Corrosion Science 77 (2013) 210–221

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Corrosion Science

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Effect of low temperature on hydrogen-assisted crack propagationin 304L/308L austenitic stainless steel fusion welds

0010-938X/$ - see front matter � 2013 Elsevier Ltd. All rights reserved.http://dx.doi.org/10.1016/j.corsci.2013.08.004

⇑ Corresponding author. Present address: Structural Integrity Associates, Inc., SanJose, CA 95138, USA. Tel.: +1 408 833 7201.

E-mail address: hjackson@structint.com (H.F. Jackson).

H.F. Jackson ⇑, C. San Marchi, D.K. Balch, B.P. SomerdaySandia National Laboratories, Livermore, CA 94550, USA

a r t i c l e i n f o

Article history:Received 5 February 2013Accepted 5 August 2013Available online 11 August 2013

Keywords:A. Stainless steelB. SEMC. Hydrogen embrittlement

a b s t r a c t

Effects of low temperature on hydrogen-assisted cracking in 304L/308L austenitic stainless steel weldswere investigated using elastic–plastic fracture mechanics methods. Thermally precharged hydrogen(140 wppm) decreased fracture toughness and altered fracture mechanisms at 293 and 223 K relativeto hydrogen-free welds. At 293 K, hydrogen increased planar deformation in austenite, and microcrackingof d-ferrite governed crack paths. At 223 K, low temperature enabled hydrogen to exacerbate localizeddeformation, and microvoid formation, at austenite deformation band intersections near phase bound-aries, dominated damage initiation; microcracking of ferrite did not contribute to crack growth.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Austenitic stainless steels are used extensively in the nuclearpower, chemical processing, and oil and gas industries, in applica-tions where they are susceptible to stress corrosion cracking (SCC)[1–3]. SCC begins with an electrochemical process, such as localanodic dissolution, however synergistic processes of anodic disso-lution and cathodic hydrogen damage are thought to contribute toSCC in stainless steels [4,5]. Hydrogen damage is typically charac-terized by a significant loss of tensile ductility. In various alloy/solution systems, interactions between dislocations at a crack tipand corrosion products, such as hydrogen, can play a key role inSCC. Detailed fractographic and microstructural studies supportthe premise that adsorbed or absorbed hydrogen promotes crackgrowth by localized plastic flow [3,6–10]. This suggests that effectsof a corrosive environment and sustained tensile stress are madeworse by the presence of hydrogen. Understanding the contribu-tion of hydrogen to deformation and fracture processes is thereforecentral to developing a mechanistic understanding of SCC.

Austenitic stainless steels are highly resistant to embrittlementpromoted either by hydrogen or low temperature, retaining signif-icant ductility and fracture toughness in these environments[4,11–15]. Austenitic stainless steel welds can be more susceptiblethan base materials to fracture, corrosion, and other degradation.The compositions of austenitic welds are typically selected to pro-mote primary solidification as d-ferrite and solid-state transforma-tion to austenite to suppress solidification cracking [16]. The as-

welded microstructure usually retains several volume percent ofthe bcc ferrite phase at room temperature, resulting in a duplexaustenitic/ferritic microstructure. Studies of duplex stainless steelsemphasize their attractive mechanical and corrosion properties,intermediate between those of austenite and ferrite. High hydro-gen fugacities at the metal surface promote hydrogen pickup, espe-cially in applications with cathodic protection or sour gas or oilcontaining significant H2S [1,2]. For hydrogen-exposed austeniticwelds and other duplex (two-phase) austenite/ferrite microstruc-tures, the detrimental influence of ferrite on room-temperatureductility and fracture toughness has been noted in several studies[17–23]. Hydrogen transport is complicated by the lower solubilitybut higher diffusivity and permeability of hydrogen in the ferritephase [24]. Austenitic steels are not as readily embrittled by hydro-gen as ferritic steels [2]. Low temperature can induce a ductile–brittle transition in ferrite and localizes deformation in austenite,yet few studies are reported for hydrogen-assisted fracture ofaustenitic stainless steel welds at low temperature, particularlystudies that quantify crack growth resistance by fracture mechan-ics methods. Mills [25] reviewed cryogenic fracture toughness ofAISI 304 and 316 base metals and their welds at temperatures aslow as 4 K, however, studies at less severe sub-ambient tempera-tures are needed.

The mechanism of environment-assisted crack propagation isknown to differ depending on the specific combination of metallur-gical, environmental, and mechanical variables [5,26]. For example,features of hydrogen damage in austenitic and duplex stainlesssteels differ when they result from exposure to high pressurehydrogen gas versus cathodic charging in an aqueous solution; orupon tensile straining after precharging with hydrogen versus insitu straining in a hydrogen environment [1,27–30]. Austenitic al-

H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221 211

loys’ susceptibility to embrittlement has been correlated with apropensity for localized deformation, which can arise due to met-allurgical factors such as low stacking fault energy (SFE) [31–33] orenvironmental factors such as low temperature [33–37] or hydro-gen exposure [4,38,39]. Hydrogen-assisted crack growth is sensi-tive to the combination of metallurgical, environmental, andmechanical variables, hence it is important to understand effectsof relevant combinations of material and environmental variables(absorbed hydrogen, low temperature, ferrite/austenite phase dis-tribution) on fracture of austenitic welds. In the present manu-script, we have isolated effects of a high-purity gaseous hydrogenenvironment. However, hydrogen derived from aqueous sources(e.g. corrosion reactions at a stress-corrosion crack tip) is thoughtto affect alloys in a similar manner [4,14,38].

The present study expands on the results of a previous investi-gation of room-temperature hydrogen-assisted crack propagationin gas-tungsten arc (GTA) welds having a 304L base metal and308L filler [21]. The objectives of the present study are to charac-terize hydrogen-assisted crack propagation of these welds at lowtemperature. Fracture mechanics specimens were thermally pre-charged in hydrogen gas, and fracture initiation toughness andcrack growth resistance curves were measured at 223 K (�50 �C).The effects of low temperature on hydrogen-assisted crack initia-tion and propagation mechanisms were assessed via electronmicroscopy of fracture surfaces and fracture profiles.

2. Experimental

The processes of weld fabrication, hydrogen precharging, andfracture testing of welds at room temperature were detailed byJackson et al. [21] and are summarized here.

2.1. 304L/308L Welds

The weld, referred to as 304L/308L throughout this work, wasfabricated from a round bar (64 mm diameter) of annealed 304Lstainless steel having a longitudinal U-shaped groove. The groovewas filled with 1.1 mm diameter 308L filler wire using 9 gas-tung-sten arc (GTA) weld passes. Bulk elemental compositions of thebar, wire, and weld fusion zone are reported in Table 1. The volumefraction of ferrite in the weld fusion zone was estimated by mag-netic measurement (Feritscope, Helmut-Fischer GmbH, Sindelfin-gen, Germany) at several distances from the weld root along thecenterline. The as-welded microstructure was imaged using opticaland scanning electron microscopy (SEM, JSM-840, JEOL, Tokyo, Ja-pan). The elemental compositions of the ferrite and austenitephases within the weld were quantified by electron probe micro-analysis (EPMA, JXA-8200, JEOL, Tokyo, Japan). Variation in tensileproperties as a function of depth in the 304L/308L weld was char-acterized by tensile tests of subsized specimens which were ex-tracted from the root, mid-height, and top of the weld. Upongoing from the root to the top of the weld, yield strength decreases

Table 1Elemental compositions (wt%) of base metal, filler, and weld fusion zone (balance Fe).

Cr Ni Mn Si Mo C

304La 19.85 10.73 1.6 0.57 0.17 0.030308La 20.5 10.3 1.56 0.5 <0.01 0.028304L/308L weld b,c 21.27 10.23 1.69 0.51 0.05 0.04304L/308L weldb,d 20.75 10.19 1.65 0.51 0.04 0.02

a Per manufacturer certification.b By emission spectroscopy.c Analysis at mid-height of weld.d Analysis at root of welde Nieq = Ni + 35C + 20N + 0.25Cu, and Creq = Cr + Mo + 0.7Nb per Kotecki and Siewert [4

from 420 MPa to 350 MPa, while strain at failure increases from40% to 50% [21].

2.2. Preparation of fracture specimens

Disk-shaped compact-tension (CT) specimens, illustrated sche-matically relative to the weld in Fig. 1, had a width W (distancefrom the load-line to the back face of the specimen) of 40.6 mm,thickness B of 6.4 mm, and net thickness Bn (between the side-grooves) of 4.8 mm, consistent with ASTM E1820 [40] and ASTME1737 [41]. Specimens tested at room temperature in the non-charged condition had the same W but with a B of 12.7 mm andBn of 9.9 mm.

In order to locate the fatigue precrack tip in a similar micro-structure in all specimens, the precrack was grown to the same dis-tance into the weld fusion zone, about 1.5 mm. To accommodatevariation in the distance from the notch tip to the root of the weldin the machined specimens, the final crack length-to-width ratiovaried between 0.63–0.68. Precracks were grown in air along theradial direction of the welded bar along the weld centerline at10 Hz under a load ratio of 0.1 and final maximum stress intensityfactor Kmax of 30 MPa m1/2.

Residual stresses in precracked specimens are minimized byconducting precracking in accordance with standard proceduresin ASTM E1820 [ref]. Specimen dimensions conformed to allowabledimensions (with respect to specimen geometry) of the starternotch, final precrack, and amount of fatigue crack growth. The Kmax

applied during precracking was kept well below the material frac-ture toughness that is measured during subsequent fracture test-ing. During the precracking procedure, the displacement is keptconstant, hence the Kmax and DK range decrease progressively asthe crack grows.

Thermal precharging of CT specimens was in 99.9999% hydro-gen gas at 138 MPa and 573 K for a minimum of 39 days and upto 61 days. The charging time and temperature were selected toachieve at the specimen midthickness a minimum hydrogen con-centration of 90% of the equilibrium hydrogen concentration atthe surface. After hydrogen precharging but before mechanicaltesting, specimens were stored at below 250 K to minimize hydro-gen outgassing. The specimen was equilibrated at the test temper-ature, either 293 or 223 K, for at least 30 min prior to mechanicaltesting. The hydrogen concentration in the fusion zone was140 wppm (�0.8 at.%), as measured by inert gas fusion (WahChang, Albany, OR). This concentration is consistent with that pre-dicted based on the thermal precharging parameters and thehydrogen solubility of 300-series alloys [42].

2.3. Fracture mechanics testing

Elastic–plastic fracture mechanics tests of the fatigue-pre-cracked disk CT specimens were conducted according to ASTME1820 [40]. Three or four replicate specimens were tested for each

S P N Cu Nb Creqe Nieq

e

0.001 0.005 0.02 – – 19.85 12.180.012 0.006 0.055 0.018 <0.005 20.51 12.380.02 0.006 0.047 0.02 0.01 – –0.011 0.005 0.048 0.02 – – –

3].

Fig. 1. Optical image of the 304L/308L weld in cross-section, with the orientationand location of disk-shaped CT specimens illustrated schematically.

212 H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221

condition. Room-temperature tests were conducted in the ambientlaboratory air. Tests at 223 K were conducted within an environ-mental chamber mounted in the load train of a servo-hydraulic testframe (MTS Corporation, Eden Prairie, MN). Evaporated liquidnitrogen was used as a coolant, with a resistive heater and temper-ature controller maintaining the test temperature of 223 K.

Testing was conducted at a constant (monotonic) displacementrate of 0.4 mm min�1. The crack-opening displacement (COD) atthe load-line was measured with a clip gauge, and crack extensionwas monitored continuously using the direct current potentialdrop (DCPD) method of ASTM E1737 [41]. The test was terminatedmanually after a greater than 10% load drop from peak load and anincrease in DCPD by at least 15% from its initial value, or whenload-line displacement exceeded the measurement capacity ofthe clip gauge (4–5 mm).

Following the test, CT specimens were heat-tinted at 623 K for60 min to mark the extent of crack growth during the test. Thespecimens were then broken apart in liquid nitrogen, and the finalcrack length and precrack length were measured from the fracturesurfaces. The calculated crack length based on DCPD data was lin-early scaled so that the final calculated and measured crack lengthswere equal. Prior to correction, the discrepancy between the finalmeasured and calculated crack lengths was less than 1.5%.

The J-integral crack growth resistance (J–R) curve was con-structed from the load vs. COD and DCPD vs. COD records. The ini-tial J vs. crack growth increment (Da) relationship was assumed tobe a linear crack blunting response following ASTM E1737. The ini-

Fig. 2. (a) Optical and (b) backscattered electron images of the fusion zone in the 304solidification direction is nominally from bottom to top of the image. The microstructur

tiation of crack extension following blunting was identified as thesecond slope change in the DCPD vs. COD record, which is consis-tent with observations from interrupted fracture mechanics testsof stainless steels [22]. The fracture initiation toughness was iden-tified from the intersection of the 0.2 mm offset blunting line withthe R curve.

2.4. Post-testing characterization

Fracture surfaces and fracture profiles parallel to and normal tothe crack growth direction were characterized by SEM. Elementalanalysis of compositional partitioning by energy dispersive X-rayspectroscopy (EDS) was used to infer phases in the weld fusionzone.

3. Results

3.1. As-welded microstructure

The microstructure of the fusion zone consisted of primary d-ferrite in a matrix of austenite (c), as seen in Fig. 2. The ferritenumber (FN) as measured at distances of 2, 6, 10, and 14 mm fromthe weld root along the centerline was 9.4, 10.1, 10.7, and 10.1,respectively. These values represent the mean of five measure-ments per location. The FN approximates the volume percent fer-rite, and the measurements are consistent with values of 8–10predicted from the WRC-1992 diagram [43], based on the calcu-lated Creq and Nieq for the weld metal (Table 1).

Table 2 shows the elemental compositions of the austenite andferrite phases within the weld, as measured by EPMA spot analysesat two locations: at mid-height of the weld and near the weld root.Table 2 also compares the compositions of the individual phases tothe average composition of the weld fusion zone, as measured byemission spectroscopy in the same locations.

In the fusion zone, the average composition by emission spec-troscopy closely matched that of the original 308L filler wire. Weldaustenite had approximately the same wt% Ni and Cr as in the 304Land 308L starting materials, while the approximately 8–10 vol% ofd-ferrite was depleted in Ni and enriched in Cr.

3.2. Fracture toughness

Compared to the non-charged condition, precharged hydrogenmarkedly decreased the fracture initiation toughness of 304L/308L welds at both 293 and 223 K, as shown by crack growth resis-

L/308L weld near the fusion zone/base metal interface at the weld centerline. Thee contains skeletal ferrite (d) in an austenite (c) matrix.

Table 2Elemental compositions (wt%) of ferrite and austenite phases within the weld.

Phase Location Cr Ni

Weld ferrite Weld mid-heighta 28.17 4.41Weld roota 27.86 4.38

Weld austenite Weld mid-heighta 20.02 10.24Weld roota 19.73 10.40

Weld average Weld mid-heightb 21.27 10.23Weld rootb 20.75 10.19308L filler wirec 20.5 10.3

Base metal Outside fusion zonea 18.44 10.21304L annealed barc 19.85 10.73

a EPMA, mean of three measurements for ferrite and nine for austenite.b Emission spectroscopy.c Per manufacturer certification.

H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221 213

tance (J–R) curves (Fig. 3) and fracture initiation toughness (Jmax orJIH) values (Table 3). Also reported in Table 3 are the correspondingstress-intensity factor KJ values, determined by:

KJ ¼ffiffiffiffiffiffiffiffiffiffiffiffiffiffi

EJ1� m2

rð1Þ

where E is Young’s modulus, J is fracture initiation toughness, and mis Poisson’s ratio [40]. The J–R curve slope over the first 0.5 mm ofcrack extension immediately following blunting is reported as crackgrowth resistance dJ/dDa.

Fig. 3. Crack growth resistance (J–R) curves for 304L/308L GTA welds with andwithout precharged hydrogen, tested at 293 K and 223 K. Up to four J–R curves percondition are plotted. (a) At longer crack extensions, welds tested at 293 K havegreater crack growth resistance than at 223 K. (b) Immediately following crackinitiation, J–R curves were indistinguishable at 293 versus 223 K, as were JIH values.

Welds not exposed to hydrogen and tested at both 293 and223 K exhibited significant crack tip blunting, and tests were ter-minated manually before the onset of noticeable crack extension,as confirmed by post-test analysis. Crack extension in non-chargedspecimens was insufficient to permit calculation of JIC, and insteadwe report as Jmax the lesser of two values, either (1) the maximum Jvalue recorded during the test or (2) the maximum J-integralcapacity of the specimen as permitted by ASTM E1820 [40], i.e.the lesser of b0rY/10 or BrY/10, where b0 is the original uncrackedligament, B is the original specimen thickness, and rY is the mate-rial flow stress. These Jmax values represent lower-bound fractureinitiation toughness measurements.

It is important to note that the difference in Jmax between non-charged welds tested at 293 and 223 K (Table 3) was primarily dueto the difference in specimen thickness and the correspondingdependence of the maximum J-integral capacity on this dimension.Non-charged specimens tested at 223 K were thinner than thosetested at 293 K, while hydrogen-precharged specimens were allof the same thickness. Regardless, crack blunting and a lack ofnoticeable crack extension were observed for non-charged speci-mens tested at both 293 and 223 K.

The measured fracture initiation toughness of hydrogen-pre-charged welds represents a reduction of greater than 59% com-pared to the lower-bound fracture initiation toughness of non-charged welds. Hydrogen-precharged specimens met the dimen-sional requirements in ASTM E1820 [40] to define conditional JQ

as JIC. To reflect the hydrogen-precharged condition, fracture initi-ation toughness is denoted JIH and stress-intensity factor is KJIH.

Fracture initiation toughness with internal hydrogen is essen-tially the same at 293 and 223 K. At the initiation of crack growth(Fig. 3b), the J–R curves for both temperatures are indistinguish-able. However, at longer crack extensions (Fig. 3a), welds testedat the lower temperature exhibit lower crack growth resistancedJ/dDa. This divergence in J–R curve slope occurs at a crack exten-sion of about 0.8–1 mm from the precrack tip. This distance wasfound to correspond to a change in orientation and morphologyof austenite grains and d-ferrite dendrites at a depth in the fusionzone where two weld passes overlapped (Fig. 4). Crack paths weretortuous in general, and any crack deflection or differences in frac-ture appearance due to microstructural differences in the weldpass overlap region were not visibly distinguished from that inother regions.

3.3. Fractography

Fracture features are significantly different for specimens frac-tured with and without hydrogen; for specimens fractured inhydrogen, a different distribution of fracture features is seen whentest temperature is decreased from 293 to 223 K. Effects of hydro-gen and test temperature on fracture features are compared inFig. 5 and in higher magnification in Fig. 6.

Non-charged welds at both 293 and 223 K exhibited significantcrack blunting followed by a slight amount of ductile crack exten-sion. These fracture surfaces featured only equiaxed and U-shapeddimples (Fig. 5a and b), indicating microvoid coalescence (MVC).Under these conditions, microvoids nucleate at fine sphericalprecipitates.

Hydrogen exposure altered the fracture mode of austeniticwelds (Fig. 5c and d). In hydrogen-precharged welds fractured at293 K, the fracture surface was dominated by features which wereelongated and aligned parallel to the crack propagation direction,which was also the weld solidification direction (Figs. 5c, 6a andb). The dendritic morphology of these features (denoted by D inFigs. 5c and 6a) resembled the underlying d-ferrite microstructure(Fig. 2), with smooth, relatively flat features interspersed with tearridges (denoted by T in Fig. 6a), i.e. vertical steps joining two flat

Table 3Fracture initiation and crack growth toughness of 304L/308L GTA welds

Hydrogen-precharged Non-chargeda,b dJ/dAac (kJ m2 mm-1)

JIH (kJ m2) KJIH (MPa m�0.5) dJ/dDa c (kJ m2 mm�1) Jmax (kJ m�2) KJ (MPa m�0.5) dJ/dDa c (kJ m2 mm�1)

223 K 144 175 171123 161 175109 152 181

Mean 125 163 176 304 254 965

293 K 110 153 209123 161 190132 167 196111 154 242

Mean 119 159 209 625 364 965

a Mean of 3 replicate tests 6 of non-charged specime.b B of non-charged specimens was 6.35 mm for tests at �50 �C and 12.7 mm for tests at 20 �C.c Slope of J–R curve, measured over 0.5 mm of crack extension immediately following blunting.

Fig. 4. A change in microstructural orientation and morphology was observed where two weld passes overlap, about 0.8–1 mm from the precrack tip and 2.3–2.5 mm fromthe base metal/weld interface. Backscattered electron image.

Fig. 5. Fracture surfaces of hydrogen precharged and non-charged welds, tested at 223 or 293 K. Without hydrogen, at (a) 293 K or (b) 223 K, fracture surfaces had onlydimpled rupture. With hydrogen at (c) 293 K, the fracture surface was dominated by elongated features with a dendritic morphology (D) and aligned with the crack growthdirection, while at (d) 223 K, the dominant features were elongated (E) and equiaxed (Q) dimples, with few dendritic-shaped features. Crack growth direction is from bottomto top in all images.

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Fig. 6. Fracture features at matching locations on both halves of hydrogen-precharged fracture surfaces. At 293 K (a and b), the dominant features were dendritic (D) withtear ridges (T) between dendritic features. At 223 K (c and d), few regions had dendritic features (D). The dominant features (e and f) were planar arrays of elongated (E)dimples and patches of equiaxed (Q) dimples. Crack growth direction is from bottom to top in all images.

H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221 215

areas of dimpled rupture, with steps also exhibiting dimples. Mostof the fracture surface exhibited features related to the underlyingdendritic structure of ferrite and the inter-dendritic regions ofaustenite.

In hydrogen-precharged welds fractured at 223 K (Figs. 5d, 6c,and d), considerably fewer regions of the fracture surface showedevidence of the dendritic microstructure (denoted by D in Fig. 6cand d). Most regions showed either elongated or equiaxed dimples(denoted by E and Q respectively in Figs. 5d and 6e). Elongateddimples occurred in nominally parallel and coplanar arrays(Fig. 6e and f). Parallel arrays of elongated dimples formed stepson the fracture surface (Figs. 5d and 6e).

Matching locations on both fracture surfaces of hydrogen-ex-posed welds (Fig. 6) were analyzed by EDS to provide a basis forinferring the phase that was present at the surface based on theapproximate Ni and Cr concentrations, which were consistent withthe values for austenite or ferrite in Table 2. EDS analysis alonemay be inconclusive if the ferrite dimensions are very thin and fer-rite evades detection by the EDS probe. By comparing elementalanalyses and fracture features on both halves of the fracture sur-face, the crack path could be identified as fracture within ferrite,at the austenite/ferrite phase boundary, or within austenite. All

three types of crack path were observed in various locations atboth temperatures.

In the present work, no distinction was made between cracks atc/c boundaries and cracks within a c grain, although both may oc-cur. For this reason, it cannot be conclusively stated that d was ab-sent at a given fracture location and that the crack path wasentirely contained within the c phase.

Planar arrays of elongated dimples, seen only at 223 K, wereassociated with fracture within the austenite phase; such crackgrowth could be either transgranular or at c/c interfaces. Previ-ously [20,21], it was established that for hydrogen-prechargedaustenitic welds at 293 K, tear ridges and equiaxed dimples wereassociated with fracture within austenite; the flat areas containingdendritic features were associated either with fracture at the aus-tenite/ferrite phase boundary or within ferrite.

At both temperatures, fracture surfaces exhibited significantout-of-plane crack growth and tortuous crack topographies. Withhydrogen at both 293 and 223 K, crack extension evolved as multi-ple discontinuous microcracks that joined up to form the maincrack path. At 293 K, the microcracks formed on planes parallelto the Mode I crack growth plane, i.e. normal to the tensile axis[21]. At 223 K, microcracks formed primarily on planes inclined

216 H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221

to the Mode I crack growth plane, as confirmed from stereo pairfractographs.

3.4. Fracture profile metallography of hydrogen-exposed welds

Fracture profiles parallel and transverse to the crack growthdirection in hydrogen-exposed welds were imaged by backscat-tered electron microscopy (Figs. 7 and 8). Atomic number contrastdistinguishes d-ferrite as the dark grey dendritic phase and austen-ite as the light grey matrix. Localized deformation bands appear asbright parallel streaks in the light grey austenite matrix.

Fracture profiles (Fig. 7) suggested that, although microvoids/microcracks initiated near ferrite at both 223 and 293 K in hydro-gen, crack growth occurred by different mechanisms at the twotemperatures.

At 293 K, microcracks were predominantly associated with d-ferrite (Fig. 7a and b). Microcracks nucleated either within the fer-rite phase or at the austenite/ferrite phase boundary (Fig. 7a) andpropagated parallel to ferrite dendrites (denoted by d in Fig. 8a).These microcracks often created coarse cracks (Fig. 7b) parallel tothe main crack, which were located up to several hundred microm-eters above or below the fracture surface (Fig. 8c).

At 223 K, there was no evidence for large microcracks awayfrom the fracture surface. Fine microvoids 1 lm or less in diameter(Fig. 7d and f) were the main type of microstructural damage ob-

Fig. 7. Fracture profiles show different mechanisms of microvoid nucleation at 293 versu223 K (c–f), microvoids nucleate where localized deformation bands intersect other deforform via the coalescence of multiple microvoids (c, d, and f). Crack growth direction is fro

served. With an equiaxed cross-section, the three-dimensionalshape of these microvoids could be either spherical or tubular.Some fine porosity was randomly distributed in the as-weldedmicrostructure, however the microvoids that were associated withdamage accumulation and fracture had a distinctive arrangement.Specifically, these microvoids tended to be aligned and to coincidewith the intersections between deformation bands (Fig. 7f). Thepredominant microvoid nucleation site was adjacent to d-ferrite(Fig. 7c–f), while a smaller fraction formed elsewhere in austenitegrains. Microcracks evolved via the nucleation and coalescence ofclosely-spaced microvoids (Fig. 7c, d, and f).

Fracture profiles reveal a second key difference between frac-ture at 293 K and at 223 K: the macroscopic crack path. At both293 and 223 K, fractures had significant out-of-plane crack growthand tortuous crack topographies (Fig. 5c and d). This is reflected incross-sections as macroscopic steps (denoted by S in Fig. 8a,c,d).

At 293 K, steps corresponded to microcracks growing along d-ferrite on various planes ahead of the crack tip that were nominallyparallel to each other and perpendicular to the tensile axis (Fig. 8aand c). The main crack grew when these microcracks joined alongplanes parallel to the tensile axis, facilitated by intense shear in theligaments between them (denoted by I in Fig. 8a). This out-of-planecrack growth corresponds to that seen on fracture surfaces(Fig. 5c).

s 223 K. At 293 K (a and b), microcracks originate at and propagate along d-ferrite. Atmation bands, and microvoids are concentrated near phase boundaries. Microcracksm left to right in all images except for a, in which crack growth is normal to the page.

Fig. 8. Fracture profiles reveal that at both temperatures, the crack path is marked by macroscopic steps (S). At 293 K (a and c), cracks propagate along ferrite dendrites (d) butcan link up through intense shear (I) in the ligaments between them. At 223 K (b and d), crack paths are primarily within austenite and do not clearly follow ferrite.

H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221 217

At 223 K, macroscopic steps and out-of-plane crack growthwere also observed. However, microcracks were observed to formon planes inclined to the tensile axis (Fig. 6c–f). In cross-section,the crack path does not clearly follow the ferrite phase (Fig. 8band d) and is primarily within the austenite phase.

4. Discussion

As shown in previous studies [20,21], one primary role ofhydrogen in degrading fracture resistance and altering fracturemechanisms in austenitic stainless steel welds at room temper-ature is enhancing planar deformation in austenite. This phe-nomenon produces stress concentrations where localizeddeformation bands intersect d-ferrite. These stress concentrationsfacilitate microcracking of the elongated d-ferrite grains, whichare aligned with the direction of crack growth. This dominantrole of ferrite in the initiation and propagation of microcracksin austenitic welds at room temperature is illustratedin Figs. 6a, 6b, 7a, andb and documented in several studies[17–21].

Extending the study on 304L/308L welds at room temperature[21], we found that low-temperature (223 K) exposure had no ef-fect on the fracture initiation toughness of hydrogen-prechargedwelds (Table 3), although the fracture appearance was clearlychanged (Fig. 5c and d). In contrast to room-temperature fracture,at 223 K fracture surfaces and fracture profiles showed a dimin-ished role of d-ferrite in facilitating crack propagation (Figs. 6and 7), relative to room temperature.

We propose that low temperature alters the competition be-tween potential fracture initiation modes in hydrogen-prechargedaustenitic welds. The two main effects of low temperature are that(1) low temperature localizes plastic deformation in austenite, ren-dering this phase more susceptible to hydrogen-enhanced local-ized plasticity, and (2) low temperature increases the hydrogenconcentration in austenite near phase boundaries through its ef-fects on the thermodynamics of hydrogen trapping. In the follow-ing sections, we discuss effects of low temperatures on plastic

deformation and hydrogen distribution in welds and how thesechanges govern the fracture mode.

4.1. Sequence of crack initiation and propagation in hydrogen-precharged welds at 223 K

In the presence of hydrogen, low temperature (223 K) produceda significantly different fracture surface than seen at room temper-ature (293 K).

At 293 K, a mixture of ferrite cleavage and phase boundaryseparation was observed. Microcracks nucleated and propagatedpreferentially at d-ferrite (Figs. 7a and b), resulting in elongated,dendritic features on fracture surfaces (Figs. 5c, 6a, and 6b)that resemble the underlying ferrite morphology (Fig. 2). Thesemicrocracks formed in the ferrite phase itself or at austenite/ferrite interfaces, facilitated by high stress concentrations at thesesites.

At 223 K, neither ferrite cleavage nor phase boundary separa-tion is an important fracture mechanism, based on the observedfracture surface features (Fig. 6c–f). Rather, the dimpled surfaceindicates that fracture was caused by MVC within the austenitephase. Microcracks originated as microvoids at intersections be-tween localized deformation bands in austenite (Fig. 7c–f). Becausedeformation bands are parallel, intersections between deformationbands occur in rows, as do the microvoids that form at these inter-sections. The non-equiaxed growth of these aligned microvoidsproduced coplanar arrays of elongated fracture dimples (Figs. 5d,6e, and 6f). In limited locations, the fracture appearance at 223 Kresembled the dendritic features associated with microcracks ind-ferrite or at austenite/ferrite boundaries at 293 K (denoted by Din Fig. 6c and d).

An important feature of the proposed low-temperature frac-ture mechanism is that, although deformation band intersec-tions occurred throughout the austenite bulk, microvoids werenot randomly distributed but were concentrated near austen-ite/ferrite phase boundaries (Fig. 7c–f). We can infer that inwelds or two-phase microstructures, two conditions promotemicrovoid nucleation in austenite: (1) deformation bands inter-

218 H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221

sect and (2) the intersection is in close proximity to an austen-ite/ferrite phase boundary. The role of hydrogen in facilitatingvoid formation near phase boundaries will be discussed in thefinal section.

Despite the differences in fracture appearance, the fracture pro-cess at 223 K had a similar sequence of steps compared to the frac-ture process at 293 K [20,21].

First, planar deformation in austenite developed as localizedbands. The formation of these deformation bands is facilitated bythe hydrogen enhanced localized plasticity (HELP) mechanism[44–46]. In austenite, HELP results in greater slip planarity[4,38,39], exacerbating the effects of metallurgical factors predis-posing an alloy to planar slip, such as low stacking fault energy(SFE) [31–33]. In metastable 300-series alloys such as 304L and308L, SFE decreases as nickel content decreases [47,48]. It followsthat the effects of low SFE, low temperature, and hydrogen wouldproduce more severely localized slip than any factor acting inde-pendently. As slip becomes more localized, other planar deforma-tion modes such as twinning and e-martensite become activatedat lower strains [33,35,49–53].

The next step in the process was the impingement of localizeddeformation bands on microstructural obstacles, thus producingstress concentrations. Austenitic stainless steels contain variousmicrostructural features that can act as obstacles to slip or twin-ning, including grain boundaries, secondary phases, and annealingtwin boundaries. At 293 K, the d-ferrite phase could serve as a bar-rier to propagation of planar deformation bands, and the resultingstress concentration facilitated microcrack initiation either ascleavage of ferrite or decohesion at the austenite/ferrite phaseboundary [21]. At 223 K, the resulting microstructural damagewas instead a series of discrete microvoids at deformation bandintersections near phase boundaries. Hence, low temperaturechanged the dominant mode of damage accumulation, leading toa different mode of microcrack initiation relative to that at roomtemperature.

The final step was crack propagation via the growth and coales-cence of microcracks. At 293 K, microcracks formed from cleavagein ferrite or decohesion of austenite/ferrite interfaces, then thesemicrocracks linked by intense shear in the remaining ligaments(Figs. 7a, 7b, 8a, and 8c). At 223 K, microvoids nucleated predomi-nantly adjacent to weld ferrite, while ferrite cleavage and austen-ite/ferrite interface decohesion were infrequently observed.Instead, under low-temperature conditions, decohesion was not adominant fracture mode; an alternative mode of crack propagationintervened, and multiple microvoid nucleation and coalescenceevents were required for microcracks to form (Fig. 7c–f). Similarto the process at room temperature, individual microcracks aheadof the crack tip at low temperature linked up via shear in the liga-ments between them, producing coarse steps on the fracture sur-faces (Fig. 8b and d).

4.2. Effect of temperature on crack initiation: localized deformation

In this study, temperature governed the mode by which hydro-gen-assisted cracks initiated and propagated. At both 293 and223 K, planar deformation and microstructural obstacles played aprominent role in the initiation of hydrogen-assisted cracking inaustenitic stainless steel welds.

We propose that low temperature altered fracture mecha-nisms in hydrogen-precharged welds in part by conditioning theaustenite for more severe planar slip that is further exacerbatedby hydrogen. The planar deformation bands that evolve increasethe severity of stress concentrations at intersecting deformationbands near the austenite/ferrite phase boundary, perhaps elevat-ing stress concentrations at these sites relative to the stress con-centrations at intersections of deformation bands with d-ferrite.

Consequently, void nucleation in austenite near phase boundariesis favored over microcracking in ferrite or at austenite/ferriteinterfaces.

Low temperature exacerbates localized deformation in severalways. First, low temperature increases the tendency toward planarslip in austenite by decreasing SFE [33–37]. More intense planarslip on [1] planes in fcc austenite gives rise to deformation twins,bcc a0-martensite, and hcp e-martensite at lower strains[33,35,49–53]. These deformation modes promote the confinementof deformation in discrete bands, leading to more pronouncedstrain incompatibility and enhanced stress concentration at obsta-cles. After deforming 301LN stainless steels in tension at tempera-tures between 233 and 353 K, Talonen and Hanninen [33] observedthat, as temperature decreased, a greater volume fraction of a0-martensite formed, and at lower levels of strain. They proposedthat low SFE at low temperature facilitates the martensitic trans-formation. In contrast, at higher temperature, the overlapping ofstacking faults is less regular, hindering the formation of martens-ite nuclei. Deformation at lower temperature produces more twinsand martensite – the results of planar deformation modes – than atroom temperature.

Second, low temperature could also change the types of defor-mation products that preferentially form. In addition to decreasingSFE, low temperature could alter the free-energy differences be-tween the fcc, bcc, and hcp phases, and thus the thermodynamicdriving force for transformation to the hcp or bcc martensitephases. During an investigation of tensile deformation of austeniticsteels containing 16-18 wt% Cr and 11–13 wt% Ni at temperaturesbetween 77 and 523 K, Lecroisey and Pineau [35] observed defor-mation bands containing both deformation twins and e-martens-ite, with a higher proportion of twins at higher deformationtemperatures. They proposed that the tendency to nucleateaustenitic twins versus e-martensite at a given temperature de-pends on the relative values of extrinsic and intrinsic stacking faultenergies as well as the volume contraction due to the fcc-to-hcptransformation. Temperature probably determines the relativeproportions of hcp or bcc martensite and fcc twins that composedeformation bands. Although insufficient data exists to evaluate ef-fects of deformation band makeup on the severity of stress concen-trations, it is reasonable to presume that deformation productssuch as twins and e-martensite confine deformation in discretebands, promoting strain incompatibility and stress concentrationat obstacles [21].

Third, more severe planar deformation has been correlated withphysically thinner deformation twins and e-martensite plates and,consequently, higher stress concentrations where they impinge onobstacles.

Based on an investigation of uniaxial tensile fracture of nitro-gen-bearing austenitic steels at temperatures between 4 and298 K, Mullner and coworkers [54,55] attributed fracture initiationto stress concentrations at deformation twin intersections. Lowertemperature decreased the thickness of twins and the dislocationdensity, and they asserted that such conditions rendered twinseffective obstacles against propagation of impinging twins, creat-ing stress concentrations high enough to initiate microcracks attwin intersections.

Based on fractographs and fracture profiles of low-SFE, high-Mn austenitic steels fractured in Charpy tests at 77 K, Takakiet al. [56] inferred that intersections between e-martensiteplates and the associated stress concentrations led to the forma-tion of parallel tunnel-like voids at these intersections. They in-voked this mechanism to explain the dimpled planes and step-like ridges they observed on fracture surfaces. It is plausible thatpropagation of both twins and e-martensite plates could beblocked by the austenite/ferrite phase boundary as well. Thinnerdeformation twins and e-martensite plates produced at lower

Table 4Hydrogen diffusivity D and diffusion distance x in ferrite and austenite phases at 293and 223 K

T (K) D a (m2 s�1) x (lm)

t = 30 min t = 3 h

Ferrite 223 1 � 10�13 30 60293 8 � 10�12 200 600

Austenite 223 2 � 10�19 0.04 0.1293 2 � 10�16 1 3

a Based on the values recommended by Perng and Altstetter [62] and San Marchi

H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221 219

temperatures would lead to more severe stress concentrationsthan at room temperature.

While the preceding discussion supports the notion that lowtemperature conditions austenitic stainless steels for more local-ized deformation, which can then be exacerbated by elevatedhydrogen, leading to elevated stress concentrations at intersectingdeformation bands, one detail still requiring attention is the obser-vation that at low temperature, microvoids form predominantlynear austenite/ferrite phase boundaries.

One possible explanation for this observation is that deforma-tion is disproportionately localized near austenite/ferrite phaseboundaries, where Ni depletion locally decreases SFE. Brookset al. [18,57] attributed the Ni concentration profile observedfor AISI 309 GTA welds to partitioning during the diffusion-con-trolled solid-state transformation from ferrite to austenite follow-ing primary ferrite solidification. The authors attributed Nienrichment at the centers of austenite grains (i.e. the boundariesof solidification cells) to secondary austenite solidification due toNi enrichment of the liquid. EDS concentration profiles of thepresent weld suggested that Ni concentration was 2.5–3 wt% low-er near the phase boundaries than in the centers of austenitegrains (i.e. the boundaries of solidification cells), where the 12–13 wt% Ni exceeded the nominal Ni concentration in weld austen-ite, about 10.2 wt% (Table 2). Such concentration gradients couldbias localized deformation toward the austenite/ferrite phaseboundaries.

In summary, low temperature enhances the formation of planardeformation bands in austenite, especially near austenite/ferritephase boundaries. One potential role of hydrogen in linking defor-mation and fracture is that hydrogen magnifies planar slip. Thisintensified planar slip or the deformation products that it stimu-lates (e.g. twinning, e-martensite) leads to stress concentrationsat obstacles such as intersecting deformation bands. The observa-tion that microvoid nucleation in austenite near phase boundariesout-competes microcrack formation in ferrite could be due in partto increased stress concentrations at deformation band intersec-tions near phase boundaries relative to those at intersections be-tween deformation bands and d-ferrite. The difference in damageinitiation modes gave rise to the significant difference in fracturesurface appearance.

While the fracture modes at 293 and 223 K differed signifi-cantly, temperature had no notable effect on fracture initiationtoughness (Table 3). We propose that the similar JIH values at293 and 223 K are a coincidence. In other words, it is coincidentalthat the dominant fracture processes at these two temperaturesare characterized by a similar magnitude of the local fracture crite-rion, e.g., a critical local strain. For example, at cryogenic tempera-tures, crack nucleation would likely occur by cleavage of d-ferritedue to its ductile–brittle transition [25]. Other fracture modesmight dominate at different sub-ambient temperatures, alteringcrack initiation toughness.

On the other hand, dJ/da, a measure of crack growth resis-tance, was slightly lower at 223 K than at 293 K. According toRitchie and Thompson [58], metallurgical factors affecting the lo-cal fracture strain have a greater effect on crack growth than oncrack initiation, owing to a weaker strain singularity ahead of astable growing crack versus a stationary crack. Metallurgical fac-tors dictating ductility include local microstructure and the de-gree of planar deformation. In the present welds, the J–Rcurves for tests at 293 and 223 K coincided near initiation butdiverged after cracks extended into a region of overlap betweenweld passes and a change in microstructural orientation andmorphology (Fig. 4). This microstructural change and the in-crease in planar deformation at 223 K had a stronger influenceon dJ/da at 223 K than at 293 K.

4.3. Effect of temperature on crack initiation: local hydrogendistribution

We propose that low temperature alters the thermodynamics ofhydrogen trapping in austenitic stainless steel welds in a way thatredistributes hydrogen from the interiors of d-ferrite grains tophase boundary trap sites. At 223 K, this could contribute to hydro-gen-assisted fracture initiation near phase boundaries rather thanwithin d-ferrite.

The total concentration of dissolved hydrogen consists ofhydrogen in interstitial lattice sites and hydrogen trapped atmicrostructural features such as grain boundaries, phase bound-aries, dislocations, and impurities [59]. Luppo et al. [60] providedevidence that austenite/ferrite phase boundaries in welds serveas hydrogen trap sites. According to Hirth [61], trap sites becomemore densely populated as temperature decreases, as expressedin following equation:

hT

1� hT¼ hL exp½� EB

kT� ð2Þ

where hT is the trap site occupancy, hL is the lattice hydrogen con-centration (in this formulation, hL � 1), Eb is the binding energyof hydrogen in the trap site, k is Boltzmann’s constant, and T is tem-perature. Because a hydrogen-precharged specimen can be consid-ered a closed system with a fixed total hydrogen concentration, thehydrogen demanded by trap sites at lower temperature must besupplied by lattice sites. Conversely, when temperature increases,trap occupancy decreases, with hydrogen migrating into latticesites.

From a kinetic standpoint, the d-ferrite phase is more likelythan austenite to supply hydrogen to trap sites at phase bound-aries. Based on the values recommended by Perng and Altstetter[62] for ferritic stainless steel 29Cr–4Mo–2Ni and those recom-mended by San Marchi et al. [42] for 300-series austenitic stainlesssteels, the diffusivity of hydrogen in ferrite and austenite weredetermined at 293 and 223 K and are summarized in Table 4.Hydrogen diffusivity is four to six orders of magnitude faster inbcc ferrite than in fcc austenite at the test temperatures. Eventhough the lattice solubility of hydrogen in austenite is about100 times that in ferrite [63], and thus the concentrations will dif-fer between the two phases, the overall diffusional flux will still begreater in d-ferrite.

The present analysis demonstrates that hydrogen diffusivity ind-ferrite is rapid enough to satisfy the increased demand for hydro-gen at traps as induced by low temperature. The hydrogen diffu-sion distance x varies with diffusivity D and time t according toEq. (3), derived from the solution to the transient diffusionequation:

x � 2ffiffiffiffiffiffiDtp

ð3Þ

Specimens were exposed to the test temperature for at least30 min and for up to several hours prior to mechanical testing.The hydrogen diffusion distance in d-ferrite during a typical expo-

et al. [42].

220 H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221

sure (Table 4) is significantly greater than the maximum requiredmigration distance, i.e. from the center of a ferrite grain to trapsites at phase boundaries, about 0.5 lm.

At low temperatures, hydrogen migration from lattice sites in d-ferrite to traps at phase boundaries both depletes hydrogen fromferrite grains and enriches phase boundaries with hydrogen. Thistilts the competition away from microcrack formation in ferriteand toward microvoid nucleation near phase boundaries.

As a result of this hydrogen enrichment at phase boundariesthat is induced by low temperature exposure, we propose thathydrogen has two roles in concentrating damage near phaseboundaries. First, hydrogen in concert with low temperature local-izes deformation and increases stress concentrations near phaseboundaries, consistent with the concepts described in the previoussection. Second, hydrogen could have a direct role in facilitatingmicrovoid nucleation near phase boundaries through several pos-sible mechanisms.

Vacancies are generated by intense localized plastic deforma-tion and stabilized by hydrogen, according to the model of hydro-gen-stabilized vacancy formation proposed by Nagumo [64].Building upon this model, Neeraj et al. [65] propose a nanovoidcoalescence (NVC) micromechanism of hydrogen embrittlementfor ferritic steels. In the NVC mechanism, nanovoids nucleate at acritical excess vacancy concentration. They proposed that NVCgave rise to the nanoscale (10–20 nm) dimples they observed onquasi-brittle facets on fracture surfaces.

In the previous section, we linked the formation of microvoidsin austenite to intersections between planar deformation bandsthat may consist of e-martensite plates [56] and deformation twins[54,55]. It is plausible that hydrogen could directly promote frac-ture by encouraging void formation at these intersections, in addi-tion to hydrogen’s indirect role in fracture, i.e. localizingdeformation.

In the context of the present work, we propose that hydrogenenrichment at phase boundaries, particularly at low temperature,makes near-phase-boundary austenite the preferred site for dam-age nucleation. At low temperature, this damage proceeds in amanner consistent with the aforementioned mechanisms, inwhich hydrogen stabilizes the vacancies that eventually formmicrovoids or hydrogen facilitates microvoid nucleation at theintersection of localized deformation bands consisting of e-mar-tensite and deformation twins. At room temperature, the damageis predominantly in the form of microcracks in ferrite or at phaseboundaries.

5. Conclusions

Elastic–plastic fracture mechanics tests were conducted to eval-uate effects of low temperature on hydrogen-assisted crack propa-gation in austenitic stainless steel welds fabricated from AISI 304Lbase metal and 308L filler metal and which had been thermallyprecharged with hydrogen gas:

1. Hydrogen degrades fracture resistance by >59% relative to thelower-bound fracture initiation toughness of non-chargedwelds. At room temperature (293 K), hydrogen alters the dom-inant mode of damage accumulation in austenitic welds,thereby altering fracture mechanisms by (1) enhancing local-ized planar deformation in austenite and (2) facilitating micro-cracking of d-ferrite or its interfaces. The dendriticmicrostructure of weld ferrite, which is aligned with the direc-tion of crack growth, governed the crack path.

2. Low temperature (223 K) exposure altered fracture mecha-nisms in hydrogen-precharged welds. At 223 K, fracture initia-tion was dominated by microvoid formation at deformationband intersections in austenite near phase boundaries, while

microcracking of ferrite and its interfaces did not play a signif-icant role in crack growth. Low temperature was presumed to(1) lower SFE and condition the austenite for more localizeddeformation, which was further exacerbated by hydrogen, and(2) enrich phase boundaries with hydrogen by increasing theoccupancy of hydrogen at trap sites relative to lattice interstitialsites in ferrite. The more localized deformation in austenite,particularly near phase boundaries, could lead to more severestress concentrations at deformation band intersections relativeto those at intersections of deformation bands with ferrite.

3. Additionally, we propose that elevated hydrogen concentrationnear phase boundaries may directly promote localized voidnucleation in austenite, consistent with the mechanisms ofhydrogen-stabilized vacancies proposed by Nagumo [64] andvacancy-induced nanovoid coalescence (NVC) proposed byNeeraj et al. [65]. Alternately, hydrogen facilitated microvoidformation at intersections between localized deformationbands, consistent with literature reports of stress concentra-tions and microvoid nucleation at intersections between e-mar-tensite plates [56] and deformation twins [54,55].

4. Exposure to 223 K had no effect on fracture initiation toughnessof precharged welds, despite the difference in the dominantmode of damage accumulation. It is likely coincidental thatthe dominant fracture modes at the two test temperatures(microvoid formation at intersecting deformation bands at223 K vs. microcracking in ferrite or its interfaces at 293 K)had similar local fracture criteria, resulting in similar fractureinitiation toughness values.

Acknowledgments

The authors are grateful to J. Campbell for hydrogen pressuresystems support, A. Gardea for metallographic preparation, and J.Chames and R. Nishimoto for SEM imaging. Sandia is a multipro-gram laboratory operated by Sandia Corporation, a wholly ownedsubsidiary of Lockheed Martin Corporation, for the US Departmentof Energy’s National Nuclear Security Administration under con-tract DE-AC04-94AL85000.

References

[1] A.A. El-Yazgi, D. Hardie, The embrittlement of a duplex stainless steel byhydrogen in a variety of environments, Corros. Sci. 38 (1996) 735–744.

[2] W. Zheng, D. Hardie, The effect of hydrogen on the fracture of a commercialduplex stainless steel, Corros. Sci. 32 (1991) 23–36.

[3] Y. Mine, T. Kimoto, Hydrogen uptake in austenitic stainless steels by exposureto gaseous hydrogen and its effect on tensile deformation, Corros. Sci. 53(2011) 2619–2629.

[4] A.W. Thompson, I.M. Bernstein, The role of metallurgical variables inhydrogen-assisted environmental fracture, in: M.G. Fontana, R.W. Staehle(Eds.), Advances in Corrosion Science and Technology, Plenum Press, NewYork, NY, 1980, pp. 53–175.

[5] D. Delafosse, T. Magnin, Hydrogen induced plasticity in stress corrosioncracking of engineering systems, Eng. Frac. Mech. 68 (2001) 693–729.

[6] S.P. Lynch, Metallographic contributions to understanding mechanisms ofenvironmentally assisted cracking, Metallography 23 (1989) 147–171.

[7] S.P. Lynch, A fractographic study of hydrogen-assisted cracking and liquid-metal embrittlement in nickel, J. Mater. Sci. 21 (1986) 692–704.

[8] S.P. Lynch, A fractographic study of gaseous hydrogen embrittlement andliquid-metal embrittlement in a tempered-martensitic steel, Acta Metall. 32(1984) 79–90.

[9] Y. Mine, K. Hirashita, M. Matsuda, M. Otsu, K. Takashima, Effect of hydrogen ontensile behaviour of micrometre-sized specimen fabricated from a metastableaustenitic stainless steel, Corros. Sci. 53 (2011) 529–533.

[10] M. Koyama, E. Akiyama, T. Sawaguchi, K. Ogawa, I.V. Kireeva, Y.I. Chumlyakov,K. Tsuzaki, Hydrogen-assisted quasi-cleavage fracture in a single crystallinetype 316 austenitic stainless steel, Corros. Sci. 75 (2013) 345–353.

[11] R.J. Walter, W.T. Chandler: effects of high-pressure hydrogen on metals atambient temperature, Report no. NASA-CR-102425, Rocketdyne, Canoga Park,CA, 1969.

H.F. Jackson et al. / Corrosion Science 77 (2013) 210–221 221

[12] H.G. Nelson, Hydrogen embrittlement, in: C.L. Briant, S.K. Banerji (Eds.),Treatise on Materials Science and Technology: Embrittlement of EngineeringAlloys, Academic Press, New York, NY, 1983, pp. 275–359.

[13] G.R. Caskey Jr., Hydrogen effects in stainless steels, in: R.A. Oriani, J.P. Hirth,M. Smialowski (Eds.), Hydrogen Degradation of Ferrous Alloys, NoyesPublications, Park Ridge, NJ, 1985, pp. 822–862.

[14] N.R. Moody, S.L. Robinson, W.M. Garrison Jr., Hydrogen effects on theproperties and fracture modes of iron-based alloys, Res. Mech. 30 (1990)143–206.

[15] C. San Marchi, B.P. Somerday: technical reference on hydrogen compatibility ofmaterials, Report no. SAND2008-1163, Sandia National Laboratories,Livermore, CA, 2008. <http://www.sandia.gov/matlsTechRef/>.

[16] J.A. Brooks, A.W. Thompson, Microstructural development and solidificationcracking susceptibility of austenitic stainless steel welds, Int. Mater. Rev. 36(1991) 16–44.

[17] J.A. Brooks, A.J. West, Hydrogen induced ductility losses in austenitic stainlesssteel welds, Metall. Trans. A 12A (1981) 213–223.

[18] J.A. Brooks, A.J. West, A.W. Thompson, Effect of weld composition andmicrostructure on hydrogen assisted fracture of austenitic stainless steels,Metall. Trans. A 14A (1983) 75–84.

[19] M.J. Morgan, G.K. Chapman, M.H. Tosten, S.L. West, Tritium effects on fracturetoughness of stainless steel weldments, in: Trends in welding research:proceedings of the 7th international conference, ASM, International, 2005, pp.743–748.

[20] B.P. Somerday, M. Dadfarnia, D.K. Balch, K.A. Nibur, C.H. Cadden, P. Sofronis,Hydrogen-assisted crack propagation in austenitic stainless steel fusion welds,Metall. Mater. Trans. A 40A (2009) 2350–2362.

[21] H.F. Jackson, K.A. Nibur, C. San Marchi, J.D. Puskar, B.P. Somerday, Hydrogen-assisted crack propagation in 304l/308l and 21cr–6ni–9mn/308l austeniticstainless steel fusion welds, Corros. Sci. 60 (2012) 136–144.

[22] K.A. Nibur, B.P. Somerday, D.K. Balch, C. San Marchi, The role of localizeddeformation in hydrogen-assisted crack propagation in 21cr–6ni–9mnstainless steel, Acta Mater. 57 (2009) 3795–3809.

[23] C. San Marchi, B.P. Somerday, J. Zelinski, X. Tang, G.H. Schiroky, Mechanicalproperties of super duplex stainless steel 2507 after gas phase thermalprecharging with hydrogen, Metall. Mater. Trans. A 38A (2007) 2763–2775.

[24] W.C. Luu, P.W. Liu, J.K. Wu, Hydrogen transport and degradation of acommercial duplex stainless steel, Corros. Sci. 44 (2002) 1783–1791.

[25] W.J. Mills, Fracture toughness of type 304 and 316 stainless steels and theirwelds, Int. Mater. Rev. 42 (1997) 45–82.

[26] F.P. Ford, The crack-tip system and its relevance to the prediction of crackingin aqueous environments, in: R.P. Gangloff, M.B. Ives (Eds.), Environment-induced Cracking of Metals, NACE, Houston, TX, 1990, pp. 139–165.

[27] A.M. Brass, J. Chene, Hydrogen uptake in 316l stainless steel: Consequences onthe tensile properties, Corros. Sci. 48 (2006) 3222–3242.

[28] T. Zakroczymski, A. Glowacka, W. Swiatnicki, Effect of hydrogen concentrationon the embrittlement of a duplex stainless steel, Corros. Sci. 47 (2005) 1403–1414.

[29] J.R. Buckley, D. Hardie, The effect of pre-straining and d-ferrite on theembrittlement of 304l stainless steel by hydrogen, Corros. Sci. 34 (1993) 93–107.

[30] C. San Marchi, T. Michler, K.A. Nibur, B.P. Somerday, On the physicaldifferences between tensile testing of type 304 and 316 austenitic stainlesssteels with internal hydrogen and in external hydrogen, Int. J. HydrogenEnergy 35 (2010) 9736–9745.

[31] P. Müllner, C. Solenthaler, P.J. Uggowitzer, M.O. Speidel, On the effect ofnitrogen on the dislocation structure of austenitic stainless steel, Mater. Sci.Eng. A A164 (1993) 164–169.

[32] L. Vitos, J.-O. Nilsson, B. Johansson, Alloying effects on the stacking fault energyin austenitic stainless steels from first-principles theory, Acta Mater. 54 (2006)3821–3826.

[33] J. Talonen, H. Hänninen, Formation of shear bands and strain-inducedmartensite during plastic deformation of metastable austenitic stainlesssteels, Acta Mater. 55 (2007) 6108–6118.

[34] R.M. Latanision, J.A.W. Ruff, The temperature dependence of stacking faultenergy in Fe–Cr–Ni alloys, Metall. Trans. 2 (1971) 505–509.

[35] F. Lecroisey, A. Pineau, Martensitic transformations induced by plasticdeformation in the Fe–Ni–Cr–C system, Metall. Trans. 3 (1972) 387–396.

[36] F. Abrassart, Stress-induced c ? a martensitic transformation in two carbonstainless steels. Application to trip steels, Metall. Trans. 4 (1973) 2205–2216.

[37] J. Singh, Influence of deformation on the transformation of austenitic stainlesssteels, J. Mater. Sci. 20 (1985) 3157–3166.

[38] D.G. Ulmer, C.J. Altstetter, Hydrogen-induced strain localization and failure ofaustenitic stainless steels at high hydrogen concentrations, Acta Metall. Mater.39 (1991) 1237–1248.

[39] K.A. Nibur, D.F. Bahr, B.P. Somerday, Hydrogen effects on dislocation activity inaustenitic stainless steel, Acta Mater. 54 (2006) 2677–2684.

[40] ASTM, Standard test method for measurement of fracture toughness, ASTMInternational, West Conshohocken, PA, 2009.

[41] ASTM, Standard test method for J-integral characterization of fracturetoughness, ASTM International, West Conshohocken, PA, 1996.

[42] C. San Marchi, B.P. Somerday, S.L. Robinson, Permeability, solubility anddiffusivity of hydrogen isotopes in stainless steels at high gas pressures, Int. J.Hydrogen Energy 32 (2007) 100–116.

[43] D.J. Kotecki, T.A. Siewert, Wrc-1992 constitution diagram for stainless steelweld metals: a modification of the wrc-1988 diagram, Weld. J. 71 (1992) 171–179.

[44] C.D. Beachem, A new model for hydrogen-assisted cracking (hydrogen‘‘embrittlement’’), Metall. Trans. 60 (1972) 437–451.

[45] H.K. Birnbaum, P. Sofronis, Hydrogen-enhanced localized plasticity - amechanism for hydrogen-related fracture, Mater. Sci. Eng. A A176 (1994)191–202.

[46] D.P. Abraham, C.J. Altstetter, Hydrogen-enhanced localization of plasticity inan austenitic stainless steel, Metall. Mater. Trans. A 26A (1995) 2859–2871.

[47] C.G. Rhodes, A.W. Thompson, The composition dependence of stacking faultenergy in austenitic stainless steels, Metall. Trans. A 8A (1977) 1901–1906.

[48] R.E. Schramm, R.P. Reed, Stacking fault energies of seven commercialaustenitic stainless steels, Metall. Trans. A 6A (1975) 1345–1351.

[49] P.L. Mangonon Jr., G. Thomas, The martensite phases in 304 stainless steel,Metall. Trans. 1 (1970) 1577–1586.

[50] J.F. Breedis, L. Kaufman, Formation of hcp and bcc phases in austenitic ironalloys, Metall. Trans. 2 (1971) 2359–2371.

[51] L. Rémy, A. Pineau, Observation of stacked layers of twins and e martensite in adeformed austenitic stainless steel, Metall. Trans. 5 (1974) 963–965.

[52] G. Olson, M. Cohen, A general mechanism of martensitic nucleation: Part i.General concepts and the fcc ? hcp transformation, Metall. Trans. A 7A (1976)1897–1904.

[53] L. Rémy, The interaction between slip and twinning systems and the influenceof twinning on the mechanical behavior of fcc metals and alloys, Metall. Trans.A 12A (1981) 387–408.

[54] P. Müllner, C. Solenthaler, P.J. Uggowitzer, M.O. Speidel, Brittle fracture inaustenitic steel, Acta Metall. Mater. 42 (1994) 2211–2217.

[55] P. Müllner, On the ductile to brittle transition in austenitic steel, Mater. Sci.Eng. A A234–236 (1997) 94–97.

[56] S. Takaki, T. Furuya, Y. Tokunaga, Effect of si and al additions on the lowtemperature toughness and fracture mode of fe–27mn alloys, ISIJ Int. 30(1990) 632–638.

[57] J. Brooks, J. Williams, A. Thompson, Microstructural origin of the skeletalferrite morphology of austenitic stainless steel welds, Metall. Trans. A 14A(1983) 1271–1281.

[58] R. Ritchie, A. Thompson, On macroscopic and microscopic analyses for crackinitiation and crack growth toughness in ductile alloys, Metall. Trans. A 16A(1985) 233–248.

[59] A. Turnbull, Hydrogen diffusion and trapping in metals, in: R.P. Gangloff, B.P.Somerday (Eds.), Gaseous Hydrogen Embrittlement of Materials in EnergyTechnologies, Woodhead Publishing, Oxford, UK, 2012, pp. 89–128.

[60] M.I. Luppo, A. Hazarabedian, J. Ovejero-Garcia, Effects of delta ferrite onhydrogen embrittlement of austenitic stainless steel welds, Corros. Sci. 41(1999) 87–103.

[61] J.P. Hirth, Effects of hydrogen on the properties of iron and steel, Metall. Trans.A 11A (1980) 861–890.

[62] T.-P. Perng, C.J. Altstetter, Effects of deformation on hydrogen permeation inaustenitic stainless steels, Acta Metall. 34 (1986) 1771–1781.

[63] M.R. Louthan Jr., R.G. Derrick, Hydrogen transport in austenitic stainless steel,Corros. Sci. 15 (1975) 565–577.

[64] M. Nagumo, Hydrogen related failure of steels – a new aspect, Mater. Sci.Technol. 20 (2004) 940–950.

[65] T. Neeraj, R. Srinivasan, J. Li, Hydrogen embrittlement of ferritic steels:observations on deformation microstructure, nanoscale dimples and failure bynanovoiding, Acta Mater. 60 (2012) 5160–5171.