A Study of Internal Hydrogen Embrittlement of Steels

13
Materials Science and Engineering A286 (2000) 269 – 281 A study of internal hydrogen embrittlement of steels G.P. Tiwari a, * , A. Bose b , J.K. Chakravartty a , S.L. Wadekar a , M.K. Totlani b , R.N. Arya a , R.K. Fotedar b a Metal Physics Section, Materials Science Di6ision, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, India b Materials Processing Di6ision, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, India Received 22 October 1999; received in revised form 2 February 2000 Abstract A novel procedure for hydrogen charging and studying the Internal Hydrogen Embrittlement (IHE) of steels is described here. A cylindrical notched tensile sample with an extended end is employed for hydrogen charging. The extended portion of the sample forms the cathode in an alkaline bath and a constant uni-axial tensile load is applied during hydrogen charging. The stress gradient set up by the notch, which is not in contact with the electrolyte, enhances the hydrogen concentration at various trapping sites of the matrix beyond the solubility limit. Subsequent to charging, the specimen is kept under the same load as that during charging, for another 24 h to stabilize the population of hydrogen within the specimen matrix. At the end of this stage, the specimen is tensile tested to failure at room temperature. Two different steels namely maraging and mild steels have been chosen to study the effect of hydrogen ingress on mechanical properties. While an increase in tangent modulus (linear portion of the stress – strain diagram), yield strength, work hardening rate and ultimate tensile stress (UTS) has been observed on hydrogenation, a decrease in total elongation has been noticed for both the steels studied. Fractographic investigation has revealed that the fracture mode is predominantly ductile dimple (failure by micro-void coalescence) in both the materials and that the hydrogen reduces the size of the dimples. The observations of this investigation are significant in two respects: firstly, it demonstrates the efficacy of a hydrogen charging method for steels which can introduce hydrogen to a level much higher than its solubility limit and secondly, it reports for the first time enhancement of modulus and work hardening by hydrogen charging. These observations have been rationalized on the basis of current understanding on the effect of hydrogen on plastic properties and hypothesis of the models of IHE. It is suggested that the trapping of hydrogen by dislocations and other structural features of the matrix and the mutual interactions of their strain fields can account for the observed effects on yield strength, tangent modulus, work hardening rate, UTS and ductility. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Internal hydrogen embrittlement; Steel; Trapping; Mechanical properties www.elsevier.com/locate/msea 1. Introduction The effects of hydrogen on the mechanical properties of iron and steels have been widely investigated [1–3]. As regards the embrittling effect of hydrogen it is well known that the chief effects are a decrease in ductility and true stress at fracture. Hydrogen embrittlement (HE) of steels can be classified into three main types [4,5]. 1. Hydrogen reaction embrittlement arises because of the generation of hydrogen on the surface as a result of a chemical reaction [6,7]. The resulting hydrogen can either form blisters in the sub-surface region or gaseous methane in the interior. Precipitation of hydrogen as hydride in hydride forming elements such as zirconium and titanium [8,9] are other ex- amples where chemical reaction aids the hydrogen- induced embrittlement of the matrix. 2. Environmental embrittlement takes place in the hy- drogen containing atmospheres through adsorption of molecular hydrogen on the surface and its ab- sorption within the lattice after dissociation into atomic form [5]. 3. Internal hydrogen embrittlement (IHE), in contrast takes place in the absence of a hydrogenated atmo- sphere and is brought about by hydrogen which has * Corresponding author. Tel.: +91-22-5505050; fax: +91-22- 5505151. E-mail address: [email protected] (G.P. Tiwari) 0921-5093/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII:S0921-5093(00)00793-0

Transcript of A Study of Internal Hydrogen Embrittlement of Steels

Page 1: A Study of Internal Hydrogen Embrittlement of Steels

Materials Science and Engineering A286 (2000) 269–281

A study of internal hydrogen embrittlement of steels

G.P. Tiwari a,* , A. Bose b, J.K. Chakravartty a, S.L. Wadekar a, M.K. Totlani b,R.N. Arya a, R.K. Fotedar b

a Metal Physics Section, Materials Science Di6ision, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, Indiab Materials Processing Di6ision, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, India

Received 22 October 1999; received in revised form 2 February 2000

Abstract

A novel procedure for hydrogen charging and studying the Internal Hydrogen Embrittlement (IHE) of steels is described here.A cylindrical notched tensile sample with an extended end is employed for hydrogen charging. The extended portion of the sampleforms the cathode in an alkaline bath and a constant uni-axial tensile load is applied during hydrogen charging. The stressgradient set up by the notch, which is not in contact with the electrolyte, enhances the hydrogen concentration at various trappingsites of the matrix beyond the solubility limit. Subsequent to charging, the specimen is kept under the same load as that duringcharging, for another 24 h to stabilize the population of hydrogen within the specimen matrix. At the end of this stage, thespecimen is tensile tested to failure at room temperature. Two different steels namely maraging and mild steels have been chosento study the effect of hydrogen ingress on mechanical properties. While an increase in tangent modulus (linear portion of thestress–strain diagram), yield strength, work hardening rate and ultimate tensile stress (UTS) has been observed on hydrogenation,a decrease in total elongation has been noticed for both the steels studied. Fractographic investigation has revealed that thefracture mode is predominantly ductile dimple (failure by micro-void coalescence) in both the materials and that the hydrogenreduces the size of the dimples. The observations of this investigation are significant in two respects: firstly, it demonstrates theefficacy of a hydrogen charging method for steels which can introduce hydrogen to a level much higher than its solubility limitand secondly, it reports for the first time enhancement of modulus and work hardening by hydrogen charging. These observationshave been rationalized on the basis of current understanding on the effect of hydrogen on plastic properties and hypothesis of themodels of IHE. It is suggested that the trapping of hydrogen by dislocations and other structural features of the matrix and themutual interactions of their strain fields can account for the observed effects on yield strength, tangent modulus, work hardeningrate, UTS and ductility. © 2000 Elsevier Science S.A. All rights reserved.

Keywords: Internal hydrogen embrittlement; Steel; Trapping; Mechanical properties

www.elsevier.com/locate/msea

1. Introduction

The effects of hydrogen on the mechanical propertiesof iron and steels have been widely investigated [1–3].As regards the embrittling effect of hydrogen it is wellknown that the chief effects are a decrease in ductilityand true stress at fracture. Hydrogen embrittlement(HE) of steels can be classified into three main types[4,5].1. Hydrogen reaction embrittlement arises because of

the generation of hydrogen on the surface as a result

of a chemical reaction [6,7]. The resulting hydrogencan either form blisters in the sub-surface region orgaseous methane in the interior. Precipitation ofhydrogen as hydride in hydride forming elementssuch as zirconium and titanium [8,9] are other ex-amples where chemical reaction aids the hydrogen-induced embrittlement of the matrix.

2. Environmental embrittlement takes place in the hy-drogen containing atmospheres through adsorptionof molecular hydrogen on the surface and its ab-sorption within the lattice after dissociation intoatomic form [5].

3. Internal hydrogen embrittlement (IHE), in contrasttakes place in the absence of a hydrogenated atmo-sphere and is brought about by hydrogen which has

* Corresponding author. Tel.: +91-22-5505050; fax: +91-22-5505151.

E-mail address: [email protected] (G.P. Tiwari)

0921-5093/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved.PII: S0921 -5093 (00 )00793 -0

Page 2: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281270

entered the lattice during processing or fabricationof steel [10]. Having entered the lattice, hydrogenembrittles the steel over a period of time which is afunction of concentration, temperature and state ofstress within the matrix.

The details of the effect of hydrogen on mechanicalproperties and the mechanism of degradation are rea-

sonably well understood in the case of hydrogen reac-tion embrittlement. However, generally acceptedmechanisms for the other two classifications of HE,have not yet been established because of the contradic-tory experimental results obtained by various re-searchers [2,3,5]. This is more so in the case of IHE andone of the main reasons for the discrepancies in thereported results has been attributed to the possiblestructural damage caused by hydrogen charging whichmasks the intrinsic effect of hydrogen on mechanicalproperties [11,12].

This paper presents the results of experiments de-signed to elucidate the intrinsic effect of hydrogen onthe mechanical properties of ferritic steels by employinga new technique of hydrogen charging. Two differentsteels namely mild steel and maraging steel with widelydifferent strength levels have been chosen for this study.

2. Experimental procedure

As received materials in tempered condition havebeen employed in this study. The nominal compositionsof these two steels are given in Table 1.

The three essential steps in the experimental proce-dure in this study on internal hydrogen embrittlementare:1. The electrolytic charging of specimen with hydrogen

under uniaxial loading conditions.2. After charging hydrogen for a specified number of

hours, the charging cell is removed and the speci-men is held in a fixture, shown in Fig. 1, under apredetermined load for a period of 24 h.

3. Finally the specimen is removed from the chargingset-up and immediately strained up to fracture at anominal strain rate of 3×10−5 s−1 in an UniversalTesting Machine.

Fig. 2 depicts the geometry as well as the dimensionsof the specimen. One end of the specimen is extended.The hydrogen charging assembly is wall mounted asshown in Fig. 1. The extended portion of the tensilespecimen forms the cathode of the electrolytic cell usedfor hydrogen charging for various periods of time. Thecharging time was selected in such way that it was wellin excess of the estimated time required for hydrogen toreach the notched section of the loaded tensile samples.An electrolytic solution of 1 wt.% NaOH with 1 g l−1

As2O3 as recombination poison was used. A currentdensity of 1.5×10−2 A cm−2 was employed for hydro-gen charging. The specimen was held rigidly by itsthreaded sections. This arrangement prevented any lat-eral movement and also helped in maintaining theproper vertical alignment of the specimen. A deadweight was applied below the elastic limit of the mate-rial. This load was arrived at from the knowledge of the

Table 1Chemical composition of the steels used

Composition (wt.%)Element

Mild steel Maraging steel grade 350

C 0.12 0.005Si 0.25 –

0.40Mn –0.04S –0.04P ––Ni 18.39–Mo 3.99

Co – 12.32–Ti 1.63

Al – 0.12Fe BalanceBalance

Fig. 1. Schematic of fixtures used for hydrogenation and loadingmechanism.

Page 3: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281 271

Fig. 2. Schematic showing details of the specimen used for chargingand testing.

neous. The main difference in the method of hydrogencharging employed by GE with that used in this studylies in the fact that the test portion of the sample in ourexperiments is fed with hydrogen through diffusion.Secondly, holding the specimen under load (for a pe-riod of 24 h) subsequent to hydrogen charging, allowssufficient time for the hydrogen population to stabiliseitself within the matrix. Hence, the distribution ofhydrogen in the matrix is expected to undergo relativelysmall changes during the tensile testing, whereas inGE’s experiments, the concentration as well as thedistribution of hydrogen remains in a state of flux tillthe end of experiments.

An optical microscope was used to characterize themicrostructures of the starting material as well as of thehydrogenated material. Fractography of the deformedspecimens was carried out using a scanning electronmicroscope (SEM).

2.1. Estimation of hydrogen diffusi6ity andconcentration

Estimation of hydrogen diffusivity in the mild steelstock used in the present investigation was carried outusing the method suggested by Devanathan andStachurski [16]. In this technique, a specimen havingrectangular cross-section (approximately 1×1 cm) and1-cm thick is cathodically charged with hydrogen up tothe saturation limit. This saturation limit is reachedonly after hydrogen has occupied all the available siteson reversible as well as irreversible traps and the con-centration in the lattice has risen up to the solubilitylimit. Fig. 3(a) shows the current versus time plot forthe cathodic charging of the 1-cm-thick mild steel spec-imen. The electrolyte used was the same as that usedfor hydrogen charging. Open circuit potential was −235 mV (w.r.t. to SCE) and the charging was performedat −400 mV (w.r.t. SCE). The concentration of dif-fusible hydrogen in the sample controls the chargingcurrent and saturation is reached around 28 mA. Thisstate is equivalent to steady state diffusion so that allthe hydrogen entering the specimen goes at the otherend. According to Deluccia and Berman [17], the diffu-sivity, ‘D ’ of hydrogen in steel is related to tmax, thetime required to reach the saturation level, as

D=L2/(4tmax) (1)

where L is the thickness of the specimen. With the useof data in Fig. 3, D is estimated to be 3×10−4 cm2

s−1, on the basis of tmax=15 min. This diffusivity valueis in very good agreement with an earlier reported value[18].

In order to validate this estimation, an anodic oxida-tion (dehydrogenation) curve of the same specimen wasalso carried out as shown in Fig. 3(b). Prior to oxida-tion, a thin oxide layer was grown on both faces of the

elastic limit of the material and stress concentrationfactor for the notch used. The presence of a notch inthe centre creates a stress gradient within the specimen.The tendency of the hydrogen to migrate up the stressgradient facilitates hydrogen charging of the specimen[3,13,14]. The procedure outlined above thus closelysimulates IHE. Keeping the bare (uncoated) specimenunder load in the absence of the source of hydrogenwill reduce out-gassing of hydrogen and allow thehydrogen to be present in the lattice in excess of thesolubility limit. At the end of this period, most of thehydrogen present in the specimen in excess of solubilityis expected to diffuse out of the ferritic lattice andresidual hydrogen exists in strongly bonded trappingsites of the steel. Hence the embrittlement effects re-ported here are primarily due to the presence of hydro-gen in tightly bonded irreversible traps.

The procedure used for hydrogen charging employedhere is somewhat similar to that used by Groeneveldand Elsea [15]. These authors (GE) used a specimenwhich was extended at both ends and held under tensileload while charging. The charging was done electrolyti-cally as is the case here. However, the charging wasdone in the middle section of the specimen which wasused for the test. The specimen was held under load inthe electrolytic bath to determine the cracking time as afunction of load and charging conditions. Thus, in theirinvestigation, the charging of hydrogen into the speci-men and its trapping within the matrix were simulta-

Page 4: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281272

specimen to minimize the interfacial reaction whichmay otherwise influence the rate of oxidation. Notethat initial current is nearly the same as the saturationcurrent in Fig. 3a. The rate of anodic reaction isdiffusion-controlled and hence the time required forsaturation and complete dehydrogenation is nearly thesame.

The concentration of hydrogen in the sample at anytime ‘t ’ can be calculated from the anodic curve of Fig.3b and the use of the following formula again fromDeluccia et al. [17]

J/nF=Co{D/(p× tmax)}1/2 (2)

with J, anodic current; n, no. of electrons, Co, concen-tration of hydrogen and F, Faraday’s constant.

The value of C0 is calculated to be 0.9 ppm which isnearly the same as the solubility limit of hydrogen in

steel [19]. This is the concentration of diffusible hydro-gen in the sample. The estimations of diffusivity andaverage hydrogen content have helped to establish theefficacy of the charging method for introduction ofhydrogen into the steel specimens.

The flux equation for the diffusion of hydrogenunder combined action of concentration and stressgradient can be written [20,21] as

Jc= − (DC/RT) · [RT(d ln C/dx)−V*/3 · (ds/dx)](3)

where, Jc is the hydrogen flux, D is the diffusivity ofhydrogen in steel, C is the concentration of hydrogen atany point X, V* is the activation volume for diffusionof hydrogen in steel, s is the tensile stress at any pointand R and T have their usual meaning.

Differentiation of Eq. (3) yields

dC/dt=D [(d2C/dx2)+K(dC/dx)] (4)

where, K= (V*/3RT) · (ds/dx)If ds/dx is constant then K is also constant and the

solution of the second-order non-linear differential Eq.(4) can be found analytically [20] or can be solved byappropriate numerical methods. A solution of thisequation with proper boundary conditions can then beutilised to estimate the hydrogen concentration at anysection of the specimen. However, this has not beendone for two reasons. Firstly, the cross-section of thetensile specimen varies from place to place and henceEq. (4) cannot be applied directly to the present case.Secondly, as mentioned earlier, the specimen is heldunder load without the presence of electrolytic chargingcell. During this period, most of the diffusible hydrogenwill escape away from the specimen in the absence ofthe concentration gradient. Tensile loading preferen-tially activates the trapping sites whose axes are at rightangles to the load [22–24]. It is believed that all thehydrogen present in the sample is held in such traps forwhich Eq. (4) is not applicable in its present form.However, it is clear from these expressions that therewill be a considerable enrichment of hydrogen at thenotched section of the specimen.

In order to obtain an estimate of the hydrogencontent in the bulk at the time of fracture, vacuumfusion analysis has been carried out on the samplestaken from regions just below the fractured notchedportion of the specimen. An average value of 5 ppmhydrogen has been obtained in such samples and thisvalue is much higher than the solubility limit of hydro-gen in ferrite matrix at ambient conditions ($1 ppm).This value may be considered as the lower boundsolubility limit of hydrogen at irreversible trapping sitesbecause out-gassing of hydrogen after fracture and theactual hydrogen concentration will be much higherwithin the loaded sample before the commencement ofthe tensile test.

Fig. 3. (a) A plot of current versus time during cathodic charging ofmild steel specimen; (b) Anodic oxidation (dehydrogenation) curvefor mild steel.

Page 5: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281 273

Fig. 4. Stress–strain plot for notched mild steel specimens chargedwith hydrogen for different durations and an un-notched specimenwithout introduction of hydrogen.

ally, such tests also yield information on the deforma-tion behaviour of the material in the absence ofhydrogen. A dead weight of 4300 N (corresponding toapproximately 65% of the yield load) was chosen forhydrogen charging of the mild steel and this load set upstresses of 220 and 150 MPa at the notch and in theregion away from the notch, respectively. The corre-sponding stress values for maraging steel samples were250 and 175 MPa and were generated using a deadweight of 5000 N which was approximately 65% of theyield stress.

An inspection of Figs. 4 and 5 reveals that thehydrogen influences the deformation behaviour of mildsteel and the maraging steels in the following way.

3.1. Effecti6e modulus

The initial parts of stress–strain curves in Figs. 4 and5 represent the elastic deformation and the slope yieldsthe value of the tangent modulus (taken here as theslope of the initial linear portion of the stress–straincurve). It is clear from Tables 2 and 3, hydrogenenhances the magnitude of the tangent or effectivemodulus in both mild steel and maraging steel. In mildsteel, for charging times up to 13 h, there was nonoticeable change in the values of the effective mod-ulus. However, a significant increase in tangent mod-ulus was observed for longer charging times (e.g. for 22and 40 h). The increase in the tangent modulus ishigher for maraging steel compared to that in mild steelfor similar charging times. In mild steel, the increasesare 33 and 80%, respectively for 22 and 40 h of charg-ing and the corresponding increase in maraging steelfor 22 h of charging is about 61%. To our knowledge,this observation of the enhancement in tangent mod-ulus of steel due to the presence of hydrogen duringtensile deformation is new and has not been reportedbefore. The slope of the linear portion of the stress–strain plot is dependent on both Young’s modulus andthe micro-yielding process. In a polycrystalline material(as in the case of these two steels), the early stage ofdeformation (micro-yielding) is associated with yieldingin the isolated individual grains which are most fa-vourably oriented with respect to the direction of stress-ing and movement of dislocations in those grains fromthe source to slip barriers such as grain boundaries,precipitates, etc. At these barriers, the piled up disloca-tions create stress concentrations to activate slip in theadjacent grain. The process repeats till all the grains inthe load bearing section of the specimen yield at themacroscopic yield stress.

3.2. Loss of ductility and enhancement of strength

The most significant effect of the presence of hydro-gen on the deformation behaviour is the loss in the

Fig. 5. Stress–strain plot at room temperature for notched maragingsteel samples with and without hydrogen.

3. Results

The results of tensile tests on mild steel and maragingsteel are given in Figs. 4 and 5. Fig. 4 shows the tensiletest results for mild steel with and without notch. Formaraging steel, the tests were performed only onnotched specimens with and without hydrogen. Thepurpose of such tests was to estimate the limiting elasticload to be applied during hydrogen charging. Addition-

Page 6: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281274

ductility (or notch ductility) which decreases from avalue of 11.2%, for the notched and uncharged speci-men of mild steel to 5.7, 5.5 and 4.1% for 13, 22 and 40h of hydrogen charging, respectively. In percentileterms, this amounts to a decrease of 49–64%. The lossin ductility in maraging steel is even greater andamounts to nearly 69% for 22 h of charging comparedto uncharged material. The detailed analyses of tensiletest data of Figs. 4 and 5 are given in Tables 2 and 3.It is observed that hydrogen charging causes a largeincrease in the yield strength (160 and 75% increase inmild steel and maraging steel, respectively) compared touncharged condition. With increasing charging times,however the yield strength remains virtually unchangedin mild steel. In contrast to the effect of hydrogen onductility, the UTS (Notched Tensile Strength) values ofthe hydrogen charged specimen show a rising trendwith increasing charging times for both mild steel andmaraging steel as found in the case of effective modulusand yield strength. However, the increases in UTS inboth these materials are not as large as in the case ofeffective modulus. It must be mentioned here that anincrease in UTS should not be taken as an indication ofimprovement of toughness also. Since toughness of anotched specimen depends strongly on the notch ductil-ity, a reduction of more than 50% in this parameter asobserved here at longer charging times (22 h andhigher) indicates a strong embrittling effect ofhydrogen.

3.3. Enhancement of work hardening rate

Tables 2 and 3 show yet another new and interestingresult on the effect of hydrogen on the work hardeningrate. The work hardening rate, which is the averagework hardening rate in the plastic portion of the

stress–strain curve, is enhanced significantly in thepresence of hydrogen. For mild steel, the work harden-ing rate is found to increase from 5.6×103 MPa forun-hydrogenated state to 5.8×103, 6.2×103 and 1.2×104 MPa for 13, 22 and 40 h of hydrogen charging,respectively. The corresponding percentile increases are4, 11 and 116% compared to un-hydrogenated mildsteel material. For maraging steel, after 22 h of charg-ing, the work hardening rate jumps to 1.3×104 MPafrom the uncharged value of 5.9×103 MPa whichamounts to an increase of more than 100%.

3.4. Fracture mode

Results of fractographic investigations are shown inFigs. 6 and 7 for mild steel and maraging steel, respec-tively. It is observed that fracture modes of both un-notched (Fig. 6(a)) and notched (Fig. 6(b)) specimens inuncharged conditions and those of hydrogen chargedfor 13 h (Fig. 6(c)) are essentially ductile and fractureoccurs by micro-void coalescence (mvc). At longercharging time (22 h), the fracture mode remains pre-dominantly ductile in nature with some quasi cleaved/cleaved area (Fig. 6(d)). Although no quantitativemeasurements have been made, it is apparent that thedimple size is larger in the uncharged samples com-pared to that in the charged samples. In addition tothis, dimple size was found to decrease with increasingcharging time. A similar general trend in fracture modeof maraging steel has also been observed as shown inFig. 7(a,b). It is noted that the failure in this material isalso by micro-void coalescence and the presence ofhydrogen decreases the dimple size. However, it mustbe mentioned that damage such as cracking (or de-co-hesion) at the particle–matrix interface and grainboundaries have not been seen even at longer charging

Table 2Effect of hydrogen on mechanical properties of mild steel

Effective modulus Total elongationWork hardening rate UTSSpecimen(MPa, ×103)(GPa) (MPa) (%)

11 3.1Without notch and without hydrogen 652 13.311.5 5.6Notched and without hydrogen 845 11.2

5.711.5 880Notched and 13 h hydrogen 5.815.3 6.2Notched and 22 h hydrogen 905 5.5

Notched and 40 h hydrogen 20.6 12 935 4.1

Table 3Effect of hydrogen on mechanical properties of maraging steel

Effective modulusSpecimen Work hardening rate Total elongationUTS(MPa, ×103) (MPa)(GPa) (%)

13.2 19.35.9Notch without hydrogen 124721.2 6.113.0Notch and 24 h hydrogen 1340

Page 7: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281 275

Fig. 6. Scanning electron micrograph of the surface of mild steel samples pulled to failure in tension at room temperature; (a) without notch andwithout hydrogen; (b) with notch and without hydrogen; (c) with notch and with hydrogen (13 h charging) and (d) with notch and with hydrogen(22 h charging).

times. The results of fractography of notched samplespresented here, correspond to the areas near the notchtip and the fracture at region away from the notch (i.e.near the centre of the specimen) occurred essentially bymvc. The effects of hydrogen on fracture mode and onthe size of the dimples are indicative of the fact thathydrogen influences the local plasticity to a consider-able extent. The results of fractographic analysis sub-stantiates the observed effects of hydrogen on ductilitygiven in Tables 2 and 3. A reduction in the dimple sizewith associated loss of ductility on hydrogen charginghas been reported earlier in the case of IHE in variousmaterials by several authors [25,26].

4. Discussion

This investigation suggests a new method of charginghydrogen and studying the intrinsic effect of hydrogenon mechanical properties of steel as during IHE. Hy-drogen is charged into the test specimen by stressassisted diffusion of hydrogen generated remotely byelectrolytic charging. Hydrogen charged in this manneris not expected to cause any damage to the microstruc-ture as the test piece is not in direct contact with theelectrolyte and is thus not exposed to high fugacityhydrogen that prevails during cathodic charging. InIHE, it is believed that hydrogen having once enteredthe lattice remains dormant till the ambient conditionsallow its transport (by diffusion) to the stressed region.

From mechanistic considerations, the two processeswhich control the influence of hydrogen on mechanicalproperties are its rate of diffusion and its interactionwith the trapping sites within the matrix [25,27,28]. Thedesign of the present experiment is such that it allowsto some extent to separate these two processes. Thepresence of stress accentuates the rate of diffusion andthe extended charging periods allow sufficient time fortrapping events to occur. Subsequently, the holding ofthe specimen under load without the presence of thesource of hydrogen, stabilises the population of hydro-gen atoms within the matrix. Hence, the diffusion ofhydrogen, per se, has little effect on the deformationcharacteristics of hydrogen charged steels reported here.The change in mechanical properties is mainly due tothe changes in the properties of the structural featuressuch as dislocation, particle–matrix interface, grainboundaries and magnetic domain walls brought aboutby their association with hydrogen.

4.1. New results on IHE and hydrogen trapping

4.1.1. Summary of the resultsThe present investigation has found that for the steels

studied, the intrinsic effect of hydrogen is to increasethe tangent modulus, strength and work hardening rateand to decrease the elongation. The results alsostrongly suggest that hydrogen causes hardening insteels. Such hardening effects in ferritic steels by hydro-gen have been reported by many workers [2,3,29,30].

Page 8: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281276

The effects of hydrogen on ductility and work harden-ing are also consistent with earlier reported research[3,11]. However, the observations made on the effect ofhydrogen on the tangent modulus and UTS are noveland significant and have not been reported before. Theincrease in UTS as found in this study is unusual and isapparently at variance with the generally observed ef-fect of hydrogen which is known to reduce the fracturestress in ferritic steels [1–3,12,14]. The most interestingobservation of this study is the significant increase ofthe tangent modulus of both the steels on hydrogena-tion. This is opposite to the observation made byBastein and Azou [31] who reported that Young’smodulus was unaffected by hydrogenation. As regardsthe mode of fracture, it is found that hydrogen does notalter the fracture mode which is predominantly trans-granular ductile failure. However, a hydrogen-inducedreduction in dimple size is noted and the origin of thesedimples is invariably the particle–matrix interfacerather than the grain boundaries as no intergranularfailure has been observed. The participation of smallercarbide/precipitate particles as void initiation sites ap-pears to be responsible for smaller dimple size and

lower ductility in the case of hydrogenated materials[5,25].

4.1.2. Hydrogen trapping in steelsThe summary of the results obtained here suggests

that the reason for the unusual effects of hydrogenstems from the accumulation of a high concentration ofhydrogen within the matrix through the new method ofhydrogen charging. It is known that a large amount ofhydrogen (much higher than the equilibrium concentra-tion) can be retained at ambient temperature in ferriticmatrix because of the presence of numerous structuraldefects which trap hydrogen[14,32,33]. Pressuoyre hasclassified trapping sites in terms of their physical naturein iron and steels and suggested that vacancy, alloyingelements, dislocations, micro-voids and interfaces arepossible trapping sites [32]. In steels, the major trappingsites are found to be grain boundaries, sub-grainboundaries, iron carbide and alloy carbides [32]. It hasbeen demonstrated by Podgurski and Oriani [33] thateven structural features like insoluble impurities cansustain hydrogen concentration far beyond the solubil-ity limit. As regards the effect of stress, in addition toenhancing local solubility at a trapping site, it maymodify trapping characteristics of a structural feature.For example, the grain boundary, in general, has amuch higher trapping efficiency compared to an inter-face in steels. However, if the grain boundaries areunfavourably oriented with respect to the tensile stressdirection, the interfaces may act as more efficient trapsfor hydrogen [22].

From the preceding section, it is clear these factorsare expected to increase the net amount of trappedhydrogen in steels containing a high density of struc-tural defects as used in this investigation and that weare dealing with samples having significantly higheramounts of hydrogen (much greater than 5 ppm). Fur-ther, it is logical also to infer that a high density oftrapped hydrogen cannot exist in isolation and hencethere must be some hydrogen present in the lattice aswell. A dynamic equilibrium of the concentration ofhydrogen should then exist between the lattice andtrapping sites. The ratio of concentration level at thesesites will depend upon the relative binding energy for aproton.

4.1.3. Effect of trapped hydrogen on mechanicalproperties

The present results are most satisfactorily interpretedif hydrogen or hydrogen–solute complex interacts withdislocations to immobilize them and interferes with thenucleation of dislocations from their possible sources.The hydrogen in the lattice, apart from interacting withdislocations by creating dilatational field [3], can mod-ify the characteristics of dislocations in a variety ofways. Earlier studies on internal friction of hydro-

Fig. 7. Scanning electron micrograph of the surface of maraging steelsamples pulled to failure in tension at room temperature; (a) notchedsamples without hydrogen and (b) notched samples with hydrogen(22 h charging).

Page 9: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281 277

genated iron and steels suggest that dislocations aretrapping sites for hydrogen and hydrogen atoms arevery effective in anchoring them at ambient tempera-tures [34–36]. The trapped hydrogen induces elasticdipoles and mutual interactions of dipoles centred ondislocations and other structural defects increases thelevel of back stresses in the lattice. Ke [37] has shownthat some hydrogen lodged at dislocations resides inmolecular form which is also expected to reduce theirmobility. Kornmuller, from magnetic relaxation studieson the interaction between hydrogen and magneticdomain walls concluded that the presence of hydrogenin domain walls enhanced the strength of the elasticstress field experienced by neighbouring dislocations[38].

In view of the factors mentioned above, the increasein tangent modulus as observed on hydrogenation isobvious as hydrogen is likely to impede glide of individ-ual dislocations during the micro-yield and also preventgeneration or unlocking of dislocations from its sourcein the adjacent grains. In transition metals like nio-bium, vanadium and titanium, hydrogen in solution hasbeen found to increase as well as decrease Young’smodulus [39]. We admit that our experimental result donot produce any evidence to include explicitly the con-tribution of hydrogen to enhance Young’s modulus.However, it is apparent from Figs. 4 and 5 that thetangent modulus (the slope of the initial portion ofstress–strain diagram) which is dependent on bothYoung’s modulus and micro-yielding has increased. Itis pertinent to note that Laurent et al. [40] observed avery large concentration (about 600 atoms per A, ofdislocation) of trapped hydrogen in the interface anddislocations of hydrogenated Armco iron by using highresolution electron autoradiography. A much largerconcentration of hydrogen than this is expected in suchtraps in stressed condition as in this investigation. Thestate of the lattice surrounding these traps and theirpossible effect on mechanical properties is not yetknown.

With regard to the effect of hydrogen on yieldstrength, it was suggested by Oriani et al. [41] thatsolute hydrogen can increase the yield or flow stress ofsteels by hydrogen drag upon moving dislocation orimpeding cross-slip. The hardening is also expectedfrom the fact that trapped hydrogen may make disloca-tion nucleation difficult from grain boundary disloca-tions and other possible sources. Further it is knownthat multiple slip will be favoured in a structure con-taining second phase/carbide particles because of elasticand plastic incompatibility effects. Hirth [3], from areview of a large number of studies, has inferred thathardening occurs when deformation is by multiple slip(at relatively large strain) and that when work harden-ing is large and the dislocations form a complex array.It is known that a relatively large strain prevails ahead

of a notch at general yield and the yield stress for thenotched sample reported in this work thus correspondsto a strain of approximately 5% (estimated from notchgeometry used). Therefore, in the investigated steelscontaining various carbides and precipitates, hardeningby hydrogen is expected.

The increased work hardening rate in the hydro-genated material relative to the hydrogen-free materialcan be explained by noting that hydrogen immobilizesdislocations and adversely affects the recovery processesby inhibiting cross slip. As a consequence, there will bebuild up of a high density of immobile dislocations inthe form of loops, tangles, piled up dislocations atcarbide/second phase particle etc. which are going toimpede glide dislocations. With increasing deformation,the mean free path of dislocations in the matrix willdecrease progressively and increasingly higher force willbe required to drive dislocations through these barriers.

AS UTS is controlled by both yield strength andwork hardening, an increase in UTS on hydrogenationis thus apparent. The observed increase in UTS, how-ever, in both the steels by hydrogen is not as effective asin the case of yield strength. This observation is com-mon in any strengthening mechanism (e.g. dynamicstrain ageing, solid solution strengthening, etc.) where arelative change in yield strength is usually larger thanthat in UTS and a consequence of this strengthening isloss of ductility. Similar observations have been madeby West et al. [27] on the effect hydrogen on strengthand ductility in 21-6-9 stainless steels. It has beenpointed out earlier that an increase in UTS should notbe taken as an indicator of increase in toughness onhydrogenation. The loss of ductility or embrittlementby hydrogen results from the attainment of criticalstress for tensile instability at lower strain compared tohydrogen free material. Similar observations were alsomade in an earlier study by West and Louthan [27].Although the steels selected for this study are verydifferent in terms of their chemical composition andmechanical properties, the changes in mechanical prop-erties brought about by hydrogen are strikingly similar.However, the effect of hydrogen in enhancing strengthand modulus and decreasing ductility are greater inmaraging steel compared to that in mild steel for thesame charging time. This is expected as maraging steelis capable of dissolving large amounts of hydrogen bytrapping because of the presence of a number of alloy-ing elements besides titanium which have a strongaffinity for hydrogen [42,43]. Further the maraging steelhas a much finer microstructure and contains a largervolume fraction of second phase particles with a muchsmaller inter-particle spacing than exists in the mildsteel of this study. Therefore a larger increase instrength and work hardening rate is expected in marag-ing steel on hydrogenation.

Page 10: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281278

We also note that the effect of hydrogen on deforma-tion characteristics is augmented directly in proportionto its concentration on dislocations and other defectsdemonstrating that the observed effects of hydrogen arereal and intrinsic. The relatively large changes instrength, tangent modulus, work hardening and ductil-ity are related to the high levels of hydrogen trapped inthe matrix by the conditions of the experiment.

It may be useful at this stage to identify the mi-crostructural features that control the plastic propertiesand fracture in hydrogenated steels. It is known thatthe binding energy of hydrogen atom to the dislocationis smaller than particle/matrix interfaces and these in-terfaces in turn provide traps or hydrogen which areless efficient than grain boundaries. This arises primar-ily from the fact that the grain boundary has a muchhigher energy than the interfacial energy of the particle/matrix interface and thus provides a deeper trap. How-ever, the observations that fracture occurred afterconsiderable macroscopic plastic deformation and theabsence of grain boundary fracture in hydrogencharged steels in this investigation suggest that thestress concentration produced by dislocations duringdeformation are not sufficient to initiate a crack atgrain boundaries. Further, the particle/matrix interfaceor dislocations acting alone cannot explain the mode offracture on hydrogenation. Therefore, it is reasonableto conclude that the interaction of dislocations or dislo-cation arrays with carbide/precipitate particle controlsplastic deformation and builds up stress concentrationslocally to initiate fracture in these steels. The efficacy ofthe stressed interface in trapping hydrogen has beendemonstrated by Ohma et al. [22].

4.2. Mechanisms of hydrogen embrittlement

To date various models have been put forward toexplain the mechanisms of hydrogen embrittlement;however no single theory can fully explain the mecha-nism of degradation of mechanical properties of steelsby hydrogen [44,45]. In the following we take up a fewrelevant mechanisms to examine the possible reasonsfor embrittlement.

4.2.1. Precipitation of hydrogen at internal defectsWhen the hydrogen atoms distributed in the lattice

precipitate as gaseous hydrogen into pre-existing micro-pores/voids, the pressure exerted by the gas adds to theexternally applied tensile load on the specimen andconsequently the effective fracture stress is reduced [46].

4.2.2. De-cohesion modelTroiano [47] suggested that hydrogen dissolved in the

regions of high stress in a lattice could reduce thestrength of inter-atomic bonds. Oriani [48] further de-veloped Troiano’s concept and suggested a mechanistic

theory of lattice de-cohesion by hydrogen. These mod-els basically propose that dissolved hydrogen in thelattice can reduce the stress needed to initiate or propa-gate a crack by lowering the cohesive strength of atoms.Gerberich and coworkers [49] examined in detail therole of the state of stress on hydrogen-induced cracking(HIC) in the region of a crack tip. In order to quantifythe effect of state of stress, they modeled the crack tipbehaviour for HIC by considering small-scale yieldingand discretised dislocation computer simulation ofstress field. They further hypothesized that the mostsevere micro mechanical conditions for HIC prevailedunder plane strain condition with mode I loading whichwas in accordance with Troiano–Oriani’s de-cohesionmodel.

4.2.3. Reduction of surface energy adsorptionAdsorption of hydrogen on the surface reduces the

surface energy, hence the total energy required for thefracture is lowered [50]. Just as in the case of theprecipitation mechanism mentioned above, this mecha-nism also requires the presence of micro-porosity.

4.2.4. Association of hydrogen with dislocationsBastien and Azou [51] suggested that embrittlement

can result from segregation of hydrogen forming Cot-terell atmosphere around dislocations during plasticdeformation. This segregated hydrogen atom in a cavitycan lead to the formation of a void containing molecu-lar hydrogen gas which under pressure would cause thematrix to rupture. Based on this concept, Tien et al.[52] developed the dislocation–transport model of hy-drogen embrittlement. An enrichment of hydrogen inplastically deformed region is predicted when disloca-tions lose their hydrogen atmosphere during straining.A slightly different model was proposed by Louthan etal. [53] who suggested that a high concentration ofhydrogen might result from dislocation–hydrogen in-teraction which could modify the processes of bothplastic deformation and fracture. In contrast, Beachem[54] proposed that the basic hydrogen–dislocation in-teraction in steels is one of easing motion of disloca-tions or generation of glide dislocations or both therebyinducing lattice softening. An accumulation of sufficientconcentration of hydrogen in the lattice just ahead of acrack tip is assumed to aid the prevailing deformationprocesses [54]. By noting similarities among hydrogenassisted cracking (HAC), stress corrosion cracking andadsorption induced liquid metal embrittlement for vari-ous alloys, Lynch [55] has suggested that HAC involvesadsorption of hydrogen at the crack tip. In Lynch’smodel adsorbed hydrogen is expected to weaken theinter-atomic bonds thereby facilitating dislocation nu-cleation at the crack tip. This model is somewhatsimilar to Beachem’s but differs in choosing the loca-tion of the effect of hydrogen. Further evidence of

Page 11: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281 279

hydrogen–dislocation interactions has been providedby Birnbaum and coworkers [56–58] by carrying out insitu deformation and fracture in TEM environmentalcell containing hydrogen gas. These results togetherwith the measurements of thermally activated deforma-tion parameters associated with dislocation motion,following the trend of macroscopic stress–strain be-haviour and theoretical treatment of hydrogen anddislocation interactions prompted these researchers tosuggest a hypothesis for rationalising the hardening–softening behaviour on hydrogenation. According tothis hypothesis, the presence of hydrogen in solid solu-tion decreases the barriers to dislocation motionthereby increasing the amount of deformation thatoccurs in a localized region adjacent to the crack sur-face. This hydrogen enhanced localized plasticity(HELP) has been attributed to the presence of a hydro-gen atmosphere around dislocations which shields itsinteraction with other elastic centres such as disloca-tions, solute atoms, etc. Although this mechanism sug-gests that the fracture process is initiated by a highlylocalized plastic failure at the microscopic level, theembrittlement manifests macroscopically by loweringductility as during a tensile test. These researchers havealso reported occurrence of HELP in a wide variety ofmaterials having different crystal structures and haveasserted that this mechanism is able to explain ingeneral the effect of hydrogen on deformation (bothelastic and plastic) and fracture.

Several investigators have reported the occurrence ofhydrogen embrittlement in hydrogen environment at apressure of 1 atm or less [6,10]. In these conditions, thepressure of hydrogen within the micro-voids or greaterthan the ambient cannot be sustained on thermody-namic considerations. In general, the decrease in UTSof steels because of hydrogen is far greater than 1 atm.When the hydrogen charged tensile specimens are testedin air, any significant amount of pressure induced bythe presence of hydrogen in the matrix cannot besustained within the micro-pores present in the matrix.Such results do not support the gaseous pressure in-duced model. Precipitation of hydrogen at internaldefects to explain IHE is inconsistent with the presentresults as well since we observe an increase in UTS onhydrogenation.

As regards the Petch–Stables model of surface ad-sorption [50], the hydrogen present in the lattice can, inprinciple, migrate to the pre-existing microvoid/pores,and lower the surface energy. Such microvoids, actingas sites for crack initiation can indeed lower the rupturestrength. The fractographic features of both hydro-genated and hydrogen-free materials of this study areessentially ductile and involve considerable plastic de-formation before fracture and thus the process of hy-drogen adsorption is not responsible for hydrogenembrittlement. Next, we consider the de-cohesion

model in which the decrease in rupture strength isbrought about by the weakening of lattice bonds. Theabsence of cleavage fracture in hydrogen charged indi-cates that hydrogen is not affecting the lattice energy.The dislocation transport theory [52] suggests that dis-locations are capable of generating a large concentra-tion of hydrogen in excess of lattice solubility in somemicrostructural features. The excess hydrogen may pre-cipitate as high pressure gaseous hydrogen in micro-voids and aid void growth which leads toembrittlement. Louthan et al. [53] also showed thatlocalized hydrogen concentration as high as 104 timesthe equilibrium value may be realised in micro cracksproduced by dislocation interactions. This high concen-tration of hydrogen may create internal gas pressure ofthe order of 106 atm or more which is large enough tocause embrittlement by pressure expansion mechanism.

If the micro-voids/cracks get activated in the pres-ence of hydrogen as postulated in both these mecha-nisms, then a reduction in the stress for initiation andgrowth of cracks is expected. In such a situation, UTSshould decrease with increasing hydrogen content. Incontrast to this, we note that UTS increases withincreasing charging time (i.e. increasing hydrogen con-tent). Thus, it appears that both these models do ex-plain the results of this study. For similar reasons, ourobservations on the effect of hydrogen on strength andwork hardening also preclude Beachem’s [54] andLynch’s [55] models as possible mechanisms of embrit-tlement. The HELP mechanism of hydrogen embrittle-ment postulates that hydrogen in solution in the latticeenhances mobility of dislocations by decreasing barriersto dislocation motion and the fracture is a highlylocalized plastic failure process rather than an embrit-tlement [56,57]. It must be mentioned here that this isthe only mechanism which attempts to rationalize theeffects of hydrogen on elastic, plastic and fractureproperties on the basis of the interaction of hydrogenwith dislocations. It has been shown that H-inducedslip localization may also lead to an increase in flowstress depending on the conditions of hydrogen charg-ing and tests variables employed during deformation.The observed effects of hydrogen on the strength, duc-tility and fracture process of the steels studied appear tobe in accordance with the proposition of HELP mecha-nism. However, the increase in the tangent modulusand the work hardening rate by hydrogen as observedin this investigation cannot be explained if hydrogendecreases the interaction energy of dislocation withother defects. A decrease in interaction energy of dislo-cations due to solute hydrogen would decrease theelastic modulus [58] and intuitively decrease the rate ofwork hardening also. Thus the suggestion that hydro-gen facilitates motion of dislocation as described in theHELP model is incongruous and against the increase intangent modulus and work hardening reported in this

Page 12: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281280

work. In addition, a decrease rather than an increase indimple size of both the steels on hydrogenation furthersubstantiates our view that lattice softening as well asHELP are not supported by our results on the effects ofhydrogen on deformation behaviour of the two com-mercial steels. It is difficult to arrive at firm conclusionsregarding the exact role of hydrogen in the embrittle-ment of these steels on the basis of the present results.However, the combination of the effect of hydrogen onmechanical properties and fractographic informationdoes imply that the intrinsic effect of hydrogen is toenhance strength and decrease ductility by impeding themotion of dislocations.

From a review of a large number of studies, Louthanet al. [28] have inferred that although there is noagreement among the investigators over the mecha-nisms of hydrogen embrittlement, all mechanisms sug-gest accumulation of a highly localized hydrogenconcentration as a prerequisite for the embrittlementprocess. A similar observation has also been made byPressouyre [32] and he proposed the trap theory whichin some respects attempted to unify various other theo-ries. According to this theory, regardless of the mecha-nism of embrittlement, a crack is initiated or its growthis assisted when the concentration of hydrogen trappedin a pre-existing stressed defect exceeds some criticalvalue. This theory can be used for rationalizing theobservations of this study. The important defects in thiscontext are dislocations, particle–matrix interfaces, thegrain boundaries and magnetic domain walls [59,60].The association of hydrogen with these structural fea-tures influences deformation behaviour and fracture asdiscussed earlier. The results of this investigation clearlysuggest that some critical stress concentration besides acritical hydrogen concentration must be achieved atsome trapping sites in order to nucleate a void for theductile fracture. In this study, such stress concentrationis achieved primarily at the particle–matrix interface bydislocation pile up. As all the defects present in themicrostructure trap a considerable amount of hydrogenbefore commencement of the tensile test, further enrich-ment by large-scale dislocation transport appears to beof only secondary importance in the present context.However, this discussion is somewhat speculative at thisstage and further work is needed to confirm this. Thusthe observed effects of hydrogen on these two steelsmay be construed as demonstrating that the increase instrength, work hardening and modulus and loss ofductility were due to the presence of a large concentra-tion of trapped hydrogen in various types of defectsintroduced by the novel method of charging. Further-more, an interval of 24 h between charging and tensiletesting allows the hydrogen to distribute itself at vari-ous trapping sites within the matrix. Hence large-scaleredistribution of hydrogen through lattice diffusionmay not be taking place during deformation andfracture.

5. Conclusions

A novel technique for charging hydrogen has beendemonstrated. This method permits hydrogenation to avery high level without causing damage to the mi-crostructure and facilitates studying the intrinsic effectof hydrogen on the mechanical properties of steel.

The intrinsic effect of hydrogen on both the steelsstudied are the enhancement of yield strength, ultimatetensile strength, tangent modulus, work hardening rateand reduction in ductility. The fracture mode was duc-tile dimple and hydrogen was found to reduce dimplesize compared to hydrogen-free materials.

The changes in mechanical properties were primarilydue to the presence of a large concentration of hydro-gen at various traps like dislocations, grain boundariesand matrix/carbide or precipitate interfaces and theinteraction of hydrogen with these traps.

The results of this investigation could be rationalizedby invoking trap theory and it appears that a criticalconcentration of hydrogen should be available at somestressed defect to initiate a micro-void/crack. In thisinvestigation particle/matrix interfaces have been iden-tified as possible locations of crack initiation.

References

[1] A.R. Troiano, Trans. Am. Soc. Metals 52 (1960) 54.[2] R.A. Oriani, Annu. Rev. Mater. Sci. 8 (1978) 327.[3] J.P. Hirth, Metal. Trans. 11A (1980) 861.[4] H.H. Gray, ASTM STP 543 (1974) 3.[5] M.R. Louthan Jr, in: I.M. Bernstein, A.W. Thompson (Eds.),

Hydrogen in Metals, ASM, Metals Park, OH, 1974, p. 53.[6] H.H. Johnson, in: I.M. Bernstein, A.W. Thompson (Eds.), Hy-

drogen in Metals, ASM, Metals Park, OH, 1974, p. 35.[7] D.A. Westphal, F.J. Worzala, in: I.M. Bernstein, A.W. Thomp-

son (Eds.), Hydrogen in Metals, ASM, Metals Park, OH, 1974,p. 79.

[8] D.O. Nothwood, U. Kosasih, Int. Metal. Rev. 28 (1985) 3323.[9] D.L. Davidson, D. Eylon, Metal. Trans. 11A (1980) 837.

[10] I.M. Bernstein, R. Garber, G.M. Pressouyre, in: A.W. Thomp-son, I.M. Bernstein (Eds.), Effect of Hydrogen on Behaviour ofMaterials, TMS-AIME, New York, 1976, p. 37.

[11] J.K. Lin, R.A. Oriani, Acta Metall. 31 (1983) 1077.[12] Y. Tobe, W.R. Tyson, Scripta Metall. 11 (1977) 849.[13] C. St. John, W.W. Gerberich, Metal. Trans. 4A (1973) 589.[14] I.M. Bernstein, Mater. Sci. Eng. 6 (1970) 1.[15] T.P. Groeneveld, A.R. Elsea, ASTM STP 543 (1974) 11.[16] M.A.V. Devanathan, Z.O.J. Stacheurski, Proc. R. Soc. A270

(1962) 90.[17] J.J. Deluccia, D.A. Berman, ASTM STP 227 (1981) 256.[18] J. Volkl, G. Alefeld, in: G. Alefeld, J. Volkl (Eds.), Hydrogen in

Metals I (Basic Properties), Springer, Berlin, 1978, p. 321.[19] M.L. Hill, E.W. Johnson, Trans. AIME 221 (1961) 622.[20] J.L. Waisman, G. Sines, L.B. Robinson, Metal. Trans. 4 (1973)

291.[21] C.E. Ells, C.J. Simpson, in: I.M. Bernstein, A.W. Thompson Jr

(Eds.), Hydrogen in Metals, ASM, Metals Park, OH, 1974, p.345.

Page 13: A Study of Internal Hydrogen Embrittlement of Steels

G.P. Tiwari et al. / Materials Science and Engineering A286 (2000) 269–281 281

[22] G. Ohma, G.P. Tiwari, Y. Injima, K. Hirano, in: Hydrogen inMetals (JIMIS–2), Minakami, Japan, 1979, p. 229.

[23] H.A. Wriedt, R.A. Oriani, Acta Metall. 18 (1970) 753.[24] J.C.M. Li, Y.T. Chou, Trans. AIME 245 (1969) 607.[25] J.M. Hyzak, D.E. Rawl Jr, M.R. Luthan Jr, Scripta Metall. 15

(1981) 937.[26] A.W. Thompson, Mater. Sci. Eng. 14 (1974) 253.[27] A.J. West Jr, M.R. Louthan Jr, Metall. Trans. 13 (1982) 2049.[28] M.R. Louthan, M.J. Morgan, J. Nondestructive Eval. 15 (1996)

113.[29] R.A. Oriani, P.H. Josephic, Metal. Trans. 11A (1980) 1809.[30] H. Kimura, H. Matsui, T. Kimura, in: Hydrogen in Metals

(JIMIS–2), Minakami, Japan, 1979, p. 533.[31] P. Bastein, P. Azou, Compd. Rond. 231 (1959) 147.[32] G.M. Pressouyre, in: C.G. Interrante, G.M. Pressuouyre (Eds.),

Current Solutions to Hydrogen Problems in Steels, ASM, MetalsPark, OH, 1982, p. 18.

[33] H.H. Podgurski, R.A. Oriani, Metal. Trans. 3 (1972) 2055.[34] A.J. Kumnick, H.H. Johnson, Acta Metall. 18 (1980) 33.[35] P. Schiller, Nuovo Cimento 276 (1976) 3313.[36] R. Gibala, Trans. AIME 239 (1967) 1574.[37] T.S. Ke, in: Hydrogen in Metals (JIMIS–2), Minakami, Japan,

1979, p. 573.[38] H. Kornmuller, in: G. Alefeld, J. Volkel (Eds.), Hydrogen in

Metals, Topics Appl. Phys., 28 (1978) 289.[39] C. Wert, O. Buck, in: Hydrogen in Metals (JIMIS–2), Mi-

nakami, Japan, 1979, p. 221.[40] J.P. Laurent, G. Lapasset, M. Aucouturier, P. Lacombe, in:

Hydrogen in Metals (JIMIS–2), Minakami, Japan, 1979, p. 559.[41] R.A. Oriani, P.A. Josephic, Metal. Trans. 11A (1980) 1809.

[42] R.D. McCright, in: Proc. Conf. Stress Corrosion Cracking andHydrogen Embrittlement of Iron Base Alloys, NACE-5 Unieux-Firminy, France, 1973, p. 306.

[43] G.M. Pressouyre, I.M. Bernstein, Acta Metall. 27 (1979) 89.[44] J. Smith, in: N.R. Moody, A.W. Thompson (Eds.), Hydrogen

Effects on Materials Behaviour, Met. Soc. AIME, Warrendale,PA, 1990.

[45] H.K. Birnbaum, Proc. Conf. Environment-Induced Cracking ofMetals, NACE, Houston, TX, 1990, p. 21.

[46] C.A. Zapffe, A.G. Moore, Trans. AIME 154 (1943) 355.[47] A.R. Troiano, Trans. ASM 52 (1960) 54.[48] R.A. Oriani, Ber. Bunsenges Phys. Chem. 76 (1972) 848.[49] X. Chen, T. Foecke, M. Lii, Y. Katz, W.W. Gerberich, Eng.

Fract. Mech. 35 (1990) 997.[50] N.J. Petch, P. Stables, Nature 169 (1952) 842.[51] P. Bastien, P. Azou, in: Proceedings of the 1st World Metallurgi-

cal Congress, ASM, Cleveland, OH, 1951.[52] J.K. Tien, A.W. Thompson, I.M. Bernstein, R.J. Richards,

Metal. Trans. 7A (1976) 821.[53] M.R. Louthan, G.R. Caskey Jr, J.A. Donovan, D.E. Rawl Jr,

Mater. Sci. Eng. 10 (1972) 357.[54] C.D. Beachem, Metall. Trans. 3 (1972) 437.[55] S.P. Lynch, Acta Metall. 36 (1988) 2639.[56] H.K. Birnbaum, P. Sofronis, Mater. Sci. Eng. A176 (1994) 191.[57] P. Sofronis, H.K. Birnbaum, J. Mech. Phys. Solids 43 (1995) 49.[58] P.J. Ferreira, I.M. Robertson, H.K. Birnbaum, Acta Mater. 46

(1998) 1749.[59] R.A. Oriani, Acta Metall. 18 (1970) 147.[60] M. Kurkela, G.S. Frankel, R.M. Latanision, S. Suresh, R.O.

Ritchie, Scripta Metal. 16 (1982) 455.

.