Corrosion Science 75 (2013) 176–183

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Corrosion Science

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Role of chlorides on pitting and hydrogen embrittlement of Mg–Mnwrought alloy

0010-938X/$ - see front matter � 2013 Elsevier Ltd. All rights reserved.http://dx.doi.org/10.1016/j.corsci.2013.05.030

⇑ Corresponding author. Tel.: +91 22 2576 7892; fax: +91 22 2572 3480.E-mail address: vsraja@iitb.ac.in (V.S. Raja).

V.S. Raja a,⇑, Bharat S. Padekar a,b,c

a Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology Bombay, Mumbai 400 076, Indiab IITB-Monash Research Academy, Indian Institute of Technology Bombay, Mumbai 400076, Indiac Department of Mechanical and Aerospace Engineering, Monash University, Melbourne, Vic. 3800, Australia

a r t i c l e i n f o

Article history:Received 25 March 2013Accepted 22 May 2013Available online 7 June 2013

Keywords:A. magnesiumB. SEMC. Hydrogen embrittlementC. Stress corrosion

a b s t r a c t

The role of chlorides on stress corrosion cracking behavior of Mg–Mn hot rolled alloy was studied inMg(OH)2 saturated, 0.01 M and 0.1 M NaCl solutions. The alloy was found to fail by hydrogen embrittle-ment mechanism both in presence and absence of chlorides. However, the role of chloride has been foundto be to damage the passive film, cause pitting and increasing hydrogen embrittlement tendency of thealloy. Crack initiation occurred through pitting and grew in a transgranular manner involving hydrogen.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

The need to use lighter engineering materials in transport andaerospace industries is growing stronger with depleting oil andgas resources and the demand to bring down environmental pollu-tion. Magnesium and its alloys being the lightest engineeringmetallic material, they are being explored for extensive applicationin these industries. One of the problems, facing these alloys is theirsusceptibility to environmentally assisted cracking (EAC) [1–14] invariety of environments namely, chlorides [1,2,4–7], carbonates[8], sulfates [9]; the alloy embrittled even in distilled water [1,5–7,10,13]. In all these cases, the EAC mechanism has been attributedto hydrogen embrittlement [1–7,12–14]. However, a detailedunderstanding of various factors causing EAC is still lacking. Forexample, pitting has been found to enhance EAC tendency of mag-nesium alloy [1,2]. Kannan and Dietzel [11] have pointed out arelation between pitting tendency and EAC in the case of AZ80 al-loy through a comparison of potentiodynamic polarization andslow strain rate (SSRT) studies. They further indicated that thepre exposed AZ80 alloy in 0.5 wt.% NaCl solution exhibited higherembrittlement when tensile tested immediately after exposurethan testing the specimen after desiccating for 7 days [11]. The lat-ter showed better ductility than the former. The role of chlorideson the stress corrosion cracking susceptibility of austenitic stain-less steel is well known. Several mechanisms such as film break-down and absorption induced dislocation emission have been

proposed in the case of austenitic stainless steel [15]. However,the role of chlorides on EAC behavior of magnesium alloys is notwell studied. Hence, the objective of this work is to investigatethe role of chlorides on EAC behavior of magnesium alloy. It shouldbe emphasized that most of the magnesium alloys on which EACstudies were carried out corresponded to as-cast structure. A castalloy, due to its inherent chemical segregation can cause localizedcorrosion, even in the absence of chlorides and hence is not ideal tostudy the effect of chlorides. Hence, in the present case a wroughtalloy of composition Mg–0.1Al–0.3Zn–1.75Mn–0.25Ce-0.3(otherrare earth), has been undertaken to investigate the role of chlo-rides. Two different chloride concentrations along with distilledwater have been chosen as environments to examine the role ofchlorides on EAC of Mg–Mn alloy. Notably, no published literatureis available on the EAC behavior of this alloy, which deals withcrack initiation.

2. Experimental

2.1. Material

Magnesium based hot rolled alloy that was annealed at 300–350 �C for 30 min and then air cooled was commercially availablein the thickness of 12.2 mm. Its chemical composition in weight%is presented in Table 1.

For microstructural examination using optical microscopy,specimens were abraded successively till 5000 grit size (4.5 lm)on silicon carbide papers and finally on cloth using diamond pasteof 0.25 lm and were etched in a solution of ethanol 100 ml, water

Table 1Chemical composition of the test alloy.

Elements Al Zn Mn Ce REa Si Be Fe Ni Cu

wt.% 0.1 0.3 1.75 0.25 0.3 0.1 0.002 <0.005 0.007 <0.005

a RE – rare earth elements (other).

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20 ml, acetic acid 5 ml and picric acid 6 g. Grain size was deter-mined using the mean linear intercept method [16]. SEM micro-graph on longitudinal section surface of plate was obtained tosee the attack near second phase particles after etching for 30 sin aqueous solution of 2% oxalic acid. Energy dispersive X-ray spec-troscopy (EDX) analysis was carried out for characterization ofbroad chemical composition of the secondary-phase particles. X-ray diffraction analysis was done on abraded specimen to analyzephases present in the alloy. PANalytical X’Pert pro MRD make ma-chine with Cu Ka radiation of 1.5406 nm wavelength with Ni filterwas used. The 2h scan was carried out from 20� to 90� with a stepsize of 0.020�.

2.2. Tests specimens and environment

Cylindrical tensile specimens having gauge length 30 mm anddiameter 6 mm were machined from the rolled plate having onelot of axis alignment in the rolling direction (longitudinal) and an-other lot in the long transverse. Specimens of both the lots wereused for SSRT and long transverse alignment was used for constantload test (CLT). The gauge surface was polished in the axial direc-tion until 5000 grit size. The specimens were cleaned using acetoneimmediately before loading on SSRT or CLT. The test environmentsfor SSRT were 0.01 M and 0.1 M NaCl solutions, both saturatedwith Mg(OH)2, and distilled water. SSRT tests were also carriedout in the inert environment of glycerol. The environment usedin CLT was 0.1 M NaCl solution saturated with Mg(OH)2. The envi-ronment cell was designed such that the gauge length remainedimmersed in the test environment during the entire test durationand the volume of corrosive environment was about �40 ml cm�2

of the immersed surface.

Fig. 1. (a) Optical micrograph, (b) SEM micrograph after 2% oxalic acid etches;arrow shows the rolling direction.

2.3. Slow strain rate testing

For magnesium alloys most of the SCC studies using SSRT aredone at 10�6 s�1 [1,2,4]. In the present study all the specimenswere tested with the crosshead travel speed of 0.0017 mm min�1,which is equivalent to a strain rate of 10�6 s�1. United CalibrationCorporation (STM-20) tensile testing machine was used for theSSRT experiments. The corrosion cell is made to fit around theshoulder of the specimen, the arrangement of test specimen, corro-sion cell, and grip on both the sides is similar to that shown inschematic for CLT rig in Elsevier [2]. The cell was filled with testsolution before starting the machine for SSRT. The load was appliedby employing the constant extension rate mode. Tests were per-formed on the specimens having longitudinal and long transverseaxis alignment. A pre-load of �20 MPa was applied to take careof the machine slack before each test. The load and the extensionwere monitored continuously by a load cell and an optical encoder.As there was no extensometer attached to the specimen gaugelength, SSRT tests produced apparent stress–strain curves. Theelongation (ef) was obtained using actual measured value on failedspecimen. The fracture surfaces were cleaned in boiling chromicacid (180 g l�1 in water), washed in distilled water and rinsed withacetone, before examination using scanning electron microscope(SEM).

2.4. Constant load testing

In the CLT three similar rigs were used. Schematic of CLT rig andarrangement for loading and determining the exact time of fail isgiven Elsevier [2]. The corrosive environment (0.1 M NaCl solutionsaturated with Mg(OH)2) was filled in the cell when the desiredload was achieved, and replaced weekly as well as any reasonableloss due to evaporation was replenished by topping up with dis-tilled water. A plot of applied stress vs. time to fail provided thethreshold stress (rSCC), below which there is no SCC failure. Testswere terminated after 1030 h. The specimens that survived1030 h were forcibly fractured. Fracture surfaces were cleanedand dried (as described in Section 2.3) and examined using SEM.

3. Results

3.1. Microstructure

A typical optical micrograph of the alloy under study is shownin Fig. 1. Equiaxed grains characteristics of wrought alloy are seen.

178 V.S. Raja, B.S. Padekar / Corrosion Science 75 (2013) 176–183

Based on the linear intercept method [16], the alloy was found toexhibit an average grain size of 12 ± 3 lm. Scanning electronmicrograph obtained in back scattered electron image mode on alongitudinal section revealed the presence of second phases, whichwere found to align along the rolling direction. Chemical composi-tion obtained through EDX analysis of these particles is summa-rized in Table 2. These particles contain Mn and Ce apart fromMg. The size of the particles is small and so an exact compositioncould not be determined except indicating the chemical constitu-ents of these particles. X-ray diffraction (XRD) patterns of the alloyshowed peaks corresponding to Mg, (Mg, Mn)12Ce and Mn (Fig. 2).(Mg, Mn)12Ce precipitates are reported to be a ternary intermetal-lic compound (Mg, Mn)12Ce having the same tetragonal structureas that of Mg12Ce [17]. The Mn content in Mg12Ce has been re-ported to very between 0.3 and 0.6 at.%, depending on alloy com-position and quenching temperature [17]. Earlier reports showedthese precipitates to contain 0.59 at.% Mn [18]. Ce addition hasbeen reported to facilitate the formation of Mg12Ce in Mg-base al-loys [19]. When Mn is added, the precipitates become (Mg, Mn)12-

Ce. Rare earth has been found to form compounds (Mg7Zn3RE)even in Zn containing alloy, ZE41 [20–22]. On the other hand,Pan and Yang [23] reported the formation of only Mg12Ce in Mg–3Sn–1Mn alloy when 0.87 wt.% Ce was added even though the al-loy had 1%Mn. Al is added to Mg–Mn alloys to form AlMn precip-itate [24]. Formation of Al8Mn5 phase was reported for AM50 alloy[25]. However in present alloy we were not able to identify AlMnphase may be because the alloy contains only 0.1 wt.% Al. The otheraspect needs to be pointed out here is that the XRD pattern showeda small peak corresponding to manganese phase, although the SEMstudy could not reveal the presence of such phase. A similar obser-vation has been made by other authors in the case of Mg–1.8Mn–5.6Ce alloy [17].

3.2. Slow strain rate testing

Apparent stress–strain curves of specimens having longitudinaland long transverse axes alignments tested at 10�6 s�1 strain ratein glycerol (considered as inert environment), distilled water and

Table 2Atomic % of the elements detected in EDX analysis on the abraded surface of Mg–Mnalloy of the two different points on second phase particles.

Point Mg Mn Ce

1 98.16 0.49 1.182 98.63 0.75 0.58

Fig. 2. X-ray diffractogram of Mg–Mn alloy surface.

Mg(OH)2 saturated 0.01 M NaCl and 0.1 M NaCl solutions are pre-sented in Fig. 3. Summary of SSRT data of both the alignments ispresented in Table 3. The reported mechanical properties namelyelongation (ef) and ultimate tensile strength (UTS) are average oftriplicate test results rounded to their respective nearest digit.The alloy showed 140 and 200 MPa respectively for yieldstrength(rYS), which is at 0.2% offset strain, and UTS when testedin air at a strain rate of 10�4 s�1 using an extensiometer in aZwick/Roell universal testing machine (UTM). Notably, the alloydid not show much variation in both rYS and UTS values betweenthe longitudinal and long transverse alignments. However, theelongation of the alloy varied marginally between the alignments.Along the longitudinal alignment the alloy showed 30% ef; whilethat along the long transverse it showed 28% ef. A 3% scatter wasfound among the data of triplicate tests.

Mg–Mn alloy tested in the longitudinal alignment in glycerol at10�6 s�1 strain rate showed a slightly higher ef (38%) and UTS

Fig. 3. Apparent stress–strain curves of (a) longitudinal, and (b) long transversealignment of specimens tested in: Gly: glycerol, Dw: distilled water, 0.01 M and0.1 M NaCl solutions saturated with Mg(OH)2, at strain rate 10�6 s�1.

Table 3Summary of SSRT results of longitudinal and long transverse alignment of axis.

Alloy ef (%) UTS (MPa)

Gly DW 0.01 M 0.1 M Gly DW 0.01 M 0.1 M

Longitudinal 38 35 16 7.5 170 170 160 150Transverse 37 34 15 6.5 165 165 155 140

Gly: glycerol, DW: distilled water, 0.01 M and 0.1 M NaCl solutions saturated withMg(OH)2.

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(170 MPa) than those obtained in the long transverse alignment(Table 3). At the same strain rate, the alloy tested in distilled watershowed a marginally lower ductility but almost the same UTS ascompared to that obtained in glycerol in both the alignments. Onthe contrary, the alloy exhibited significantly lower values in0.01 M NaCl saturated with Mg(OH)2. When the chloride levelwere increased to 0.1 M NaCl a further significant reduction in ef

and UTS values were observed. In chlorides, a significant influencewith respect to SCC has been found in the long transverse align-ment than along the longitudinal direction (Table 3).

3.3. Constant load testing

Constant load tests were carried out using the most aggressiveof the four test environments employed in the SSRT study, namely,0.1 M NaCl solution saturated with Mg(OH)2. Similarly the speci-mens were tested in the long transverse alignment, as the alloywas found to be more susceptible to SCC along this alignment than

Fig. 4. Plot of stress vs. time to fail for Mg–Mn alloy in CLT condition in 0.1 M NaClsolution saturated with Mg(OH)2 (arrows indicate no failure).

Fig. 5. SSRT specimen tested in 0.1 M NaCl saturated Mg(OH)2: (a) side, and (b) transversepits and initiation of cracks from them, and (c) longitudinal section view of transverse axphase particles.

in the longitudinal alignment. The specimens were loaded at differ-ent initial stress levels and each test was conducted for a maxi-mum duration of 1030 h and discontinued thereafter even if thespecimen did not fail. It should be mentioned that this test dura-tion was higher than the exposed test periods reported in the liter-ature for similar studies carried out on Mg-alloys [26–32]. Stressvs. time to fail plot obtained for Mg–Mn alloy using CLT is shownin Fig. 4. The Mg–Mn alloy did not fail when the applied stresseswere6126 MPa (Fig. 4). Based on the data an empirical relation be-tween the time to fail and applied stress as given in Eq. (1) isobtained.

tf ¼ �263 lnðr� 12674

Þ ð1Þ

where tf is time to fail (in h) and r is applied stress (in MPa).Fitting the data in Eq. (1) yielded threshold stress as 126 MPa

(which is 90% of the rYS). The 200 MPa value for instant failure(in Fig. 4) corresponds to the ultimate tensile strength obtainedby a tensile test.

4. Fractography

Scanning electron microscope was employed to analyze theoverall cracking tendency of the alloy. For illustration, low magni-fication images of a failed specimen of long transverse axis align-ment exposed to Mg(OH)2 saturated 0.1 M NaCl solution andtested under SSRT are shown in Fig. 5. The side (Fig. 5a) and trans-verse (Fig. 5b) images of the specimen revealed the presence ofsurface pits where the cracks originated. The depth and the num-ber of pits were found to increase with chloride content, thoughmild pitting was noticed even on specimens exposed to distilledwater. Stampella et al. [9] have reported that the cracks originatedfrom pits in the case of pure magnesium even when tested in chlo-ride free Na2SO4 solution. To understand SCC initiation and growth,longitudinal section of the long transverse axis alignment of sev-eral specimens were examined. As shown in Fig. 5c, though SCCis initiated by pits the second phase particles aligned along therolling direction (also in Fig. 6a) have become selective path for

macrographs of the specimen along the transverse axis showing localized corrosionis alignment shows SCC crack that had initiated from the pit and grew along second

Fig. 6. CLT specimen loaded to 140 MPa and failed after 460 h. (a) shows pits on the edge (shown by arrows), (b) evidence of dissolution and selective dissolution of globularprecipitates (Detail A), (c) beach mark (Detail B), and (d) SCC propagation in TG manner (Detail C). The in set in d shows typical ductile failure of a specimen failed in SSRT inglycerol revealing dimple features.

180 V.S. Raja, B.S. Padekar / Corrosion Science 75 (2013) 176–183

the stress corrosion cracks to grow, there by indicating the influ-ence of aligned second phase particles ((Mg, Mn)12Ce) on SCC sus-ceptibility of the alloy. As a consequence, cracks in longitudinalspecimens subjected to SSRT were found to be shorter than thatfound in long transverse specimens.

Fractographic features of failed (after 460 h at 140 MPa) CLTspecimens shown in Fig. 6 further bring out microscopic crackgrowth mechanism of EAC in the alloy. As seen from Fig. 6a, pitson the edge were deep that terminated into transgranular stresscorrosion cracks. Also the fractograph (Fig. 6a) exhibited finger likepatterns mostly aligned in the direction of rolling. This aligned fea-tures support the earlier observation (Fig. 5c) that aligned secondphase particles (Fig. 1b) facilitate crack growth. Though the crackinitiation region (Fig. 6a) looked like a simple pit, examination ofthis area at high magnification revealed the presence of secondarycracks (Fig. 6b) within the so called pit indicating a possible syn-ergy between pitting and crack growth. This implies that eventhe pit growth might be assisted by the applied stress. The pit sur-face still remained smooth. However, the fracture features weremore evident in region B, as brought in Fig. 6c. Here the crackswere predominantly secondary in nature and even beach marksrepresenting discontinuous crack growth were seen (Fig. 6c). Nota-bly, the region shown in Fig. 6c had less number of micropits thanthat observed in Fig. 6b. This might be possibly because of the factthat the region close to the specimen edge, as shown in Fig. 6a,might be exposed for longer duration to environment than thespecimens towards the centre (Fig. 6c). Hence, the former couldsuffer more localized attack such as selective dissolution and pit-ting. In contrast to the regions A and B, the region C, as broughtout in Fig. 6d, showed a distinctly different fracture feature. Thefeature seen in region C (Fig. 6d) is similar to those generally re-ported in the literature [3,33] and is attributed to hydrogenembrittlement. It appears that under constant load condition, thestress concentration at a growing crack tip increases with time.As a consequence the crack velocity increases with time, whilethe electrochemical dissolution suffered by the specimen de-creases. Therefore, the so-called beach marks characteristics ofslow growth rate as seen in region B (Fig. 6c) did not appear. Itshould however be noted that this region suffered hydrogen

embrittlement, similar to that reported earlier [3,4,33] and notover load failure. A typical fractograph of the alloy failed in SSRTin glycerol environment revealing dimple features characterizingductile failure is shown as inset of Fig. 6d for comparison.

5. Discussion

In order to quantify SCC susceptibility of the magnesium alloysubjected to SSRT, the data presented in Table 3 for ef and UTSwas quantitatively evaluated to obtain the susceptibility index(ISCC) for each case as per the following equation [1,2]:

ISCC ¼ðeGly:%� eSCC%Þ

eGly:%� 100 ð2Þ

where eGly and escc are the elongation measured at the end of thetest in glycerol and the environment of interest respectively. (Sim-ilar calculations were also done using UTS values). Here glycerol isconsidered as an inert environment for comparison purpose. SCCsusceptibility index, ISCC, expressed here is percentage loss of prop-erty. Higher ISCC indicates higher SCC susceptibility of alloy in a par-ticular environment with respect to inert environment.

The SCC susceptibility indices for Mg–Mn alloy in the longitudi-nal and the long transverse alignments of the specimens obtainedat 10�6 s�1 strain rate for the test environment are compared inFig. 7. The data showed interesting results on the effect of speci-men orientation and the nature of environment on the SCC suscep-tibility of the alloy.

(a) Overall, the long transverse specimens exhibited high sus-ceptibility to SCC than that of the longitudinal specimen.

(b) SCC susceptibility increased with an increasing severity ofthe environment. It can be assumed that the increasingseverity of the environment increased with chlorideconcentration.

(c) There was a synergy between (1) and (2). That is, the differ-ence in the SCC susceptibility between the long transverseand the longitudinal specimens increased with an increasingseverity of the environment.

Fig. 7. Comparison of SCC susceptibility index in different environments based onmeasured values (Table 3) of three identical specimens: (a) ef (%), and (b) UTS.Scatter of <3% was observed in the experimental data.

V.S. Raja, B.S. Padekar / Corrosion Science 75 (2013) 176–183 181

The possible reason as why the long transverse alignment spec-imens suffered higher SCC susceptibility than the longitudinalalignment specimens can be explained as follows. The precipitateswere aligned along the rolling direction (see Fig. 1b) which couldfacilitate growth of cracks along the matrix/precipitate interfaceas the interface suffered selective corrosive attack. Second phasesaffect localized corrosion/selective attack of magnesium alloys[34,35]. As the extent of selective attack could increase with theseverity (chloride content) of the environment, SCC susceptibility,as indicated by ISCC, is expected to increase, especially on the longtransverse specimens.

Accordingly, the alignment of (Mg,Mn)12Ce phases along thecrack growth direction in long transverse specimen has causedmore lowering of UTS (140 MPa) than longitudinal specimens(150 MPa) in Mg(OH)2 saturated 0.1 M NaCl solution. On the con-trary, Song et al. [36] reported ZE41 alloy to be more susceptiblealong the longitudinal alignment than along the long transversealignment in 0.01 M NaCl solution. Notably, the above suggestionwas made based on the reduction in area values. However, their re-ported UTS data indicated the fact that even ZE41 alloy was moresusceptible along long transverse alignment than along the longi-tudinal alignment. It is also important to point out the fact thatZE41 showed a steep reduction in elongation from 25.3% to 1.4%in 0.01 M NaCl solution, where as the present alloy showed onlya smaller drop in elongation from 38% in glycerol to 16% in0.01 M NaCl. The present environment however may be considered

as relatively milder in relation to what employed by Song et al.[36], as the present environment is buffering with Mg(OH)2.

The residual tensile strength (as UTS) of the specimens survivedafter 1030 h of CLT in 0.1 M NaCl solution saturated with Mg(OH)2

were determined by a tensile test carried out at 10�4 s�1 strainrate. All the specimens showed UTS value of 195 ± 1 MPa, whichis only marginally (2.5%) lower than 200 MPa, to the later being ob-tained on a specimen tested in air. However, the specimensshowed only 13 ± 0.2% ef, which is far lower than 28% ef obtainedon specimens tested in air.

This can be attributed to stress concentration because of pitsand/or hydrogen absorption of the alloy as a result of 1030 himmersion in 0.1 M NaCl solution in CLT condition (Fig. 7b). Thefact that the residual UTS values of the exposed specimen wereonly marginally lower than that of the unexposed specimen whileshowing a significant loss in ductility, as measured by elongation,suggests that pits did not exert as much notch effect to reduceUTS, as hydrogen did to embrittle the alloy.

The role of chloride on stress corrosion cracking behavior of thepresent alloy is examined. Examination of the role of chlorides onthe hydrogen evolution kinetics on this alloy is interesting. It indi-rectly accelerates the hydrogen evolution kinetics as has been ex-plained in the ensuing discussion. The potential differencebetween Mg2+ + 2e� = Mg (E� = �2.363 V standard hydrogen elec-trode (SHE)) and 2H+ + 2e� = H2 (E� = 0.000 V SHE) is large enoughto drive the hydrogen evolution reaction. However, the oxide/hydroxide passive film formed on Mg alloys seems to lower thehydrogen reduction kinetics (H+ + e� = ½ H2). Therefore passivefilm/oxide breakdown by either of the following means, namely(1) electrical breakdown, (2) chemical breakdown, and (3)mechanical breakdown and in any combination of the above canaccelerate hydrogen evolution reaction (cathodic reaction) onmagnesium and its alloys. A raise in potential (in anodic direction)causes film break down and as a consequence more hydrogen hasbeen found to liberate on Mg alloys. The ‘‘negative difference ef-fect’’ observed on Mg alloys [34], during anodic polarization has in-deed been attributed to passive film breakdown under theinfluence of potential and the consequent increase in hydrogenevolution [37]. Such observations are also made on aluminum al-loys as well [38].

In the present case the chlorides were responsible for oxide/hydroxide film break down/pitting which in-turn facilitate H2 evo-lution. An increase in chloride concentration can (a) raise the pit-ting susceptibility of the alloy (b) stifle the repassivationtendency of a pit or crack. Creation of bare metal surface due toany of this phenomenon can facilitate hydrogen evolution and thenhydrogen embrittlement.

In order to examine the above proposition, tests were con-ducted to determine the role of Cl� on (a) the repassivation kineticsand (b) H2 evolution kinetic. Our studies on the present alloyexposing it to different chlorides indeed support the first point thatchlorides raise the pitting susceptibility of the alloy by the fact thatan increase in chloride content of the medium was found to in-crease both the number and depth of pits. Since, such observationsare on the expected lines this data is not presented.

For repassivation studies, specimens were held at �1.48 V(standard calomel electrode (SCE)), which lies in the passive regionof the alloy in Mg(OH)2 saturated solution of 0.01 M NaCl. Afterreaching a stable passive current, a scratch was made using aglass-tip. The variation in current was measured and the same isshown in (Fig. 8). Comparison of the data shows that chloridesnot only significantly increased peak current but also loweredthe rate of current decay (rate of repassivation). These experimentssupport the above proposition that Cl� damages the film. Filmdamage in turn is expected to enhance H2 evolution kinetics.

Fig. 8. Potentiostatic tests at �1.48 V (SCE) compare variation of the current due toscratch on alloy exposed to Dw-distilled water (insets) and 0.1 M and 0.01 M NaClsolutions saturated with Mg(OH)2. Note that the chlorides significantly increasedthe peak current due to scratching and lowered the rate of current decay(repassivation rate).

Fig. 10. The plot of pit size vs. applied stress along with time to fail vs. Appliedstress in CLT specimens tested in 0.1 M NaCl saturated with Mg(OH)2 (arrowindicates no failure).

182 V.S. Raja, B.S. Padekar / Corrosion Science 75 (2013) 176–183

The influence of chlorides on H2 evolution is examined directly.Hydrogen evolution test was done using fishing line specimen [39]for which the procedure was followed as mentioned by Venugopalet al. [38], Shi et al. [40], and Liu et al. [41]. Fig. 9 shows the cumu-lative hydrogen evolution occurred on Mg–Mn alloy with immer-sion time for two different chloride levels, namely, 0.01 M and0.1 M NaCl saturated with Mg(OH)2. It is quite clear, that an in-crease in chloride concentration caused a significant increase inhydrogen evolution kinetics.

The role of applied stress is interesting. The applied stress cannot only cause hydrogen embrittlement, but also promote oxidefilm breakdown. Thus, the applied stress and chlorides can act to-gether to enhance the pitting tendency of the alloy.

Examination of low magnification images (Fig. 5) and fracto-graphs obtained by SEM (Fig. 6) showed pitting to be a precursorevent for initiation of stress corrosion cracking. Hence, the deepestpit in each of the CLT specimen was measured for radial depth onfractured surface and plotted as shown in Fig. 10 to see if any cor-relation existed between pitting, and the applied stress. For thosespecimens survived in CLT, they were first fractured by tensileloading and the pits present on the fractured surface were exam-

Fig. 9. Hydrogen evolution measurements for Mg–Mn alloy in 0.01 M and 0.1 MNaCl solution saturated with Mg(OH)2.

ined. It is also assumed that the specimens failed on those sectionswhere pitting is predominant. Though pitting is a stochastic phe-nomenon, stress corrosion cracking is deterministic in nature ascrack will preferentially originate from the deepest pit; the latterbeing the biggest stress raiser. As the present data correspondsto the actual pits responsible for crack initiation, it brings out theeffect of pit size to SCC susceptibility of the alloy.

The plot shows that the time to fail decreased with appliedstress as expected. However, the pit depth does not follow any sin-gle relation with applied stress. The pit depth remained same from75–126 MPa. There after it increased up to an applied stress levelof 155 MPa and then declined on further increases the stress level.There after it fails with increase in stress level. Non-variation in pitdepth (deepest pit) over 70–126 MPa can be attributed to fact thatall these specimens were exposed for 1000 h in the environmentand pitting is more dependent on exposure time than appliedstress. However, what is interesting is that the pit depth increaseswith applied stress thereafter, although time to fail decreases withapplied stress; meaning the pit growth is increased by the appliedstress. It is possible that at these stress level the strain on the spec-imen could be high enough to damage the film so as to acceleratethe pit growth. In fact Winzer et al. [10] using direct current poten-tial difference (DCPD) showed that film break down increases withstrain even within the elastic region, supporting our propositionthat pit grows under the influence of stress. Thus, synergy betweenpitting and applied stress has been found in the present work. Thereduction in the pit depth at 170 MPa can be explained by the factthat at high stress levels even smaller pits are sufficient to initiatestress corrosion cracks and the failure is dominated more by stressthan pitting. The forgoing discussion has established the role ofchloride on pitting and hydrogen evolution on one hand and thepit growth on the other.

Based on the results obtained in this study and the analysismade above, it is possible to propose a phenomenological modelto explain a possible SCC mechanism and the role of chlorides ininfluencing SCC in Mg alloys. The same is provided in Fig. 11. Pit-ting is precursor event in SCC, weather or not the environmentcontains chlorides (Fig. 11a), as specimens failed even in distilledwater exhibited pitting. Pits are initiation sites for SCC (Fig. 11b),their nucleation and growth are influenced by chlorides and ap-plied stress. These two variables also increase hydrogen evolutionkinetics and thereby cause hydrogen embrittlement in Mg–Mn al-loy. During the initial stages of pitting/SCC, the growth of crack isslow and progresses discontinuously by with smaller step(Fig. 11b). During the later stage, high stress concentration leads

Stage I

Precursor event

Stage II

Initiation event

Stage III

Growth event

a b c

Fig. 11. Diagrammatic representation of SCC phenomena: (a) pitting or selectiveattack at the precipitate/ matrix interface, (b) crack initiation at pits, and (c) crackgrowth in discontinuous manner from beach mark.

V.S. Raja, B.S. Padekar / Corrosion Science 75 (2013) 176–183 183

to larger steps and faster growth rate and the crack follows alsoeasy path of precipitate alignment along the rolling direction(Fig. 11c). At higher stress level, it is possible that the alloy doesnot require high level of hydrogen (produced due to anodic dis-solution) for sustained crack growth and the growth becomesfaster in relation to anodic dissolution of the alloy. Besides this,the high stress levels can prevent the alloy from repassivationand maintain active dissolution. The fact that pitting or filmbreakdown is an important precursor event in hydrogen embrit-tlement mechanism of Mg alloys, efforts should be directed indeveloping pitting resistant Mg alloys. It is worth mentioningthe fact that EV31A alloy [2] did not undergo SCC even whenit was subjected to stress level beyond its rYS as it was foundto be resistant to pitting. It is in this context segregation ofalloying elements in cast structures, presence of hydrogen gener-ating impurities and absence of passivating elements can con-tribute one way or other to SCC susceptibility of the alloy viz-a-viz pitting and hydrogen evolution.

6. Conclusions

The main findings of this study are as follows.

(1) Mg–Mn alloy was susceptible to hydrogen embrittlement inMg(OH)2 saturated solutions of 0.01 M and 0.1 M NaCl andmuch less susceptible in distilled water.

(2) The rSCC determined in Mg(OH)2 saturated solutions of 0.1 MNaCl using CLT was 126 MPa which is 90% rYS (measured at0.2% offset strain) of Mg–Mn alloy.

(3) Pitting was the main precursor event in the SCC initiationprocess. However, its main role lies in enhancing hydrogenevolution kinetics rather than to act as stress raisers. Themain role of chloride was to damage the passive film (pit-ting) to create a bare surface and there by enhance hydrogenevolution kinetics.

(4) A phenomological model was proposed to explain the inter-relation among chlorides-pitting-hydrogen evolution andhydrogen embrittlement.

(5) It is believed that enhancing the pitting resistance canimprove the resistance of Mg-alloys to hydrogen-embrittlingenvironments.

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