anopportunity to win it.

ACKNOWLEDGMENTS

Much of the information in this paper was assembled from Babcock & Wilcoxinformation, and was developed by P.E. Ralston and L.M. Shepard. I am grateful for theirpermission to use it. The opinions and suggestions concerning issues associated with theASME Boiler and Pressure Vessel Code or ASTM material specifications are entirely myown and do not necessary represent the views of ASME or ASTM.

Author Index

A Kim, K-C., 18Kirschner, W., 104

Antos, D. J., 129 Kusuhashi, M., 56Azuma, T., 267, 330 Kwon, H.-K, 18

B LBocquet, P., 3, 65 Leap, M. J., 160

Levacher, R, 259C Li, Y, 249

Lin, M., 33Cheviet, A, 3, 65 Lueg, J., 104Coudreuse, L., 3

MD

Malas, J. C., 116Dumont, R, 3 Manning, M. P., 259

Morin, F., 65F Mozden, C. A, 160

Murley, P. F., 259Fielding, J. E., 93 Murphy, E. L., 241, 344Fischer, C. E., 116Focht, R B., 33, 93 N

G Nelson, T. D., 33Ni, Z., 249

Gold, M., 353 Nisbett, E. G., 129Gunasekera, J. S., 116

PH

Price, S., 79, 196Hansen, S. S., 33Hebel, T. E., 148Honeyman, G. A, 79 RHuang, K, 224

Reese, W., II, 224I Reppert, K F., 93

Ronemus, D. C., 224Ikeda, Y,267,330Ishii, R, 267, 317Ishiguro, T., 330 SHoh, A, 305

Sato, I., 56K Schonfeld, K H., 259

Stein, G., 104Kim, J.-M., 18 Steiner, J. E., 241, 344Kim, J.-T., 18 Suzuki, K, 56

363


T W

Walsh, M. A., 196Takenouchi, To, 305 Williams, R. D., 344Tanaka, Y., 267, 330 Wingert, Jo c., 160Tihansky, E. L., 93Tsuda, Yo,267, 317 YTsukada, H., 56

Yamada, H., 305Yamada, M., 267, 317

V Yang, K., 249Yoshida, H., 330

Veerabadran, To, 224 ZViertl, J. R. Mo, 213Viswanathan, R., 280 Zimmers, E. W., Jr., 224

Subject Index

A D

Deformation, 249Aerospace industry, 104 Density difference, 305Aging, 196, 330 Deoxidation, 18, 56resistance, 65 Detection and characterization,

Aluminum, 18,56,79 213American Society of Mechanical Die-forging hammer, 249

Engineers code, 3, 353 Disk forgings, 280Anvil block, 249 Disk manufacturing, processASTM standards, 353 model development, 116A 707, 196

Austenite ~rain size, 160 EAustenitizmg, 33Aviation industry, 104 Electric Power Research

Institute, 280B Electric utility applications, 280

Embrittlement, 3, 280Bearing journal wear, 129Blow force, maximum, 249Bottom pouring, 93 FBursts, forging, 241

Fatiguecrankshaft, 129

C life, 148low cycle, 330

Carbon, 79, 104, 196 FEM modeling, 224carbon monoxide reaction, 56 Ferrite, 65

Chromium, 33, 330 Ferritic steel, 267, 317Chromium ferritic steel, 267, Finite element method, 224

317 Flake-like defect, 241Chromium-manganese-nickel Flaking, 241, 344

alloy, 79 Flasks, transport, 79Chromium-molybdenum-nickel- Fracture, 241

niobium steel, 259 brittle, 160Chromium-molybdenum vanadium ductile, 160

steel, 3, 241, 305 toughness, 330Copper precipitates, 196 toughness, plane strain, 18Corrosion resistance, 104, 259Corrosion, stress, 280Cracking, premature, 148 GCrack propagation, unstable,

160 Gear component forging, 241Crankshafts, grain flow, 129 Generators, 129Creep, 3, 280 rotors, 213rupture strength, 267, 317 Geothermal power station, 259

CRONIDUR,104 Grain flow crankshaft, 129

365


Grain refining, 56 Molybdenum, 3precipitates, 160 carbides, 33

Grain size, 79 effect on creep ruptureaustenite, 18 strength, 267, 317

effect on segregation, 305H in martensitic stainless steels,

104,259Hammer, die-forging, 249 in pressure vessel steels, 3, 18Hardness, aviation stainless in steel forging bursts, 241

steels, 104 in superclean steels, 330Head-forming, 224Hydrogen, 344 NHydrogen damage, 3Hydrogen flakes, 241 Nickel, 305, 330

in pressure vessel steel, 18,33I in steel forging bursts, 241, 259

in transport casks, 79Impact toughness, 160 Nickel chromium molybdenumInclusions, effect on toughness, vanadium, 241

305 Nil-ductile transition temperature,Indication sets, 213 18,33Induction hardening, 129 Niobium, 104,259,305,317,330Ingot cogging,224 Nitriding, 129Ingot production, 93 Nitrogen, 104,330International business, 353 Nitrogen-alloyed stainless steels,

104J Nozzle integration, 56

Nuclear power plant forgingsJoint ventures, 353 pipings, 65

pressure vessels, 18,33, 93L steam generators, 56

transport casks, 79Ladle refining, 93, 344Laves phase, 317 0Lead lmer, 79Locomotive engine crankshafts, Open die forging process, 224

129 Optimization, forged diskmanufacturing processes, 116

M

Machinin~, 116 PdistortIon following, 148

Manganese, 18,79, 305, 330 Piping, primary, 65Martensite, 160 Power generation diesel engineMechanical stress relief, 148 crankshafts, 129Models and modeling Precipitates, copper, 196computer assisted design, 224 Pressure vesselsdevelopment, simplified nuclear, 18,33, 93

forging, 116 oil industry, 3FEM, 224 Pressurized slag remelting, 104methodologies, 224 Process model development, 116prediction, 305 Punch forming, 224

INDEX 367

R T

Tensile strength, 259Rhenium, 317 Tension leg platform, 196Rolls, 93 Thermal efficiencies, 267Rotors, 93, 213, 259, 267, 344 Thermal flakes,·241superclean steels, 280, 305 Transport flasks, spent nuclear

fuel, 79Tungsten, 317

S Turbine disks or blades, 104,317Turbine rotors, 213, 259, 267,

Sesregation prediction, 305 280Silicon, 33, 56 Turbine, steam, 317, 330Silicon killing, 18Steam drum head integration, 56 USteel and steel alloysA 508, Class 3, 18 Ultrasonic testing, 65, 241A 508, Grade 1,33 rotor, 213AlSI 304L, 65 Upper-nose tempercarbon-manganese-nickel embrittlement, 33

alloy, 79chromium ferritic, 267, 317 Vchromium-molybdenum

(vanadium), 3, 241, 305 Vacuum carbon deoxidation,cop'per-bearing, 196 18,56ferntic,267 Vacuum degassing, 344forged, 79, 344 Vacuum stream degassing, 93high strength, 160 Vanadium, 3, 104,241,305,330low alloy, 196,213 Vibration stress relief, 148manufacturing, 93SAS08,Class 3, 56 Wstainless, austenitic, 65stainless, martensitic, 104, Welding, 148

160,241,259 Welds, 65superclean, 280, 305, 330 distortion and cracking, 148

Stress corrosion, 280 Worldwide business, 353Stress relief, 148Structure property relationships, Y

196Sulf~r flaking, 344 Yield strength, 259

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