Aalto- Hydrogen effects on DD austenitic stainless steels

9HSTFMG*agcafb+

ISBN 978-952-60-6205-1 (printed) ISBN 978-952-60-6204-4 (pdf) ISSN-L 1799-4934 ISSN 1799-4934 (printed) ISSN 1799-4942 (pdf) Aalto University School of Engineering Department of Engineering Design and Production www.aalto.fi

BUSINESS + ECONOMY ART + DESIGN + ARCHITECTURE SCIENCE + TECHNOLOGY CROSSOVER DOCTORAL DISSERTATIONS

Aalto-D

D 6

8/2

015

Hydrogen embrittlement resistance is an important parameter for the development of new steel grades for different applications. Hydrogen embrittlement is a phenomenon that causes loss of ductility in a metal, making it brittle and enhancing cracking and failure. The thesis describes studies of hydrogen effects in two classes of steels – austenitic stainless steels and advanced high-strength carbon steels. One of the used methods is thermal desorption spectroscopy, which allows to analyse hydrogen diffusion and trapping in metals, applying mathematical modeling of the processes of hydrogen uptake and desorption. The parameters of hydrogen diffusing and trapping were estimated for the different grades of austenitic stainless steels. The role of non-metallic inclusions in hydrogen embrittlement and hydrogen-induced fracture was studied for high-strength carbon steels with different micro-alloying. Mechanisms of hydrogen interaction with non-metallic inclusions are discussed.

Olga T

odoshchenko H

ydrogen effects on austenitic stainless steels and high-strength carbon steels A

alto U

nive

rsity

Department of Engineering Design and Production

Hydrogen effects on austenitic stainless steels and high-strength carbon steels

Olga Madelen Ingrid Todoshchenko

DOCTORAL DISSERTATIONS

Aalto University publication series DOCTORAL DISSERTATIONS 68/2015

Hydrogen Effects on Austenitic Stainless Steels and High-Strength Carbon Steels


A doctoral dissertation completed for the degree of Doctor of Science in Technology to be defended, with the permission of the Aalto University School of Engineering for public examenation and debate in Auditorium K215 at the Aalto University School of Engineering (Espoo, Finalnd) on the 12th of June 2015 at noon (at 12 o'clock).

Aalto University School of Engineering Department of Engineering Design and Production Engineering Materials

Supervising professor Hannu Hänninen Thesis advisor Yuriy Yagodzinskyy Preliminary examiners Prof. David Porter, University of Oulu, Finland Prof. Thomas Böllinghaus, Federal Institute for Materials Research and Testing (BAM), Berlin, Germany Opponent Yukitaka Murakami, Kyushu University, Japan

Aalto University publication series DOCTORAL DISSERTATIONS 68/2015 © Olga Madelen Ingrid Todoshchenko ISBN 978-952-60-6205-1 (printed) ISBN 978-952-60-6204-4 (pdf) ISSN-L 1799-4934 ISSN 1799-4934 (printed) ISSN 1799-4942 (pdf) http://urn.fi/URN:ISBN:978-952-60-6204-4 Unigrafia Oy Helsinki 2015 Finland Publication orders (printed book): [email protected]

Abstract Aalto University, P.O. Box 11000, FI-00076 Aalto www.aalto.fi

Author Olga Madelen Ingrid Todoshchenko Name of the doctoral dissertation Hydrogen Effects on Austenitic Stainless Steels and High-Strength Carbon Steels Publisher School of Engineering Unit Deparment of Engineering Design and Production

Series Aalto University publication series DOCTORAL DISSERTATIONS 68/2015

Field of research Engineering Materials

Manuscript submitted 9 February 2015 Date of the defence 12 June 2015

Permission to publish granted (date) 20 April 2015 Language English

Monograph Article dissertation (summary + original articles)

Abstract The resistance to hydrogen embrittlement is an important factor in the development of new

steel grades for a variety of applications. The thesis describes investigations on hydrogen effects on two classes of steels - austenitic stainless steels and advanced high-strength carbon steels.

Hydrogen solubility and diffusion in metastable austenitic stainless steels are studied with thermal desorption spectroscopy (TDS). This method, together with the mathematical modeling of the processes of hydrogen uptake and desorption, allows to analyse hydrogen diffusion and trapping in a metal.

Temperature dependencies of hydrogen desorption for the studied steels manifest a complex peak. Specific features of hydrogen uptake and desorption for a multi-component alloy in comparison with that for pure metals are analysed by the proposed model. It was found that plastic strain affects the shape of the TDS peak for all the studied materials. In metastable austenitic stainless steels the parameters of hydrogen diffusion and trapping in martensite phase, which is forming under pre-straining, were estimated by the proposed model.

High-strength carbon steels of the strength level from 1000 to 1400 MPa with different contents of Ti were studied. It was found that there is an optimal content and size distribution of Ti-based non-metallic inclusions (NMI) when the steel is most resistant to hydrogen embrittlement.

Fractography of the studied steels shows that the fracture mechanism depends on the chemical composition of the studied steels and hydrogen-induced cracking exhibits intergranular or transgranular character occurring often in the form of hydrogen flakes. It was found that hydrogen-induced cracks of the hydrogen flakes initiate at Ti-based NMIs. TDS analysis evidences that the main trapping sites for hydrogen in the high-strength carbon steels are NMI interfaces. The mechanisms of hydrogen interaction with NMIs are discussed.

All the studied high-strength carbon steels are sensitive to hydrogen under slow strain rate tensile tests. Constant load tests show that different processes influence on hydrogen-induced fracture at low and high applied stresses for the high-strength carbon steels with high Ti-alloying. Tempered high-strength carbon steels are less sensitive to hydrogen embrittlement than their non-tempered counterparts.

Keywords hydrogen embrittlement, austenitic stainless steels, high-strength carbon steels, thermal desorption spectroscopy, fracture, non-metallic inclusions

ISBN (printed) 978-952-60-6205-1 ISBN (pdf) 978-952-60-6204-4

ISSN-L 1799-4934 ISSN (printed) 1799-4934 ISSN (pdf) 1799-4942

Location of publisher Helsinki Location of printing Espoo Year 2015

Pages 144 urn http://urn.fi/URN:ISBN:978-952-60-6204-4

Tiivistelmä Aalto-yliopisto, PL 11000, 00076 Aalto www.aalto.fi

Tekijä Olga Madelen Ingrid Todoshchenko Väitöskirjan nimi Vedyn vaikutus austeniittisiin ruostumattomiin teräksiin ja ultra-lujiin hiiliteräksiin Julkaisija Insinööritieteiden korkeakoulu Yksikkö Koneenrakennustekniikan laitos

Sarja Aalto University publication series DOCTORAL DISSERTATIONS 68/2015

Tutkimusala Koneenrakennuksen materiaalitekniikka

Käsikirjoituksen pvm 09.02.2015 Väitöspäivä 12.06.2015

Julkaisuluvan myöntämispäivä 20.04.2015 Kieli Englanti

Monografia Yhdistelmäväitöskirja (yhteenveto-osa + erillisartikkelit)

Tiivistelmä Vetyhaurauden estäminen on tärkeää erilaisten uusien terästen suunnittelussa ja

kehityksessä. Väitöskirjassa esitetään vedyn vaikutuksen tutkimustulokset kahdella terästyypillä: austeniittisilla ruostumattomilla teräksillä ja ultra-lujilla hiiliteräksillä.

Vedyn liukoisuutta ja diffuusiota metastabiileissa austeniittisissa ruostumattomissa teräksissä tutkittiin termisen desorption spektroskopian (TDS) avulla. Tällä menetelmällä voidaan analysoida vedyn diffuusiota ja loukkuuntumista metalleissa käyttäen matemaattista mallinnusta.

TDS käyrissä esiintyy monimutkainen piikki kaikilla tutkituilla teräksillä. Vedyn loukkuuntumista ja desorptiota on verrattu monikomponenttisissa seoksissa ja puhtaissa metalleissa ja analysoitu kehitetyn mallin mukaan. Todettiin, että plastinen deformaatio vaikuttaa TDS piikin muotoon kaikilla tutkituilla teräksillä. Martensiitin kiderakenteen vaikutukset vedyn diffuusioon ja loukkuuntumiseen on arvioitu metastabiileissa austeniittisissa ruostumattomissa teräksissä, joissa martensiitti muodostuu plastisessa deformaatiossa.

Väitöskirjassa tutkittiin myös ultra-lujia hiiliteräksiä lujuusluokassa 1000...1400 MPa varioimalla titaani-seostusta. Tulosten perusteella todettiin, että on optimaalinen TiN(C) partikkelien määrä ja kokojakauma, jolloin teräs kestää parhaiten vetyhaurautta.

Tutkittujen terästen murtopintatutkimus osoitti, että murtumismekanismi riippuu teräksen kemiallisesta koostumuksesta. Vedyn aiheuttama murtuma voi esiintyä raerajamurtumana, rakeiden läpi etenevänä murtumana tai vetyläikkinä, jotka ydintyvät epämetallista partikkeleista ja ovat aina rakeiden läpi eteneviä murtumia. Vetyläikät ydintyvät yleensä TiN(C) epämetallisista partikkeleista. TDS analyysi vahvisti, että ultra-lujissa hiiliteräksissä tärkeimpiä vedyn loukkuuntumispaikkoja ovat näiden partikkelien pinnat. Vedyn ja epämetallisten partikkeleiden vuorovaikutuksen mekanismeja arvioidaan.

Kaikki hidasvetokokeissa tutkitut ultra-lujat hiiliteräkset ovat alttiita vedyn vaikutukselle. Vakio jännityksen alaisena tehdyt kokeet osoittivat, että eri mekanismit vaikuttavat vedyn aiheuttamaan murtumiseen matalilla ja korkeilla jännityksillä kaikissa ultra-lujissa hiiliteräksissä, joissa on korkea Ti pitoisuus. Lämpökäsitellyt päästetyt ultra-lujat hiiliteräkset ovat vähemmän alttiita vetyhaurauteen kuin lämpökäsittele-mättömät sammutetut teräkset.

Avainsanat Vetyhauraus, austeniittiset ruostumattomat teräkset, ultra-lujat hiiliteräkset, termisen desorption spektroskopia, murtuminen, epämetalliset partikkelit

ISBN (painettu) 978-952-60-6205-1 ISBN (pdf) 978-952-60-6204-4

ISSN-L 1799-4934 ISSN (painettu) 1799-4934 ISSN (pdf) 1799-4942

Julkaisupaikka Helsinki Painopaikka Espoo Vuosi 2015

Sivumäärä 144 urn http://urn.fi/URN:ISBN:978-952-60-6204-4

Sammandrag Aalto-universitetet, PB 11000, 00076 Aalto www.aalto.fi

Författare Olga Madelen Ingrid Todoshchenko Doktorsavhandlingens titel Väte effekter på austenitiska rostfria stål och höghållfasta kolstål Utgivare Högskolan för ingenjörsvetenskaper Enhet Institutionen för produktutveckling och produktion

Seriens namn Aalto University publication series DOCTORAL DISSERTATIONS 68/2015

Forskningsområde Konstruktionsmaterialier

Inlämningsdatum för manuskript 09.02.2015 Datum för disputation 12.06.2015

Beviljande av publiceringstillstånd (datum) 20.04.2015 Språk Engelska

Monografi Sammanläggningsavhandling (sammandrag plus separata artiklar)

Sammandrag Motståndet för väteförsprödning är en viktig faktor för utveckling av nya stålkvaliteter i en

mängd tillämpningar. Tesen beskriver undersökningar på väte effekter i två klasser av stål - austenitiska rostfria stål och avancerade höghållfasta kolstål.

Väte löslighet och diffusion i de metastabila austenitiska stålen studerades med termisk desorption spektroskopi (TDS). Denna metod tillsammans med den matematiska modeleringen av processer väteupptag och desorption gör det möjligt att analysera diffusion och vätefällor i en metall.

Temperaturberoende av väte desorption för de undersökta stålen bevisar en komplicerad topp. Specifika särdrag av väteupptag och desorption för en multikomponent legering i jämförelse med det för rena metaller är analyserade med den föreslagna modellen. Det upptäckets att plastisk deformation påverkar formen av TDS toppen på alla studerade materialer. För metastabila austenitiska rostfria stålen är parametrar av väte diffusion och upptag i den martensitiska fasen, som formas under plastisk deformation, beräknade med den föreslagna modellen.

Höghållfasta kolstål med draghållfasthetsegenskaperna från 1000 MPa till 1400 MPa med olika mängder av Ti blev undersökta. Det framkom att där finns en optimal mängd och storleksfördelning av Ti-baserade icke-metalliska inneslutningar när stål är mest resistent mot väteförsprödning.

Fraktografi av de undersökta stålen påvisar att brott-mekanismen beror på den kemiska sammansättningen av det undersökta stålet och väte-inducerad sprickbildning visar både interkristallin eller transkristallin karaktär och förekommar ofta i en form av väte flakes. Det konstaterades att väte-inducerade sprickor initieras på Ti-baserade icke-metalliska inneslutningar. TDS analyser visar att väte fällorna är fasgränser av inneslutningar och stål. Mekanismen av väte interaktion med icke-metalliska inneslutningar har diskuterats.

Alla studerade höghållfasta kolstål är känsliga för väte under dragprovning. Provning med konstant belastning påvisar att olika processer inverkar på väte-inducerad sprickning med låg eller hög mekanisk belastning av stål med hög Ti-legering. Seghärdade höghållfasta kolstål är mindre sensitiva för väte försprödning som stål utan efterföljande anlöpning.

Nyckelord väteförsprödning, austenitiska rostfria stål, höghållfasta kolstål, termisk desorption spektroskopi, brott, icke-metalliska inneslutningar

ISBN (tryckt) 978-952-60-6205-1 ISBN (pdf) 978-952-60-6204-4

ISSN-L 1799-4934 ISSN (tryckt) 1799-4934 ISSN (pdf) 1799-4942

Utgivningsort Helsingfors Tryckort Esbo År 2015

Sidantal 144 urn http://urn.fi/URN:ISBN:978-952-60-6204-4

Preface

The present thesis is based on the work done at Aalto University School

of Engineering by the research group of Engineering Materials at the De-

partment of Engineering Design and Production during 2010 - 2014. The

study was funded by RFCS project "Hydrogen Sensitivity of Different Ad-

vanced High Strength Microstructures" (HYDRAMICROS), by FIMECC’s

Light and efficient solutions program (LIGHT) project 1: "Production and

properties of breakthrough materials" (TEKES, Finland) and by the project

of the Academy of Finland "Interstitial Effects on Plastic Strain Localiza-

tion in Austenitic Alloys" (IEPLA).

This thesis would have not been possible without the help of many peo-

ple who worked with me inspiring me to these studies. I wish to express

my gratitude to the supervisor of my thesis, professor Hannu Hänninen

for giving me opportunity to study engineering materials in Aalto Univer-

sity, for his guidance and advise. I would like to thank my thesis advisor

Dr. Yuriy Yagodzinskyy for his theoretical support and practical advices,

for his scientific discussions, which encouraged me in the studies.

It is my pleasure to acknowledge kind assistance of the laboratory per-

sonnel. Special thanks to Tapio Saukkonen for SEM observations. Also I

would like to thank Kim Widell, Jari Hellgren, Janne Peuraniemi, Seppo

Nurmi and Valentina Yagodzinska for their kind support in preparation

and running of the experiments and Johanna Salmela for her help in ad-

ministrative issues.

I am grateful to my husband Dr. Igor Todoshchenko from O.V. Lounas-

maa Laboratory for his discussions and comments, helping me in physical

understanding of the studied phenomena. I wish also to thank my chil-

dren Valentin, Seva and Jelena for their patience and understanding.

1

Preface

Espoo, May 3, 2015,


2

Contents

Preface 1

Contents 3

List of Publications 5

Author’s Contribution 7

Original Features 9

List of Abbreviations and Symbols 11

1. Introduction 13

1.1 Mechanisms of hydrogen embrittlement . . . . . . . . . . . . 13

1.2 Inclusions in steels . . . . . . . . . . . . . . . . . . . . . . . . 15

1.2.1 The origin of non-metallic inclusions in steels . . . . . 16

1.2.2 Non-metallic inclusions and hydrogen embrittlement 16

1.3 Thermal desorption spectroscopy in the studies of hydrogen

in metals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

1.4 Aims of the study . . . . . . . . . . . . . . . . . . . . . . . . . 20

2. Experimental 23

2.1 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

2.1.1 High-strength carbon steels . . . . . . . . . . . . . . . 23

2.1.2 Austenitic stainless steels . . . . . . . . . . . . . . . . 24

2.2 Electrochemical charging . . . . . . . . . . . . . . . . . . . . . 25

2.3 Thermal desorption spectroscopy . . . . . . . . . . . . . . . . 27

2.4 Mechanical tests . . . . . . . . . . . . . . . . . . . . . . . . . . 28

2.5 Fractography . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29

2.6 Optical observations . . . . . . . . . . . . . . . . . . . . . . . 30

3

Contents

3. Studies of hydrogen diffusion, trapping and solubility in

austenitic stainless steels 31

3.1 Model and simulations of thermal desorption of hydrogen . . 31

3.2 Hydrogen diffusion, trapping and solubility in different grades

of austenitic stainless steels . . . . . . . . . . . . . . . . . . . 32

4. Studies of hydrogen effects on mechanical properties and

fracture of high-strength carbon steels 39

4.1 Constant extension rate tests (CERT) . . . . . . . . . . . . . 39

4.2 Constant load tests (CLT) . . . . . . . . . . . . . . . . . . . . 39

4.3 Fractography and non-metallic inclusions . . . . . . . . . . . 42

4.4 Hydrogen - NMI interactions and crack initiation . . . . . . 45

4.5 Thermal desorption analysis of hydrogen in high-strength

carbon steels . . . . . . . . . . . . . . . . . . . . . . . . . . . . 46

4.6 Strain localization . . . . . . . . . . . . . . . . . . . . . . . . . 48

4.6.1 EBSD mapping . . . . . . . . . . . . . . . . . . . . . . 48

4.6.2 Optical observations using LaVision system . . . . . . 50

5. Conclusions 53

Bibliography 55

Errata 63

Publications 65

4

List of Publications

This thesis consists of an overview and of the following publications which

are referred to in the text by their Roman numerals.

I Y. Yagodzinskyy, O. Todoshchenko, S. Papula, H. Hänninen. Hydro-

gen Solubility and Diffusion in Austenitic Stainless Steels Studied with

Thermal Desorption Spectroscopy. Steel Research Int., Vol. 82, No 1, pp.

20-25, 2011.

II O. Todoshchenko, Y. Yagodzinskyy, S. Papula, H. Hänninen. Hydro-

gen Solubility and Diffusion in Metastable Austenitic Stainless Steels

Studied with Thermal Desorption Spectroscopy. In Hydrogen-Materials

Interactions. Proceedings of the 2012 International Hydrogen Confer-

ence, edited by B.P. Sommerday and P. Sofronis, USA, ASME Press, pp.

615-623, 2014.

III O. Todoshchenko, Y. Yagodzinskyy, H. Hänninen. Thermal Desorption

of Hydrogen from AISI 316L Stainless Steel and Pure Nickel. Defect and

Diffusion Forum, Vol. 344, No Properties of Lanthanum Hexaboride - A

Compilation, pp. 71-77, 2013.

IV M. Ganchenkova, O. Todoshchenko, Y. Yagodzinskyy, H. Hänninen.

Hydrogen Solubility and Trapping in AISI 316L Austenitic Stainless

Steel. In Hydrogen-Materials Interactions. Proceedings of the 2012 In-

ternational Hydrogen Conference, edited by B.P. Sommerday and P. Sofro-

nis, USA, ASME Press, pp. 605-613, 2014.

5

List of Publications

V O. Todoshchenko, Y. Yagoszinskyy, T. Saukkonen, H. Hänninen. Role of

Non-metallic Inclusions in Hydrogen Embrittlement of High-strength

Carbon Steels with Different Micro-alloying. Metallurgical and Materi-

als Transactions A, Vol. 45, No 11, pp. 4742-4747, 2014.

VI O. Todoshchenko, Y. Yagodzinskyy, V. Yagodzinska, T. Saukkonen, H.

Hänninen. Hydrogen Effects on Fracture of High-Strength Steels with

Different Micro-alloying. Accepted for publication in Corrosion Reviews,

13 p., 2 March 2015.

VII Y. Yagodzinskyy, O. Todoshchenko, T. Saukkonen, H. Hänninen. Role

of Non-metallic Inclusions in Hydrogen-induced Fracture of High-strength

Carbon Steels. In Proceedings of the 4th International Symposium on

Steel Science (ISSS-2014), edited by K. Tsuzaki, T. Ohmura and N. Tsuji,

ISIJ, pp. 147 - 150, 2014.

6

Author’s Contribution

Publication I: “Hydrogen Solubility and Diffusion in AusteniticStainless Steels Studied with Thermal Desorption Spectroscopy”

The author has prepared and run TDS measurements and carried out hy-

drogen charging of the specimens. The author has also contributed in the

analysis of TDS results. The author was responsible for the development

of the model of hydrogen diffusion in the presence of traps and for simu-

lations of TDS spectra and has developed original computer program for

the fitting of TDS peaks using MatLab software.

Publication II: “Hydrogen Solubility and Diffusion in MetastableAustenitic Stainless Steels Studied with Thermal DesorptionSpectroscopy”

The author has performed all TDS measurements and analysed obtained

results. The author also provided computer modelling and simulations,

using developed computer programmes. The manuscript of the article

was prepared and written by the author.

Publication III: “Thermal Desorption of Hydrogen from AISI 316LStainless Steel and Pure Nickel”

The author has prepared and run TDS measurements and carried out hy-

drogen charging of the specimens. The author worked out matematical

model and made computer simulations of TDS spectra, using developed

and MatLab software. The author has analysed experimental and calcu-

lated data and has written the text of the article.

7

Author’s Contribution

Publication IV: “Hydrogen Solubility and Trapping in AISI 316LAustenitic Stainless Steel”

The author has carried out TDS measurements, computer simulations of

TDS spectra and participated in the discussion of the results.

Publication V: “Role of Non-metallic Inclusions in HydrogenEmbrittlement of High-strength Carbon Steels with DifferentMicro-alloying”

The author has prepared and run CERT and CLT tests and made SEM

microscopy of the samples. The author has analysed the experimental

data, contributed in the analysis of the results and has written the text of

the article.

Publication VI: “Hydrogen Effects on Fracture of High-StrengthSteels with Different Micro-alloying”


microscopy. The author has contributed in the analysis of the experimen-

tal data and has written the text of the article.

Publication VII: “Role of Non-metallic Inclusions inHydrogen-induced Fracture of High-strength Carbon Steels”


microscopy. The author has contributed in the analysis of the experimen-

tal data and in the discussion of the results.

8

Original Features

Thermal desorption flux of hydrogen from multi-component alloys such

as austenitic stainless steels cannot be described by a simple diffusion

model. The model of TDS, taking into consideration the diffusion of hy-

drogen into the metal crystal lattice and its trapping by the trapping sites

with different binding energies of hydrogen atoms and different concen-

trations of trapping sites was developed. New data for hydrogen solubility,

trapping energies and diffusion parameters were obtained for industrially

important stainless steels.

The analysis of experimental TDS curves and accurate simulation of

the processes by the developed model allow to estimate the parameters

of hydrogen diffusion and trapping for the studied austenitic stainless

steels. The proposed model allows also to estimate the hydrogen diffusion

parameters of martensite phase of the steel, which is difficult to measure

experimentally.

A characteristic feature of the fracture surface of high-strength carbon

steels with high Ti-alloying after CERT and CLT tests under continuous

hydrogen charging is hydrogen flakes, which form transgranularly around

titanium-based particles. For the steels with low content of Ti, high con-

tent of Mn and containing Nb hydrogen-induced cracking has intergran-

ular character and these steels are more sensitive to hydrogen embrittle-

ment.

Hydrogen sensitivity factors and times to fracture were obtained for the

different high-strength steel grades.

Maximum resistivity to hydrogen embrittlement of high-strength car-

bon steels was found in the steels with optimal Ti-alloying, where Ti-

based non-metallic particles are small.

Using the original approach to microstructural and TDS analyses it was

found the main trapping sites for hydrogen in the studied high-strength

9

Original Features

carbon steels are situated on the NMI interfaces.

It was found that high-strength carbon steels with tempered martensite

are less sensitive to hydrogen than their martensitic counterparts.

10

List of Abbreviations and Symbols

Abbreviations

CERT - constant extension rate test

CLLM - crystal lattice local misorientation map

CLT - constant load test

DIC - digital image correlation

EBSD - electron backscatter diffraction

EDS - energy-dispersive X-ray spectroscopy

fcc - face-centered cubic lattice structure

HE - hydrogen embrittlement

HELP - hydrogen-enhanced localized plasticity

HEDE - hydrogen enhanced decohesion

IPF - inverse pole figure

NMI - non-metallic inclusions

SEM - scanning electron microscopy

TDS - thermal desorption spectroscopy

TRIP - transformation-induced plasticity

UHV - ultra-high vacuum

Symbols

aH - hydrogen activity in metal

C0 - hydrogen concentration at the specimen surface

δH - hydrogen sensitivity factor

D - diffusion coefficient

DL - the diffusion constant in a crystal lattice

Ea - activation energy of detrapping of hydrogen

Eb - binding energy of hydrogen

Ed - activation energy of diffusion of hydrogen

E′ - saddle point energy of hydrogen

11

List of Abbreviations and Symbols

E - applied electrochemical potential

E0 - standard electrode potential

εxx, εyy - strain tensor components

F - Faraday’s constant

KX - density of the trapping site

kB - Boltzmann’s constant, eV· K−1

L0 - elongation to fracture for hydrogen-free specimens

LH - elongation to fracture for hydrogen-charged specimens

R - gas constant

Rm - ultimate tensile strength

Rp0.2 - yield stress

t - time

T - temperature

ts - time of hydrogen charging

Ts - temperature of hydrogen charging

td - time of hydrogen desorption

Td - temperature of hydrogen desorption

Tc - peak temperature of thermal desorption spectra

φ - heating rate of thermal desorption process

12

1. Introduction

The deleterious effect of hydrogen on the mechanical properties of metals

and alloys has been studied since the beginning of 20th century. At the

first time the phenomenon of hydrogen embrittlement of iron was men-

tioned in 1875 by W.H. Johnson [1]. Various mechanisms of hydrogen

embrittlement have been widely discussed and reviewed in the literature

[2], [3], [4], [5], [6], [7], but there is no clear understanding of this phe-

nomenon and mechanisms of hydrogen embrittlement of different mate-

rials are still under discussion.

It is widely accepted that increasing of the metallic material strength

leads to a higher sensitivity to hydrogen in a variety of engineering mate-

rials such as carbon and stainless steels, aluminum-base alloys, etc. as a

result of hydrogen embrittlement (HE) [8], [9].

1.1 Mechanisms of hydrogen embrittlement

Hydrogen embrittlement is a phenomenon that causes loss of ductility in

a metal, making it brittle and enhancing cracking and failure.

The mechanism of hydrogen embrittlement of high-strength carbon steels

has been widely discussed during the last decades [2], [10], [11], [12]. The

presence of a few wt.ppm of hydrogen, for instance in carbon steels may

result in a dramatic reduction of elongation to fracture, in a loss of duc-

tility or even initiate a brittle fracture without macroscopic plastic strain

[8], [9]. Another effect of hydrogen on steels is so-called hydrogen-induced

delayed fracture, when the material cracks under constant load which can

be even less than the load corresponding to the yield stress of the material

[13].

Different models were proposed to explain the phenomenon of hydro-

gen embrittlement. Enhanced ductile processes due to hydrogen interac-

13

Introduction

tion was suggested by Beachem [14]. He suggested that hydrogen stim-

ulates dislocation processes that localize plastic deformation sufficiently

to result in subcritical crack growth with brittle characteristics on the

macroscopic scale. The mechanism of hydrogen-enhanced localized plas-

ticity (HELP) was widely discussed [8], [15], [16]. According to the HELP

mechanism, hydrogen induced premature failures result from hydrogen

induced plastic instability.

The hydrogen enhanced decohesion (HEDE) mechanism was first sug-

gested by Troiano [17], [18], and developed in details in the works [19],

[20], [21], [22], [23]. In this model hydrogen accumulates in the lattice

and reduces the cohesive bonding strength between atoms of the host

metal. According to HEDE hydrogen damage occurs in the crack tip frac-

ture process zone when the local crack tip opening tensile stress exceeds

the maximum local atomic cohesion strength, lowered by the presence of

hydrogen [19].

Also the other models such as hydride-induced embrittlement [24], [25],

where the processes of hydride nucleation and growth, together with the

brittle nature of hydrides seems to be the main cause of embrittlement of

typical hydride forming elements.

It is widely discussed today that the vacancies enhanced by hydrogen

take the major part in the hydrogen embrittlement [7], [10], [29]. The

hydrogen - enhanced strain - induced vacancies (HESIV) mechanism [7]

which considers vacancy clusters rather than hydrogen itself as the pri-

mary factor of HE was developed. Excessive vacancies, as assumed, are

formed due to hydrogen - dislocation interactions and operate as effective

trapping sites for hydrogen. The TDS peaks observed in the high-strength

carbon steels are often referred to the vacancy or vacancy-type trapping

sites [29]. The amount of trapped hydrogen in vacancies increases with

deformation. Hydrogen content and applied stress, time of the formation

and accumulation of vacancies have been concluded to be important fac-

tors causing hydrogen embrittlement [10].

All these models can explain some aspects of hydrogen embrittlement,

but none of them has reached the complete explanation. Despite the dif-

ference in details all the models assume diffusion and segregation of hy-

drogen atoms in metals [26]. The processes of hydrogen diffusion and

trapping in metallic materials are presumed to be important for under-

standing mechanisms of hydrogen embrittlement [27], [28].

Austenitic stainless steels are considered to be resistant to hydrogen

14

Introduction

embrittlement because of low diffusivity and high solubility of hydrogen

in the FCC lattice structures. Brittle fracture associated with hydrogen

embrittlement (HE) is, however, observed for stainless steels in severe

environmental conditions such as cathodic charging [30], [31]. Some of

the steel grades designed recently, namely, metastable and Mn-alloyed

austenitic stainless steels manifest, however, a remarkable sensitivity to

hydrogen. The effects of martensitic structure on the hydrogen content of

austenitic stainless steels have been studied in a number of works [32],

[33], [34], [35]. There are many industrial applications of austenitic stain-

less steels, which provide the environmental conditions resulting in the

hydrogen uptake of electrochemical origin, for instance, under conditions

of cathodic protection of the steel constructions or under corrosion reac-

tions. The diffusion of hydrogen introduced into the steel electrochem-

ically, which is typically hydrogen of high fugacity, may differ from the

corresponding parameter in the case hydrogen is introduced from the low

fugacity gas atmosphere.

1.2 Inclusions in steels

There always exists a number of inclusions in steels which can have de-

grading or beneficial effects on their properties. Despite of small content

of non-metallic inclusions they have a significant effect on such steel prop-

erties as tensile strength, fatigue properties [36], corrosion resistance and

hydrogen embrittlement resistance, ductility, machinability and weldabil-

ity [37]. Nowadays the inclusion problem is generally associated with the

high-strength steels, while the steel industry manufactures rather clean

steels.

Non-metallic inclusion generates stress concentrations around it and in-

side the inclusion. During the hot rolling of the steel the stresses between

inclusions and matrix are relaxed. During the cooling that follows hot

rolling, tensile residual stresses are generated around an inclusion due to

the differences in the thermal expansion coefficients of the inclusion and

the matrix [36]. The value and the location of the maximum stress de-

pend on the values of Young’s modulus and Poisson’s ratio of the inclusion

and of the matrix and on the shape of the inclusion. Usually inclusions

are harder than the surrounding matrix at room temperature. It leads

to the stress and strain concentration during matrix deformation which

can produce voids by matrix-inclusion decohesion or fracture of inclusion

15

Introduction

particles [38]. Stress concentrations at a spherical and elliptical holes are

analyzed in [39]. The problem of determination of the elastic field inside

and outside of an ellipsoidal inclusion was analyzed in the works [40],

[41], [42].

1.2.1 The origin of non-metallic inclusions in steels

Non-metallic inclusions in steels can be divided into two groups, those of

indigenous and those of exogenous origin [37]. Non-metallic inclusions

of exogenous origin appear as a result of reactions taking place in the

molten or solidifying steel bath, also they can appear as a result of me-

chanical incorporation of slags, refractories or other materials with which

the molten steel comes in contact. The characteristic features of these

inclusions are generally larger size, random occurrence, irregular shape

and complex structure. The indigenous inclusions [43] are those that form

by precipitation as a result of homogeneous reactions in the steel. The re-

actions that form them may be induced either by the additions to the steel

or by changes in the solubility during the cooling and freezing of the steel

[37]. These indigenous inclusions (precipitates) have usually smaller size.

When inclusions are formed, it is not important whether this occurs from

homogeneous nucleation or from exogenous sources. Ti-based NMIs of dif-

ferent sizes as well as TiN/TiC fine precipitates are present in the studied

high-strength carbon steels.

1.2.2 Non-metallic inclusions and hydrogen embrittlement

It has been found that non-metallic inclusions (NMI) act as a strong trap-

ping site for hydrogen [13] and play the major role in hydrogen embrittle-

ment of high-strength carbon steels. Hydrogen trapping by TiC particles

was analysed and trap activation energies of hydrogen were estimated in

[44], [45], [46].

Hydrogen was also assumed to enhance non-metallic inclusion decohe-

sion or cracking [7]. Hydrogen-induced fracture initiating on TiN and

Al2O3·(CaO)x particles in steel is studied by [47]. In the study it is sup-

posed, that hydrogen is concentrated in voids, which are forming around

inclusions. Under the tensile loading voids grow, initiating the crack.

The presence of hydrogen increases the density of strain-induced trap-

ping sites and reduces ductile crack growth resistance. In the work of

Nagumo et al. [13] the delayed fracture was examined with respect to

16

Introduction

Figure 1.1. Energy states profile around a hydrogen trapping site: Ea, Eb and E′ are thedetrapping activation energy, the binding energy and the saddle point energy,respectively. Ed is the diffusion activation energy [52].

microstructures in Mo-V martensitic steels tempered at 550◦C and 650◦C

of a similar strength level. It was found that non-metallic precipitates act

as a strong trapping site for hydrogen with a uniform distribution in the

matrix for the samples tempered at 650◦C.

1.3 Thermal desorption spectroscopy in the studies of hydrogen inmetals

Thermal desorption spectroscopy (TDS) of hydrogen is an unique exper-

imental method to study hydrogen uptake and trapping in metals and

alloys [48], [Publication I], what is necessary to analyse and control HE.

In the TDS, by measuring the amount of hydrogen released from a spec-

imen that is heated with a constant rate, the hydrogen desorption rate

as a function of specimen temperature or the hydrogen desorption spec-

trum is obtained. The trapping state of hydrogen in the specimen is esti-

mated from the desorption peaks of hydrogen [58]. The energy states pro-

file of hydrogen in the crystal lattice around a trapping site is presented

schematically in Fig. 1.1. [52].

TDS spectrum provides detailed information about hydrogen trapping

states based on determination of the peak temperatures for hydrogen des-

orption during heating and applying numerical modeling of hydrogen up-

take and desorption [52], [53]. However, the distinction of trapping sites

by TDS is not straightforward because of overlapping of temperature re-

gions of the desorption peaks of different trapping sites and the possibility

of detrapping of hydrogen during TDS measurement. It has to be taken

into consideration during the calculations. In [54] it has been reported

that desorption from such defects as grain boundaries, dislocations and

microvoids take place in the same temperature region, making difficult

17

Introduction

the identification of trapping site from TDS analysis.

The simple model of the thermal-detrapping-determining desorption of

hydrogen by Choo and Lee [54], [55] is used to determine the detrapping

activation energy of the defect site corresponding to the desorption peak

of hydrogen. It is formulated as follows [46]:

dX

dt= A(1−X) exp[−Ea/RT ], (1.1)

where X = C0−CC0

, C0 and C are the amount of hydrogen in the trapping

site at t = 0 and t > 0, respectively, Ea is the detrapping energy of hydro-

gen, R is the gas constant, T is the absolute temperature. An arbitrary

constant A is related to the rate constant of the detrapping process. This

model is used to determine the detrapping activation energy of the de-

fect site corresponding to the desorption peak of hydrogen. The following

equation is derived by using the condition that the derivative of Eq. (1.1)

is zero at the peak temperature Tc:

∂ ln(φ/T 2c )

∂(1/Tc)= −Ea/R, (1.2)

where φ is the heating rate. φ and Tc are measured in TDS and detrapping

energy Ea needed to escape from the trap site to the normal lattice site

is calculated from the slope of ln(φ/T 2c ) vs (1/Tc) plot by measuring the

change of peak temperatures with the heating rates [44].

This model works well enough only for the metals with low density of

trapping sites and high hydrogen diffusivity, but it is widely used for esti-

mating the binding energy of defects in different steels.

The effective hydrogen diffusion constant, D, under an assumption of

a local equilibrium between trapped and lattice hydrogen, was given by

Oriani [19] as:

D = DL[1 +KX exp(Eb

RT)]−1, (1.3)

where DL is the diffusion constant in a crystal lattice and Eb and KX

are the hydrogen-binding energy and the density of the trap, respectively.

The Oriani’s model requires that the diffusion process is slower than the

detrapping processes and amount of hydrogen at diffusion sites around

trapping sites cannot decrease rapidly. So, the local hydrogen equilibri-

ums hold between diffusion sites and trapping sites. The Oriani’s model is

limited only for the simulation of the desorption profile with a single peak

with the adoption of the effective diffusion coefficient including the trap-

ping effect. This model was used for analyzing the hydrogen detrapping

characteristics of dislocations and grain boundaries of iron [56]. More

18

Introduction

complicated model, based on the same principles, where the hydrogen lo-

cal equilibrium postulate is numerically imposed at each step in the tem-

perature raising process is proposed by Ebihara et al. [57]. The more

rigorous model of TDS, which allows to determine the binding energies

and density of trapping sites by numerical solution was proposed by Turn-

bull et al. [58]. The traps are considered to be independent of each other

and sample is assumed to be charged uniformly. Baskes et al. have pro-

posed the embedded atom method for numerical simulation of metallic

systems where fractures, surfaces, impurities and alloying additions can

be included [50], [51].

Reproduction of the total thermal desorption spectra of hydrogen re-

quires the application of general model of the thermal hydrogen desorp-

tion process, which is based on the mass conservation of hydrogen in the

specimen with the activation energy law for the diffusion, trapping and

detrapping processes [59], [60] is described as follows [52]:

∂C

∂t+∑i

∂Ci

∂t= D

∂2C

∂x2, (1.4)

∂Ci

∂T= kC(1− Ci

Ni)− piCi, (1.5)

where i = 1, n, C is the concentration of diffusive hydrogen in lattice, Ci,

Ni are concentrations of hydrogen and trapping site of type i, respectively,

pi and k - the rate constants of trapping and detrapping processes, kB

is Boltzmann’s constant, D is the diffusion coefficient of hydrogen in the

crystal lattice, which is dependent on temperature T:

D = D0 · exp −Ed

kBT, (1.6)

pi = p0 · exp −(Ebi + Ed)

kBT, (1.7)

k = k0 · exp −Ed

kBT, (1.8)

where Ed is the diffusion activation energy, Ebi is the binding energy of

trapping site i. D0, p0 and k0 are the pre-exponential factors of diffusion

rate constants of trapping and detrapping reactions. This model is more

complicated in mathematical realization and requires numerical solution.

Nagumo et al. [29] have studied the nature of defects acting as trapping

sites of hydrogen in ferritic steels deformed to various degrees and an-

nealed following with thermal desorption spectroscopy (TDS). They have

found that for ferritic steels vacancy clusters, which themselves annihi-

late during TDS measurement, act as a trapping site of hydrogen. Hy-

drogen trapping sites and binding energies for ultra-high strength steel

19

Introduction

AERMET 100 were determined using thermal desorption methods [61].

Three major hydrogen desorption peaks were indentified and associated

with three trapping states: precipitates, interfaces and grain boundaries

and undissolved metal carbides.

Hydrogen trapping in cold deformed multi-phase high strength transfor-

mation-induced plasticity (TRIP) steels using thermal desorption tech-

nique was investigated by [62]. The microstructure of these steels is

mainly ferrite, together with bainite, retained austenite and martensite.

In order to understand the interaction of these different phases with hy-

drogen it is important to understand the behavior of each phase when

exposed to hydrogen. Because of the complex microstructure of the TRIP

steel the study was done with reference to a pure iron material. Both ma-

terials were studied before and after cold deformation in order to under-

stand the effect of deformation-induced crystal defects and phase trans-

formation in the TRIP steel. Significant impact of the degree of cold defor-

mation on hydrogen uptake was observed for TRIP steels [62]. TDS peaks

correspond to strain induced microstructural changes and martensite for-

mation, grain boundaries and retained austenite.

Traditionally hydrogen solubility and diffusion in austenitic stainless

steels are measured with the hydrogen permeation method, which con-

sists of the accurate measurement of hydrogen diffusion flux through a

metallic membrane at temperatures above 200◦C [9], [63]. The electro-

chemical methods widely used in studies of hydrogen diffusion and trap-

ping in carbon steels meet a number of difficulties in the case of austenitic

stainless steels, due to the relatively slow hydrogen diffusion in austenite

as compared to that in carbon steels. TDS method is successfully used for

the study of residual hydrogen in stainless steels [64], [65]. TDS measure-

ments together with accurate mathematical modeling of the processes

of thermal hydrogen desorption allow to obtain not only parameters of

diffusion and solubility, but also activation energies of trapping sites in

austenitic stainless steels.

1.4 Aims of the study

The aim of the research work was to determine the most important factors

influencing the hydrogen embrittlement resistance of the steels of differ-

ent grades. The influence of micro-alloying on the character of hydrogen-

induced cracking and fracture modes of high-strength carbon steels have

20

Introduction

to be analysed, comparing different grades of high-strength steels.

It was also aimed to develop the model, allowing to predict hydrogen

diffusion and trapping parameters of multi-component alloys, and to com-

pare these parameters for different grades of austenitic stainless steels.

In austenitic stainless steels the parameters of hydrogen diffusion in mar-

tensite phase have to be estimated.

Another aim was to clarify the role of NMIs of different chemical compo-

sition in hydrogen embrittlement, hydrogen-induced crack initiation and

fracture of high-strength carbon steels. Mechanisms of hydrogen - NMI

interactions have to be analysed.

21

Introduction

22

2. Experimental

In this chapter studied materials and used experimental methods and

techniques are described.

2.1 Materials

2.1.1 High-strength carbon steels

High-strength steels are important materials, for example, for automo-

bile industry, where the reduction of the weight of the structure is crit-

ical. Three high-strength carbon steel grades with different contents of

Mn, Cr, Ti and Nb were chosen for the study. Chemical composition of the

steels is shown in Table 2.1. Steels in the form of sheets with thickness

of 1.0 to 1.5 mm were supplied by ThyssenKrupp Steel Europe, Germany,

Voestalpine Stahl GmbH, Austria and ArcelorMittal, Belgium, marked

below as S1, S2 and S3, respectively. The steels were treated by suppliers

to obtain the strength level from 1000 to 1400 MPa and they consist of fer-

rite and martensite, (marked below as M02 with 20% of martensite, M04

with 40% of martensite, etc.); ferrite and tempered martensite (marked

below as TM02 with 20% of tempered martensite and TM09 with 90%

of tempered martensite); martensite, tempered martensite, bainite and

ferrite (marked below as MTM) and tempered fully martensitic (TM) mi-

crostructures. Typical microstructures of the studied steels are presented

in Fig. 2.1 and Fig. 2.2. Microstructure of the steel S1_1000_M02 observed

with SEM is shown in Fig. 2.1 a. Ferrite and martensite grains manifest a

markedly different contrast. Dark islands of martensite have a fine struc-

ture which is revealed in the figure with black lines representing grain

boundaries. There are no variations of the contrast within the ferrite

grains evidencing that ferrite is free of local strain. Microstructure of the

23

Experimental

Figure 2.1. Microstructures of the ferrite - martensite steels S1_1000_M02 (a) andS2_1200_M05 (b) observed with SEM.

Figure 2.2. Microstructures of martensitic steels S1_1400_M (a) and S1_1400_TM (b)observed with SEM.

steel S2_1200_M05 observed with SEM is shown in Fig. 2.1 b. Ferrite and

martensite grains manifest a markedly different contrast. Dark fields of

martensite have a fine structure which is revealed with black lines repre-

senting grain boundaries. There are also variations of the contrast within

the martensite islands evidencing on the fine structure as well as a high

level of the martensite crystal lattice local strain [73].

Fully martensitic structures are presented in Fig. 2.2. Microstructure

of the steel S1_1400_M is shown in Fig. 2.2 a. The microstructure looks

homogeneous with remarkable variation of the contrast within all grains.

The grain size distribution seems to be wide ranging from sub-micrometer

size to a few micrometers. Microstructure of tempered fully martensitic

steel S1_1400_TM is shown in Fig. 2.2 b. It is homogeneous and rather

similar to that of S1_1400_M steel. The grain size distribution seems also

to be wide ranging from sub-micrometer size to a few micrometers [73].

2.1.2 Austenitic stainless steels

Industrial austenitic stainless steel grades produced at Outokumpu Stain-

less (Finland), namely, metastable AISI 301LN, Mn-alloyed AISI 201 and

24

Experimental

AISI 204Cu steels were chosen to study. AISI 310 stable austenitic steel

and AISI 304 austenitic steels were used as reference materials. The

steels were shipped as solid-solution treated in the sheet form of 1.5 mm

thickness.

Austenitic stainless steel AISI 316L supplied by GoodFellow in a sheet

form of 0.4 mm and 0.05 mm was used as a model material for TDS sim-

ulations in [Publication III, Publication IV]. Chemical composition of the

stainless steels is shown in Table 2.2.

2.2 Electrochemical charging

The key idea of the hydrogen charging procedure applied in the study was

to provide the same chemical potential to the hydrogen ion in the course of

its transfer from the electrolyte into the steel. In agreement with Nernst

equation, activity in metal aH depends on the applied potential E:

E = E0 +RT

Fln aH , (2.1)

where E0 is the standard electrode potential, R the gas constant and F

the Faraday’s constant.

Thus, applying the constant cathodic potential in the potentiostatic hy-

drogen charging was assumed to provide the same activity of hydrogen for

all the studied steels.

Environmental cell used for electrochemical charging is shown in Fig. 2.3.

It consists of a double-wall glass compartment with a Luggin-probe of

Hg / Hg2SO4 reference and Pt counter electrodes. The electrolyte was

pumped and circulated between the cell and an additional compartment

used for de-aeration by nitrogen gas bubbling as it is shown on Fig. 2.3

b. Widely used electrochemical charging from 1N H2SO4 solution with

addition of 20 mg/L thiourea (CH4N2S) as a "poison" of hydrogen recom-

bination at the metal surface was chosen to introduce hydrogen into the

studied steels.

Due to high diffusivity of hydrogen in iron at room temperature the hy-

drogen charging time for high-strength carbon steels was chosen to be 1

hour to provide a homogeneous hydrogen distribution over the specimen

cross-section before the loading.

Relatively slow diffusion of hydrogen in austenitic stainless steels re-

sults in a requirement of a longer time of hydrogen charging. There is no

way to provide a homogeneous hydrogen distribution by electrochemical

25

Experimental

Figure 2.3. The general view of environmental cell (a) and the scheme of environmentalcell for electrochemical hydrogen charging (b).

Table 2.1. Chemical composition of the studied high-strength carbon steels (wt. %)

C Si Mn Al Ti Mo Cr Nb N

S1 0.164 0.72 1.90 0.04 0.114 0.01 0.37 0.001 0.0048

S2 0.157 0.19 2.24 0.053 0.002 0.004 0.46 0.022 0.006

S3 0.17 0.18 0.41 0.049 0.035 0.004 0.024 0.002 0.0056

Table 2.2. Chemical composition of the studied austenitic stainless steels (wt. %)

C Si Mn Cr Ni Cu N Mo

AISI 310 0.049 0.56 0.88 25.54 19.09 - 0.036 0.24

AISI 304 0.049 0.43 1.49 18.2 8.2 0.43 0.047 -

AISI 301LN 0.030 0.50 1.23 17.4 6.6 0.168 0.168 0.19

AISI 201 0.055 0.31 7.1 16.91 4.38 - 0.222 0.10

AISI 201A 0.045 0.35 6.97 17.6 4.5 0.25 0.198 -

AISI 204Cu 0.079 0.4 9.0 15.2 1.1 1.68 0.115 -

AISI 316L 0.022 0.55 0.98 16.4 10.5 - - 2.0

26

Experimental

Figure 2.4. Hydrogen concentration profile after electrochemical charging of 1 mm thickAISI 310 stainless steel plate for 100 h and dwelling at room temperature for0.5 h.

hydrogen charging in the millimeter scale for a reasonable charging time

of several days. Assuming that the hydrogen concentration at the speci-

men surface C0 is kept constant and solving the corresponding boundary

problem for diffusion equation one can obtain an expression for the hydro-

gen concentration profile [Publication I]:

C(x) =4C0

π

∞∑n=0

(−1)n

(2n+ 1)cos

(2n+ 1)πx

h

(1− e−

π2(2n+1)2D(Ts)tsh2

)·e−

π2(2n+1)2D(Td)tdh2 ,

(2.2)

where C0 is the concentration of hydrogen on the surface, ts and Ts are

time and temperature of hydrogen charging, td and Td are time and tem-

perature of hydrogen desorption, respectively, h is the thickness of speci-

men, and D is the diffusion coefficient.

Typical hydrogen concentration profile after electrochemical charging of

1 mm thick AISI 310 stainless steel plate for 100 h and dwelling at room

temperature for 0.5 h is shown in Fig. 2.4. The profile calculation was

performed with the following parameters: Ts = 323 K, Td = 293 K, ts =

3.6× 105 s, td = 1.8× 103 s and D(T ) = 7.16× 10−7 exp (−53.0kJ/mol/RT )

m2/s [74].

Typical hydrogen charging time for austenitic steels was chosen to be 18

h at the temperature of 40 ◦C.

2.3 Thermal desorption spectroscopy

TDS measurements were carried out with a thermal desorption apparatus

[Publication I] to analyze the uptake and trapping of hydrogen. The gen-

27

Experimental

Figure 2.5. The general view (a) and the scheme (b) of TDS apparatus.

eral view and the scheme of the apparatus is presented on Fig. 2.5. The

apparatus consists of the UHV-vacuum chamber equipped with a small

vacuum furnace, air-lock vacuum chamber for specimen supply, the spec-

imen transportation mechanism with magnetic support of the specimen

basket, the pumping system with an effective pumping rate of 6.6 × 10−2

m3/s and PC using Lab View based software for controlling the mass spec-

trometer unit and furnace heating rate. The basic vacuum in UHV cham-

ber is kept at the level of 7 × 10−9 mbar. The heating system provides a

direct control of the specimen temperature in the temperature range from

RT to 1273 K with heating rate from 1 to 10 K/min. The time for the spec-

imen installation before measurements and pumping the vacuum level of

2 × 10−8 mbar in the UHV chamber did not exceed 15 min. Specimens

for TDS measurements were cleaned before measurements with acetone

in US-bath for 1 min and dried in the He-gas flow to remove any water

residuals from the specimen surface.

Typical size of TDS specimens is 0.9× 4.0× 14 mm. TDS measurements

for the steels after mechanical testing were done at the specimens which

were cut from the gauge section of the samples after CERTs and CLTs.

The time between hydrogen charging and TDS measurement was chosen

to be one hour.

2.4 Mechanical tests

Mechanical tests were performed for high-strength carbon steels. Flat

specimens of sub-sized shape for tensile testing were cut from metal sheets

by electro-discharge machining.

Specimens were cut transverse to the rolling direction in the form of

plate of 250 x 15 mm and gauge section of 5 x 35 mm (Fig. 2.6). The gauge

28

Experimental

Figure 2.6. High-strength carbon steel specimen for tensile testing.

section of the tensile specimens was mechanically polished with emery

paper, finishing with 6 μm diamond paste. Series of smooth and notched

specimens were used. Notches were done at the center of gauge section,

the distance between notches is about 3.7 mm, the radius of the notch is

about 0.3 mm.

All the tensile tests were performed using a 35 kN MTS desktop ma-

chine equipped with a thermostatic environmental cell (see Fig. 2.7). Elec-

trochemical hydrogen charging was carried out under the constant poten-

tial of -1.2 V vs. Hg/Hg2SO4 reference electrode at the room temperature.

The value of -1.2 V was chosen as optimal for high-strength carbon steels

in 1N H2SO4 charging solution, when the specimen is under cathodic po-

tential and the processes of corrosion and dissolution of the metal are

minimal. The reference specimens were tested in air.

Constant extension rate tests (CERT) and constant load tensile tests

(CLT) were performed after hydrogen pre-charging for 1 hour. The speci-

men was loaded with a constant strain rate of 10−4 s−1 to fracture (CERT)

or to the certain load in CLT. The load was kept constant during CLT test

until the specimen fracture or for 100 h if the fracture did not occur.

2.5 Fractography

Scanning electron microscopy (SEM) and electron backscattering diffrac-

tion (EBSD) measurements were done with Zeiss Ultra 55 FEG-SEM

equipped with Nordlys F+ camera and Channel 5 software from Oxford

Instruments. In EBSD measurements misorientation angle of 10◦ was

used for defining the grain boundaries. Chemical composition of non-

metallic inclusions was analyzed by energy dispersive X-ray spectroscopy

(EDS).

Crack path and strain localization along the crack path in high-strength

carbon steels were analysed by EBSD and crystal lattice local misorienta-

tion mapping (CLLM).

29

Experimental

Figure 2.7. MTS desktop machine equipped with the thermostatic environmental cell.

2.6 Optical observations

LaVision system based on the digital image correlation (DIC) technique

was used for measuring of deformations of a sample and imaging it during

the tensile tests. The system consists of high-speed Basler acA2000-340

camera, PIXCI-E8 video capture boards and Epix XCAP software, used

for imaging and DaVis 8.1 software, used for calculations of components

of the strain tensor. This method implies that some visual information is

present on a sample to follow the deformation. The surface pattern in the

form of randomly dispersed marker spots were put on the side surface of

the sample before the tensile test.

Notched specimens of the studied high-strength steels were tested in air

and after electrochemical hydrogen charging from 1N H2SO4 solution for

1 h. The time between electrochemical hydrogen charging and mechanical

testing with optical monitoring, including the spraying and drying of the

markers did not exceed 15 min. The tests were synchronized with optical

monitoring of the marked specimen surface area between notches. The

average distance between marker spots is about 0.1 mm.

30

3. Studies of hydrogen diffusion,trapping and solubility in austeniticstainless steels

In this chapter models for TDS simulation which allow to obtain the dif-

fusion and solubility parameters of the material and calculations for the

different grades of austenitic stainless steels are presented.

3.1 Model and simulations of thermal desorption of hydrogen

TDS measurements together with mathematical modeling of the processes

of hydrogen uptake and desorption allow to distinguish different hydro-

gen trapping sites in the material and to find their binding energies.

Simple diffusion model

First, the diffusion model, where trapping is not taken into considera-

tion was applied to simulate the TDS curves. Assuming that the hydrogen

concentration at the specimen surface C0 is constant and solving the cor-

responding boundary problem for diffusion equation one can obtain an ex-

pression for the hydrogen concentration profile (Eq. 2.2). Supposing that

TDS peak is provided by the mobile hydrogen in the lattice and taking into

account the hydrogen concentration profile across the specimen thickness

(Fig. 2.2) as an initial value of the corresponding boundary problem it was

found that in the case of linear heating the hydrogen desorption flux from

the thin specimen plate as a function of temperature is [Publication I]:

J(Ti) =8C0sD(Ti)

h

∞∑n=0

(1− e−

π2(2n+1)2D(Ts)tsh2

)· e

−π2(2n+1)2

h2

ti∫

0

dτD(Ti), (3.1)

where C0 is the concentration of hydrogen on the surface of the specimen,

s is the specimen surface area, ts and Ts are time and temperature of

hydrogen charging, td and Td are time and temperature of hydrogen des-

orption, respectively, h is the thickness of specimen, and D is the diffusion

coefficient.

31

Studies of hydrogen diffusion, trapping and solubility in austenitic stainless steels

General model of TDS

TDS curves calculated with the simple diffusion model do not reproduce

the details of experimental peaks. The general model developed to simu-

late hydrogen absorption in the course of hydrogen charging and hydrogen

desorption during TDS measurement takes into consideration the diffu-

sion of hydrogen into the metal lattice and its trapping and detrapping

by trapping sites with different binding energies and concentrations. The

model is described by Eq. (1.4) - (1.8) [52]. The values of p0 and k0 were

estimated according to the work [52].

3.2 Hydrogen diffusion, trapping and solubility in different gradesof austenitic stainless steels

The behavior of hydrogen in metals and alloys are complicated and can-

not be described by a simple model. The unhomogeneous profile of hydro-

gen concentration in specimen after charging provides the diffusion peak

of complicated shape. For the pure metal model calculations with one

type of trapping site give satisfactory correspondence to the experimental

data, but for multi-component alloys several types of trapping sites have

to be taken into consideration to simulate hydrogen uptake and desorp-

tion [Publication III].

In [Publication I] diffusion parameters of three stainless steels, namely

AISI 310, AISI 301LN and AISI 210 were obtained by the proposed ana-

lytical model based on Eq. (3.1). The analytical solution of the problem al-

lows the application of a fitting procedure to obtain diffusion parameters.

However, the result of the fitting does not match well with the experimen-

tal data. It evidences clearly that the simplified model of the hydrogen dif-

fusion in the case of austenitic stainless steel, which is a multi-component

alloy with a spectrum of different trapping sites for hydrogen, does not

describe reliably the hydrogen desorption flux. Typical experimental and

simulated curves for AISI 301LN and AISI 201 are presented in Fig. 3.1.

In [Publication II] the general model of TDS was applied. Average hy-

drogen concentration, CH , diffusion coefficient, D, and concentration of

hydrogen at the specimen surface during the hydrogen charging process

C0, which is considered as a measure of hydrogen solubility, are shown for

the studied materials in as-supplied state in Table 3. It is found from the

thermal desorption spectra of hydrogen for AISI 304, AISI 301LN, AISI

201A and AISI 204Cu steels that the lowest activation energy of hydrogen

32


Figure 3.1. Typical fitting result for TDS spectrum of AISI 301LN steel (a) and AISI 201steel (b) with Eq. (3.1). Heating rate is 6 K/min.

Table 3.1. Calculated diffusion parameters of the studied austenitic stainles steels.

Steel CH , Diffusion coefficient, D, m2/s C0,

wt.ppm at.%

AISI 301LN 36 3.1× 107 × exp(−0.54 eV/kBT ) 24

AISI 201A 57 3.2× 107 × exp(−0.507 eV/kBT ) 19

AISI 204Cu 39 5.7× 107 × exp(−0.573 eV/kBT ) 59

AISI 304 40 4.4× 107 × exp(−0.555 eV/kBT ) [69] 44

diffusion equal to 0.507 eV is in AISI 201A steel and the highest activation

energy of 0.573 eV is obtained for AISI 204Cu steel. Hydrogen diffusion in

metastable Cr-Ni austenitic stainless steel AISI 301LN and in Cr-Mn-Ni

austenitic AISI 201 steel is faster than that in AISI 304 and AISI 204Cu

steels due to lower diffusion activation energy.

The apparent hydrogen solubility in metastable AISI 301LN and Mn-

alloyed AISI 201 austenitic stainless steels is less than that in the refer-

ence AISI 304 austenitic stainless steel, as well as in Mn- and Cu-alloyed

AISI 204Cu steel.

The microstructure of the steels was analysed by optical and SEM mi-

croscopy. α′-martensite phase, formed under deformation, was analysed

by EBSD mapping. The increasing of α′-martensite increases the hydro-

gen diffusivity in the deformed material. Similar observations for AISI

301LN were done in [67].

In [Publication II] the model, proposed in [Publication I] was extended

to the case of complex material, consisting of austenite and martensite

phases with effective diffusion coefficient to analyse the pre-strained states

of the materials. Martensite phase was modeled as spherical inclusions in

austenitic matrix. Using the diffusion coefficient obtained by fitting the

33


Figure 3.2. Hydrogen diffusivity for austenite and martensite phases and effective diffu-sivity for AISI 204Cu steel (a) and AISI 301LN steel (b).

TDS curves for pre-strained state as the effective diffusion coefficient of

complex material and the diffusion coefficient obtained for non-deformed

state (Table 3.1), where material is free of martensite as the diffusion co-

efficient of austenite matrix, the diffusivity in martensitic phase can be

calculated using the equation [68]:

2(Deff )2 + (D2 − 2D1 − 3Φ(D2 −D1))Deff −D1D2 = 0, (3.2)

where D1 and D2 - diffusion coefficients of austenite and martensite phases,

Deff - effective diffusion coefficient of complex material, Φ - the volume

fraction of matrix, 0 � Φ � 1. Eq. (3.2) corresponds to the continuum me-

dia consisting of small particles of martensite phase, which can provide

paths for fast diffusion of hydrogen.

Hydrogen diffusivity in the deformed austenitic stainless steels is higher,

than in non-deformed ones, what is in accordance with the experimen-

tal data for Fe-Cr-Ni alloys [69] and can be attributed to the presence of

α′-martensite phase with high diffusivity of hydrogen. The result of cal-

culations of hydrogen diffusivity in martensite phase using Eq. (3.2) are

shown on Fig. 3.2. Activation energies of hydrogen diffusion in marten-

site phase were estimated to be 0.547 and 0.540 eV and pre-exponential

factors are 6.12 ×10−7 m2/s and 6.48×10−7 m2/s for AISI 204Cu and AISI

301LN steels, respectively.

Specific features of hydrogen uptake and desorption of multi-component

alloy were studied in [Publication III]. Austenitic stainless steel 316L and

pure nickel (99,999%, GoodFellow) were chosen for TDS peak simulations.

Experimental thermal desorption curves for pure nickel and for austenitic

stainless steel AISI 316L manifest two peaks. The main peak is situated

at the temperature from 420 to 450 K for both materials and the second

34


Figure 3.3. Simulated and experimental thermal desorption spectra of hydrogen releasefor AISI 316L stainless steel (a) and for pure nickel (b).

Figure 3.4. Hydrogen concentration profile of H-charged sample during TDS measure-ment (a) and dependences of the calculated thermal desorption spectra with-out trapping on the thickness (h) of a specimen for pure nickel (b).

one is at the temperatures from 550 to 650 K. To analyze the processes

of hydrogen diffusion and trapping numerical simulations were made by

the simple diffusion and general models of TDS. The experimental values

of the diffusion parameters used for simulations for pure nickel are D0 =

6.7 · 10−7m2s−1 and Ed = 0.4 eV [70] and for AISI 316L steel D0 = 6.2 ·10−7m2s−1 and Ed = 0.557 eV [71].

For AISI 316L steel the simple diffusion model does not give correspon-

dence to the experimental data (Fig. 3.3 a), while for pure nickel the dif-

fusion peak fits well with the experimental data by both of the applied

models (Fig. 3.3 b). The second peak on the desorption curve cannot be

described by the simple diffusion model, because the trapping processes

during hydrogen charging and desorption have to be taken into consid-

eration. The general model of TDS, based on Eq. (1.4) - Eq. (1.8), as

mentioned above, was used to consider diffusion and trapping processes

during hydrogen charging and desorption. The initial distribution of hy-

drogen was obtained by simulating of the process of hydrogen charging

35


by the above mentioned model and realized using MatLab software. Then

the process of hydrogen desorption was simulated, using the same model.

Hydrogen concentration profile of hydrogen-charged sample of pure nickel

during TDS measurement is presented in Fig. 3.4 a. At the beginning of

the measurement hydrogen is concentrated near the surfaces of the sam-

ple, during the heating hydrogen re-distributes and releases. It was found

that only one type of trapping sites for pure nickel is enough to reproduce

the shape of the experimental curve. In this case, however, the homoge-

neous distribution of trapping sites inside the specimen does not give sat-

isfactory correspondence between the experimental and calculated curves.

It was assumed, that the traps are concentrated near the surfaces of the

specimen, because near the surfaces of a specimen there is more defects

due to mechanical treatment and hydrogen charging. The best correspon-

dence of the experimental and model curves was found when the trap con-

centration is described by the exponential decay when the concentration

of traps near the surfaces was 0.22 at.% and at the center of the specimen

it is close to zero (see Fig. 3.3 b). Binding energy of the trapping sites for

nickel was found to be 0.22 eV, which is close to the energy of hydrogen

dislocations interaction [72]. The dislocations are, probably, introduced in

the surface layer during the electrochemical charging of the specimen.

For AISI 316L steel two types of trapping sites with binding energies of

0.16 eV and 0.1 eV, respectively, were considered. The values of the bind-

ing energies of hydrogen atom with trapping site were obtained from ab

initio calculations [Publication IV], considering the most favorable posi-

tions of hydrogen atom in the fcc crystal lattice of AISI 316L steel. Wide

peak at 650 K for AISI 316L steel is mainly provided by the hydrogen re-

lease from the trapping sites with the binding energy of 0.16 eV, which

corresponds to the binding energy of hydrogen in vacancy in stainless

steels [72]. Concentration of one type of trapping sites, corresponding to

vacancy-type defects was found to be about 0.07 at.% and concentration

of the second type of trapping sites corresponding to interstitial positions

of hydrogen about 0.001 at.%, respectively [Publication IV].

During the analysis of the experimental and numerically simulated data

it became obvious that the differences in the profile of hydrogen concen-

tration in the specimen after charging lead to the differences in the TDS

peak shape, even if it is provided by the mobile diffusive hydrogen in the

metal lattice without traps (Fig. 3.4 b). The shape of the main diffusion

peak depends markedly on the thickness of the specimen, as it was ob-

36


served by the experimental TDS curves for both materials.

The dependences of the calculated thermal desorption spectra on the

thickness of the specimen for pure nickel without trapping of hydrogen

are shown in Fig. 3.4 b. The diffusion peak manifests a complex shape. It

can be explained by the specific hydrogen release when the distribution

profile of hydrogen in the specimen is inhomogeneous and caused by re-

distribution of hydrogen in the metal lattice and between trapping sites

during the TDS measurement.

37


38

4. Studies of hydrogen effects onmechanical properties and fractureof high-strength carbon steels

4.1 Constant extension rate tests (CERT)

Constant extension rate tests (CERT) were performed for all studied high-

strength carbon steels in as-supplied and hydrogen-charged states to ob-

tain mechanical properties and hydrogen sensitivity factors of the steels.

Typical CERT stress-strain curves for some of the studied steels with and

without hydrogen are shown in Fig. 4.1. CERTs show that elongation to

fracture decreases markedly with hydrogen charging for all the studied

materials. Mechanical properties, hydrogen content CH and hydrogen

sensitivity factors δH = L0−LHL0

, where L0 and LH are elongations to frac-

ture for non-charged and hydrogen-charged materials, respectively, are

shown in Table 4.1 for S1 steels and in Table 4.2 for S2 and Table 4.3 for

S3 steels.

The values of hydrogen sensitivity factor δH are smaller for tempered

steels comparing with martensitic ones of the same chemical composition

(see Table 4.1, Table 4.2 and Table 4.3), thus, the tempered steels are less

sensitive to hydrogen effects in CERT than the martensitic counterparts.

The value Rp0.2/Rm is higher for the steels which exhibit yield point, for

the tempered martensitic steels it is higher than for untempered ones (see

Table 4.1, Table 4.2 and Table 4.3).

4.2 Constant load tests (CLT)

One of the important parameters obtained by CLT under continuous hy-

drogen charging is time to fracture. CLTs show a remarkable hydrogen

effect on all the studied steel grades. Specimens tested in air did not

fracture during 100 h under the load corresponding to the yield stress.

39

Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels

Table 4.1. Mechanical properties, sensitivity to hydrogen δH , time to fracture tf in CLTsat applied stress of 0.3×Rm and hydrogen concentration CH of S1 steels. [73]

Steel S1_1000 S1_1000 S1_1200 S1_1200 S1_1400 S1_1400

grade M02 TM02 M04 TM09 M TM

Rp0.2, MPa 493 755 602 1230 1223 1416

Rm, MPa 1050 988 1185 1250 1572 1416

Rp0.2/Rm 0.47 0.76 0.51 0.98 0.78 1.0

δH 0.768 0.647 0.74 0.47 0.783 0.524

tf , min 1367 4299 367 311 1 25

CH , wt.ppm 0.032 0.193 0.361 0.155 0.041 0.041

as-supplied

CH , wt.ppm 2.09 2.92 1.95 2.03 2.36 1.95

H-charged


Steel S2_1000 S2_1200 S2_1200 S2_1400 S2_1400

grade TM05 MTM M05 MTM TM

Rp0.2, MPa 916 976 752 914 1330

Rm, MPa 994 1213 1298 1349 1423

Rp0.2/Rm 0.92 0.81 0.58 0.68 0.94

δH 0.525 0.659 0.633 0.795 0.739

tf , min 608 312 105 17 255

CH , wt.ppm 0.174 0.046 0.075 0.042 0.122

as-supplied

CH , wt.ppm 1.21 0.90 0.78 1.90 0.87

H-charged

40



Steel S3_1000 S3_1200 S3_1400

grade TM TM M

Rp0.2, MPa 1019 1202 1199

Rm, MPa 1051 1252 1402

Rp0.2/Rm 0.97 0.96 0.85

δH 0.48 0.43 0.574

tf , min >3000 634 157

CH , wt.ppm 0.247 0.154 0.328

as-supplied

CH , wt.ppm 1.89 3.14 2.98

H-charged

Hydrogen-charged specimens manifest fracture after loading in less than

100 h, even though the applied stress is less than the yield stress. Times

to fracture tf for all the studied steels under continuous hydrogen charg-

ing at applied stress of 0.3Rm are presented in Fig. 4.2 for smooth (a) and

notched (b) series of specimens and in Tables 4.1, 4.2 and 4.3 for smooth

specimens. All the steels with higher strength level are more sensitive to

hydrogen than the steels of the same chemical composition, but with lower

strength level. Steel S3 exhibits the longest time to fracture, comparing

with the other steels of the same strength level.

Tempered steels exhibit longer times to fracture than their untempered

counterparts (see Fig. 4.2, Tables 4.1, 4.2 and 4.3). It also confirms the

fact that tempered steels are less sensitive to hydrogen.

Dependencies of time to fracture on the applied stress for S1_1200_M04,

S3_1200_TM, S1_1000_M02 and S1_1000_TM02 are shown in Fig. 4.3 and

Fig. 4.4 a. It is seen that different processes dominate at low stresses and

at high stresses. Probably after some critical time, the local corrosion at

the particles and hydrogen-induced decohesion of the particles dominate,

initiating the fracture.

For the steels S1_1000_M02 and S1_1000_TM02 there is no dependence

of time to fracture on the applied stress when time to fracture is longer

than 1000 min. Variations in time to fracture may be caused also by the

presence or absence of large NMIs in the gage length of different samples.

For the steels with the small amount of NMIs (S2_1000_TM05 and

41


Figure 4.1. Stress-strain curves of CERT for martensite (a) and tempered martensite (b)steels with the strength level of 1200 MPa.

Figure 4.2. Times to fracture for smooth (a) and notched (b) specimens at applied stressof 0.3 of the ultimate tensile strength. Specimens, which exhibit fracturebefore the applied stress of 0.3 of the ultimate tensile strength are markedas CERT.

S2_1000_M05) the dependence of time to fracture on the applied stress

has a linear character in logarithmic coordinates (see Fig. 4.4 b).

4.3 Fractography and non-metallic inclusions

Fracture surfaces of smooth and notched samples after CERT and CLT

under continuous hydrogen charging were analysed by SEM imaging.

A lot of TiC/TiN partcles are observed on the fracture surfaces of S1

steels. Analysis of the fracture surfaces of these steels shows, that the

characteristic feature of the fracture formed in testing under continuous

hydrogen charging is hydrogen flakes or "fish eyes" (Fig. 4.5 a). The hy-

drogen flakes form transgranularly around large, about 5 to 20 μm tita-

nium carbide/nitride particles. Some particles itself are broken during

testing. Transgranular hydrogen-induced cracking was observed in the

zones around TiC/TiN particles (Fig. 4.5 b).

Steel S2 contains mainly Al2O3, CaO and MnS particles. Hydrogen-

42


Figure 4.3. Time to fracture in CLT under continuous hydrogen charging at differentapplied stresses for the steels S1_1200_M04 (a) and S3_1200_TM (b).

Figure 4.4. Time to fracture in CLT under continuous hydrogen charging at differentapplied stresses for the steels S1_1000 (a) and S2_1000 (b) martensitic andtempered martensitic steels.

induced brittle area is localized in the middle of the CLT sample (Fig. 4.6

a). Hydrogen-induced cracking has intergranular character (Fig. 4.6 b).

In steel S3 hydrogen embrittlement is a mixture of transgranular and

intergranular fracture modes. A lot of secondary cracks were observed on

the fracture surface (Fig. 4.7 a). SiO2 and TiC/TiN non-metallic inclusions

are seen on the fracture surface. Hydrogen-induced cracking is initiated

at small titanium carbide/nitride particles, about 1 to 2 μm, situated in

the middle of the small brittle transgranular areas (Fig. 4.7 b).

Small TiN, Ti(CN), TiC precipitates were analysed in [73] using the

carbon-replica technique in order to study the distribution and size of pre-

cipitates in the material. In case of the steel S3 grades, the main type of

observed precipitates are TiN precipitates in the size range of 70 nm up to

several micrometers. Very small Ti(CN) precipitates are observed in the

size range of 3-15 nm. In case of the steels S1 mainly TiC particles are

finely distributed throughout the microstructure. Typical TEM images of

precipitate distribution are presented in Fig. 4.8.

43


Figure 4.5. Micrographs of fracture surface after CLT test with continuous hydrogencharging of S1_1200_TM09 steel.

Figure 4.6. Micrographs of fracture surface after CLT tests with continuous hydrogencharging of S2_1200_M04 steel.

The size, shape and the amount of NMIs were analysed by SEM and

EDS observing the side surfaces of the tensile specimens. The size of

NMIs for all materials varies between 100 nm and 20 μm. The amount of

NMIs in S1_TM steel is 102/mm2 of the size 1 to 2 μm and 80/mm2 of the

size 3 to 10 μm. For the steel S1_M they are 43/mm2 and 68/mm2, for the

steel S2 65/mm2 and 11/mm2 and for the steel S3 157/mm2 and 72/mm2,

respectively. In S3 steel about 36000/mm2 small particles of the size 100

to 300 nm were observed. In S1_TM steel the amount of the same size

particles is about 15000/mm2, but in S2 steel only few particles of the size

100 to 300 nm were observed. NMIs distribution by size in the studied

steel grades is presented in Fig. 4.9.

In S1 and S3 steels large particles are mainly TiC/TiN particles of rect-

angular shape, which are usually broken in the mechanical testing and

in these steels cracks initiate on Ti-based particles (Fig. 4.10 a) in the

specimens which are continuously hydrogen charged during CERT and

CLT. In the specimens tested in air no cracks, but only voids around Ti-

based particles were found near the particles (Fig. 4.10 b). Non-metallic

particles themselves are usually cracked during the tensile testing under

hydrogen charging (Fig. 4.11 a, b). No hydrogen-induced crack initiation

on non-metallic particles was found in steel S2.

44


Figure 4.7. Micrographs of fracture surface after CLT with continuous hydrogen charg-ing of S3_1200_TM steel.

Figure 4.8. Transmission electron micrographs - a carbon-replica of the steel S1_1400_M(a) and of the steel S3_1400_M [73].

4.4 Hydrogen - NMI interactions and crack initiation

The mechanism of the hydrogen-induced crack formation at NMIs has to

take into consideration the elastic and plastic strain distribution around

the particle under applied external load [74]. The elastic stress forming

around the NMIs under the applied tensile loading is shown in Fig. 4.12

[42]. NMIs itself are under high tensile stress while the tensile external

load is applied to the material.

When the external load is applied to the material, hydrogen probably

accumulates in the area of the hydrostatic dilatation close to the central

plane of the particle, which is normal to the external load direction. The

hydrogen atmosphere forming close to the central plane may enhance the

local vacancy concentration due to the high hydrogen-vacancy binding en-

ergy in iron [Publication VII]. Vacancies are forming voids where hydro-

gen is accumulated, initiating hydrogen-induced cracks (see Fig. 4.10 and

4.11). The scheme of hydrogen atoms and vacancies accumulation near

NMI is presented in Fig. 4.13.

45


Figure 4.9. Distribution of non-metallic inclusions of 1 to 40 μm by size.

Figure 4.10. Micrographs of the side surface after CLT under continuous hydrogencharging (a) and in air (b) for the steel S1_1200_TM09.

4.5 Thermal desorption analysis of hydrogen in high-strengthcarbon steels

TDS measurements were carried out for as-supplied and hydrogen-charged

specimens. While the solubility of hydrogen in carbon steels is low, most

of the observed hydrogen is concentrated in trapping sites, which can be

various lattice defects, such as vacancies, dislocations, voids, grain bound-

aries and phase interfaces. Small precipitates can also influence on the

diffusion properties of the steel.

Any industrial metallic material in as-supplied state contains some amo-

unt of hydrogen collected in different trapping sites. TDS measurements

of the studied steels in as-supplied state allow preliminary evaluation of

their ability to hydrogen absorption and define the temperature locations

of peaks of the hydrogen evolution, which correspond to the trapping sites

46


Figure 4.11. Micrographs of the side surface after CLT under continuous hydrogencharging for the steels S1_1200_TM09 (a) and S1_1400_TM (b).

Figure 4.12. Elastic stress forming around hard NMI in the steel under uniaxial externalstress σ0 [42].

of different origin.

Hydrogen charging of the steels results in a remarkable increase of

the total hydrogen content, which becomes about one order of magnitude

higher than that in as-supplied state. It varies between 2 and 3 wt.ppm

for S1 and S3 steels, while it is only about 0.9 wt.ppm in S2 steels (see

Table 4.1, Table 4.2 and Table 4.3). Thermal desorption curves for the

studied steels manifest at least two peaks. The main peak is situated

at the temperature from 420 to 460 K and small second peak is at the

temperatures from 580 to 650 K.

The assumption that NMIs act as the main trapping sites is based on

the analysis of TDS peaks of different materials. S1 and S3 steels contain

more particles than S2 steels (see Fig. 4.9) and TDS peak amplitudes for

S1 and S3 steels are much higher than those for S2 steels (Fig. 4.14 a). For

the steels of the same chemical composition TDS peaks are quite similar

for tempered and untempered materials. The lowest hydrogen content

was found in S2 steel which contains the lowest amount of NMIs.

47


Figure 4.13. Scheme of hydrogen-induced crack initiation at NMI.

Figure 4.14. TDS curves after hydrogen charging for S1, S2 and S3 steels (a) and thedependence of hydrogen content of the surface fraction and volume frac-tion (top axis) of non-metallic particles (b). Points correspond to S2, S1_M,S1_TM and S3 steels.

Particle surface fraction in the volume unit and the volume fraction of

the particles were calculated for the studied steels taking into consider-

ation particle sizes and quantity on the side surfaces of the tensile spec-

imens. The analysis of the dependence of hydrogen content on the par-

ticle surface fraction and volume fraction (Fig. 4.14 b) allows to assume,

that hydrogen accumulates preferably at the NMI interfaces, because the

amount of trapped hydrogen is proportional to the NMI surface fraction

in volume, as it is seen in Fig. 4.14 b.

4.6 Strain localization

Strain localization is one of the important factors in fracture mechanisms

and it can be analyzed at different levels beginning from atomic and mi-

crostructural levels to macroscopic scale.

4.6.1 EBSD mapping

EBSD inverse pole figure (IPF) and crystal lattice local misorientation

(CLLT) maps were obtained for all studied high-strength carbon steels at

the areas of the side surfaces of tensile specimens after CLT under con-

tinuous hydrogen charging where microcracks were observed. EBSD IPF

48


Figure 4.15. EBSD IPF (a) and CLLM (b) maps of the side surface crack in S1_1000_M02steel after CLT under continuous hydrogen charging.

Figure 4.16. EBSD IPF (a) and CLLM (b) maps of the side surface crack inS1_1000_TM02 steel after CLT under continuous hydrogen charging.

maps indicate transgranular character of hydrogen-induced crack propa-

gation for the most of the studied high-strength steels. The maps for the

steels S1_1000_M02 and S1_1000_TM02 with clear transgranular charac-

ter of hydrogen-induced cracking is presented in Fig. 4.15 a and Fig. 4.16

a. In S2 steels cracks propagate mainly intergranularly, for example, steel

S2_1400_M exhibits intergranular hydrogen-induced cracking (Fig. 4.17

a).

CLLT maps show strain localization along crack path for all studied

steels (see Fig. 4.16 b and Fig. 4.17 b), except of the steels S1_1000_M02

(Fig. 4.15 b), S1_1400_M, S1_1400_TM and S2_1200_M05, where no strain

localization along crack path were observed. For example, for the steel

S1_1000_M02 there is no strain localization along the crack path, but for

its tempered counterpart S1_1000_TM02 one can see clear strain local-

ization along the crack path and at the crack tip (compare Fig. 4.15 b and

Fig 4.16 b).

49


Figure 4.17. EBSD IPF (a) and CLLM (b) maps of the side surface crack in S2_1400_Msteel after CLT under continuous hydrogen charging.

4.6.2 Optical observations using LaVision system

Deformations in the notched samples of the studied high-strength carbon

steels were studied by LaVision system, using digital image correlation

technique. Specimens were tensile tested in air in as-supplied state and

after electrochemical hydrogen charging from 1N H2SO4 solution.

About 2500 snapshots of the movies made for each specimen before frac-

ture were treated with DaVis software in order to calculate components

of the 2D strain tensor as functions of testing time or nominal strain. Hy-

drogen effect on macroscopic strain localization in notched specimens was

analysed by the comparison the value of 2D strain tensor component Eyy

(y axis is parallel to the applied load direction) just before crack initiation

for the studied high-strength steels [73]. The snapshots of the movies for

S2_1400_TM and S3_1000_TM steels for hydrogen-free (a) and hydrogen-

charged specimens (b) are shown in Fig. 4.18 and 4.19. The snapshots of

the movies, next to the presented ones correspond to complete fracture

of the specimen into two parts. Comparison of the snapshots allows to

conclude that hydrogen reduces markedly the nominal strain and stress

values (see the snapshots at Fig. 4.18 and 4.19), which correspond to the

crack initiation.

Most of the studied steels (except for S1_1000_TM02 steel) manifest a

reduction of the local strain in the presence of hydrogen as compared to

that in hydrogen-free specimen (see Fig. 4.18 and Fig. 4.19) and exhibit

a brittle fracture without visible macroscopic strain localization. It indi-

cates that the local plastic strain preceding the crack initiation is localized

in the presence of hydrogen at the scale less than the average distance be-

tween marker spots. It is seen also that in the presence of hydrogen stain

localization may occur only at one notch as, for example, in S3_1000_TM

50


Figure 4.18. Comparison of the macroscopic distributions of Eyy strain tensor compo-nent in hydrogen-free (a) and hydrogen-charged (b) notched specimens ofS2_1400_TM steel. Nominal strain and stress, which correspond to thespecimen fracture are shown below snapshots.

Figure 4.19. Comparison of the macroscopic distributions of Eyy strain tensor compo-nent in hydrogen-free (a) and hydrogen-charged (b) notched specimens ofS3_1000_TM steel. Nominal strain and stress, which correspond to thespecimen fracture are shown below snapshots.

steel (Fig. 4.19 b), that means, that the first microcrack immediately leads

to the brittle fracture of the specimen.

51


52

5. Conclusions

In the study the model for the interpretation of thermal desorption peaks

was developed. The behavior of hydrogen in metals and alloys is compli-

cated and cannot be described by a simple model. The processes of hy-

drogen trapping, re-trapping and diffusion in multi-component alloys are

defined by the specific atomic distribution of hydrogen over the trapping

sites and that has to be taken into consideration.

The unhomogeneous profile of hydrogen concentration in a specimen of

austenitic stainless steel after cathodic charging provides the diffusion

peak of complicated shape. For the pure metal, model calculations with

one type of trapping site give satisfactory correspondence to the experi-

mental data, while the alloys with the larger number of different trap-

ping sites represent complicated processes during the hydrogen release

and more parameters have to be taken into consideration to simulate it.

The detailed analysis of experimental TDS curves and simulation of the

processes by the proposed model allow to find the parameters of hydrogen

diffusion and trapping for different materials.

The activation energies of hydrogen diffusion for AISI 304, AISI 301LN,

AISI 201 and AISI 204Cu steels are found from the thermal desorption

spectra, applying mathematical modeling. The lowest activation energy of

hydrogen diffusion equal to 0.507 eV is exhibited by Cr-Mn-Ni austenitic

AISI 201A steel and the highest activation energy of 0.573 eV is obtained

for AISI 204Cu steel.

The apparent hydrogen solubility in metastable AISI 301LN and Mn-

alloyed AISI 201 austenitic stainless steels is less than that in the ref-

erence AISI 304 austenitic stainless steel, as well as in Cu-alloyed AISI

204Cu steel.

The formation of α′-martensite increases the hydrogen diffusivity in the

deformed austenitic stainless steels. Activation energies of hydrogen dif-

53

Conclusions

fusion in martensite phase were estimated to be 0.547 and 0.54 eV and

pre-exponential factors are 6.12 ×10−7 m2/s and 6.48×10−7 m2/s for AISI

204Cu and AISI 301LN steels, respectively.

For all the studied high-strength carbon steels elongation to fracture

and time to fracture are reduced markedly after hydrogen charging. All

the steels with higher strength level are more sensitive to hydrogen than

the steels of the same chemical composition, but with lower strength level.

Tempered martensite steels are less sensitive to hydrogen than their mar-

tensitic counterparts.

Non-metallic inclusions play the major role in crack initiation of hydro-

gen-charged high-strength steels. In the steels with high Ti-alloying hydro-

gen-induced cracks are initiated mainly at TiC/TiN particles. A character-

istic feature of the fracture surface after continuous hydrogen charging is

hydrogen flakes which form transgranularly around titanium-based par-

ticles, which effectively trap hydrogen and enhance the sensitivity to hy-

drogen embrittlement. The steels with low Ti-alloying, where Nb-alloying

is present show a remarkable reduction of the NMI size and density, ex-

hibit intergranular fracture mode and higher sensitivity to hydrogen em-

brittlement. The steel of medium Ti-alloying with high density of fine

TiC/TiN inclusions shows a mixed trans- and intergranular character of

hydrogen-induced fracture and is more resistant to hydrogen embrittle-

ment.

Microstructural and TDS analyses allow to assume, that the main trap-

ping sites for hydrogen in the studied high-strength carbon steels are NMI

interfaces.

54

Bibliography

[1] W. H. Johnson: On Some Remarkable Change Produced in Iron and

Steel by the Action of Hydrogen and Acids, Proc. R. Soc. London, Vol.

23 (1875), 168-179.

[2] R.P. Gangloff: Hydrogen Assisted Cracking of High Strength Al-

loys, in Comprehensive Structural Integrity. Elsevier Science, Vol.

6 (2003), New York, NY, 31-101.

[3] R.P. Gangloff: Critical Issues in Hydrogen Assisted Cracking of

Structural Alloys, in S.A. Shipilov, R.H. Jones, J.-M. Olive and R.B.

Rebak. (Eds.): Environmental Induced Cracking of Metals, Elsevier

Science, Vol. 1 (2005) Oxford, UK, 141-165.

[4] R.A. Oriani, J.P. Hirth, M. Smialowski: Hydrogen Degradation of

Ferrous Alloys, William Andrew Publishing, Noyes (1985).

[5] S. Lynch: Mechanisms of Hydrogen Assisted Cracking - a Review, in

International Conference on Hydrogen Effects on Materials Behavior

and Corrosion Deformation Interactions, Moran, WY, USA (2002),

449-466.

[6] M. Nagumo: Conformity between Mechanics and Microscopic Func-

tions of Hydrogen Failure, ISIJ International, Vol. 52 (2012), 168-

173.

[7] M. Nagumo: Hydrogen Related Failure of Steels - a New Aspect,

Mater. Sci. and Tech., Vol. 20 (2004), 940-950.

[8] J.P. Hirth: Effects of Hydrogen on the Properties of Iron and Steel,

Met. Trans., Vol. 11A (1980), 861-890.

55

Bibliography

[9] G.R. Caskey, Jr: Hydrogen Effects in Stainless Steels, in R. A. Oriani,

J. P. Hirth, M. Smialowski (Eds.): Hydrogen Degradation of Ferrous

Alloys, Noyes Publ. (1985), 822-862.

[10] T. Doshida, H. Suzuki, K. Takai, N. Oshima and T. Hirade: Enhanced

Lattice Defect Formation Associated with Hydrogen and Hydrogen

Embrittlement under Elastic Stress of a Tempered Martensitic Steel,

ISIJ Int., Vol. 52 (2012), 198-207.

[11] S. Takagi, Y. Toji, M. Yoshino and K. Hasegawa: Hydrogen Embrit-

tlement Resistance Evaluation of Ultra High Strength Steel Sheets

for Automobiles, ISIJ Int., Vol. 52 (2012), 316-322.

[12] S. Takagi, Y. Toji, S. Takagi, K. Hasegawa, Y. Tanaka, N. Roessler, B.

Hammer and T. Heller: Hydrogen Embrittlement Evaluation Meth-

ods for Ultra-high Strength Steel Sheets for Automobiles, Int. J. Au-

tom. Eng., Vol. 2 (2010), 7-13.

[13] M. Nagumo, T. Tamaoki and T. Sugawara: Functions of Microstruc-

tures in Delayed Fracture of Martensitic Steels, in Hydrogen Ef-

fects on Material Behaviour and Corrosion Deformation Interactions,

TMS (2003), 999-1008.

[14] C. Beachem: New Model for Hydrogen-assisted Cracking (Hydrogen

"Embrittlement"), Metall. Trans., Vol. 3 (2) (1972), 441-455.

[15] H.K. Birnbaum, P. Sofronis: Hydrogen-enhanced Localized Plasticity

- a Mechanism for Hydrogen-related Fracture, Mater. Sci. and Eng.

A, 176 (1994), 191-202.

[16] P. Sofronis, Y. Lianga and N.Aravas: Hydrogen Induced Shear Lo-

calization of the Plastic Flow in Metals and Alloys, Eur. J. Mech. A,

Solids, Vol. 20 (2001), 857-872.

[17] A. Troiano: The Role of Hydrogen and Other Interstitials in the Me-

chanical Behavior of Metals, Trans. ASM, Vol. 52 (1960), 54-80.

[18] J. Morlet, H. Johnson and A. Troiano: A New Concept of Hydrogen

Embrittlement in Steel, Journal of the Iron and Steel Institute, Vol.

189 (1958), 37-44.

[19] R. A. Oriani: The Diffusion and Trapping of Hydrogen in Steel, Acta

Metall., Vol. 18 (1970), 147-157.

56

Bibliography

[20] H.A. Wriedt, R.A. Oriani: Effect of Tensile and Compressive Elastic

Stress on Equilibrium Hydrogen Solubility in a Solid, Acta Metall.,

Vol. 18 (1970), 753-760.

[21] R. Oriani, P. Josephic: Equilibrium Aspects of Hydrogen-induced

Cracking of Steels, Acta Metall., Vol. 22 (1974), 1065-1074.

[22] C.J. McMahon, Jr.: Hydrogen-induced Intergranular Fracture of

Steels, Eng. Fract. Mech., Vol. 68 (2001), 773-788.

[23] S. Serebrinsky, E. A. Carter and M. Ortiz: A Quantum-mechanically

Informed Continuum Model of Hydrogen Embrittlement, J. Mech.

Phys. Solids, Vol. 52 (2004), 2403-2430.

[24] I.M. Robertson, D. Teter: Controlled Environment Transmission

Electron Microscopy, J. Microsc. Res. Tech., Vol. 42 (1998), 260-269.

[25] D.F. Teter, I.M. Robertson and H.K. Birnbaum: The Effects of Hydro-

gen on the Deformation and Fracture of β-titanium, Acta Mater., Vol

49 (2001), 671-692.

[26] H.J.Grabke, E. Riecke: Absorption and Diffusion of Hydrogen in

Steels, Mater. Tehnol., Vol. 34, (2000), 331-342.

[27] A.H.M. Krom, A.D. Bakker: Hydrogen Trapping Models in Steel,

Met. and Mater. Transactions B, Vol. 31B (2000), 1475-1482.

[28] S.-M. Lee, J.-Y. Lee: The Trapping and Transport Phenomena of Hy-

drogen in Nickel, Metallurgical Transactions A, Vol. 17A (1986), 181-

187.

[29] M. Nagumo, K. Takai and N. Okuda: Nature of Hydrogen Trapping

Sites in Steels Induced by Plastic Deformation, J. of Alloys and Com-

pounds, Vol. 293-295 (1999), 310-316.

[30] J. Chêne, in: C.G. Interrante and G.M. Pressouyre (Eds.): Current

Solutions to Hydrogen Problems in Steels, American Society for Met-

als, (1981) 263–271.

[31] E. Herms, J.M. Olive and M. Puiggali: Hydrogen Embrittlement of

316L Type Stainless Steel, Mater. Sci. Eng., Vol. A272 (1999), 279-

283.

[32] A.R. Marder, A.O. Benscoter: Microcracking in Fe-C Acicular

Martensite, Trans ASME, Vol. 61 (1968), 293-299.

57

Bibliography

[33] C.S. Robert: Effect of Carbon on the Volume Fractions and Lat-

tice Parameters of Retained Austenite and Martensite, Trans. TMS

AIME, Vol. 197 (1953), 203-204.

[34] Y. Sakamoto, T. Mantani: Effect of Quenching and Tempering on

Diffusion of Hydrogen in Carbon Steel, Trans. Jpn. Inst. Met., Vol.

17 (1976), 743-748.

[35] S.L.I. Chan, H.L. Lee and J.R. Yang: Effect of Retained Austenite on

the Hydrogen Content and Effective Diffusivity of Martensitic Struc-

ture, Met. Trans. A., Vol. 22A (1990), 2579-2586.

[36] P. Juvonen: Effects of Non-metallic Inclusions on Fatigue Properties

of Calcium Treated Steels, Dissertation for the degree of Doctor of

Science in Technology, Otamedia OY, Espoo, Finland (2004), 102 p.

[37] R. Kiessling: Non-metallic Inclusions in Steel. The Institute of Ma-

terials, London (1997), 400 p.

[38] H.V. Atkinson, G. Shi: Characterization of Inclusions in Clean Steels:

a Review Including the Statistics of Extremes Methods, Progress in

Mater. Sci., Vol. 48 (2003), 457-520.

[39] Y. Murakami: Metal Fatigue: Effects of Small Defects and Nonmetal-

lic Inclusions. Elsevier Science Ltd., Oxford (2002), 369 p.

[40] J.D. Eshelby: Determination of the Elastic Field of an Ellipsoidal

Inclusion, and Related Problems, Proc. R. Soc. London, A, Vol. 241,

No. 1226 (1957), 376-396.

[41] J.D. Eshelby: The Elastic Field Outside an Ellipsoidal Inclusion,

Proc. R. Soc. London, A, Vol. 252, No. 1271 (1959), 561-569.

[42] Z.Y. Ou, G.F. Wang, T.J. Wang: An Analytical Solution for the Elas-

tic Fields near Spheroidal Nano-inclusions, Acta Mech. Sin., Vol. 25

(2009), 821-830.

[43] C.E. Sims: Transactions, American Institute of Mining, Metallurgi-

cal and Petroleum Engineers, Vol. 215 (1959), pp.367 - 393

[44] H.G. Lee, J.-Y. Lee: Hydrogen Trapping by TiC Particles in Iron, Acta

Metall., Vol. 32, No.1 (1984), 131-136.

[45] F.G. Wei, K. Tsuzaki: Quantative Analysis on Hydrogen Trapping

of TiC Particles in Steel, Met. and Mater. Trans. A, Vol. 37A (2006),

331-353.

58

Bibliography

[46] S.M. Lee, J.Y. Lee: The Effect of the Interface Character of TiC Par-

ticles on Hydrogen Trapping in Steel, Acta Metall., Vol. 35, No.11

(1987), 2695-2700.

[47] S. Fujita, Y. Murakami: A New Nonmetallic Inclusion Rating Method

by Positive Use of Hydrogen Embrittlement Phenomenon, Metal.

and Mater. Trans. A, Vol. 44A, No.1 (2012), 303-322.

[48] S.W. Smith, J.R. Scully: The Identification of Hydrogen Trapping

States in an Al-Li-Cu-Zn Alloy Using Thermal Desorption Spec-

troscopy, Metall. Mater. Trans., Vol. 31A, No.1 (2000), 179-193.

[49] A. Turnbull, R.B. Hutchings and D.H. Ferriss: Modelling of Thermal

Desorption of Hydrogen from Metals, Mater. Sci. Eng. A, Vol. A238

(1997), 317-328.

[50] M.S. Daw, M.I. Baskes: Embedded-atom Method: Derivation and Ap-

plication to Impurities, Surfaces, and Other Defects in Metals, Phys.

Rev. B, Vol. 29 (1984), 6443-6453.

[51] J.E. Angelo, N.R. Moody and M.I. Baskes: Trapping of Hydrogen to

Lattice Defects in Nickel, Modelling Simul. Mater. Sci. Eng., Vol. 3

(1995), 289-307.

[52] K. Ebihara, H. Kaburaki: Numerical Modeling of Thermal Deso-

prtion Spectra of Hydrogen: a Review of Thermal Desorption Models,

ISIJ Int., Vol. 52 (2012), 181-186.

[53] T. Yamaguchi, M. Nagumo: Simulation of Hydrogen Thermal Des-

orption under Reversible Trapping by Lattice Defects, ISIJ Interna-

tional, Vol. 43, No.4 (2003), 514-519.

[54] W. Y. Choo, I.-Y. Lee: Thermal Analysis of Trapped Hydrogen in Pure

Iron, Metall. Trans. A, Vol. 13A (1982), 135-140.

[55] S.-M. Lee, I.-Y. Lee: The Trapping and Transport Phenomena of Hy-

drogen in Nickel, Metall. Trans. A, Vol. 17A (1986), 181-187.

[56] K. Ono, M. Meshii: Hydrogen Detrapping from Grain Boundaries

and Dislocations in High Purity Iron, Acta Metall. Mater., Vol. 40

(1992), 1357-1364.

[57] K. Ebihara, T. Suzudo, H. Kaburaki, K. Takai and S. Takebayashi:

Modeling of Hydrogen Thermal Desorption Profile of Pure Iron and

Eutectoid Steel, ISIJ Int., Vol. 47 (2007), 1131-1140.

59

Bibliography

[58] A. Turnbull, R.B. Hutchings and D.H. Ferriss: Modelling of Thermal

Desorption of Hydrogen from Metals, Mater. Sci. and Eng., Vol. A238

(1997), 317-328.

[59] K. Ebihara, H. Kaburaki, T. Suzudo and K. Takai: A Numerical

Study on the Validity of the Local Equilibrium Hypothesis in Model-

ing Hydrogen Thermal Desorption Spectra, ISIJ Int., Vol. 49 (2009),

1907-1913.

[60] M.S. Daw, M.I. Baskes: Semiempirical, Quantum Mechanical Calcu-

lation of Hydrogen Embrittlement in Metals, Phys. Rev. Lett., Vol. 50

(1983), 1285-1288.

[61] D. Li, R. P. Gangloff, J. R. Scully: Hydrogen Trap States in Ultrahigh-

Strength AERMET 100 Steel, Metall. and Mater. Trans. A, Vol. 35A

(2004), 849-864.

[62] D. Pérez Escobar, T. Depover, L. Duprez, K. Verbeken and M.

Verhaege: Combined Thermal Desorption Spectroscopy, Differen-

tial Scanning Calorimetry, Scanning Electron Microscopy and X-ray

Diffraction Study of Hydrogen Trapping in Cold Deformed TRIP

Steel, Acta Mater., Vol. 60 (2012), 2593-2605.

[63] T. Tanabe, Y. Yamanishi, K. Saweda and Sh. Imoto: Hydrogen Trans-

port in Stainless Steels, J. Nucl. Mater., Vol. 122-123 (1984), 1568-

1572.

[64] S. Papula, J. Talonen, O. Todoshchenko and H. Hänninen: Effect

of Internal Hydrogen on Delayed Fracture of Metastable Low-Ni

Austenitic Stainless Steels, Metall. and Mater. Trans. A, Vol. 45,

No.11 (2014), 5270-5279.

[65] J.-P. Bacher, C. Benvenuti, P. Chiggiato, M.-P. Reinert, S. Sgobba and

A.-M. Brass: Thermal Desorption Study of Selected Austenitic Stain-

less Steels, J. Vac. Sci. Technol., Vol. A 21 (2003), 167-174.

[66] W. Beck, J.O’M. Bockris, L. Nanis and J. McBreen: Hydrogen Perme-

ation in Metals as a Function of Stress, Temperature and Dissolved

Hydrogen, Proc. R. Soc., Vol. A290 (1966), 220-235.

[67] J. Talonen, P. Nenonen, G. Pape and H. Hänninen: Effect of Strain

Rate on the Strain-Induced γ → α′- Martensite Transformation and

Mechanical Properties of Austenitic Stainless Steels. Metall. and

Mater. Trans. A, Vol. 36, No. 2 (2005), 421-431.

60

Bibliography

[68] I.N Beckman: Phenomenological Description of Diffusion in the

Medias with Defects. Hydrogen Interactions with Metals, Nauka,

Moscow (1987), 143-177.

[69] D.J. Fisher: Hydrogen Diffusion in Metals. A 30-Year Retrospective,

Scitec Publications, Switzerland (1999), 59-126.

[70] W. Eichenauer, W. Löser und H. Witte: Löslichkeit und Diffusion-

sgeschwindigkeit von Wasserstoff und Deuterium in Einkristallen

aus Nickel und Kupfer, Z. Metalk., Vol. 56 (1965), 287-293.

[71] A.M. Brass, J. Chêne: Hydrogen Uptake in 316L Stainless Steel:

Consequences on the Tensile Properties, Corros. Sci., Vol. 48 (2006),

3222-3242.

[72] S.M. Myers, W.R. Wampler and F. Besenbacher: Trapping and Sur-

face Recombination of Ion-implanted Deuterium in Stainless Steel,

J. Appl. Phys., Vol. 56 (1984), 1561-1571.

[73] T. Hickel: Hydrogen Sensitivity of Different Advanced High-

Strength Microstructures. EU RFCS Publication RFSR-CT-2010-

00020. HYDRAMICROS Project Report (2014).

[74] T. P. Perng, C.J. Altstetter: Effects of Deformation on Hydrogen Per-

meation in Austenitic Stainless Steels, Acta Metall., Vol. 34, No. 9

(1986), 1771-1781.

61

Bibliography

62

Errata

Publication III

In Eq. (2), page 73,∫ t10 dτD(Ts) have to be

∫ ti0 dτD(Ti)

63

9HSTFMG*agcafb+

ISBN 978-952-60-6205-1 (printed) ISBN 978-952-60-6204-4 (pdf) ISSN-L 1799-4934 ISSN 1799-4934 (printed) ISSN 1799-4942 (pdf) Aalto University School of Engineering Department of Engineering Design and Production www.aalto.fi

BUSINESS + ECONOMY ART + DESIGN + ARCHITECTURE SCIENCE + TECHNOLOGY CROSSOVER DOCTORAL DISSERTATIONS

Aalto-D

D 6

8/2

015

Hydrogen embrittlement resistance is an important parameter for the development of new steel grades for different applications. Hydrogen embrittlement is a phenomenon that causes loss of ductility in a metal, making it brittle and enhancing cracking and failure. The thesis describes studies of hydrogen effects in two classes of steels – austenitic stainless steels and advanced high-strength carbon steels. One of the used methods is thermal desorption spectroscopy, which allows to analyse hydrogen diffusion and trapping in metals, applying mathematical modeling of the processes of hydrogen uptake and desorption. The parameters of hydrogen diffusing and trapping were estimated for the different grades of austenitic stainless steels. The role of non-metallic inclusions in hydrogen embrittlement and hydrogen-induced fracture was studied for high-strength carbon steels with different micro-alloying. Mechanisms of hydrogen interaction with non-metallic inclusions are discussed.

Olga T

odoshchenko H

ydrogen effects on austenitic stainless steels and high-strength carbon steels A

alto U

nive

rsity

Department of Engineering Design and Production

Hydrogen effects on austenitic stainless steels and high-strength carbon steels


DOCTORAL DISSERTATIONS

Documents

Aalto- Hydrogen effects on DD austenitic stainless steels