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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
Olga Madelen Ingrid Todoshchenko
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,
Olga Madelen Ingrid Todoshchenko
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”
The author has prepared and run CERT and CLT tests and made SEM
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”
The author has prepared and run CERT and CLT tests and made SEM
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
Studies of hydrogen diffusion, trapping and solubility in austenitic stainless steels
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
Studies of hydrogen diffusion, trapping and solubility in austenitic stainless steels
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
Studies of hydrogen diffusion, trapping and solubility in austenitic stainless steels
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
Studies of hydrogen diffusion, trapping and solubility in austenitic stainless steels
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
Studies of hydrogen diffusion, trapping and solubility in austenitic stainless steels
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
Studies of hydrogen diffusion, trapping and solubility in austenitic stainless steels
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
Table 4.2. Mechanical properties, sensitivity to hydrogen δH , time to fracture tf in CLTsat applied stress of 0.3×Rm and hydrogen concentration CH of S2 steels. [73]
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
Table 4.3. Mechanical properties, sensitivity to hydrogen δH , time to fracture tf in CLTsat applied stress of 0.3×Rm and hydrogen concentration CH of S3 steels. [73]
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
Studies of hydrogen effects on mechanical properties and fracture of high-strength carbon steels
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
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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
Olga Madelen Ingrid Todoshchenko
DOCTORAL DISSERTATIONS