The Effect of Pre-Strain on the Resistance to Hydrogen Embrittlementin 316L Austenitic Stainless Steel

Il-Jeong Park1, Jae-Gil Jung2, Seo Yeon Jo1, Sang-Min Lee1 and Young-Kook Lee1,+

1Department of Materials Science and Engineering, Yonsei University, Seoul 120-749, Korea2Light Metal Division, Korea Institute of Materials Science, Changwon 642-831, Korea

The effect of pre-strain before hydrogen charging on the resistance to hydrogen embrittlement (HE) in the 316L austenitic stainless steelwas investigated through the slow strain rate tensile test (SSRT), transmission electron microscopy, and thermal desorption analysis (TDA). Thepre-strain suppressed mechanical twinning during the SSRT, regardless of hydrogen charging. However, it accelerated the ¾-martensitictransformation in hydrogen-charged specimens. The TDA revealed that whereas hydrogen atoms migrated from grain boundaries anddislocations mainly to the austenite (£)/¾ interfaces in pre-strained specimens during the SSRTs, they moved to the boundaries of freshmechanical twins, which newly formed during the SSRTs, in the annealed specimen. The elongation loss by hydrogen charging became greaterwith increasing the pre-strain, indicating that pre-straining deteriorated the resistance to HE. This elongation loss by pre-strain resulted from boththe increase in fraction of ¾-martensite with pre-strain and the segregation of hydrogen atoms to the £/¾ interfaces.[doi:10.2320/matertrans.M2014036]

(Received January 31, 2014; Accepted March 13, 2014; Published April 25, 2014)

Keywords: ¾-martensitic transformation, pre-strain, metastable austenitic stainless steel, mechanical twins, hydrogen embrittlement

1. Introduction

The Al6061 alloy has been used as a cylinder liner forhydrogen storage of a hydrogen fuel-cell vehicle (HFCV)because of its light weight and high resistance to hydrogenembrittlement (HE).1­3) However, the Al cylinder liner mustbe fully wrapped with strong carbon fiber reinforced plastic(CFRP) to endure the high pressure of the hydrogen gas.2)

Therefore, in recent high strength ferrous materials, partic-ularly austenitic stainless steel, are being considered as analternative liner for reducing the usage of expensive CFRP.The 316L austenitic stainless steel (STS 316L)4) can bea candidate material due to the higher resistance to HE,compared to the 304L austenitic stainless steel (STS304L).5­7) When slow strain rate tensile tests (SSRTs) wereperformed with hydrogen-charged annealed austenitic stain-less steels, STS 304L exhibited a brittle fracture due to thestrain-induced ¾- or ¡A-martensitic transformation.8­10) How-ever, STS 316L had neither the martensitic transformationnor the transition in fracture mode from ductile to brittle.5,11)

The effects of cold working on both the martensitictransformation and the resistance to HE in austenitic stainlesssteels have been also investigated. Unstable austeniticstainless steels, such as STS 301, 304, and 304L, underwentthe strain-induced ¾- or ¡A-martensitic transformation duringcold working for cylinder forming,8,10,11) which deterioratedthe resistance to HE after hydrogen charging.8,9) STS 316Ldid not undergo the ¾- or ¡A-martensitic transformation, butmechanical twinning during cold working.6,11) Mechanicaltwins also provide both nucleation sites for hydrogen-inducedcracking at their intersections and propagation paths alongtheir boundaries in high Mn austenitic steel.12) Therefore,deformed STS 316L is expected to have the lower resistanceto HE, compared to the annealed STS 316L.

However, the influence of cold working before hydrogencharging (hereafter, it is called pre-strain) on the resistance to

HE in STS 316L is yet to be systematically investigated,although this study is inevitable to understand the effect ofcold deformation for cylinder fabrication on the HE.

Therefore, the objective of the present study was toexamine the effect of pre-strain on the resistance to HE inSTS 316L through the SSRT and thermal desorption analysis(TDA).

2. Experimental Procedure

A plate of STS 316L with the chemical composition of Fe­17.5Cr­11.5Ni­4.6Mn­2.2Mo­0.008C (mass%) was used inthe present study. The 4-mm thick plate was reheated at1373K for 30min and hot-rolled at approximately 1273K to1.50, 1.67, and 2.15-mm plates, respectively. The hot-rolledplates were annealed at 1373K for 20min, followed by waterquenching. In the present study, the pre-strain was controlledby the thickness reduction during the room-temperature coldrolling of annealed plates. The annealed plates were cold-rolled to 1.50-mm thick sheets with different thicknessreductions of 0% (as-annealed), 10%, and 30%. The meangrain size of annealed specimens before cold rolling wasapproximately 70 µm.

The cold-rolled specimens were mechanically polished andetched with a mixed solution of HCl 50ml, HNO3 25ml, andH2O 25ml to observe the microstructure using an opticalmicroscope (OM, Olympus, BX41M). Changes in themicrostructure by heat treatment, deformation, and hydrogencharging were investigated using an X-ray diffractometer(XRD, RIGAKU, D/MAX-RINT 2700) with a Cu­K¡ target(­ = 0.15405 nm).

Tensile specimens with a gauge portion of 6.0mm wide,25.0mm long, and 1.5mm thick were machined from thecold-rolled sheets for the SSRTs. The tensile specimens wereelectrochemically charged with hydrogen in a 3% NaClaqueous solution containing 0.3% NH4SCN with a currentdensity of 50A/m2 for 48 h at room temperature. A platinumwire was used as a counter-electrode.+Corresponding author, E-mail: yklee@yonsei.ac.kr

Materials Transactions, Vol. 55, No. 6 (2014) pp. 964 to 970©2014 The Japan Institute of Metals and Materials EXPRESS REGULAR ARTICLE

The SSRTs were conducted with hydrogen-charged and-uncharged tensile specimens at a crosshead speed of 0.1mm/min, corresponding to an initial strain rate of 4.8 ©10¹5 s¹1 at room temperature using a servo hydraulicuniversal tensile testing machine (Instron, 3382). Theelongation loss, which indicates the degree of the reductionin total elongation before and after hydrogen charging, wasexpressed by the following equation:

El:loss ð%Þ ¼ ðEl:ðH-unchargedÞ � El:ðH-chargedÞÞ=El:ðH-unchargedÞ� 100 ð1Þ

where El.(H-uncharged) and El.(H-charged) are the total elongationof the hydrogen-uncharged and -charged specimens, respec-tively.

Mechanical twins and ¾-martensite were observed using afield-emission transmission electron microscope (FE-TEM,JEOL, JEM-2000EX) operated at 200 kV. Thin foils for TEMobservation were prepared by twin-jet polishing, which wasperformed at 288K with an applied potential of 20V in amixed solution of 90% glacial acetic acid and 10% perchloricacid. TEM samples were taken at a depth of approximately50 µm from the surface, where hydrogen atoms wereconcentrated, in the hydrogen-charged specimens afterSSRTs.

The amounts of diffusible hydrogen in the hydrogen-charged specimens were measured before and after theSSRTs through TDA during continuous heating from roomtemperature to 600K at a constant rate of 100K/h. Theamount of hydrogen, which was released from a hydrogen-charged specimen, was quantitatively analyzed at an intervalof 5min using helium as a carrier gas.

3. Results and Discussion

All STS 316L specimens before the SSRTs had anaustenite single phase without ¾- or ¡A-martensite regardlessof pre-strain and hydrogen charging, which was confirmed byXRD and OM analyses. However, mechanical twins wereobserved in pre-strained specimens and their amount wasproportional to the pre-strain. This result shows goodagreement with the previous result in which the density oftwins was higher in the strain-hardened STS 316 specimen,compared to the annealed one.11)

Figure 1(a) shows the engineering stress­strain curvesmeasured during the SSRTs with both hydrogen-charged and-uncharged specimens, which were annealed and pre-strainedwith different thickness reductions. Independent of hydrogencharging, both yield and tensile strengths of the specimenswere increased at the expense of total elongation withincreasing the pre-strain. The increase in strength with pre-strain was most likely due to the increase in densities ofdislocations13) and mechanical twins14) introduced by pre-strain before the SSRT.

Regarding the effect of hydrogen charging on tensileproperties, whereas both yield and tensile strengths of theannealed specimen (0%) were almost unchanged before andafter hydrogen charging, those of the pre-strained specimenwere significantly increased by hydrogen charging. The totalelongation was decreased in all specimens after hydrogencharging, regardless of the pre-strain. However, the elonga-

tion loss was increased from 14.8 to 54.6% with increasingthe pre-strain from zero to 30%. This result indicates that theresistance to HE was deteriorated by pre-strain.

To investigate the cause for the variation in elongation withhydrogen charging in both annealed and pre-strained speci-mens, the XRD tests of fractured specimens after the SSRTswere conducted. The XRD tests showed that all hydrogen-uncharged specimens had an austenite single phase without ¾-or ¡A-martensite, regardless of pre-straining. However, hydro-

(a)

(b)

(c)

Fig. 1 (a) Engineering stress­strain curves and (b) XRD patterns of bothannealed (0%) and pre-strained (10 and 30%) specimens with and withouthydrogen charging after the SSRTs. (c) The volume fraction of¾-martensite in hydrogen-charged specimens after the SSRTs. Thepercentage in figures indicates the amount of pre-strain.

The Effect of Pre-Strain on the Resistance to Hydrogen Embrittlement in 316L Austenitic Stainless Steel 965

gen-charged specimens displayed ¾-martensite (Fig. 1(b)),whose volume fraction was increased with the pre-strain(Fig. 1(c)).

To confirm the hydrogen distributions of both annealedand pre-strained specimens, the surfaces of fractured tensilespecimens were observed under a low magnification usinga SEM (Fig. 2). Whereas the edge part of the specimensshowed the brittle fractured surface, the center part exhibitedwell-developed dimples, a feature of the ductile fracturedsurface. The reason why the edge part of the specimensshowed the brittle fractured surface is because hydrogen wasconcentrated at the subsurface of specimens after elec-trochemical hydrogen charging. The area fraction of thebrittle fractured region was increased with increasing the pre-strain, as shown in Fig. 2(d). The reason that the increase inelongation loss with increasing the pre-strain (Fig. 1(a)) wasdue to the easy transition of the fracture mode from ductile tobrittle with pre-strain in the edge part of the specimen, wherehydrogen atoms were enriched.

To investigate the effect of the pre-strain on the ¾-martensitic transformation during the SSRT in hydrogen-charged specimens, the strain hardening behaviors of bothannealed and pre-strained specimens were examined becauseextra-strain hardening is expected if the ¾-martensitic trans-formation occurs during the SSRT. Figure 3(a) shows thestrain hardening rate (SHR, d·/d¾) - true strain (¾) curves ofhydrogen-charged and -uncharged specimens. The SHR waslowered with increasing the pre-strain, regardless of hydro-

gen charging. To analyze strain hardening behaviors more indetail, modified Crussard­Jaoul (C­J) plots are drawn inFigs. 3(b) to 3(d), which have been often used for theanalysis of deformation behaviors of TWIP steels.15­18)

The ln(d·/d¾) - ln · plots of both hydrogen-charged and-uncharged annealed specimens can be divided into fivestages (A to E) based on the slope change (Fig. 3(b)). Thisvariation in SHR in STS316L was similar to that in high MnTWIP steels. According to previous studies on the strainhardening behaviors of TWIP steels,17,18) the SHR rapidlydecreases at the early stage of plastic deformation due to thedynamic recovery of dislocations (stage A). The decrease inSHR is greatly slowed with further strain so that the SHRbecomes almost constant owing to active primary mechanicaltwinning (stage B). The SHR is lowered again due to lessactive primary mechanical twinning (stage C), increased byactive secondary mechanical twinning (stage D), and finallydecreases with the formation of thicker twin bundles(stage E).

To confirm that the variation of the SHR in the annealedSTS 316L specimen was caused by primary and secondarytwinning like that in TWIP steels, the hydrogen-chargedannealed specimen was tensile strained by 0.4, correspondingto the stage D in Fig. 3(b), and then its microstructure wasobserved using a TEM. Both primary (TW1) and secondarytwins (TW2) were simultaneously observed, as was expected.The thickness of mechanical twins was between approx-imately 5 and 30 nm, as shown in Fig. 4(a). This result means

Fig. 2 SEM fractographs of the hydrogen-charged (a) annealed, (b) 10% pre-strained, and (c) 30% pre-strained specimens. Dashed linesin the specimens indicate the boundaries between brittle and ductile fractured regions. (d) The change in area fraction of the brittlefractured region with pre-strain in the hydrogen-charged specimens.

I.-J. Park, J.-G. Jung, S. Y. Jo, S.-M. Lee and Y.-K. Lee966

that the SHR in the hydrogen-charged annealed STS 316Lspecimen was greatly dependent upon mechanical twinninglike that in annealed TWIP steels.

The 10% pre-strained specimen also exhibited five stagesbut with shortened stages B and D, regardless of hydrogencharging. It is noteworthy that the slope of stage D in thehydrogen-charged specimen was increased (Fig. 3(c)). Toexamine the cause for the increase in slope at stage D,after the 10% pre-strained specimen was strained by 0.2,corresponding to the stage D, its microstructure wasobserved using the TEM.

Not only primary (TW1) and secondary twins (TW2) butalso ¾-martensite plates were observed (Fig. 4(b)). Accord-ingly, the increased slope at stage D in the 10% pre-strainedspecimen is thought to be due to the formation of ¾-martensite as well as mechanical twinning.

Meanwhile, the 30% pre-strained specimen without hydro-gen charging revealed only three stages (A, C, and E), asshown in Fig. 3(d). This implies that stages B and D wereshortened with increasing the pre-strain. This result isprobably because many mechanical twins and dislocations,which were already introduced by pre-strain, hindered theformation of additional mechanical twins during the SSRTs.

Regarding the 30% pre-strained and hydrogen-chargedspecimen, four stages (A, C, D, and E) appeared. The slope atstage D was increased due to the ¾-martensitic transformation(Fig. 3(d)), which was similar to that in the hydrogen-

charged 10% pre-strained specimen (Fig. 3(c)). From theseresults, we realized that the ¾-martensitic transformationoccurred only in pre-strained and hydrogen-charged speci-mens during the SSRT and that primary mechanical twinningoccurred prior to the ¾-martensitic transformation.

Therefore, to elucidate the interrelationship between the¾-martensitic transformation, pre-strain, and hydrogen migra-tion, the TDAwas conducted using hydrogen-charged tensilespecimens with different pre-strains before and after theSSRTs. Figure 5(a) shows the change in the hydrogendesorption rate with temperature in the hydrogen-chargedspecimens before the SSRTs. The hydrogen-charged annealedspecimen exhibited only a peak (1st peak) between 275 and475K, which is well-known to form by the desorption ofhydrogen atoms mainly from grain boundaries and disloca-tions.19,20) The size of the 1st peak was almost the same,regardless of the pre-strain. This indicates that dislocations onthe specimen surface, which were introduced by pre-strain,did not accelerate the hydrogen permeation.

However, the hydrogen-charged pre-strained specimensshowed not only the 1st peak but also another small peak(2nd peak) between 510 and 560K. The 2nd peak is knownto be generated by the detrapping of hydrogen atoms from theboundaries of mechanical twins.19) The size of the 2nd peakwas increased with pre-strain unlike that of the 1st peak,which was due to the fraction of mechanical twins increasedby pre-strain.

(a)

(b)

(c)

(d)

Fig. 3 (a) Strain hardening rate (d·/d¾) curves of annealed (0%) and pre-strained (10 and 30%) specimens. ln(d·/d¾) - ln · plots for themodified Crussard­Jaoul analysis of (b) annealed, (c) 10% pre-strained, and (d) 30% pre-strained specimens. All the data were generatedusing hydrogen-uncharged and -charged specimens.

The Effect of Pre-Strain on the Resistance to Hydrogen Embrittlement in 316L Austenitic Stainless Steel 967

The total amount of hydrogen released from the specimen,which was calculated from both the 1st and 2nd peaks, was4.45, 4.68, and 5.12mass ppm in the annealed, 10%, and30% pre-strained specimens, respectively. This increase intotal amount of hydrogen with pre-strain was caused by thesize of the 2nd peak increased with pre-strain.

Figure 5(b) shows the change in the hydrogen desorptionrate with temperature in the hydrogen-charged specimens,which were fractured after the SSRTs. Whereas the width ofthe 1st peak was broadened in all specimens, its height wasdecreased, compared to the counterparts measured beforethe SSRTs (Fig. 5(a)). According to Chun et al.,21­24) theactivation energy (22 kJ/mol) for detrapping of hydrogenatoms from austenite (£)/¾ interfaces was lower than thosefrom grain boundaries and dislocations (between 26 and35 kJ/mol).19,25) Therefore, it is thought that the reason forpeak broadening is probably due to the release of hydrogenatoms which were migrated to the £/¾ interfaces during theSSRTs. The peak broadening before and after the SSRTs wasproportional to the amount of pre-strain, which shows goodagreement with the increase in fraction of ¾-martensite withthe pre-strain.

To separate the amount of hydrogen detrapped from the£/¾ interfaces from that of hydrogen desorpted from bothgrain boundaries and dislocations in tensile specimensfractured after the SSRTs, the 1st peaks between 275 and460K were fitted by the Gaussian function using themaximum temperatures of two different peaks (Figs. 5(c) to5(e)). Whereas the peak drawn by the circle marker resultedfrom the desorption of hydrogen atoms from grain boundariesand dislocations, the peak drawn by the square marker wascaused by detrapping of hydrogen atoms from the £/¾interfaces. The areas of the two peaks were converted to theamounts of hydrogen, which are listed in Table 1. The amountof hydrogen desorpted from the £/¾ interfaces more greatlyincreased with increasing the pre-strain, compared to that ofhydrogen detrapped from grain boundaries and dislocations.

Meanwhile, the decrease in the height of the 1st peak canbe explained by the migration of hydrogen atoms from grainboundaries and dislocations to both £/¾ interfaces andmechanical twins, which formed during the SSRTs. Themigration of hydrogen atoms to £/¾ interfaces caused thebroadening of the 1st peak and that to mechanical twinsincreased the size of the 2nd peak.

The size of the 2nd peak was increased with decreasingthe pre-strain, which was completely opposite to the resultobtained before the SSRTs (Fig. 5(a)). This result can beexplained as follows: When the pre-strain was large,hydrogen atoms migrated from grain boundaries anddislocations toward the £/¾ interfaces (that is, to the 1st¾peak in Fig. 5(c)). When the pre-strain was small, hydrogenatoms migrated mainly to mechanical twins (that is, to the2nd peak), which newly formed during the SSRTs (hereafter,they are called fresh mechanical twins), not to mechanicaltwins introduced by pre-strain before hydrogen charging(hereafter, they are called pre-existing mechanical twins).The reason for the migration of hydrogen atoms to freshmechanical twins is probably because the boundaries of freshmechanical twins were hydrogen-free, whereas the bounda-ries of pre-existing mechanical twins were already hydrogen-enriched.

Consequently, the reason why the elongation loss of thehydrogen-charged pre-strained specimens was great wasbecause the ¾-martensitic transformation was accelerated inthese specimens and hydrogen atoms migrated to the £/¾interfaces during the SSRTs, resulting in the HE.

4. Conclusions

(1) The elongation loss became greater with increasing the

Fig. 4 The bright-field TEM images and SAED patterns of (a) the annealedspecimen strained by 0.4 and (b) the 10% pre-strained specimen strainedby 0.2. The applied strain for each specimen corresponds to the stage Don the ln(d·/d¾) - ln · plot in Fig. 2.

Table 1 The hydrogen amount (ppm) converted from each peak on thehydrogen desorption rate curves (Figs. 4(c) to 4(e)) measured usingannealed (0%) and pre-strained (10 and 30%) specimens after slow strainrate tensile tests.

(ppm)

Pre-strainHydrogen amount of

1st¾ peakHydrogen amount of

1stgb peakTotal hydrogen

amount

0% 0.15 2.57 2.72

10% 0.38 2.76 3.14

30% 1.21 2.84 4.05

I.-J. Park, J.-G. Jung, S. Y. Jo, S.-M. Lee and Y.-K. Lee968

pre-strain because the fraction of ¾-martensite wasincreased with pre-strain and hydrogen atoms weresegregated to the £/¾ interfaces, leading to the hydrogenembrittlement.

(2) The strain hardening rate curves of both hydrogen-uncharged and -charged annealed specimens, whichwere measured during the slow strain rate tensile tests(SSRTs), were divided into five stages based on themechanical twinning behavior without the ¾-martensitictransformation. The pre-strain made mechanical twin-ning during the SSRT less active, and insteadaccelerated the ¾-martensitic transformation.

(3) The hydrogen-charged annealed specimen before theSSRT exhibited only a peak (1st peak) between 275 and475K by the desorption of hydrogen atoms from bothgrain boundaries and dislocations. However, hydrogen-charged pre-strained specimens showed not only the 1stpeak but also the 2nd peak between 510 and 560K bythe detrapping of hydrogen atoms from the boundariesof mechanical twins. The total amount of hydrogenstored inside the specimen was increased with pre-strainprimarily due to the increased fraction of mechanicaltwins with pre-strain.

(4) After hydrogen-charged annealed and pre-strainedspecimens underwent the SSRTs, the width of the 1stpeak was broadened, compared to that measured beforethe SSRTs. The peak broadening before and after theSSRTs was proportional to the pre-strain. These resultswere caused by both the increase in fraction of ¾-martensite with pre-strain and the segregation of

hydrogen atoms to the £/¾ interfaces. Meanwhile, theheight of the 1st peak was decreased, compared to thatmeasured before the SSRTs, due to the migration ofhydrogen atoms from grain boundaries and dislocationsto both £/¾ interfaces and the boundaries of freshmechanical twins.

(5) The size of the 2nd peak after SSRTs was increasedwith decreasing the pre-strain, which was completelyopposite to the result obtained before the SSRTs. Whenthe pre-strain was small, hydrogen atoms migrated fromgrain boundaries and dislocations to the boundaries offresh mechanical twins. When the pre-strain was large,they moved mainly toward the £/¾ interfaces.

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