HYDROGEN EMBRITTLEMENT OF FERROUS MATERIALS · The stepwise repeated slow strain rate test (SW...

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Université Libre de Bruxelles Faculty of Applied Sciences Department of Matter and Materials Materials Science and Electrochemistry Group

Supervisor Jean – Luc Delplancke Co Supervisor Antoine Pourbaix

Academic year 2005 – 2006

HYDROGEN EMBRITTLEMENT OF

FERROUS MATERIALS

Mioara Elvira Stroe

Supervisor Dr. Jean Luc Delplancke

Co – supervisor Antoine Pourbaix

Local Members of the Jury Dr. Luc Segers

Dr. Marie–Paule Delplancke

Dr. Marc Degrez

Invited Members of the Jury Dr. Patrick Toussaint Industeel Belgium, Group Arcelor

Dr. Mihai Popa Institute of Physical Chemistry “I.G. Murgulescu”, Romania

Belgium, January 2006

To Luiza, Cris, Anastasis and Anouka

Preface

This thesis is submitted to the Faculty of Applied Sciences at the Université Libre

de Bruxelles, Belgium, in order to fulfil the requirements for obtaining the Ph.D. degree in

Applied Sciences.

This study was financed by Industeel Belgium, in the frame of a broader research

program between Industeel Belgium and CEBELCOR

Acknowledgements

I am deeply indebt to Antoine Pourbaix, co – supervisor of this thesis, for guiding

me during all these years. His clarity of spirit and enthusiasm continuously amazed me and

inspired me. I thank him for constantly encouraging me and for the extensive discussions that

helped me to find the way.

By his support I could visit some laboratories abroad (NPL in UK, Statoil Norway,

ECP France) and to participate to several conferences.

I would like to thank Jean Luc Delplancke, the head of Department of Materials

Science and Electrochemistry for hosting me in his Department and for all support he granted

me in the past four years.

I like to acknowledge Industeel Belgium, and particularly Patrick Toussaint and

Jean Jacques Dufrane for financial and scientific support.

I am very glad that I found in Renée Scherer, Jacques Kissel and Suzanne de

Kegel not only very kind colleagues but also friends. My work and my stay here was more

pleasant due to their presence.

I enjoyed collaborating with Jean Dille, Roger D’Haens, Lionel Canet, Olivier van

de Vyver and Victor Wertz from the Department of Materials Science and Electrochemistry.

I would like to thank to Catherine Dagbert from Ecole Centrale Paris. Part of

fractional thermal degassing tests presented in Section 7.2. were performed with her

assistance.

I am pleased to mention the good collaboration with Vincent Ligier to whom I thank

for helping me with measurements by fractional thermal degassing at CRMC Industeel Creusot

and also with many useful advices.

I thank to Liane Smith (Intetech Ltd. UK) and Stein Olsen (Statoil Norway) for the

insights on the industrial aspects of the problem of hydrogen embrittlement.

I like to thank to my sister and my parents for encouraging me all this time. The

holydays and their phone calls cheered me up whenever was needed.

My gratitude goes to my friend Tarik Bouali whose determined character and keen

sense of detail were an example for me.

ABSTRACT

This work deals with the damage due to the simultaneous presence of hydrogen in

atomic form and stress – straining.

The aim of this work is twofold: to better understand the hydrogen embrittlement

mechanisms and to translate the acquired knowledge into a more appropriate qualification test.

The phenomena of hydrogen entry and transport inside the metals, together with

the different types of damages due to the presence of hydrogen, are presented.

The analysis of the most important models proposed up to now for hydrogen

embrittlement (HE) indicated that the slow dynamic plastic straining is a key factor for the

embritteling process. There is a synergistic effect of hydrogen – dislocations interactions: on

one hand hydrogen facilitates the dislocations movement (according to the HELP mechanism)

and on the other hand dislocations transport hydrogen during their movement when their

velocity is lower than a critical value.

This work is focused on supermartensitic stainless steels, base and welded

materials. The interest on these materials is due to their broad use in offshore oil production.

First, the material’s characterisation with regards to hydrogen content and

localisation was performed. This was conducted in charging conditions that are representative

of industrial applications.

Because of previous industrial experience it was necessary to find a more

appropriate qualification test method to asses the risk of HE.

In this work we proposed the stepwise repeated slow strain rate test (SW R –

SSRT) as a qualification test method for supermartensitic stainless steels.

This test method combines hydrogen charging, test duration, plastic, dynamic and

slow strains. Thus, this test method is coherent with both the model HELP proposed for

hydrogen embrittlement and the observations of industrial failures.

The stepwise repeated slow strain rate test (SW RSSRT) is interesting not only as

a qualification test of martensitic stainless steels, but also as a qualification test of conditions for

using these materials (type of straining, range of strain and stress, strain rate, hydrogen

charging conditions, etc.).

RESUME

Ce travail se rapporte à l’endommagement provoqué par la présence simultanée

de l’hydrogène sous forme atomique et une contrainte (appliquée où résiduelle).

Ce travail a comme but une meilleure compréhension du mécanisme de la

fragilisation par l’hydrogène (FPH) et la recherche d’un essai de qualification qui soit cohérent

avec ce mécanisme.

Les phénomènes liés à l’entrée et au transport de l’hydrogène au sein des métaux,

ensemble avec les différents types d’endommagements dus à la présence de l’hydrogène, sont

présentés.

L’analyse des modèles proposés jusqu’au présent pour la fragilisation par

l’hydrogène (FPH) suggère que la déformation lente plastique dynamique est le facteur clé pour

le processus de la fragilisation. Il y a un effet synergétique des interactions entre l’hydrogène et

les dislocations: d’un coté l’hydrogène facilite le mouvement des dislocations (d’après le modèle

HELP) et d’un autre coté les dislocations transportent l’hydrogène pendant leur mouvement,

pourvu que leur vitesse soit en dessous d’une valeur critique.

Le travail a été conduit sur des aciers supermartensitiques, matériau de base et

soudé. L’intérêt pour ces matériaux réside de leur large utilisation dans la production du pétrole

en offshore.

D’abord, le matériau a été caractérisé du point de vu de la teneur et de la

localisation de l’hydrogène. Les essais ont été conduits dans des conditions représentatives

pour les cas industriels.

L’expérience industrielle d’auparavant indique qu’il est nécessaire de trouver un

test de qualification plus approprié pour estimer la susceptibilité à la fragilisation par

l’hydrogène.

Dans ce travail on propose un essai de traction lente incrémentée (SW R – SSRT)

comme méthode de qualification pour les aciers supermartensitiques.

L’essai combine le chargement en hydrogène, la durée d’essai, la déformation

lente, plastique et dynamique. Donc, cette méthode d’essai est cohérente avec le modèle HELP

proposé pour FPH et les observations des accidents industriels.

Cet essai est intéressant pas seulement comme essai de qualification pour les

aciers supermartensitiques, mais aussi comme essai de qualification pour les conditions

d’utilisation des ces matériaux (type de déformation, niveau de déformation et contrainte,

vitesse de déformation, conditions de chargement en hydrogène, etc.).

TABLE OF CONTENT NOTATION 1. INTRODUCTION …………………………………………………………..

1.1. ORIGIN AND AIMS OF THE WORK 1.2. STRUCTURE OF THE WORK

2. HYDROGEN ENTRY AND TRANSPORT IN METALS ………………

2.1. HYDROGEN ADSORPTION 2.1.1. ELECTROCHEMICAL CHARGING

2.1.2. CHARGING FROM GASEOUS HYDROGEN ATMOSPHERE

2.1.3. THE INFLUENCE OF SURFACE STATE ON ADSORPTION

2.2. HYDROGEN ABSORPTION 2.2.1. ATOMIC HYDROGEN TRANSFER (CLASSICAL

MECHANISM)

2.2.2. DIRECT TRANSFER OF H+ THROUGH THE INTERFACE

(MODEL PROPOSED BY CROLET ET AL.)

2.3. HYDROGEN TRANSPORT IN MATERIAL 2.3.1. DIFFUSION

2.3.1.1. IDEAL DIFFUSION, FICK’S LAW

2.3.1.2. TEMPERATURE EFFECTS

2.3.1.3. DIFFUSION UNDER STRESS GRADIENT

2.3.2. TRAPPING

2.3.3. TRANSPORT OF HYDROGEN BY MOVING DISLOCATIONS

2.4. SUMMARY OF HYDROGEN IN METALS

3. TYPES OF DAMAGES DUE TO HYDROGEN ………………………..

3.1. HYDROGEN INDUCED CRACKING (HIC) AND STEPWISE CRACKING (SWC)

1.

3.5.

7.

10.10.

15.

18.

18.

19.

22.

23.23.

23.

25.

26.

26.

28.

30.

33.

36.

3.2. STRESS ORIENTED HYDROGEN INDUCED CRACKING (SOHIC) 3.3. HYDROGEN REACTION WITH THE METAL MATRIX (HYDRIDE

FORMATION) 3.4. HYDROGEN REACTIONS WITH NON METALLIC PHASES 3.5. HYDROGEN EMBRITTLEMENT (HE ) OR HYDROGEN STRESS

CRACKING (HSC) 3.6. SULPHIDE STRESS CRACKING (SSC)

4. HYDROGEN EMBRITTLEMENT ………………………………………..

4.1. DECOHESION MECHANISM 4.1.1. THERMODYNAMIC ASPECTS OF INTERFACIAL

DECOHESION

4.1.2. ELECTRONIC DISTRIBUTION IN H – METAL SYSTEMS

4.1.3. EXPERIMENTAL OBSERVATIONS

4.1.4. CONCLUSIONS FOR THE DECOHESION MECHANISM

4.2. PLASTICITY MODEL: HYDROGEN ENHANCED LOCALISED

PLASTICITY “HELP” 4.2.1. HYDROGEN SHIELDING EFFECT

4.2.1.1. HYDROGEN EFFECT ON THE INTERACTIONS

BETWEEN DISLOCATIONS

4.2.1.2. INTERACTIONS BETWEEN DISLOCATIONS AND

AN IMPURITY ATOM IN THE PRESENCE OF

HYDROGEN

4.2.2. MICROSCOPIC OBSERVATIONS

4.2.3. MACROSCOPIC OBSERVATIONS

4.2.4. CONCLUSIONS FOR THE HELP MECHANISM

4.3. CONCLUSIONS ON THE PARAMETERS AND ON THE MODELS PROPOSED FOR HYDROGEN EMBRITTLEMENT

5. INDUSTRIAL ASPECTS AND NEED FOR A HE TEST METHOD ...

5.1. HISTORIC OF SUPERMARTENSITIC STAINLESS STEEL USE IN OFF SHORE

37. 38.

39. 39.

40.

41.

44.

44.

48.

48.

49.

50.

51.

52.

55.

55.

60.

61.

61.

65.

67.

5.2. FAILURES DUE TO HYDROGEN EMBRITTLEMENT 5.3. NEED FOR A HE TEST METHOD

6. MATERIALS AND EXPERIMENTAL TECHNIQUES …………………

6.1. MATERIAL DESCRIPTION 6.1.1. OVERVIEW OF GENERAL MECHANICAL PROPERTIES

6.1.2. CHEMICAL COMPOSITION AND MICROSTRUCTURE

6.1.3. WELD DESCRIPTION

6.2. EXPERIMENTAL METHODS 6.2.1. PERMEATION METHOD

6.2.2. THERMAL DEGASSING

6.2.2.1. FRACTIONAL THERMAL DEGASSING

6.2.2.2. TOTAL DEGASSING

6.2.3. NANOINDENTATION METHOD

6.2.4. MECHANICAL TESTS

6.2.4.1. CONSTANT LOAD TEST

6.2.4.2. SLOW STRAIN RATE TEST (SSRT)

6.2.4.3. REPEATED SLOW STRAIN RATE TEST (RSSRT)

6.2.4.4. STEPWISE REPEATED SLOW STRAIN RATE TEST

(SW RSSRT)

7. EXPERIMENTAL RESULTS …………………………………………….

7.1. PERMEATION TESTS RESULTS 7.2. RESULTS FOR THERMAL DEGASSING TESTS

7.2.1. FRACTIONAL THERMAL DEGASSING

7.2.2. TOTAL DEGASSING

7.2.3. CONCLUSIONS OF THERMAL DEGASSING TESTS

7.3. RESULTS OF NANOINDENTATION TESTS 7.4. RESULTS OF MECHANICAL TESTS

7.4.1. RESULTS FOR CONSTANT LOAD TESTS

69.71.

73.

75.75.

77.

78.

81.81.

88.

88.

89.

91.

94.

94.

96.

99.

100.

103.

105.

118.118.

121.

121.

122.

127.127.

7.4.2. RESULTS OF SLOW STRAIN RATE TESTS

7.4.3. REPEATED SLOW TRAIN RATE TESTS RESULTS

7.4.4. STEPWISE REPEATED SLOW STRAIN RATE TESTS

RESULTS

8. DISCUSSION ………………………………………………………………

8.1. HYDROGEN DISLOCATIONS INTERACTIONS 8.2. THE EMBRITTLING PHENOMENON: FROM HELP TO

EMBRITTLEMENT 8.3. BEHAVIOUR OF VARIOUS FERROUS MATERIALS IN THE

PRESENCE OF HYDROGEN 8.4. PROPOSAL FOR SPECIFIC HE TESTS FOR MARTENSITIC

MATERIALS 8.5. RESULTS OF HE TESTS USED IN THIS WORK 8.6. SW RSSRT AS A QUALIFICATION TEST FOR MARTENSITIC STEELS

AND FOR OPERATING CONDITIONS

9. CONCLUSIONS …………………………………………………………... 10. REFERENCES ……………………………………………………………. 11. APPENDICES ……………………………………………………………..

APPENDIX 1. HYDROGEN EVOLUTION REACTIONS …………………………. APPENDIX 2. THERMODYNAMIC AND KINETIC ASPECTS OF THE

INTERFACIAL DECOHESION …………………………………….. APPENDIX 3. ATOM SUPERPOSITION AND ELECTRON DELOCALISATION

MOLECULAR ORBITAL METHOD (ASED – MO) ……………… APPENDIX 4. HYDROGEN SHIELDING EFFECT ………………………………..

128.

148.

161.

171.

173. 175.

175.

177.

178.

181.

183.

189.

201.

203.

213.

227. 229.

NOTATION

Roman letters

A

A.R.

AFM

b

C0

CE

CL

D

Dok

E

Ea

Er

F

2Fk*

G

GTAW

h

Hnano

[H]

[H+]

HAZ

HE

HELP

HIC

HID

i

i0

J

Jabs

Jads

Jdes

Jdsb

Jss

Surface

Area reduction

Atomic force microscopy

Burgers vector

Subsurface concentration

Counter electrode

Constant load test

Diffusion coefficient

Preexponential term for kink diffusion

Potential

Activation energy

Reduced elastic modulus

Faraday constant

Free energy of formation of a double – kink on a dislocation

Shear modulus

Gas tungsten arc welding

Depth

Nanohardness

Concentration of hydrogen atoms on the metallic surface

Concentration of hydrogen cations in solution

Heat affected zone

Hydrogen embrittlement

Hydrogen enhanced localised plasticity

Hydrogen induced cracking

Hydrogen induced decohesion

Current density

Exchange current density

Flux

Flux of absorbed hydrogen atoms

Flux of adsorbed hydrogen atoms

Flux of desorbed hydrogen atoms

Flux of hydrogen atoms diffusing out from the bulk of material

Steady state flux

k

ki

l

l

L

m

ni

Ni

No

Nr

Ns

P

Pnano

PGMAW

Qk

r

R

RE

R SSRT

S

Snano

SCE

SCC

SEM

SOHIC

SSC

SSRT

SW M SSRT

SW R SSRT

t

T

TEM

TTF

UTS

vcr

V

VH

Boltzmann’s constant

Rate constant for reaction i

Dislocation length

Strain increment

Membrane thickness

Mass

Stoichiometric coefficient of species i

Density of irreversible traps

Number of interstitial sites occupied by hydrogen

Density of reversible traps

Number of interstitial sites in the matrix

Hydrogen partial pressure

Indentation load

Pulse gas metal arc welding

Activation energy for kink diffusion

Distance in polar coordinates

Gas constant

Reference electrode

Repeated slow strain rate test

Sieverts constant

Stiffness

Saturated calomel electrode

Stress corrosion cracking

Scanning electron microscopy

Stress oriented hydrogen induced cracking

Sulphide stress cracking

Slow strain rate test

Stepwise monotonic slow strain rate test

Stepwise repeated slow strain rate test

Time

Temperature

Transmission electron microscopy

Time to failure

Ultimate tensile strength

Critical velocity of dislocation

Mean molar volume

Molar volume of hydrogen

Greek letters

x

W

WB

WE

YS

YS0.2%

Distance

Activation energy

Bonding energy between hydrogen and trap

Working electrode

Yield strength

Offset yield strength

αi Transfer coefficient of reaction i β Proportionality constant γ Proportionality constant δ Separation distance between two layers of the interface δc Critical separation distance

∆l Strain to failure ε Strain η Overpotential θ Coverage degree θ0 Equilibrium coverage degree

θL Occupancy of the interstitial sites

θr Occupancy of the reversible sites

µ Bulk modulus ν Poisson’s coefficient σ Stress σh Hydrostatic stress τ Shear stress ϕ (χ, δ) Cohesive function

φ Angle in polar coordinates

1. INTRODUCTION

1. INTRODUCTION

2

1. INTRODUCTION

3

1.1. ORIGIN AND AIMS OF THE WORK Several cracking accidents occurred recently in offshore exploitations that were

attributed to hydrogen embrittlement (HE).

The structures involved are flowlines and equipments installed on the bottom of the

sea.

These flowlines are protected against corrosion by cathodic protection with

aluminium – indium (AlIn) or aluminium – zinc – indium (AlZnIn) sacrificial anodes and by heavy

duty coatings.

The flowlines are subjected to elastic and plastic deformations during laying and

during operation: reeling and dereeling, laying, movements on the sea bed due to marine

streams, thermal expansion, pressure tests, or accidental interference with fishing activities.

The materials of the pipelines that experienced failure are martensitic and

supermartensitic stainless steels (SMSS) and duplex stainless steels (DSS).

Of course, these materials are known to be sensitive to hydrogen embrittlement.

However, extensive design analysis and qualification testing [1] indicate that they are

appropriate for this application. In particular extensive long-term testing showed no cracking

under constant load and under constant plastic deformation (four point bend tests) at – 1050

mV to – 1200 mV versus saturated calomel electrode (SCE).

As these materials are much more cost effective than carbon steel, their use has

markedly increased in the recent past [2] and is planned to increase further in the future.

These cracking accidents came as a surprise to many operators, because the

materials passed all the qualification tests and because, currently, there are about 2000 km of

such SMSS and DSS pipelines in operation and only few accidents occurred.

This work has two main purposes:

a better understanding of the hydrogen embrittlement mechanisms

involved in the failures,

1. INTRODUCTION

4

in case the HE tests used appear non appropriate, find better qualification

tests that better reflect the mechanisms and parameters involved in HE.

This study is focused on martensitic stainless steels, due notably to a marked trend

for a broader use of this material.

In principle two approaches can be considered:

a detailed fundamental analysis of the HE phenomena, to see the factors

that may cause failures,

a comprehensive analysis of the accidents.

Since a comprehensive analysis of the industrial accidents is not currently

available, this study was conducted mostly on the basis of an analysis of the fundamental

aspects of hydrogen embrittlement.

Extensive work was conducted and several models were proposed for hydrogen

embrittlement. But, none of the models proposed up to now can explain all the observed

phenomena or the role of all factors.

For example, the hydrogen content was often considered as the main factor (as in

the decohesion and the adsorption models, see section 4). The stress level is also considered

in most studies. But the loading mode, the local H accumulation and the local stress

concentration in the bulk of the material were not always given due consideration.

It was thus felt necessary to revisit the existing models with the aim of identifying

and analysing the influencing parameters.

Hydrogen embrittlement is a loss of mechanical properties due to the presence of

hydrogen in atomic form and stress. A significant decrease of ductility and /or fracture strength,

delayed fracture and absence of metal loss are typical features of HE.

In fact, this work showed that the fundamental approach combined with the real life

test parameters and with the analysis of the actual failures proved to be interestingly coherent.

1. INTRODUCTION

5

1.2. STRUCTURE OF THE WORK

After the presentation of the origin and the aims of this work (Section 1.1.), the

most important aspects of the hydrogen entry and behaviour inside the materials are presented

in Section 2.

In Section 3. a summary of different damages due to hydrogen is presented. This

chapter aims to identify the characteristics of Hydrogen Embrittlement (HE) in comparison with

other types of hydrogen damages. HE is the consequence of simultaneous presence of

hydrogen in atomic form and stress.

Several models were proposed up to now for hydrogen embrittlement. An analysis of

the most important is presented in Section 4. The factors involved in these models and that

must be considered for testing materials, were identified.

The industrial context at the origin of this work and the need for specific tests are

described in Section 5.

This study is focused on supermartensitic stainless steels. The description of the

materials and the experimental methods selected and used to characterise the materials for HE

is given in Section 6.

The results are presented in Section 7.

The discussion (Section 8) addresses the test methods that are coherent with the

mechanisms for hydrogen embrittlement and analyses the results, with a comparison with the

real conditions of failures.

Section 9 (Conclusions) summarises the important factors for hydrogen embrittlement

that are derived from the HELP model, how these factors can be included in a specific test

procedure for HE, how the results of a test proposed here are coherent with the models and

with the industrial experience.

1. INTRODUCTION

6

2. HYDROGEN ENTRY AND

TRANSPORT IN METALS

2. HYDROGEN IN METALS

8


9

In this section the different steps involved in hydrogen entry and transport through

the material are presented.

The first step is hydrogen adsorption on the material surface. As source of hydrogen, electrochemical evolution of hydrogen by cathodic

polarisation, corrosion reaction or gaseous hydrogen atmosphere can be mentioned.

In this work two cases are analysed: cathodic protection in aqueous solution

(Section 2.1.1.) and adsorption from gaseous hydrogen atmosphere (Section 2.1.2.).

The influence of metallic surface state on the hydrogen adsorption is analysed in

Section 2.1.3.

The adsorbed atoms could then undergo absorption, passing through the metallic interface. This step leads to accumulation of a subsurface hydrogen concentration, C0. The C0

dependency on applied potential or current (for electrochemical charging) and pressure (for

charging from gaseous hydrogen atmosphere) is presented in Section 2.2.

The next process is the transport of hydrogen through the material. The different aspects, like diffusion, trapping and the hydrogen transport by moving dislocations are

presented in Section 2.3.


10

2.1. HYDROGEN ADSORPTION

The adsorbed hydrogen atoms can be formed by different paths, according to the

source of hydrogen.

By cathodic polarisation or corrosion reaction, electroadsorbed species are formed

on the metallic surface whereas in the presence of a gaseous atmosphere the molecular

hydrogen can undergo physisorption or chemisorption.

2.1.1. ELECTROCHEMICAL CHARGING

In solution under cathodic polarisation, the hydrated hydrogen cations, H3O+, are

transported by diffusion / migration towards the cathode. There, the cation undergoes reduction

and becomes atomic hydrogen, H.. The atomic hydrogen can recombine to form molecular

hydrogen, H2 that leaves the metallic surface.

For the reduction of hydrogen ions two different reaction mechanisms are possible,

depending on the nature of the metal:

A. Volmer – Tafel mechanism (electrochemical reduction followed by chemical recombination)

adsorbed

k

khydrated HeH ↔

−

−+ +

1

1

Volmer reaction .)1.2(

2

2

2

HHHk

kadsorbedadsorbed ↔

−

+ Tafel reaction .)2.2(

B. Volmer – Heyrovsky mechanism (electrochemical reduction followed by electrochemical recombination)

adsorbed

k

khydrated HeH ↔

−

−+ +

1

1


2

3

3

HeHHk

kadsorbedhydrated ↔

−

−+ ++ Heyrovsky reaction .)3.2(


11

Depending on the nature of the metal, the different mechanisms of hydrogen

reduction could take place. The paths followed by hydrogen reduction reaction on different

metals are given in Table 2.1.

For iron and steels the most probable mechanisms are coupled reduction –

chemical combination or slow reduction – fast electrochemical as noted in Table 2.1.

Table 2.1. Mechanisms followed by hydrogen reduction on different metals [3]

Metal Mechanism

Fe A : Coupled reduction, recombination

or

B :Slow reduction, fast electrochemical

Ti B: Fast reduction, slow electrochemical

Pd A : Fast reduction, slow recombination

Pt A : Fast reduction, slow recombination

Ni A : Slow reduction, fast recombination

For each of the two mechanisms, the reactions and the dependency of the degree

of coverage, θ, on the current or potential will be described.

θ is necessary to calculate the subsurface concentration C0.

The subsurface concentration, C0, is an important factor as it determines the

hydrogen content in the material (see Section 2.2. and 2.3.). It also determines the filling of

reversible traps, as will be presented in Section 2.3.

A. Volmer – Tafel mechanism The first reaction (Volmer reaction) is the cathodic reduction of hydrated cations to

form atomic hydrogen that remains adsorbed on the metallic surface.

The second reaction is the recombination of atomic hydrogen to form molecular

hydrogen and is a purely chemical reaction.


12

The general expression for the rate of an electrochemical reaction is:

RT

FrevEE

einiC

RTW

kerate

))( −−

Π−

=

α

.)4.2(

Assuming that the Volmer reaction is the rate determining reaction, the expression

for the potential – current density dependency can be written as:

[ ]

−−

+−

= − η

αη

α

RT

FVHkRT

FVHki)1(

expexp 111 .)5.2(

The concentration of atomic hydrogen adsorbed on the surface, [H], is proportional

to the degree of coverage, θ. The reduction of hydrogen cation H+ occurs only on the sites that

are not covered by adsorbed hydrogen atoms. This part is equal to (1 – θ), so the second term

on the right side of equation (2.5.) is directly proportional to (1 – θ).

The current – potential relationship becomes:

−−−−

= − η

αθη

αθ

RT

FVkRT

FVki)1(

exp)1(exp 111 .)6.2(

By introducing the exchange current density for Volmer reaction, i0,V, and the

overvoltage expressions, η, the equation (2.6.) becomes:

−−

−−

−

= ηα

θθηα

θθ

RTF

RTFii VVV

)1(exp11exp

00,01 .)7.2(


13

The equation (2.7.) expresses the dependence of the degree of coverage on the

current density i1 .

The Tafel reaction is a pure chemical reaction, so its rate constant does not depend on the potential. The reaction rate is given by:

[ ] ²)1( 2222 θθ −−−=−= kkdtHdFi .)8.2(

At equilibrium, i.e. when the overall rate is zero, the degree of coverage reaches

the equilibrium value, θ0, given by:

²)1( 022

022,0 θθ −=−= kki .)9.2(

For the Volmer – Tafel mechanism when the rate is determined by the Volmer

reaction, the degree of coverage is (See Appendix 1 ):

2,0

10 1 i

i−= θθ .)10.2(

where i0,2 is the reaction exchange current density for the equilibrium value of the

potential.

B. Volmer – Heyrovsky mechanism (electrochemical reduction followed by electrochemical recombination (Figure 2.2.)

adsorbed

k

khydrated HeH ↔

−

−+ +

1

1


2

3

3

HeHHk

kadsorbedhydrated ↔

−

−+ ++ Heyrovsky reaction .)3.2(


14

The Heyrovsky reaction (2.3.) consists of the reduction of a hydrated hydrogen

cation with a hydrogen atom that is adsorbed on the metallic surface, with the formation of

molecular hydrogen. This equation is also an electrochemical reaction, as well as the Volmer

reaction (2.1.). The rate of the reaction is potential dependent.

The rate of the cathodic partial reaction is proportional to the degree of coverage,

θ, and the hydrogen cations concentration, [H+]. The reversal reaction is proportional to the

molecular hydrogen concentration, [H2], and the free part of the surface (1 – θ). The rate of

reaction (2.3.) is given by:

[ ] [ ]

−−−

−= +− η

αθη

αθ

RT

FHHkRT

FHHki)1(

expexp)1( 3233 .)11.2(

For the Volmer – Heyrovsky mechanism, θ depends on the current density of each

partial reaction involved, as both of them are charge transfer reactions.

Considering the equations (2.5.) and (2.11.) and for large overvoltage, the degree

of coverage becomes independent of the potential. E.g. for high cathodic overvoltage, θ is

expressed as:

H

V

ii

,0

,0

0

0

11

1

θθ

θ

−+

= .)12.2(

where i0,V and i0,H are the exchange current densities for the Volmer and

Heyrovsky reactions respectively.

The hydrogen coverage on the metal substrate depends both on the mechanism of

hydrogen evolution on the metal (A or B) and on the charging conditions (current density or

potential).


15

2.1.2. CHARGING FROM GASEOUS HYDROGEN ATMOSPHERE

The model proposed by Wang in 1936 [4] and then improved by other authors [5 –

10] for the dissociative chemisorption of gaseous hydrogen is presented in Figure 2.1.

The process involves several steps:

a. The gas molecule strikes the material surface and splits into atoms that adhere

there.

adsorbedHH 22 ↔ .)13.2(

The flux of adsorbing atoms is second order in (1 – θ) and could be expressed by:

PkJ adsads )²1( θ−= .)14.2(

b. There is a reversal reaction of recombination of adsorbed atoms with the

formation of molecular hydrogen, which leaves the metallic surface. The desorbing flux is

second order in θ, since two adjacent atoms are needed to recombine:

²θdesdes kJ = .)15.2(

c. The adsorbed atoms could traverse the metallic interface and become absorbed

atoms. The corresponding flux is:

γθ=absJ .)16.2(


16

d. The atoms inside the material could diffuse out to the surface and the flux is

proportional to the subsurface hydrogen concentration, C0, and the fraction of unoccupied

surface sites, through which hydrogen can diffuse out:

0)1( CJdsb θβ −= .)17.2(

The net flux has the following form:

²)²1( θθ desads kPkJ −−= .)18.2(

0)1( CJ θβγθ −−= .)19.2(

At equilibrium, there is no net flux. The degree of coverage, θ0, and subsurface

hydrogen concentration, C0, can be deduced as:

Pkk

des

ads=− 0

0

1 θθ

.)20.2(

PkkCdes

adseq β

γ=,0 .)21.2(

The expression for the concentration is known as the Sieverts’ law:

PSC eq =,0 .)22.2(

where S is Sieverts’ constant that depends on the kinetics of adsorption and

desorption processes:

βγ

des

ads

kkS = .)23.2(


17

This model explains quantitatively how the net flux changes with the external gas

pressure.

The Sieverts’ law shows that the subsurface hydrogen concentration is

proportional to the square root of the hydrogen atmosphere pressure.

In Figure 2.2. an example of this dependency is presented for carbon steel at

different temperatures [11]. On the ordinate the flux units are presented, as the flux is

proportional to the subsurface concentration (see section 6. 1. Permeation test method).

J ads =kads(1–θ )²P

J dsb = β(1–θ )C0

J des = kdesθ²

J abs = γ θ

gas interface metal bulk

Figure 2.1. Fluxes involved in the hydrogen adsorption and absorption from gaseous atmosphere

0

5

10

15

20

25

0 50 100 150 200 250 300

square root of hydrogen pressure (Pa1/2)

Flux

(in

arbi

trar

y un

its)

Figure 2.2. Hydrogen flux – hydrogen pressure dependency at different temperatures for carbon steel [11]

413 K 483 K

533 K

633 K

698 K


18

This type of charging can take place in the same time with electrochemical

charging: (for instance: pipelines at great depths, cathodically protected, with simultaneous

electrochemical charging and charging from gaseous H2 from bubbles at high hydrostatic

pressure).

2.1.3. THE INFLUENCE OF SURFACE STATE ON THE ADSORPTION

Presence of promoters for hydrogen entry An important aspect of hydrogen behaviour is the substantial enhancement of

absorption in the presence of specific compounds. These compounds such as S2-, HS –, H2S,

As etc., hinder the recombination of hydrogen atoms on the metallic surface and therefore

enhance the absorption reaction.

These compounds are usually referred to as poisons.

Even small quantities of poison strongly increase the hydrogen uptake. In Table

2.2. the hydrogen amount penetrating the material for constant charging current, when the

amount of sulphide in solution increases, is presented [12].

Table 2.2. Hydrogen uptake function of sulphide content in solution [12]

[S2–], ppm Hydrogen penetrating the steel (%)

3.5 x 10 – 3 1.6

1.3 x 10 – 2 2.4

2.75 x 10 – 2 5.8

3.6 x 10 – 1 25.6

Oxide films formed on the metallic surface are barriers for H absorption and are

hindering the hydrogen passage through the interface [13 – 17].

2.2. HYDROGEN ABSORPTION

With regard to the passage of hydrogen through the metallic interface, with the

accumulation of a subsurface concentration, C0, two mechanisms were proposed up to now:


19

one mechanism considers that the same species (atomic H) adsorbed on

the surface lead to molecular and absorbed hydrogen (classical mechanism

presented in Section 2.2.1.)

Crolet et al. recently suggested that hydrogen is passing in ionic form (H+)

directly through the metallic interface to form a solid solution (Section 2.2.2.)

2.2.1. ATOMIC H TRANSFER (CLASSICAL MECHANISM)

After the reduction of hydrogen cations, a part of the hydrogen atoms that are

adsorbed on the metallic surface will recombine to form molecular hydrogen that leaves the

surface.

Another part of adsorbed hydrogen atoms will undergo an absorption reaction inside the material, according to the equilibrium:

absorbedadsorbed HH ↔ .)24.2(

The direct reaction, of passage of atomic hydrogen through the interface, depends

on the surface coverage (θ) and on the number of available sites in the subsurface that

hydrogen can occupy.

The consequence is the accumulation of hydrogen under the metallic surface,

leading to a concentration C0.

The reversal reaction, of hydrogen passage form the subsurface towards the

surface can take place.

The rate for the reversal reaction is proportional to the subsurface concentration

and to the concentration of empty sites on the surface through which hydrogen could be

desorbed (1 – θ ).

The rate of the overall reaction is:

)1(1 00 θθ −−

−= CkNN

ki dess

absabs .)25.2(


20

At equilibrium (i abs = 0) and for small coverage degree θ and small degree of

saturation, the subsurface concentration of hydrogen C0 is proportional to the degree of

coverage:

0Ckk desabs =θ .)26.2(

or:

θKC =0 .)27.2(

where K is the ratio of the rate constants for the direct and reversal reaction of the

hydrogen passage through the interface (2.24.).

The equation (2.27.) shows that the subsurface concentration of hydrogen, C0, is a

function of the degree of coverage, θ.

We remind here that in the previous section we saw that θ depends on the

charging conditions (current density and potential) and also on the temperature.

Thus, the concentration of hydrogen in the material will also be depending on these

parameters.

Assuming that the passage of hydrogen through the interface (2.24.) is the slowest

step, then, for the Volmer - Tafel mechanism the relationship between the subsurface

concentration and the charging current can be found considering that reactions (2.1.) and (2.2.)

are at equilibrium, thus:

[ ] ²exp)1( 21 θηαθ kRTFHki =

−= + .)28.2(

Replacing θ from (2.28.) into (2.27.) leads to the value of C0 :

ik

KC2

01

= .)29.2(


21

The subsurface concentration C0 depends not only on the charging current, but

also on the kinetics of hydrogen evolution (through k2).

The square root dependence on the charging current is followed up to certain value

of the current density. Above a limiting value, the subsurface concentration of hydrogen

becomes independent of the charging current (Figure 2.3.). This corresponds to the saturation

of the surface ( θ = 1).

In the presence of promoters (or poisons) this square root dependency is followed

up to high values of charging current density as presented in Figure 2.3 [18].

Thus, the subsurface concentration, C0, depends on the charging conditions

(applied current or potential), on temperature, pH, state of the surface (presence of poisons,

presence of oxides, other species than hydrogen adsorbed on the surface).

1,0E+00

1,0E+01

1,0E+02

1,0E+03

1,0E+00 1,0E+01 1,0E+02 1,0E+03 1,0E+04i1/2 (µA/cm²)

Ct (

ppm

wt)

1,0E-05

1,0E-04

1,0E-03

1,0E-02

C0

(ppm

wt)

Figure 2.3. Example of experimental hydrogen content – charging current density dependency. Tests were conducted on 22Cr duplex stainless steel in a chloride solution without H2S (1) and

saturated with H2S (2) [18].

(1)

(2)


22

2.2.2. DIRECT TRANSFER OF H+ THROUGH THE INTERFACE (MODEL

PROPOSED BY CROLET ET AL.)

In the classical mechanism presented above, it is considered that the same

adsorbed atoms lead either to molecular hydrogen evolution or to atomic hydrogen absorbed in

the material.

Crolet [19 – 21] considers that a hydrated hydrogen cation looses its water

atmosphere and will proceed directly through the interface. Thus, the subsurface concentration

is given by the passage of H+ (in ionic form) through the interface.

In this new model for hydrogen absorption the species that are leading to

molecular hydrogen (atomic hydrogen adsorbed on the surface) are different from the ones that

are undergoing absorption.

The equilibrium (2.24.) is written as:

++ ↔ metalhydrated HH .)30.2(

The direct transfer of the proton H+ could take place not only from H+ from water,

but also by deprotonation of other complexes adsorbed on the surface: H2S becomes HS- ,and

similarly for H3As, H3P, HSCN, HF.

The severity of these species with regard to cracking increases with decreasing

stability of these complexes. Thus, the adsorbed HS-ads, is not an inhibitor for recombination

reaction (2.2.) or (2.3.) but a catalyst for the direct transfer of H+ through the interface (reaction

2.30.):

+−

−+

+→

→+

metaladsorbedadsorbed

adsorbedadsorbedhydrated

HHSSH

SHHSH

2

2

.)31.2(


23

This mechanism gives a new explanation of the damaging effects of H2S and other

poisons: these species favour the direct transfer of H+ in the metal without decreasing the

recombination of atomic H.

This new direct transfer mechanism does not necessarily exclude the classical

mechanism, where atomic hydrogen is passing through the interface.

Apparently more studies are still conducted on these interesting questions.

2.3. HYDROGEN TRANSPORT WITHIN THE MATERIAL

In the following, the diffusion laws are presented for the ideal case of diffusion

without trapping. Then the different types of traps that can be encountered in a real material are

presented, with the expression for diffusion in the presence of traps. In the last subsection

hydrogen diffusion inside the material under other gradients than composition gradient is

presented.

2.3.1. DIFFUSION

2.3.1.1. Ideal diffusion, Fick’s laws

Due to its small volume, a hydrogen atom can diffuse and occupy the interstitial

sites inside the material (Figure 2.4.).

(a) (b) Figure 2.4. Octahedric interstitial sites of face centred cubic (a) and body

centred cubic (b) matrix [22]


24

The face-centred cubic (f.c.c.) lattice has one octahedral interstitial site per metal

atom and two tetrahedral interstitial sites per metal atom in the unit cell.

The body-centred cubic (b.c.c.) lattice has three octahedral interstitial sites per

metal atom and six tetrahedral interstitial sites per metal atom in the unit cell. In the f.c.c. lattice

the octahedral positions (O) have the largest free volume, whereas in the b.c.c. lattice the

tetrahedral sites (T) are the largest. In Table 2.3. several metals with crystallographic structure

and preferred occupied sites are presented [22].

Table 2.3. Interstitial sites occupied by hydrogen in different metals [22]

Host lattice Crystallographic structure Occupied sites

α – Fe b.c.c. T

γ – Fe f.c.c. O

Pd f.c.c. O

Ta b.c.c. T

V b.c.c. T

Nb Rhomb. T

where T are tetrahedral sites and O octahedral sites .

The subsurface concentration, C0, determines a concentration gradient that is the

driving force for the diffusion.

The diffusion obeys the Fick’s laws:

Fick’s first law for diffusion: xCDJ∂∂

−= .)32.2(

and

Fick’s second law for diffusion: ²

²xCD

tC

∂∂

=∂∂

.)33.2(

where D is the diffusion coefficient for the ideal case , where diffusion takes place

without trapping.


25

For a given material the diffusion coefficient is constant. Some values of the

diffusion coefficient for different materials are presented in Table 2.4.

Table 2.4. Diffusion coefficient of hydrogen in different materials at room temperature

Material D (cm²/s) Ref.

Carbon steel 2.5 x 10 – 6 [11]

Ferritic stainless steel 10 – 7 [23]

Austenitic stainless steel 2.15 x 10 – 12 [33]

Martensitic stainless steel 2 x 10 – 9 [24 – 28]

Duplex stainless steel 10 –9 – 10 – 10 (depending on

the ferrite / austenite ratio)

[23, 29 – 32 ]

2.3.1.2. Temperature effect

The diffusion coefficient has an exponential expression dependence with the

temperature:

−=kTEDD aexp0 .)34.2(

In Figure 2.5. the diffusion coefficient – temperature dependency for different steels

is presented [33].

1,00E-14

1,00E-10

1,00E-06

1,00E-02

0 100 200 300 400

temperature (°C)

D (m

²/s)

C steelduplexaustenitic

Figure 2.5. Evolution of diffusion coefficient (D) with the

temperature for different steels [33]


26

The interstitial hydrogen is reversible, therefore at room temperature it can diffuse

out from the metal.

The diffusion coefficients (apparent and real) and the subsurface concentration, C0,

can be easily measured and calculated from permeation experiments (See Section 6.).

2.3.1.3. Diffusion under stress gradient

When a stress is applied, hydrogen diffuses under stress gradient toward the

places of high stress. The diffusion flux depends not only on the concentration gradient, but also

on the stress gradient according to equation:

−−=

−

hgradRTVcgradcDJ σ .)35.2(

Stress – induced hydrogen diffusion takes place whether the inhomogeneous

stress is caused by applied forces or residual stresses.

Due tot the stress gradient, the diffusion of hydrogen can take place even when

hydrogen distribution is uniform inside the material (grad c = 0).

2.3.2. TRAPPING

In real cases, the hydrogen atoms are not located only in the interstitial positions,

but they are trapped by the different defects inside the material.

Any metallurgical defect inside the material can act as a trap for hydrogen.

According to their energy, traps are divided into reversible, for which the energy is low and

hydrogen can leave easily the trap, and irreversible (or deep traps) where more energy has to

be provided for hydrogen release.

Examples of traps are presented in Table 2.5.


27

The presence of traps in the material will hinder the hydrogen diffusion. Therefore,

the diffusion will be apparently slower than for an ideal crystal.

Table 2.5. Types of traps existing in ferrous materials

Traps Binding energy

(kJ / mol)

Degassing temperature

(°C) Material Ref.

Matrix 6.9 Room

temperature Fe [34]

Grain boundaries 17.15 112 Fe [35]

215 Fe [35]

200 Fe [36]

Dislocations 20 – 26

272

AISI 4340

(carbon steel

0.39%C)

[37]

338 AISI 4340 [37]

305 Fe [35] Microvoids 35 – 48

480 Carbon steel

(0.47%C) [38]

MnS inclusions 72 495 AISI 4340 [37]

Carbides

interfaces 96.6 723 Fe [36]

The diffusion laws have to take into account the presence of reversible and

irreversible traps.

The energy of reversible traps is low and comparable to the energy of interstitial

sites.

According to Oriani’s theory [39], the amount of hydrogen trapped into the

reversible traps is always in equilibrium with the lattice hydrogen.


28

−=

− RTWB

L

L

r

r exp11 θθ

θθ

.)36.2(

The diffusion law is expressed by the equations:

tN

tN

xCD

tC i

ir

r ∂∂

−∂∂

−∂∂

=∂∂ θθ

²²

.)37.2(

where :

rrrr pCkt

θθθ −−=∂∂ )1( .)38.2(

)1( irii Cktθθ −=

∂∂

.)39.2(

The Fick’s law can be written as:

²²xCD

tC

eff ∂∂

=∂∂

.)40.2(

where Deff is the effective diffusion coefficient including the effect of traps.

2.3.3. TRANSPORT OF HYDROGEN BY MOVING DISLOCATIONS

The plasticity model proposed for HE suggests that hydrogen enhances

dislocations mobility (See Section 4.2. Plasticity model). In the same time, moving dislocations

could transport hydrogen during their movement.

Therefore, the transport of hydrogen by dislocation is an important aspect in

understanding the hydrogen effects on the mechanical properties of the material.


29

Under plastic deformation the number of dislocations is increased and, as

dislocations offer more space for hydrogen, there is more hydrogen in the metal. Also, as

plastic deformation imposes a movement of dislocations, the transport of hydrogen is facilitated

[40].

There are indeed experiments that show an increased solubility and a decreased

apparent diffusivity of hydrogen for samples that have been cold worked prior to permeation

experiment. The latter is due to the trapping effect of dislocations.

Considerable controversy has arisen with respect to the exact role that dislocation

transport plays in the embrittlement process.

Experiments conducted on Ni [41] and Ni – Cr alloys [42] showed that the

diffusivity increases very much under plastic deformation during hydrogen charging. The

diffusivity appears to increase by approximately 5 orders of magnitude, from about 10-10cm²/sec

to 10-5 cm²/sec.

Bastien and Azou suggested that moving dislocations could carry hydrogen [43]

and that a local high hydrogen concentration can develop when dislocations annihilate or

intersect a void.

Franck [44] observed that deformation enhances hydrogen outgassing of mild

steel.

Two types of transport behaviour are generally discussed:

a. All the hydrogen is introduced into the material prior to plastic deformation.

Initially, the hydrogen is uniformly distributed and the total hydrogen content remains constant.

When the system is mechanically loaded, new dislocations are formed. A

redistribution of the hydrogen from the lattice and old traps towards the newly formed

dislocations occurs.

At the same time, the dislocation motion starts. According to their velocity, the

moving dislocations transport or loose the trapped hydrogen. Authors generally consider that if

the velocity is sufficiently high, all the hydrogen will be left behind as dispersed through the

lattice and perhaps partially absorbed by static traps.

For steels, this dumping process may double the lattice hydrogen concentration if

the deformation rate is very high.


30

Dislocations will act as hydrogen source inside the material.

For velocities near or less than a critical value (case of slow strain rate tests) the

hydrogen is transported by moving dislocations.

If the dislocation slows down, e.g. after intersection with an obstacle, its hydrogen

concentration may increase again to a level higher than that of fast moving dislocations.

Hence, dislocations which are stopped or slowed down are sinks for hydrogen.

b. The second case is the case where hydrogen charging and plastic deformation

occur simultaneously. The dislocations moving from the surface can be assumed to transport

hydrogen at a concentration equal to that of the external charging medium. In this case

dislocations will be sources of hydrogen inside the material because:

– fast moving dislocations loose their hydrogen, as in a.

– slow moving dislocations may loose hydrogen when their initial concentration is

higher than inside the material

The critical velocity of dislocation to transport or loose hydrogen was calculated as

[45]:

−

−=

RTF

RTQ

RTblDv kkokcr

*2expexpσ .)41.2(

For iron all these parameters are known [45] : l = 1 µm, b = 0.248 nm, ν = 1013 s–1 ,

G = 86 GPa (shear modulus), Qk = 3.4 kJ / mol and 2Fk* = 28.4 kJ / mol, and the critical velocity

equals:

vcr = 8.33 x 103 (σ / G), in m/s (2.42.)

For material testing, the most severe conditions will be those where hydrogen is

transported more easily by dislocations, thus at a strain rate below the critical value given by

equation (2.42.).


31

2.4. SUMMARY OF HYDROGEN IN THE METAL

The hydrogen entry and transport through the material is influenced by the

charging conditions (through C0), the material structure, the level of stress and the strain.

The charging conditions (potential or current density, hydrogen pressure,

temperature, solution pH) lead to a given subsurface concentration, C0.

The subsurface concentration, C0, was for a long time considered as the main

factor for the susceptibility to hydrogen embrittlement (See Section 3. and Section 4.)

The subsurface concentration, C0, is the driving force for diffusion of hydrogen

inside the material.

It is pointed out that there is a limit value for C0 corresponding to saturation in given

charging conditions. This limit value is very much increased in the presence of promoters for

hydrogen entry.

The total hydrogen content is given by the reticular hydrogen (hydrogen located in

interstitial positions of the matrix) and by the hydrogen trapped in metallurgical defects

(hydrogen reversibly or irreversibly trapped).

Thus the total hydrogen content depends not only on the charging conditions but

also on the history of the material: for instance, a previous plastic deformation increases the

density of dislocations, thus there are more available hydrogen trapping sites in the material.

The transport by diffusion takes place under concentration gradient established by

C0 and also by stress gradient, due to externally applied forces or residual stresses. As a

consequence, the total content of hydrogen inside the material is increased when C0 increases,

but also when a stress is applied to the sample.

Under dynamic plastic deformation, dislocations can transport hydrogen during

their movement. For certain materials, like Ni [41, 42], this type of transport is more rapid than

by simple diffusion.


32

3. TYPES OF DAMAGES DUE TO

HYDROGEN

3. TYPES OF DAMAGES DUE TO HYDROGEN

34


35

In this chapter a summary of the most important damages caused by the presence

of hydrogen in the material is presented.

The purpose is to identify where the problem studied in this work is situated among

the series of hydrogen damages.

These damages are generally classified according to various factors:

- the form of hydrogen that produces the damage: atomic hydrogen, gaseous

hydrogen, other gases (CH4, H2O), metal hydrides ;

- the source of hydrogen: electrochemical hydrogen (cathodic protection,

electroplating, corrosion, pickling), gaseous hydrogen atmosphere, H2S and other poisons;

- presence or absence of applied stress.

There are inevitable interferences between these three classification modes.

It is possible that a better understanding of the mechanisms of hydrogen damages

might in the future suggest a more logical classification.

The problem studied in this work is related to damages due to the presence of hydrogen in atomic form, under applied stresses.


36

3.1. HYDROGEN INDUCED CRACKING (HIC) AND STEPWISE

CRACKING (SWC)

Atomic hydrogen H diffuses inside the material and recombines as molecular

hydrogen, H2, at specific sites. This H2 can develop high pressure at these sites.

In more or less ductile materials (yield strength below about 540 MPa [46]), these

high pressures deform the material and produce blisters of gaseous molecular hydrogen H2.

HIC blisters are typical of low strength steels in environments containing H2S.

For less ductile material cracking can occur once the pressure exceeds a critical

value.

Atomic hydrogen is usually produced by a corrosion reaction in H2S containing

environments.

This type of failure is typical for medium or high strength carbon steel pipe

transporting H2S containing products.

Figure 3.1. Hydrogen induced cracking (HIC) of carbon steel in H2S environment [47]: blisters are associated

to MnS inclusions


37

The recombination of H into H2 generally occurs on elongated MnS inclusions [47]

(Figure 3.1.). Counter measures are reducing the S content in steel to very low levels (0.001 %)

or adding calcium (Ca) or rare earth elements (Cerium for example) to build spherical

inclusions.

This type of damage occurs in the absence of applied stress. The stress is

produced by the internal hydrogen pressure.

HIC may propagate in a stepwise manner (stepwise cracking SWC), which is

shown in Figure 3.2.

3.2. STRESS ORIENTED HYDROGEN INDUCED CRACKING

(SOHIC)

SOHIC is a combination of HIC and sulphide stress cracking SSC or hydrogen

embrittlement (see below). It involves blisters formed by H2, as in HIC, but it occurs in the

Figure 3.2. Typical aspect for cracking by stepwise cracking (SWC) [46]


38

presence of an applied or residual stress. The joint presence of atomic H and stress produces

thin cracks than can interconnect HIC cracks (see arrow on Figure 3.3.).

.

These cracks develop perpendicular to the applied stress direction [47]. Figure 3.3.

presents a typical aspect of cracking by stress oriented hydrogen induced cracking (SOHIC).

3.3. HYDROGEN REACTION WITH THE METAL MATRIX (HYDRIDE FORMATION)

Hydrogen can form brittle hydrides with a number of metals: Ti, Zr, V, Nb, Ta, Mg,

U, Th and their alloys.

Failures of Ti tubes were observed in condensers near tube plates that were

cathodically protected.

Hydride formation occurs above a critical hydrogen concentration. Pre-existing

cracks may be the initiation site, since H tends to diffuse towards the region of high stress at the

Figure 3.3. Stress Oriented Hydrogen Induced Cracking (SOHIC) [47]

stress direction stress direction


39

crack tip. The hydride phase has a higher volume than the parent metal and this creates local

elastic or plastic deformation.

This type of failure does not occur on ferrous based materials.

3.4. HYDROGEN REACTIONS WITH NON METALLIC PHASES Hydrogen can react with iron carbides in the steel to form methane CH44. The

reaction occurs at temperature of 200 to 300°C.

These decarburising damages affect the material in two ways: by reduction of

carbon content (and a reduction of the strength of the material) and by high internal pressures

produced by methane that may lead to cracking.

Also, hydrogen can react at high temperature with certain non-metallic inclusions. One example is the reaction of hydrogen with copper oxide inclusions during annealing of copper in hydrogen atmosphere.

The reaction is

)(222 2 gOHCuHOCu +→+

with a weakening of the copper matrix.

3.5. HYDROGEN EMBRITTLEMENT (HE) (OR HYDROGEN

STRESS CRACKING HSC)

This type of failure occurs in the presence of atomic hydrogen and under residual or applied stress.

This phenomenon is of concern for high strength steels (yield stress YS > 600 N /

mm²). In the presence of hydrogen, their mechanical resistance can decrease considerably,

below the yield strength.

For many steels, there is a threshold stress below which HE does not occur.

However, this threshold stress is a function of the strength of the material and of the


40

environment. Generally, the higher the yield or tensile strength, the lower the threshold stress.

HE is definitely linked to absorption of H, and there is often an incubation time for charging and

transport of H, resulting in delayed fracture.

Hydrogen embrittlement (or hydrogen stress cracking, as stress is a prerequisite)

generally produces clear single cracks (intergranular or transgranular), as opposed to often

branching cracks for stress corrosion cracking SCC.

An interesting manner to differentiate HE (or HSC) and SCC was presented by

Brown [48] who showed a decreasing time to failure when the electrochemical potential is either

increased or decreased: shorter time to failures at higher potentials are attributed to SCC, with

anodic dissolution, while early cracking at lower potentials are attributed to hydrogen effects,

possibly in the absence of any corrosion.

The sources of atomic hydrogen can be cathodic protection, electroplating, pickling

or just corrosion.

This work is focused on this type of damage; the source of hydrogen is cathodic protection, there is no corrosion and stress and strain are present.

3.6. SULPHIDE STRESS CRACKING SSC

Sulphide stress cracking, SSC, is a special case of HE or HSC where sulphide

compounds (or other poisons) increase the entry of atomic hydrogen in the metal lattice. The

name was given in relation with this type of environment.

Sulphide stress cracking may be particularly rapid and catastrophic.

Hard microstructures are more sensitive to sulphide stress cracking. Improper heat

treatments or heat affected zone of low heat input welds are frequent causes of sulphide stress

cracking.

The occurrence of this type of damage in sour environment is often prevented by

limiting the hardness of the material.

The problem dealt with in this work, does not consider H2S or other

“poisons”.

4. HYDROGEN EMBRITTLEMENT


42


43

From a fundamental point of view, HE is not completely understood, despite the

large amount of studies conducted on this topic.

Several models were proposed for hydrogen embrittlement and a brief description

of them is presented in the following.

Adsorption model One of the first mechanisms proposed for HE is the one proposed by Petch [49].

This model considers that hydrogen adsorbed on the surface decreases the surface energy.

The reduction of surface energy is proportional to the concentration of adsorbed hydrogen.

When hydrogen is adsorbed on the internal surfaces of cracks or voids the surface energy

decreases, reaching values below a critical stress and thus leading to fracture.

This model does not explain why other gases, like O2 or N2, with higher adsorption

energies, have not the same embritteling effects.

Decohesion model [50] (see Section 4.1.)

According to this model, hydrogen inside the material decreases the cohesion

forces that exist between atoms.

This model is applicable to brittle fracture (cleavage or intergranular) but it cannot

explain the fractures accompanied by plasticity.

This model also considers that there is a critical hydrogen concentration for the

occurrence of brittle fracture.

Plasticity models

Regarding the effect of hydrogen on the plasticity of the materials, two contradictory theories are developed:

Stroh [51] considers that hydrogen blocks dislocations and thus increases the flow

stress.

This theory is contradicted by the observations of Beachem [52] who proposes that

hydrogen facilitates the dislocations movement, favouring the plasticity of the material.


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These latter observations are the basis for the HELP model (Hydrogen Enhanced Localised Plasticity) that will be detailed in Section 4.2.

Another mechanism that considers the effects of hydrogen on the plasticity of

material is proposed by Lynch [53 – 55]. According to this mechanism, hydrogen adsorbed on

crack tip facilitates the dislocations injection and crack advancement. This mechanism is similar

to the mechanism for embrittlement by liquid metals and is taking into account only the effects

of hydrogen on the surface of the material.

Among all these models, the most elaborated and well established models are the

decohesion mechanism and the plasticity model HELP.

A detailed analysis of these two models is presented in the following sections.

4.1. DECOHESION MECHANISM

The decohesion mechanism is one of the earliest mechanisms proposed for HE. It

was initiated by Troiano [50], and improved by Oriani [56 – 58].

This mechanism is based on the postulate that solute hydrogen decreases the forces required to separate the crystal along a crystallographic plane, i.e. the cohesive force and correspondingly the energy to form cleavage surface (Figure 4.1.).

Decohesion is supported by the fact that HE seems to occur in the absence of

significant local deformation, by thermodynamic arguments and by theoretical calculations of

electronic distribution of systems in the presence of hydrogen.

4.1.1. THERMODYNAMIC ASPECTS OF INTERFACIAL DECOHESION [59] The interfacial layer was modelled in order to obtain a relationship between the

hydrogen content and the cohesive strength and the work of decohesion.

The thermodynamic of decohering interface was well established [59 – 62] with the

emphasis of two limiting cases: infinitely fast separation and infinitely slow separation.


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Figure 4.1. Hydrogen effect on the cohesive energy (U) and on the cohesive

stress (σ) of a material. U°cohesion is the cohesive energy (the energy

required to separate the two half solids along the cleavage plane to a separation larger than the critical distance (r) in the absence of hydrogen; UHcohesion is the cohesive energy in the presence of

hydrogen in solid solution, σ°cohesion is the cohesive stress (the

stress to disrupt the atomic bonds) in the absence of hydrogen;

σHcohesion is the cohesive stress in the presence of hydrogen in solid

solution; a is the lattice parameter; εH is the strain induced by

hydrogen in solid solution [50, 56 – 58].

Ucohesion , σcohesion

U°cohesion

UHcohesion

σ°cohesion

σHcohesion r

HH

H HH

HH

HHH H

a = a0 + εH a0

a = a0 + εH a0

rH

(a) (b)


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For slow separation, the diffusion of impurity towards the interface is much faster than the separation; the chemical potential of impurity remains constant while its interfacial

concentration varies.

In the case of fast separation, where the diffusion is slow compared to the separation time, the interface remains isolated from the impurity source, the concentration in the

interface remains constant, and the chemical potential changes during the separation process.

Recently, the intermediate situation, where the two processes (separation and

diffusion) are comparable to each other, was modelled [57, 62]. This situation is closely related

to the diffusion – controlled brittle fracture taking place in many materials.

The three regimes can be represented in concentration – separation distance

coordinates, (c, δ), as in Figure 4.2. There is a critical state for which the separation reaches the

critical value, δ = δc. On the right of the critical separation line the fracture occurs, whereas on

the left side the interface is subjected to stress but the fracture does not occur yet.

For fast separation, the H concentration remains constant during the evolution of

system up to fracture and this regime is represented by a horizontal line.

For slow separation, the evolution of system is represented by a curve µ(c, δ) =µ0.

The intermediate regime lies between these two curves.

The cohesive strength and work of decohesion, that are concentration dependent,

where calculated.

The mathematical development of the model is presented in Appendix 2.

The cohesion of the interface is described by a function that depends on the

concentration of hydrogen (c) and the separation δ between the layers that form the interface.

This function is called the cohesive function, ϕ (c, δ). The cohesive function depends linearly on

hydrogen concentration and by a more complex function (cubic polynomials) on the separation

δ. For the critical separation δc, the cohesive function becomes zero.

The interfacial stress depends on c and δ through ϕ.

The calculations conducted by this model (see Appendix 2) show that the

cohesive stress of the interface decreases with an increase of the hydrogen concentration (Figure 4.3.)


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critical separation line

slow

fast

Figure 4.2. Schematic trajectories of interfacial separation in coordinates

hydrogen concentration c versus separation distance δ. The

initial state is characterised by c0 and δ0. The interface fails

when the trajectory crosses the critical separation (δc) line

[59]

Normalised separation, δ/δc

Nor

mal

ised

coh

esiv

e st

ress

σ∗

Figure 4.3. Normalised cohesive stress for different values of hydrogen concentration [59]


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4.1.2. ELECTRONIC DISTRIBUTION IN H – METAL SYSTEMS

The decohesion mechanism is supported by calculation of the electronic

distribution of the metal – hydrogen systems.

There is an increasing amount of studies on to the hydrogen effects on metallic

bonds [62 – 71].

The atom superposition and electron delocalisation molecular orbital method

(ASED – MO) (see Appendix 3.) is a semi – empirical method that can predict the molecular

structure from atomic data (atomic wave functions and ionisation potentials).

The calculations conducted using this method show that hydrogen inside the

transition metals determines an electronic delocalisation of the surrounding metals atoms.

The bond between hydrogen atom and the metallic atoms is formed at the expense

of metal – metal bond of neighbouring atoms.

The calculations conducted for H – Fe system show that hydrogen diminishes the

strength of Fe – Fe bond for the first neighbour atoms by more than 40% of the initial value.

4.1.3. EXPERIMENTAL OBSERVATIONS

Intergranular fracture consistent with decohesion has been observed in some particular systems where other embritteling species are segregated at the boundaries: Ni – S

system with low levels of S [72] and the systems Ni3Al containing B, where solute H did appear

to decrease the grain boundary cohesion [73].

The decohesion mechanism seems to be operative in the case of β – Ti – H

system [62], where the strain to failure and the nature of the fracture surfaces depend on the

hydrogen content (Figure 4.4.). The strain to failure depends on the hydrogen concentration and

it has a sharp transition in ductility for a hydrogen concentration (expressed as atomic ratio) H /

M = 0.28. For H / M < 0.21 the fracture surfaces are characterised by ductile microvoid

coalescence with decreasing microvoid sizes as the H concentration increases. At H / M > 0.27

the fracture surfaces is cleavage in nature.

For systems with low solubility of hydrogen, such as steels, the hydrogen

distribution at the crack tip is dominated by the trapped hydrogen and the concentration at the

crack tip exceeds that of the bulk by a factor of 102. Since hydrogen embrittlement occurs in


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steels at an average concentration of about 10 atomic ppm, the enhanced concentration is of

the order of 103 atomic ppm. This is a low concentration to account for low macroscopic strain

fracture on the basis of decohesion.

4.1.4. CONCLUSIONS FOR DECOHESION MECHANISM

The decohesion mechanism considers that there is a critical concentration of

hydrogen atoms for which brittle fracture occurs.

This mechanism could be applied for intergranular fracture, where high

concentration of hydrogen accumulates at grain boundaries (and thus reaching the critical

concentration for brittle fracture)

The plasticity associated with the fracture is not compatible with the decohesion

mechanism.

Figure 4.4. Strain to failure dependency on the hydrogen concentration

of β – Ti alloys. A sharp transition on ductility is observed

for hydrogen concentration of about 0.28 [62]


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There is little definitive evidence of the hydrogen decohesion in systems with low

solubility of hydrogen, such as steels.

The available evidence suggests that solute hydrogen does not weaken atomic

bond. The elastic constant measurements [74 – 76] and phonon dispersion curves [77] show a

stiffening of the lattice bonds, rather than a softening.

4.2. PLASTICITY MODEL: HYDROGEN ENHANCED LOCALISED PLASTICITY – HELP

The basis of this model was established by Beachem [52], who was the first to

propose that in the case of hydrogen embrittlement the failure occurs by locally ductile

processes.

The HELP (Hydrogen Enhanced Localised Plasticity) mechanism is based on consistent observations that the presence of hydrogen in solid solution increases the mobility of

dislocations (Figure 4.5.) and creates localised high deformation regions [78 – 88].

The reason of this increased mobility is attributed to reduction of interactions

between dislocations and between dislocations and other obstacles (such as C atoms, grain

boundaries) when hydrogen is present in system.

Dislocations thus move closer to each other, and closer to obstacles, and produce

denser or more compact pile-ups when H is present.

The result of this is that microscopic regions of high deformation (where H

increases the mobility of dislocations) are surrounding less ductile zones where dislocations are

closely packed. The applied stress is then concentrated on these hard zones that represent only

a small portion of the cross section. When the tensile stress in these small portions is higher

than the ultimate tensile strength, failure occurs.

Even if, at very localised (microscopic) level, plasticity is enhanced by hydrogen, at

macroscopic level the material exhibits a brittle behaviour (reduced strain to failure, lower

fracture strength) [78].


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The assumption for the HELP mechanism is that hydrogen creates an atmosphere

around dislocations that shields the elastic interactions between dislocations, and between

dislocations and other obstacles.

This model is well established with respect to the influence of hydrogen on the

dislocations behaviour.

This model is supported by calculations of elastic shielding due to hydrogen

atmosphere (see 4.2.1) and by in situ transmission electron microscopy (T.E.M.) observations

of the influence of hydrogen on the dislocation behaviour (see 4.2.2, ).

4.2.1. HYDROGEN SHIELDING EFFECT

The hydrogen shielding effect was modelled [89 – 95]. Analytical and finite element

Figure 4.5. Hydrogen effect on dislocations velocity: dislocations velocity increases when hydrogen pressure increases. In the figure the ratio dislocation velocity in hydrogen atmosphere to velocity in vacuum is represented. Curve 1 shows the effect of introducing hydrogen for the first time, and curve 2 the effect after hydrogen was removed and reintroduced. Tests conducted on alfa -Titanium


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calculations were conducted to estimate the effects of the presence of hydrogen on the

interactions between two edge dislocations and between a dislocation and an impurity atom

(like interstitial C atoms).

When hydrogen diffuses into the lattice, despite its small volume, it introduces an

elastic strain in the matrix.

In this model, first the distribution of hydrogen in equilibrium with an applied stress

is calculated.

The stress varies with the position inside the material, thus, the hydrogen

distribution evolves similarly.

Then the stress exerted by the hydrogen atmosphere on a dislocation is calculated.

The total shear stress exerted on a dislocation is the sum of the shear stress due

to the hydrogen stress field and the shear stress due to the other dislocations.

4.2.1.1. Hydrogen effect on the interactions between dislocations The hydrogen shielding effect was modelled taking into account the strain induced

by the presence of hydrogen inside the matrix, the hydrogen segregation around dislocations

and the interactions between parallel dislocations with the same sign and Burgers vectors.

(a) Analytical calculation

The entire mathematical development is presented in Appendix 4.

The shear stress exerted by the H atmosphere is calculated and is determined as a

function of hydrogen concentration, CH, and of material characteristic through µ (bulk modulus):

φφφνπ

τπ

drdr

rCNVµ R

rH

A

HH

2sin),()1(2

2

0 2∫ ∫−−= .)1.4(

The interaction between two parallel dislocations of the same sign , in the absence

of hydrogen, is calculated and is established as being function of the material (through the bulk

modulus µ and Poison’s coefficients v), of the Burgers vector b and the separation distance l ;


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lµb

Dωω

νπτ 2coscos

)1(21

−−= .)2.4(

The total stress exerted on dislocation is the sum τH + τD.

(b) Numerical integration results

Numerical integration was calculated for the niobium system [96].

(For iron and steels the available data are still dispersed.)

According to this model, the hydrogen atmosphere in equilibrium with a stress field

of a single dislocation is symmetric with respect to the dislocation plane (Figure 4.6.).

When two dislocations approach one another, the symmetry of hydrogen

atmosphere changes, due to the superposition of the stress fields of the two dislocations.

For dislocations of the same sign (Figure 4.7.), the hydrostatic stress field is

reinforced positively below the slip plane and negatively above the slip plane. This

reinforcement increases as the dislocations approach each other. Consequently, the hydrogen

concentration increases in the regions of positive stress enhancement and its value becomes

larger than the concentration of the corresponding region in the atmosphere of a single

dislocation.

For dislocations of opposite sign (Figure 4.8.), that attract each other, the positive

hydrostatic stress field of each dislocation is weakened and the hydrogen concentration is lower

than that of a single dislocation.

The shear stress is diminished in the presence of hydrogen and this reduction

depends on the separation distance between the dislocations. The effect of hydrogen shielding

is higher when the distance between the dislocations is smaller. For high dislocations

separation (> 10 Burgers vectors) the magnitude of the effect is almost independent of the

distance between dislocations.


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Figure 4.6. Symmetrical contours of normalised hydrogen concentration C/C0 around a single edge dislocation at a normal hydrogen concentration C0 = 0.1 and temperature of 300K [96]

Figure 4.7. Contours of normalised hydrogen concentration C/C0 around two parallel edge dislocations of equal Burgers vectors b, on the same slip plane, at a normal hydrogen concentration C0 = 0.1 and temperature of 300K [96]

Figure 4.8. Contours of normalised hydrogen concentration C/C0 around two parallel edge dislocations of opposite and equal Burgers vectors b, on the same slip plane, at a normal hydrogen concentration C0 = 0.1 and temperature of 300K [96]


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A direct consequence of the shielding effect of hydrogen on the dislocation behaviour is that in the presence of hydrogen the spacing between dislocations in a pile – up decreases. This effect is higher for the dislocations at the tip of the pile – up, where the separation distance between dislocations is smaller.

The result of this is that microscopic regions of high deformation (where H

increases the mobility of dislocations) are surrounded by hard zones, where more compact pile-

ups are formed. The stress distribution is not uniform through the whole section of the

specimen and spots of stress concentration occur.

Maybe, for high hydrogen concentration, dislocations of same sign can join causing

fracture by coalescence.

4.2.1.2. Interaction between dislocations and an impurity atom in the

presence of hydrogen

A model similar to the one presented above was used to calculate the interactions

between dislocations and a defect and the influence of hydrogen on these interactions [89].

Interstitial atoms (like carbon) are barriers

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