Corrosion Science 71 (2013) 1–9

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Corrosion Science

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Stress corrosion cracking of a recent rare-earth containing magnesium alloy,EV31A, and a common Al-containing alloy, AZ91E

Bharat S. Padekar a,b,d,⇑, R.K. Singh Raman b,c, V.S. Raja d, Lyon Paul e

a IITB-Monash Research Academy, Indiab Department of Mechanical and Aerospace Engineering, Monash University (Melbourne), Vic. 3800, Australiac Department of Chemical Engineering, Monash University (Melbourne), Vic. 3800, Australiad Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology Bombay, Mumbai 400 076, Indiae Magnesium Elektron Ltd., Manchester, United Kingdom

a r t i c l e i n f o

Article history:Received 14 June 2012Accepted 1 January 2013Available online 23 January 2013

Keywords:A. MagnesiumB. SEMC. Stress corrosion

0010-938X/$ - see front matter � 2013 Elsevier Ltd. Ahttp://dx.doi.org/10.1016/j.corsci.2013.01.001

⇑ Corresponding author. Tel.: +91 22 25764607.E-mail address: bspadekar@iitb.ac.in (B.S. Padekar

a b s t r a c t

Stress corrosion cracking of the magnesium alloy Elektron 21 (ASTM–EV31A) and AZ91E was studiedusing constant load test in 0.1 M NaCl solution (saturated with Mg(OH)2), and slow strain rate test usingglycerol, distilled water and Mg(OH)2 saturated, 0.01 M and 0.1 M NaCl solutions. Slow strain rate testindicated that EV31A was less susceptible to stress corrosion cracking than AZ91E. Under less intenseloading of constant load, EV31A was found to be resistant to stress corrosion cracking. Fractography ofEV31A specimens showed little evidence of hydrogen embrittlement. The superior resistance of EV31Ais attributed to a more robust oxide/hydroxide layer.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Magnesium alloys are the lightest metallic engineering materi-als, and hence they are extremely attractive for light weight appli-cations, such as aerospace and automotive industries. Magnesiumalso offers several other advantages, such as castability, dampingcapacity, electromagnetic interference shielding, availability, rela-tively high thermal and electrical conductivities, and recyclability[1]. In spite of these attractive properties, their use has historicallybeen restricted largely due to their chemical reactivity, causingpoor resistance to corrosion [1–4] and corrosion-assisted cracking[5–8], besides their other limitations, viz., low elastic modulus andlimited strength and creep resistance at elevated temperatures[1,9]. Newer corrosion resistant alloys, such as EV31A and highpurity versions of AZ91(D & E) ,have addressed general corrosionconcerns, and rare earth elements (RE) containing alloys such asEV31A provide good elevated temperature performance up to200 �C.

In most of the traditional magnesium alloys, aluminum hasbeen one of the common alloying elements [1–9] for the purposeof both corrosion resistance and precipitation strengthening. How-ever, there has been a rapid increase recently in research into Al-free magnesium alloys, particularly for the specific requirementof such alloys in biological applications [10,11]. Most Al-free alloysgenerally contain one or more rare earth (RE) elements, again for

ll rights reserved.

).

the purposes of precipitation strengthening and corrosion resis-tance [12]. It may be imperative now to divide magnesium alloysinto two broad categories: Mg-RE containing alloys (Al-free), andMg–Al containing alloys (RE-free). Zirconium is not compatiblewith Mg–Al containing alloys. In Al-free magnesium alloy, how-ever, zirconium has two beneficial effects: (a) it reacts with thedeleterious impurities (viz., Fe, Ni and Co that can otherwise formhighly cathodic intermetallic compounds and accelerate corrosion[1,12]) and precipitates them at the alloy melting stage, therebyimproving corrosion resistance [1], and (b) it is an effective grainrefiner, causing strengthening [1,12]. Rare earth additions offer fur-ther benefits including improved strength at elevated tempera-tures and good creep performance.

It is well established in the literature that Al-containing magne-sium alloys (AZxx) are susceptible to stress corrosion cracking(SCC) in distilled water and chloride solutions [5–8]. However,SCC of the family of Al-free and RE-containing alloys has receivedvery limited research attention [6,12–14]. In fact, even the limitedliterature on the SCC of rare-earth containing magnesium alloyspresents contradictory views. In their comprehensive review ofthe SCC of magnesium alloys, Winzer et al. [5] have suggested thatREs such as Nd, Zr and Th have little influence on SCC susceptibil-ity, whereas Rokhlin [12] has reported that Nd or Cd additions toincrease the SCC resistance of an Mg–Zn–Zr alloy. Bobby Kannanet al. [6] reported each of the alloys used in their study (viz.,AZ80, ZE41, QE22 and EV31A (the latter three are RE containingalloys)) was susceptible to SCC during slow strain rate (SSRT)testing in a chloride solution. Though the authors have reported

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the SCC resistance of the Zr- and RE-containing alloy, EV31A, to besuperior to other alloys tested in their study, they have providedlittle concrete explanation for this comparative behavior. It isnoted that European Cooperation Space Standardization acknowl-edges the superior SCC resistance of EV31A T6 alloy [15].

EV31A is a relatively recent rare-earth containing and Zr-grainrefined magnesium alloy that has a considerably superior strengthto the traditional magnesium alloys [16]. Because of its Nd and Gdcontents, this alloy develops a robust and stable film of mixed oxidesof Zr, Nd and Gd (as opposed to the chemically and mechanically lessstable oxide/hydroxide films that develop on Al-rich alloys such asAZ91). The nature and stability of the oxide/hydroxide films that de-velop on the surface of metals and alloys in aqueous solutions areknown to profoundly influence the initiation and propagation ofstress corrosion cracks [17]. SCC susceptibility also depends on theintensity of loading that governs the mechanical stability of theoxide/hydroxide films [18]. SCC of EV31A has been investigated bySSRT [6]. The continuously increasing dynamic strain in SSRT invari-ably subjects the material (and the oxide/hydroxide films) to anexcessive loading that will disrupt any oxide/hydroxide film, andhence may not be effective in distinguishing the role of a morestable/robust oxide film that develops on EV31A alloy as opposedto the less stable oxide films on AZ91. This study develops acomparative mechanistic understanding of the SCC susceptibilitiesof an Al-free and RE-containing magnesium alloy (EV31A) and anAl-containing and RE-free alloy (AZ91E) in response to the role of arelatively less intense loading that will be experienced by thedistinctly different oxide/hydroxide films that develop on the twoalloys during constant load testing (CLT). This mechanistic under-standing will be compared with the SCC susceptibilities under thecondition of intense loading experienced during SSRT.

2. Experimental procedures

2.1. Test materials and microstructure

The chemical compositions of the two magnesium alloys usedin this study, EV31A and AZ91E, are presented in Table 1. The al-loys were received from Magnesium Elektron UK as sand castplates (250 � 150 � 22.5 mm) with T6 heat treatments. ForEV31A, T6 heat treatment corresponds to 8 h at 520 �C, hot waterquench and then 16 h at 200 �C; whereas T6 for AZ91 was 24 hat 413 �C, air cooled and then 16 h at 170 �C. The mechanical prop-erties of the two alloys were determined by tensile testing in air ata strain rate of 10�4 s�1, using a Zwick/Roell universal testing ma-chine (UTM). Common mechanical properties of the two alloyswere determined to be: (a) AZ91E, elongation (El): 4.5%, yieldstrength (YS): 170 MPa and ultimate tensile strength (UTS):250 MPa, and (b) EV31A, El: 6%, YS: 170 MPa, and UTS: 280 MPa.

Microstructures of the two alloys were examined by opticalmicroscopy of specimens prepared by standard metallographicprocedures and etched in a solution of ethanol (100 ml), water(20 ml), acetic acid (7.5 ml) and picric acid (7 g). Energy dispersiveX-ray (EDX) analysis was carried out for the characterization of thesecondary phase particles, which included X-ray mapping of pre-cipitates and surrounding areas.

Table 1Chemical compositiona of the test alloys (wt.%).

Alloy Al Zn Mn Zr Nd Gd

EV31A 0.28 0.6 2.9 1.5AZ91E 8.5 0.6 0.2

a Chemical analysis was done at Magnesium Elektron Ltd. UK.b RE – Rare earth (other).

2.2. Test environment and tests specimens

Cylindrical tensile specimens having gauge length 30 mm anddiameter 6 mm were used for CLT and SSRT. The gauge surfacewas abraded in the axial direction successively till 2500 grit sizeof silicon carbide abrasive papers. The specimens were cleanedusing acetone immediately before loading. 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 (as shown in Fig. 1 indicated by b) suchthat the gauge length remained immersed in the test environmentduring the entire test duration.

2.3. Slow strain rate testing

In the present study all the specimens were tested with thecrosshead travel speed of 0.0017 mm min�1, which is equivalentto a strain rate of 10�6 s�1. United Calibration Corporation (STM-20) tensile testing machine was used for the SSRT experiments.The corrosion cell is made to fit around the shoulder of the speci-men. The arrangement of test specimen, corrosion cell, and grip onboth the sides are similar to that shown for schematic of CLT rig inFig. 1. The cell was filled with test solution before starting the ma-chine. The load was applied by employing the constant extensionrate mode. A pre-load of �20 MPa was applied to take care of themachine slack. The load and the extension were monitored contin-uously by a load cell and an optical encoder. As there was no exten-sion meter attached to the specimen gauge length, SSRT testsproduced apparent stress–strain curves. The El was obtained bymeasuring actual value on failed specimens. The fracture surfaceswere cleaned in boiling chromic acid (180 g l�1 in water), washedin distilled water and rinsed with acetone, before examinationusing scanning electron microscope (SEM).

2.4. Constant load testing

Schematic of constant load test (CLT) rig is shown in Fig. 1.Three such rigs were used during this work. Load was applied onspecimen by the reactive force of helical compression spring.Springs used were having stiffness as 350, 300, and 250 N mm�1

and maximum allowable deflection as 50, 75 and 100 mm respec-tively. Spring with upper and lower guiding cup for pull rod wereloaded in alignment and calibrated on a UTM. Pull rod (indicatedby d in Fig.1) is provided with threads of 2.5 mm pitch. Load wasapplied by tightening the nut manually at the rate of half revolu-tion per minute. Distance between upper and lower guiding cupwas measured using a vernier calliper (least count: 0.02 mm).The applied load was maintained by adjusting the compressionof the spring (which included adjustment for any slack). In orderto avoid rotation of the pull rod during loading and or adjustingthe slack, a flat locator is provided with loading frame. A clock hav-ing external battery connection with banana clip was fixed to theupper and lower cups of loading spring. When specimen failed,the banana clip disconnects and clock stops, enabling to determinethe exact time of failure. The corrosive environment (0.1 M NaCl

REb Ag Fe Ni Cu

0.14 <0.001 0.003 <0.0001 <0.0001

Fig. 1. Schematic of constant load test rig showing: (a) test specimen, (b) corrosioncell, (c) grip member, (d) pull rod, (e) flat locator, (f) loading spring, (g) banana clip,and (h) clock.

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solution saturated with Mg(OH)2) was filled in the cell when thedesired load was achieved, and replaced weekly as well as any rea-sonable loss due to evaporation was replenished by topping upwith distilled water. A plot of applied stress vs. time to fail pro-vided the threshold stress for SCC (rSCC). Tests were terminatedafter 1008 h. The specimens that survived 1008 h were forciblyfractured. Fracture surfaces were cleaned and dried (as describedin Section 2.3) and examined using SEM.

3. Results

3.1. Microstructure

The microstructure of AZ91E alloy possesses a a-phase matrixand second phase precipitates, predominantly at the grain bound-aries (Fig. 2a). These precipitates have been reported to be Mg17-

Al12 [1,3]. High magnification SEM/BSE image, shown as an insetin Fig. 2a, reveals the presence of dendritic eutectic (with lamellarmorphology) appearing adjacent to the grain boundary precipi-tates, b (Mg17Al12).

Microstructure of EV31A appears to have a greater density offine second phase particles within the grain as well as a few coarseprecipitates at the grain boundaries (Fig. 2b). The high magnifica-tion SEM/BSE image shown as an insert in Fig. 2b reveals both grainboundary precipitates and fine precipitates. Images of X-raymapping (Fig. 3) suggest the grain boundary particles to be richin Nd, Gd, Zr and Zn. This is further associated with a notable Zrdepletion surrounding the precipitate. Notably; no publishedliterature is available on the nature of phases formed in EV31A

Fig. 2. Reflected light images with SEM/BSE images in the insets: (a) AZ91E showing bzirconium rich precipitate (shown by arrow).

magnesium alloy. Hence the effects of Gd, Nd and Zn in formingvarious phases in magnesium alloys are presented here. The workby Vostry et al. [19] indicates that Mg–5, 10 and 15 wt.% Gd binaryalloys form Mg5Gd intermetallic phase. Birbilis et al. [20] reportedthe formation of Mg3Nd in Mg–3.53 wt.%Nd alloys. Similarly Zn hasbeen reported to form intermetallic phases such as MgZn [21] withMg in Mg–1, 5, 7 wt.% Zn alloys. These studies indicate the ten-dency of Gd, Nd and Zn to form intermetallic phases. InMg–7%Gd–2.25%Nd–0.6%Zr alloy, having similar alloying elementsas in the EV31A, Apps et al. [22] have reported formation ofMg5(Nd0.5Gd0.5) phase. This indicates that Nd and Gd can alloytogether to form a complex magnesium intermetallic phase.However, contrary to our studies, they have not reported the pres-ence of Zr in the phase even though their alloy contained 0.6 wt.%Zr. While Apps et al. results do not support our observation; Coyet al. [23] have reported the grain boundary precipitate to containZr in ZE41 magnesium alloy. Study on Mg–3Nd–0.2Zn–0.4Zr (wt.%)alloy by Chang et al. [24] is noteworthy to mention. They havereported formation of Mg12Nd, which does not seem to containeither Zn or Zr, possibly because of the fact that these alloyingelements are low in content. This suggestion is supported by thefact that Drits et al. [25], reported formation of binary phases,Mg9Nd, Mg7Zn3, MgZn and ternary phases MgNd4Zn5, Mg6 Nd2Zn7,Mg2Nd2Zn9 in Mg–Zn–Nd alloys containing up to 8 wt.% Zn and3 wt.% Nd. Here, Nd and Zn seem to alloy together in forming inter-metallic phase because of the fact that their Zn content is higherthan the alloy studied by Chang et al. Thus, it looks logical thatthe grain boundary precipitate in the present case contains Gd,Nd, Zn and Zr.

3.2. Slow strain rate testing and fractography

Mechanical properties UTS and El for EV31A and AZ91E testedat a strain rate of 10�6 s�1 in glycerol, distilled water and thetwo chloride solutions (0.01 M NaCl and 0.1 M NaCl solutions,(each saturated with Mg(OH)2) are compared in Fig. 4. Mechanicalproperties reported here, UTS and El are average of triplicate testresults (deviated <3%) and are rounded to their respective nearestdigit. The extent of loss in mechanical properties of the two alloys,as a result of their exposure to the corrosive solutions, is repre-sented by the comparison of UTS and El in various test environ-ments, as shown in Table 2.

SCC susceptibility can be quantitatively represented in terms ofsusceptibility index (ISCC) [5]:

ISCC ¼ðElGly: � ElSCCÞ

ElGly:� 100 ð1Þ

where ElGly and Elscc are the elongations measured at the end of thetests in glycerol and the corrosive solutions respectively. ISCC can be

(Mg17Al12) phase, and (b) EV31A showing heavy grain boundary precipitate and

Fig. 3. SEM/BSE image of grain boundary precipitate (top left) and corresponding X-ray mapping for different elements of alloy EV31A.

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described also on the basis of UTS data. The El and UTS data shownin Table 2 were used for generating the ISCC data shown in Fig. 5 forthe two alloys in different corrosive media.

The fractographs of EV31A and AZ91E alloys tested in the inertenvironment of glycerol are shown in Fig. 6. The fracture featuresof EV31A were found to be transgranular with some localized duc-tility revealing as dimples (Fig. 6a). Fracture surface of AZ91E re-vealed cracking by a mix of a transgranular and a lathlike mode(Fig. 6b). These features could be attributed to the presence oflamellar eutectic areas adjacent to the grain boundary precipitatesin the microstructure of the alloy (as seen in Fig. 2a). Accordingly,an area of brittle fracture (presumably, of grain boundary precipi-tates) was often found adjacent to such lamellar feature.

In order to examine for the evidence of SCC of alloy EV31A andAZ91E after SSRT, transverse images of fracture specimens wereobtained. Only a selected few are presented to bringout the sailentpoints. Fractured surfaces of SSRT specimens tested in 0.1 M testsolution were observed using a stereo microscope. Typical imagesare shown in Figs. 7a and 8a. SCC region is not distantly visible inthe case of EV31A (Fig. 7a) but it was possible to distinguish anddemarcate the region in case of AZ91E (Fig. 8a). The area within

the white line demarcation was found to undergo mechanical over-load failure.

The longitudinal section views of the specimens are shown inFigs. 7b and 8b. Localized corrosion attack is clearly seen on thespecimen edge in SEM image of the longitudinal section of EV31(Fig. 7b). Such attack occurs along the grainboundary near the largegrain boundary precipitates. Alloy AZ91E also suffered localized at-tack on specimen surface, which becomes precursor for stress cor-rosion cracks (Fig. 8b). In this case pits were found to occur evenwithin the grains. Crack length is shorter for EV31A (Fig. 7b) thanfor AZ91E (Fig. 8b), for the identical SSRT condition. As a result,the SCC region was not as clearly distinguishable in the case ofEV31A in transverse micrograph (Fig. 7a). Initiation of cracks dueto pits and their propogation as stress corrosion cracks have alsobeen reported by Stampella et al. [26] for pure magnesium duringSSRT in Na2SO4. Winzer et al. [27] suggested a crack initiationmechanism involving slip induced film rupture and localized disso-lution in weld, heat affected zone and base metal of AZ31 alloy. Inalloy EV31A tested in distilled water and 0.1 M NaCl, crack initia-tion site is the bulky grain boundary (GB) precipitates and propo-gation is generally intergranular (IG), as shown in (Fig. 7b). IGSCC

Fig. 4. Comparison of EV31A and AZ91E based on: (a) El, and (b) UTS, SSRT datatested in Gly: glycerol, Dw: distilled water, 0.01 M and 0.1 M: NaCl solutionssaturated with Mg(OH)2.

Table 2Summary for SSRT results of EV31A and AZ91E.

Alloy El (%) UTS (MPa)

Gly DW 0.01 N 0.1 N Gly Dw 0.01 N 0.1 N

EV31A 7 6 5 4 270 255 240 230AZ91E 6.5 5 4 3 235 200 185 180

Note: Gly: glycerol, Dw: distilled water, 0.01 M and 0.1 M: NaCl solutions saturatedwith Mg(OH)2.

Fig. 5. Comparison of SCC susceptibility index based on: (a) El, and (b) UTS data (inTable 2), Dw: distilled water, 0.01 M and 0.1 M: NaCl solutions saturated withMg(OH)2.

B.S. Padekar et al. / Corrosion Science 71 (2013) 1–9 5

caused by the second phase particle along grain boundries was alsoreported by Ben-Hamu et al. [28] for ZSM6X0 alloy (SSRT in3.5 wt.% NaCl saturated with Mg(OH)2).

Fractographic features of alloy EV31A after SSRT at a strain rateof 10�6 s�1 in distilled water and Mg(OH)2 saturated 0.01 M and0.1 M NaCl solutions were found to be similar. Localized corrosionat the circumference facilitated crack initiation as shown by darkarrows in the representative fractograph of specimen tested in0.1 M NaCl solution saturated with Mg(OH)2 (Fig. 7c). Intergranularcracking was facilitated by the bulky precipitates at grain bound-ary (Fig. 7c). But, the predominant feature was a mix of intergarnu-lar and transgranular cracking (Fig. 7d) whereas the central area ofthe fracture surface consisted predominantly of the feature of over-load ductile failure (similar to that shown in Fig. 6a).

A comparison of the fractographic features of alloy EV31A afterSSRT in distilled water and Mg(OH)2 saturated 0.01 M and 0.1 MNaCl solutions suggests an increase in localized corrosion andsecondary cracking at the circumference with the increasing

corrosivity of the aqueous solutions. This trend is consistent withthe trend of increasing SCC susceptibility of EV31A alloy with thechange in corrosive medium from distilled water to 0.01 M NaCl,and then to 0.1 M NaCl (as shown in Fig. 5). With the change in testenvironment from 0.01 M NaCl to 0.1 M NaCl, the extent of inter-granular stress corrosion cracking (IGSCC) was found to increase.With very little IGSCC in distilled water, this also suggests a trendof increasing IGSCC with change in corrosive medium from dis-tilled water to 0.01 M NaCl, and then to 0.1 M NaCl. In each case,discontinuous crack propagation was evident by the parallel facet-ing. The central part in each case had the feature of overload duc-tile failures similar to Fig. 6a.

Fractographic features of alloy AZ91E tested in the three testmediums were similar in nature, i.e., crack initiated from preferen-tially corroded /pit area and propagated in transgranular (TG) man-ner. The representative fractographs of a specimen tested in 0.1 MNaCl solution saturated with Mg(OH)2 are shown in Fig. 8c. Cracksinitiated from preferentially corroded regions or pits on the speci-men surface (shown by arrows in Fig. 8c), and propogated as trans-granular stress corrosion cracks in a discontinuous manner (shownin Fig. 8c Detail A). However, the extent of SCC region (shown inFig. 8c Detail A) was found to increase with the change in corrosivemedium from distilled water to 0.01 M NaCl, and then to 0.1 MNaCl. Beach marks or parallel facets shown in Fig. 8c (Detail A)have been attributed in the literature to the hydrogen assistedcracking for the AZ series alloys [29–34]. This trend of increase inhydrogen assisted cracking is consistent with the trend of increas-ing SCC susceptibility of AZ91E alloy with the change in corrosive

Fig. 6. SEM fractographs of SSRT specimen tested in glycerol: (a) EV31A, and (b) AZ91E.

Fig. 7. Stereo micrograph, and SEM micrograph and fractographs of alloy EV31A strained to fracture in 0.1 M test solution: (a) transverse stereo micrograph, (b) longitudinalsection showing localize corrosion/pit and crack, (c) features of localized corrosion on the specimen circumference (shown by dark arrows), intergranular cracking alongbulky precipitates at grain boundary (shown by bright arrows), and (d) mix of intergarnular (shown by dark arrows) and transgranular cracking.

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medium, as shown in Fig. 5. The central area of the fracture surfaceconsisted predominantly of the feature of overload ductile failuressimilar to Fig. 6b. Fairman and Bray [35] mentioned that TGSCC isdiscontinuous and involves a brittle fracture step which is proba-bly crystallographic in Mg–6.5%Al–1.1%Zn single phase alloy testedin chromate chloride solution and IGSCC is also discontinuousexhibiting alternating fast and slow stages in two phase alloy. Intheir review, Meletis and Hochman [36] mention that fracture sur-face produced by hydrogen embrittlement of Mg-7.5Al alloy con-sisted of flat facets that follow f2; �2;0;3g planes.

3.3. Constant load testing and fractography

The alloys, EV31A and AZ91E, were subjected to CLT, for inves-tigating under less stringent loading conditions of CLT (than inSSRT) the role of mechanical stability of the considerably differentoxide scales that develop on the two alloys, and thereby examiningdifference, if any, in SCC susceptibility under CLT condition. CLTtests were carried out using the most aggressive of the test envi-ronments employed in this study (i.e., 0.1 M NaCl saturated withMg(OH)2) at different initial loads for a maximum duration of

1008 h. Stress vs. time to fail plots for EV31A and AZ91E are shownin Fig. 9. AZ91E did not fail when applied stress was 6102 MPa(Fig. 9a), indicating the rSCC of the alloy in 0.1 M NaCl to be102 MPa (which is 60% of yield stress).

The equation of exponential decay curve that best fit the exper-imental data of AZ91

r ¼ 148 exp�tf

289

� �þ 102 ð2Þ

where r is applied stress (in MPa) and tf is time to fail (in h).The Eq. (2) that best fits experimental data of CLT of AZ91E

yields applied stress equal to 250 MPa at time of failure equating0 h which is UTS of AZ91E and at time of failure equating infinity(i.e. 102.05 MPa) which is rSCC. Fairman and Bray [35] obtainedthreshold stress as 100 MPa in chromate-chloride solution and70 MPa in dichromate–chloride solution under constant load testcondition for Mg–6.5%Al–1.1%Zn single phase alloy using speci-mens of 0.5 mm diameter. Dependence of susceptibility on Al con-tent was also studied by Fairman and Bray [37] using constant loadtest (specimen diameter of 0.76 mm) in solution of 4% NaCl + 4%Na2CrO4 and obtained threshold stress of 85, 83, 77, and 65 MPa

Fig. 8. Stereo micrograph, and SEM micrograph and fractographs of alloy AZ91E strained to fracture in 0.1 M test solution: (a) transverse stereo micrograph (white linedemarcating area of SCC), (b) longitudinal section showing localized corrosion/pit and crack, and (c) crack initiation from the edge (shown by arrows), transgranular anddiscontinuous crack propagation (Detail A).

Fig. 9. Stress vs. time to fail plots: (a) AZ91E, and (b) EV31A, in 0.1 M NaCl solutionsaturated with Mg(OH)2 (arrows indicate no failure).

B.S. Padekar et al. / Corrosion Science 71 (2013) 1–9 7

for the alloys Mg–1Al, Mg–3Al, AZ61, and AZ80 respectively. How-ever the study [35,37] was carried out for maximum time durationof 28 h mainly in chromate containing solution which is inhibitor/passivating and there is also effect of specimen size. So, compari-son of these data to those in the present study will not bereasonable.

EV31A did not fracture during 1008 h CLT even when the ap-plied stresses were considerably greater than the yield strength.As described in Section 2.4, the specimens that survived 1008 hwere forced to fracture to determine the retained mechanical prop-erties, and fracture surfaces thus produced were examined. The re-tained mechanical properties of a few critical specimens arepresented in Table 3.

Longitudinal section view of the specimen of AZ91E alloy testedat 160 MPa that failed after 245 h shows localized corrosion/pitand cracks (Fig. 10a). Fractographs of AZ91E subjected to CLT, at160 MPa in 0.1 M NaCl (Fig. 10b) which invariably showed featuresconfirming SCC/hydrogen embrittlement, viz., localized corrosionat the circumference, transgranular cracking, secondary crackingadjacent to localized corrosion, and discontinuous crack propaga-tion in Fig. 10 Detail B (parallel facets). Fracture surfaces of thespecimens tested at or below rSCC (102 MPa) and then forced tofracture were not having features of hydrogen embrittlement (nodiscontinuous crack propagation). The major fractographic fea-tures were the fine dimples similar to those shown in Fig. 6b. How-ever, this specimen had suffered localized corrosion (similar to asshown in Fig. 10b) at the circumference, during their long exposureto the corrosive environment, which accounts for the loss in its re-tained mechanical properties (shown in Table 3). It should benoted that for same environment, a SSRT specimen spent muchshorter duration in the test solution than a CLT specimens. Forexample, a test on AZ91E in 0.1 M NaCl solution saturated withMg(OH)2, ended within 18–20 h (attaining a maximum stress180 MPa) whereas CLT specimen loaded at 160 MPa took 245 hto fail. Therefore, it is common that the failed SSRT specimenshowed less visible evidence of localize corrosion on circumfer-ence. But, due to a longer exposure, the CLT specimen loaded at

160 MPa suffered localized corrosion at its circumference (asshown in Figs. 10a and b) which assisted SCC initiation andpropogation.

Table 3Tensile test (strain rate, 10�4 s�1) results of EV31A and AZ91E that survived CLT atdifferent stresses in 0.1 M NaCl for 1008 h, showing retained mechanical properties.

Specimens (CLT stresses) El (%) UTS (MPa)

EV31A (153–230 MPa) 4 ± 0.2 235 ± 5AZ91E (68–102 MPa) 2.5 ± 0.2 165 ± 5

Fig. 10. SEM micrograph and fractographs of alloy AZ91E fractured during CLT, loaded to 160 MPa: (a) longitudinal section showing localized corrosion/pit and crack, and (b)localized corrosion at circumference (shown by bright arrows), transgranular and secondary cracking (shown by dark arrows in Detail A), and discontinuous crackpropagation (Detail B).

8 B.S. Padekar et al. / Corrosion Science 71 (2013) 1–9

Fractographic features of EV31A subjected to CLT at variousstresses and then forced to fracture were similar to those for thealloy tested in inert environment (Fig. 6a), i.e., ductile dimplesand only occasional transgranular features, confirming that the al-loy did not suffer SCC in 0.1 M NaCl solution under static stressesas high as 230 MPa. Also, these specimens suffered very little lossof cross sections due to localized corrosion at the circumferenceduring CLT in 0.1 M NaCl solution. This explains the limited lossin the retained mechanical properties of this alloy (Table 3).

Fig. 11. XRD analyses of the oxide/hydroxide films on plain abraded specimensafter immersing in distilled water for 175 h: (a) EV31A and (b) AZ91E.

4. Discussion

Though a recent study has suggested a combined effect ofhydrogen induced cracking and anodic dissolution [38], hydro-gen-assisted embrittlement is believed to be the most commonmechanism for SCC of magnesium alloys in aqueous solutions[5,6]. In this mechanism, the applied stress is suggested to beresponsible for disruption in the oxide/hydroxide films that devel-op on magnesium alloys, thereby allowing entry of hydrogen intothe alloy matrix to cause embrittlement [5,39]. Therefore, the nat-ure and stability of the oxide/hydroxide films that develop on mag-nesium alloys is the most important parameter for their SCC.

EV31A and AZ91E develop distinctly different oxide/hydroxidescales in 0.1 M NaCl solution saturated with Mg(OH)2. EV31Adevelops a dark (black colored) passive layer that appeared robustand protective. This is attributed in the literature [6] to the forma-tion of robust layer of the mixed oxides of Zr, Nd and Gd that formon this alloy in the aqueous solution. Al-rich alloys (such as AZ91E)develop an Al-rich hydroxide film of Mg(OH)2 that is entirely

different in appearance. Ambat et al. [40] have reported oxide filmcontaining Mg(OH)2, MgO, and Al2O3 for AZ91D ingot immersed in

B.S. Padekar et al. / Corrosion Science 71 (2013) 1–9 9

3.5% NaCl solution. X-ray diffraction (XRD) of the corrosion filmsshown in Fig. 11 confirms the formation of chemically differentoxides on the two alloys. XRD was obtained on flat specimensabraded successively until 5000 grit size and then immersed in dis-tilled water for 175 h. Ben-Hamu et al. [39] have reported the rare-earth rich alloys to develop an oxide film of rare-earth elements(Mg–Me–O, where Me is Zr and Nd) that are more stable thanMg(OH)2 passive layer observed in the case of Al-rich alloys (e.g.,AZ91). The disruptions seen in the oxide/hydroxide films that de-velop on AZ91E suggest a less robust film on this alloy whichcauses a greater hydrogen evolution.

The remarkable resistance of EV31A to SCC during CLT (as seenin Fig. 9b) can be attributed to the robust oxide/hydroxide layerthat develops on this alloy. It is presumed that the relatively lessintense straining experienced during CLT allowed the sustenanceof the film. But, the oxide/hydroxide films that develop on AZ91Ealloy was able to sustain the less intense loading during CLT onlywhen the stress was below 102 MPa. But, at higher stresses, thefilm got disrupted and the alloy suffered SCC (as seen in Fig. 9a).

Both EV31A and AZ91E alloys suffered SCC during SSRT. SCC ofAZ91E is attributed to the ease of disruption of the oxide/hydroxidelayer that develops on this alloy [37]. It is suggested in the case ofEV31A that the more intense and continuous straining experiencedduring SSRT could disrupt even the robust oxide/hydroxide layerthat develops on this alloy, causing SCC. Another observation wasthe greater propensity for intergranular SCC of EV31A during SSRT.In fact, as described in Section 3.2, this propensity increased withthe increasing NaCl content. This behavior can again be explainedon the basis of the reported nature of oxide/hydroxide that devel-ops at microscopic level on this alloy. Zirconium is reported[23,39,41] to enrich in areas away from the grain boundaries, leav-ing the areas adjacent to grain boundaries lean in Zr. Electrochem-ical corrosion/dissolution resistance of the Zr-rich central area isreported [41,42] to be superior to that of the Zr-lean areas. Coyet al. [23] reported that the precipitate containing Zr to have po-tential differences in excess of +170 mV with respect to the matrixfor the alloys ZE41, WE43, and WE54, in 3.5% NaCl solution. Theless corrosion resistant oxide/hydroxide layer that developed lo-cally in the Zr-lean areas adjacent to the grain boundaries will besusceptible to disruption due to aggressive anions such as chloride.The increase in the susceptibility of localized disruption in areasadjacent to grain boundaries, with increasing chloride content, willalso increase the opportunity for hydrogen entry into the EV31A al-loy at/around grain boundaries, which explains the increase inIGSCC susceptibility with increasing chloride content.

5. Conclusion

During constant load testing, AZ91E was found to be susceptibleto SCC but EV31A showed resistance to SCC. The resistance ofEV31A to SCC is attributed to the formation of a robust oxide layeron the alloy and sustenance of this film during CLT.

The slow strain rate testing, which is meant to test until speci-men fails, indicates EV31A to possess higher tensile strength and

ductility than AZ91E in the tested environments. Based on theseresults it is suggested that EV31A alloy possesses a superior resis-tance to SCC than AZ91E.

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