- Email: [email protected]
Stress corrosion cracking behavior of magnesium alloys EV31A and AZ91E
Materials Science & Engineering A 583 (2013) 169–176
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Stress corrosion cracking behavior of magnesium alloys EV31A and AZ91E Bharat S. Padekar a,b,c, V.S. Raja b,n, R.K. Singh Raman c,d, P. Lyon e a
IITB-Monash Research Academy, India Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology Bombay, Mumbai 400076, India c Department of Mechanical and Aerospace Engineering, Monash University, Vic. 3800, Australia d Department of Chemical Engineering, Monash University, Vic. 3800, Australia e Magnesium Elektron Ltd, Manchester, United Kingdom b
art ic l e i nf o
a b s t r a c t
Article history: Received 11 March 2013 Received in revised form 26 June 2013 Accepted 29 June 2013 Available online 8 July 2013
The stress corrosion cracking of the sand cast magnesium alloys, Elektron 21 (ASTM-EV31A) and AZ91E, was studied using compact tension specimens in distilled water and in Mg(OH)2 saturated 0.1 M NaCl solution under constant load. EV31A showed a higher KISCC than AZ91E, but exhibited a higher stage II crack growth rate than the latter. Fractography showed transgranular cracking for AZ91E and EV31A in distilled water, whereas mixed intergranular and transgranular cracking was observed for EV31A in the chloride environment. The crack propagation involved hydrogen embrittlement. & 2013 Elsevier B.V. All rights reserved.
Keywords: Magnesium alloys Casting Fracture Hydrogen embrittlement
1. Introduction Magnesium alloys are attractive for light structural applications, such as automobile and aerospace, because their density is only 2/3rd of aluminum. A range of alloys, from magnesium– aluminum alloys (i.e., AZ91) to the high strength and high temperature magnesium–zirconium alloys (i.e., WE43), are commercially available. Rare-earth addition to magnesium alloys has a significant effect on their creep resistance, which is primarily attributed to the formation of rare-earth containing phases along the grain boundaries of such alloys [1–4]. Controlled addition of the individual rare-earth elements in magnesium alloys also influences their castability as well as age-hardening response [5]. Magnesium alloys can be divided into two broad categories [1,5]: (1) Zr-containing and (2) Zr-free alloys. Zr offers several advantages over magnesium alloys. It reacts with the impurity elements (that are extremely detrimental for corrosion resistance viz, Fe, Ni and Co), to form their intermetallics that are removed during melting, thus producing intrinsically “high purity” alloys with good corrosion resistance (note poor corrosion resistance is a major concern for Mg-alloys) [3]. Zr is also an effective grain refiner [1,3,6]. However, as Zr reacts with Al to form detrimental intermetallics, magnesium alloys would rarely contain both Al and Zr. Further, the Zr
n
Corresponding author. Tel.: +91 22 2576 7892; fax: +91 22 2572 3480. E-mail address: [email protected] (V.S. Raja).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.06.085
containing alloys often contain alloying elements like rare earths (RE), Ag, Th, etc. Zinc is added in most of the magnesium alloys in sufficient quantities to achieve precipitation hardening. In combination with Zn and Zr, addition of neodymium imparts hardness and strength upon artificial aging [5], as a result of precipitation both within the grains and at grain boundaries. Neodymium also improves castability of the magnesium alloys. Al containing magnesium alloys are susceptible to stress corrosion cracking (SCC) [2,7]. The literature on SCC of rare-earth containing magnesium alloys is very limited. Even the limited literature [2,8–10] reports contradictory views on SCC of such alloys. In a critical review of the SCC behavior of magnesium alloys Winzer et al. [10] have inferred that neodymium and zirconium have very little or no influence on SCC. In contrast, Rokhlin [2] reported that neodymium addition to Mg–Zn–Zr alloy increases the SCC resistance. Al-free magnesium–zinc alloys that are alloyed with either zirconium (i.e., ZK60) or rare earths (i.e., ZE10) have intermediate SCC resistance [8]. Kannan et al. [9] reported that rare-earth containing magnesium alloys such as ZE41, QE22 and EV31A possess superior SCC resistance than a RE free alloy AZ80. European Cooperation Space Standardization (ECSS) has also identified the RE-containing alloy EV31A T6 to have a considerably high resistance to SCC in comparison with the RE-free alloys [11]. It is noteworthy that the limited published literature on SCC behavior of RE containing Mg alloys is mainly concerned with studying smooth tensile specimens [9,12], which cannot address the potential role of surface defects. The fracture mechanics based
170
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
SCC tests that are designed specifically to account for such pre existing surface defects are in general limited [13–19], and such studies on rare-earth containing alloys [18] are even less. The fracture mechanics approach to SCC concerns the susceptibility of an existing flaw/crack to grow under the influence of corrosive environment and stress intensity at crack-tip. One of the common modes of quantifying this susceptibility is the crack growth rate (da/dt) vs. stress intensity (KI) curve [20,21]. Critical parameters in designing damage tolerant components susceptible to SCC are threshold stress intensity for SCC (KISCC), slope for stage I regime and stage II crack growth rate [20,22]. The slope (m) of curve for stage I crack growth is particularly important for life prediction but such detailed studies have not been carried out for Mg alloys. EV31A is a relatively new alloy developed for elevated temperature creep resistance and is suitable for investment casting and sand casting of relatively thicker sections (up to 75 mm) [23]. However, there is little reported on stress corrosion crack growth behavior of such alloy, and hence, the present study. This study also compares the stress corrosion crack growth behavior of EV31A with that of a relatively known alloy, AZ91E (sand cast and T6 heat treated). The employed testing procedures complied with the plane strain condition, which is a material property.
2. Investigation procedures 2.1. Materials Two magnesium alloys used in this study (EV31A and AZ91E) were received from Magnesium Electron UK as sand cast plates (250 150 22.5 mm) with respective T6 heat treatments. For EV31A, T6 heat treatment corresponded to 8 h hold at 520 1C, hot water quenching and then 16 h aging at 200 1C, whereas T6 for AZ91 it was 24 h hold at 413 1C, air cooling and then 16 h aging at 170 1C. The chemical composition in wt% of EV31A was 0.28 Zn–0.6Zr–2.9Nd–1.5Gd–0.14other RE (each o0.01%)– o0.001Ag– o0.0001Cu–0.003Fe– o0.0001Ni, and that of AZ91E was 8.5Al– 0.6Zn and 0.2Mn. Reflected light microscopy was used to observe the microstructure of the specimens developed upon standard metallography followed by etching with a solution of ethanol 100 ml, water 20 ml, acetic acid 7.5 ml and picric acid 7 g. Energy dispersive X-ray (EDX) analysis was carried out for broad chemical characterization of the secondary-phase particles. To observe preferential corrosion if any, due to microstructural features, abraded specimens were exposed to an aqueous solution of 2 wt% oxalic acid and examined under a scanning electron microscope (SEM) for localized corrosion.
2.2. Specimen design ISO 7539-6 [24], ASTM E399 [25] and ASTM E 1681[26] standards were followed to design compact tension (C(T)) specimens. Based on the available material size, specimen width (W) of 38.1 mm and specimen thickness (B) of 22 mm were designed. Straight-through notch was produced by EDM wire cut with 1 mm end slot of 0.16 mm radius. The notch was abraded manually using 1000 grit silicon carbide paper to remove EDM marks. The specimen sides were abraded till 5000 grit silicon carbide paper, with a final finish using a diamond paste of 0.25 mm grit. Marker lines to limit the fatigue pre-crack length were made on both the sides of specimens. Smooth tensile specimens of 6 mm diameter and 30 mm gage length were prepared as per ASTM E 8 [27] and were abraded in the axial direction to a final finish of 5000 grit size.
Fig. 1. Schematic of CLT test rig with CMOD measuring arrangement using dial indicator: (a) CT specimen, (b) clevis, (c) CMOD extension, (d) dial indicator, (e) load cell, (f) loading frame, (g) corrosive drip line, (h) loading spring, (i) banana clip, and (j) clock.
2.3. Pre-cracking The fatigue pre-cracking of C(T) specimens was carried out (till W/a 0.5) at the laboratory temperature with 4 Hz frequency and load ratio (R) as 0.1, using a load control sine wave on a servo hydraulic controlled Zwick/Roell UTM. Side grooves using an end mill cutter of 601 nose angle and end radius of 0.15 mm were milled to a depth of 5% of the specimen thickness ( 1.1 mm) per side [24,25] on both the sides of the specimens. Holes with M3 tapping were produced to fit the extension of PTFE for crack mouth opening displacement (CMOD) measurement using a dial test indicator (Fig. 1d). Deburring and cleaning were done before loading the specimen on a constant load test rig. 2.4. Constant load test setup A rig with horizontal loading axis was used so that drip of the corrosive solution could easily access the fatigue pre-crack and then drain down the specimen. A schematic of the test rig is shown in Fig. 1. Two identical rigs were used. Each rig is equipped with a load cell of 2000 kg capacity with a least count of 1 kg to measure the applied load. The load was applied through a compression spring having a spring index of about 400 N mm 1 such that the load did not decrease by greater than 1% of the applied load in the event of the change in load line deflection before resetting time (normally 12 h). The rotary motion of the pull rod was restricted by a locator to avoid twisting of specimen during loading and while adjusting the load whenever required. A clock having external battery connection with a banana clip was fixed to the end cups of the loading spring. When the specimen failed, the banana clip disconnected and the clock stopped, enabling determination of the exact time to fail. The pulling grips were provided with the alumina coated loading pins and wrapped with PTFE so as to facilitate the adjustment of the specimen to proper loading line and allow rotation of the specimen during testing [25]; also PTFE washers (0.5 mm thick) on both sides were provided to avoid friction and galvanic couple of specimen with the pulling grips. All the tests were done at the laboratory temperature (22 73 1C). The load was applied slowly so that the rate of the applied stress intensity (KI) will not exceed 100 MPa m1/2/min [26]. The drip of the corrosive solution was turned on when the applied load reached 100 N (10 kg), which provided some opening for the solution to reach the crack tip. The flow of solution was adjusted to 1.5 ml/min. Mg (OH)2 saturated solution of 0.1 M NaCl with initial pH as 10.5 and distilled water were employed as the test solutions. Each test started with a fresh solution, which was replaced with another fresh solution after every week of use. A schematic of the CMOD arrangement is shown in Fig. 1c–d. A mechanical dial indicator with a least count of 1 mm was used at a movement arm of 41 mm (above load line) to detect the crack growth. Deflection of the dial
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
increased as the crack grew under constant load test conditions. During the loading of C(T) specimen, the dial indicator of CMOD arrangement was set to 0 mm when there was no load. Load and CMOD were measured at every hour for the first 6 h, and subsequently, at least twice a day throughout the remaining test period. Tests were terminated after 1008 h if the specimen did not fail under the applied constant load.
2.5. Measurement of crack size Crack length and crack growth were measured on the fractured surface using a traveling microscope. Detailed procedures of measurement of crack size by the nine line method are mentioned in our recent publication [18]. The average crack length was obtained from the nine measurements using the formula [18,28] ( ) 8 1 a1 þ a9 a¼ þ ∑ ai ð1Þ 8 2 i¼2
2.6. Determination of KI In constant load test (CLT), the load was applied assuming an approximate crack size from the experience of specimens that were tested for determination of KIC. Parameters like B, thickness at the root of side groove (Bn), and W were measured before loading. After failure of the specimen or after the survived specimen was forced to fracture, the initial crack size (ai) and final crack size (af) were measured from the fractured surface. From the applied load and measured dimensions, KI was calculated using standard formula [24–26,28].
171
3. Results 3.1. Microstructure and corrosion resistance Microstructure of EV31A appears to have a greater density of fine second-phase particles within the grains as well as a few coarse precipitates at the grain boundaries (Fig. 2a). The average grain size of the alloy determined by the linear intercept method [29] was found to be 50 75 μm. Back scattered electron images of the etched microstructure reveal the fine second phase particles to have rod/plate like morphologies (Fig. 2a inset). EDX analysis revealed these precipitates to have Zn (1.54 at%), Zr (2.14at%), Nd (0.24 at%), Gd (0.43 at%) and Mg (94.64 at%). Here Zn and Zr contents are nearly in proportion to Zn2Zr3. Intergranular phases of Zn2Zr3 and Mg4Zn7 have been reported by Li et al. [30] for Mg– 5Zn–0.6Zr alloy. A more detailed analysis is made in our recent publication [12]. Microstructure of AZ91E alloy possesses α-phase matrix and second-phase precipitates, predominantly at the grain boundaries (Fig. 2b). The average grain size of the alloy was determined to be 320 715 μm. EDX analysis of the grain boundary precipitates, identified by SEM/BSE, was found to be rich in Mg and Al. The main precipitates in AZ91E alloy have been reported to be β-phase of composition Mg17Al12 [7,10]. SEM/BSE image (Fig. 2b) also reveals areas of eutectic (with lamellar morphology) occurring
2.7. Validity check Every C(T) specimen used in CLT was checked for validity of plane strain conditions to satisfy such parameters as B, a, and W–a, according to the criteria mentioned in the standards [24–26]. Also none of the nine dimensions measured for crack size differ by more than 0.05 B ( 1 mm) from the average value of crack, a [28].
2.8. Determination of crack growth rate Average crack growth rate (da/dt) of SCC was obtained from the difference of final and initial crack length (af ai) that was measured on the fractured surface and the time to fail or time of testing (for survived specimen) as [18,24] da SCC length ðmÞ ¼ dt Time of testing ðsÞ
ð2Þ
2.9. Fractography Stereo micrographs were obtained on fractured surfaces to observe broad features after subjecting them to cleaning in distilled water, drying, and finally ultrasonic cleaning in acetone and drying. However, for a detailed fractographic analysis, using SEM (Hitachi S-3400N), required regions were cut and chemical cleaning was carried out to remove surface oxides. Sections were cleaned in boiling solution of 20 wt% chromic acid in distilled water by immersing for 20 s and then rinsed in distilled water, dried, and ultrasonically cleaned with acetone before fractographic examination.
Fig. 2. SEM micrographs for (a) EV31A, showing coarse grain boundary precipitate, and precipitates in the interior of grain which are rod like and plate like (shown in the insets), and (b) AZ91E, showing heavy grain boundary precipitate (β), surrounded by eutectic, and α phase.
172
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
adjacent to the grain boundary precipitates. The etching in aqueous solution of 2 wt% oxalic acid produced galvanic attack adjacent to grain boundary precipitates which is consistent with the reported highly cathodic ( 300 mV) nature of β precipitate (Mg17Al12) as compared to α matrix [31]. Corrosion rate of magnesium alloys can be easily assessed by the rate of hydrogen evolution [32]. The hydrogen evolved during the immersion test was collected, following the procedure of Venugopal et al. [33] and Song et al. [34], and the results are presented in Fig. 3. The corrosion rates for the alloys EV31A and AZ91E were calculated to be 0.53 mg cm 2 day 1 (1.12 mm yr 1) and 2.16 mg cm 2 day 1 (4.54 mm yr 1) respectively in 0.1 M NaCl solution. AZ91E has exhibited higher corrosion rate than EV31A despite higher Al content of the former. The higher corrosion rate of AZ91E than EV31A can be attributed to the presence of cathodic phases namely Mg17Al12 in AZ91E. 3.2. Mechanical behavior
Fig. 4. Plot of CMOD vs. time to fail (arrows indicate specimens survived for 1008 h) for alloy AZ91E tested in 0.1 M NaCl solution saturated with Mg(OH)2.
Various mechanical properties obtained in the present study are summarized in Table 1. Tensile properties reported here are the average of triplicate test results obtained at 10 4 s 1 strain rate tested in air. The tensile properties such as ultimate tensile strength (UTS), yield strength (YS) (obtained at 0.2% offset strain), Young's modulus (E) and fracture toughness (KIC) values were in agreement with the material data sheet [35]. Fracture toughness tests provided a clear idea about the maximum load that may be employed in CLT. For the present geometry of C(T) specimens (i.e., for W and a), it was established that the specimens satisfied the plane strain criterion as long as the initial KI for the CLT specimens was below 14.8 MPa m1/2. A representative set of CMOD vs. time to fail plots of AZ91E tested in 0.1 M NaCl is presented in Fig. 4 for a maximum duration of 150 h (arrow indicates the specimens that survived in 1008 h). These plots suggest that there is no or very less incubation period to start SCC crack from fatigue pre-crack after attaining the applied
Fig. 5. Initial stress intensity vs. time to fail plots of EV31A and AZ91E in distilled water and Mg(OH)2 saturated 0.1 M NaCl solution, generated using CLT data (arrows indicate no failure).
load in CLT. Increasing CMOD indicates crack growth, and slope approaching infinity indicates fracture, i.e., when growing crack approached the critical size (acrit). Plots of initial KI vs. time to fail are presented in Fig. 5. When a specimen did not fail in 1008 h, the test was terminated. Such specimens are identified by arrows in Fig. 6. The initial KI approaches the threshold because of an exponential decay in the KI vs. time to fail relationship [13,15,16,18]. A generalized equation (Eq. (3)) that best fits the data of KI vs. time to fail in distilled water (Dw) and 0.1 M NaCl solution saturated with Mg (OH)2 for EV31A as well as AZ91E is given as follows: Fig. 3. Hydrogen evolution measurements for EV31A and AZ91E alloys in 0.1 M NaCl solution saturated with Mg(OH)2 with immersion time.
Table 1 Summary of the mechanical properties of the two alloys. Alloy
Elongation (El %)
YS (MPa)
UTS (MPa)
E (GPa)
KIC (MPa m1/2)
EV31A AZ91E
6 4.5
170 170
280 250
44 44
15 13.2
K I ¼ a exp
t f þc b
ð3Þ
where KI is the stress intensity (MPa m1/2), and tf is the time to fail (h). KISCC, a, b, and c are summarized in Table 2. KISCC (KI corresponding to time to fail tending to infinity) as calculated from Eq. (3) was determined to be 10 MPa m1/2 (in Dw) and 8 MPa m1/2 (in NaCl) for EV31A. At tf ¼0, KI is 15 MPa m1/2 which is the fracture toughness (KIC) of EV31A. Similarly, using Eq. (3), KISCC was determined for AZ91E to be 8 MPa m1/2 (in Dw)
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
173
and 6 MPa m1/2 (in NaCl), and KI (at tf ¼0) was found to be 13 MPa m1/2 which is the KIC of AZ91E. Plots of log crack growth rate vs. initial KI for the two alloys tested in distilled water and Mg(OH)2 saturated 0.1 M NaCl are presented in Fig. 6. Stage I data points of the alloys in the two environments fit to an empirical linear relation given by da ¼ mK I þ n log ð4Þ dt
Fig. 6. Plot of log crack growth rate vs. initial stress intensity factor for alloy EV31A and AZ91E in distilled water (Dw) and Mg(OH)2 saturated 0.1 M NaCl (0.1 M) solution using CLT.
Table 2 Summary of fracture mechanics data of two alloys obtained by fitting to the data of Fig. 5 using Eq. (3) and of Fig. 6 using Eq. (4). KISCC corresponds to crack growth rate ¼10 10 m s 1 obtained using Eq. (4). Alloy
EV31A AZ91E
Env.
Dw NaCl Dw NaCl
Eq. (3)
Eq. (4)
a
b
c
m
n
5 7 5 7
30 7 220 37
10 8 8 6
5.5 7 0.8 2
65 66 16.4 22
KISCC (MPa m1/2)
Stage II da/dt (m s 1)
10 8 8 6
7.05 10 9 7.25 10 8 2.69 10 9 3.98 10 8
Env, environment; Dw, distilled water; NaCl, 0.1 M NaCl saturated with Mg(OH)2.
where crack growth rate (da/dt) is in m s 1, KI in MPa m1/2, m is the slope and n is the intercept on the axis of log crack growth rate. These data are summarized in Table 2 for each alloy– environment combination. Table 2 also presents stage II crack growth rates. The KISCC values in the table correspond to a crack growth rate of 10 10 m s 1 [18,28]. These KISCC data (determined using Eq. (4)) are consistent with those determined from KI vs. time to fail plot (using Eq. (3)). KISCC data of EV31A alloy presented here are among the first such data which show higher KISCC (i.e., greater resistance to SCC crack propagation) than alloy AZ91E in both the test environments. 3.3. Fractography Representative stereo micrographs tested in each alloy–environment combination are shown in Fig. 7. Various regions, viz., machined notch region, fatigue pre-cracked region (that had typical parabolic crack front), the SCC region (demarcated by white lines), and mechanically failed region, are delineated on the fractographs. Fracture surface of C(T) specimen of EV31A tested in distilled water is intentionally presented in Fig. 7(a) to show the maximum extent of variation in the crack depth over the nine locations that was measured on fracture surface of specimen that met the validity requirements for variation in crack dimensions [28]. Side grooves are clearly visible on all the micrographs (Fig. 7). There is no visible evidence of deformation along the thickness (Fig. 7) which indicates that the specimens satisfied the plane strain condition.
Fig. 7. Stereo micrographs of the fractured surfaces of the two magnesium alloys in distilled water (Dw) and Mg(OH)2 saturated solution of 0.1 M NaCl (0.1 M): (a) EV31A—Dw, (b) EV31A—0.1 M, (c) AZ91E—Dw, and (d) AZ91E—0.1 M, showing different regions: (1) machined notch, (2) fatigue-pre-crack, (3) SCC region (demarcated by white line), and (4) mechanical failed region.
174
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
Fractographic features of the different broad regions of the fracture surfaces shown in Fig. 7 were observed at higher magnifications, as seen in Figs. 8–10. Fatigue pre-crack regions of both
the alloys have transgranular striations, as the general feature (Fig. 8a and c); however, there are areas of dissolution along grain boundary precipitates, particularly when tested in the 0.1 M NaCl
Fig. 8. Fractographs of EV31A (a) fatigue region, (b) mechanical failed region, and of AZ91E (c) fatigue region and (d) mechanical failed region. Note – the dissolution near heavy grain boundary precipitates is shown by arrows (in a and c).
Fig. 9. Fractographs of EV31A for SCC region (a) showing TGSCC in distilled water and parallel facets (Detail A), (b) predominantly TGSCC with isolated occurrence of IGSCC (shown by arrows) and parallel facets (Detail B) in Mg(OH)2 saturated 0.1 M NaCl.
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
175
Fig. 10. Fractographs of AZ91E for SCC region: (a) TGSCC region in distilled water reveals secondary cracks (shown by arrows) and parallel facets (shown in Detail A), and (b) TGSCC in 0.1 M NaCl saturated with Mg(OH)2 (showing secondary cracks marked by arrows) and parallel facets (Detail B). Fractograph shown in B reveals localized attack which could affect the otherwise sharp transgranular crack front.
solution (presumably that occurred during the long CLT in 1008 h). The mechanically failed regions have fine dimples (Fig. 8b and d), sign of ductile failure, similar to the fracture feature of the two alloys that failed under slow strain rate tensile (SSRT) tests in glycerol [36]. EV31A tested in distilled water showed mostly transgranular feature (Fig. 9a). The transgranular SCC (TGSCC) in distilled water propagated in a discontinuous manner having parallel facets (Fig. 9a—Detail A). The SCC region of EV31A tested in Mg(OH)2 saturated 0.1 M NaCl showed mainly TGSCC as shown in Fig. 9b and b—Detail B. There are a few evidences of dissolution near grain boundary precipitates (shown by arrows in Fig. 9b). Similar fractographic features were reported by Kannan et al. [9] for the failed SSRT specimens of EV31A that shows TGSCC when tested in distilled water, and mixed intergranular SCC (IGSCC) and TGSCC in 0.5 wt% NaCl. They attributed the IGSCC feature to the electrochemical dissolution along the continuous grain boundary precipitates in the rare-earth containing alloys (i.e. ZE41, QE22 and EV31A). On the other hand, the TGSCC feature is discontinuous and involves alternating mechanical and electrochemical processes. Similar to EV31A, the SCC region of AZ91E tested in distilled water showed mostly transgranular feature with parallel facets (Fig. 10a and Detail A). SCC region of specimen tested in Mg(OH)2 saturated 0.1 M NaCl also showed mostly the TGSCC feature (Fig. 10b and Detail B) with evidence of secondary cracks shown by arrows in Fig. 10a and b. TGSCC features having parallel facets and evidence of dissolution are observed at higher magnification (shown in Detail B). The fractographic features observed in the SCC region of C(T) specimens of alloy EV31A and AZ91 tested in distilled water and 0.1 M NaCl saturated with Mg(OH)2 have similar features to the author's earlier work on SSRT specimens tested in the same solution [12,36]. The transgranular features having flat parallel facets (Figs. 9 and 10) indicate discontinuous crack growth and the possible role of hydrogen in cracking.
Winzer et al. [37] reported TGSCC feature for AZ91 alloy tested in distilled water under constant extension rate test using cylindrical tensile specimens. Chakrapani and Pugh [38] suggested SCC of an Mg–Al alloy in NaCl–K2CrO4 solution to be a form of hydrogen embrittlement and the fracture to be cleavage like. For pure Mg in NaCl–K2CrO4 solution, Meletis et al. [39] have reported the cracking to be TGSCC, having a series of parallel crystallographic facets separated by steps and jogs, similar to the features observed in the present study.
4. Discussion EV31A and AZ91E alloys have shown the presence of stage I and stage II crack growth characteristics and the absence of stage III crack growth (Fig. 6), much the same as reported in the literature for magnesium alloys [10,17–19], and aluminum alloys [21]. The log (da/dt) vs. KI plots bring out the difference in crack growth tendency of the two alloys. Notably, EV31A exhibits higher crack growth rate than AZ91E in stage II, even though its KISCC is greater. Several factors seem to contribute to the SCC resistance of the alloys. The addition of chloride has caused a significant decrease in KISCC of both the alloys. Chloride ions increase the tendency for IGSCC as against TGSCC found in distilled water for EV31A alloy (Fig. 9). Between the two alloys, AZ91E liberates more hydrogen than EV31A (Fig. 3), which could be one of the reasons why the former exhibits lower KISCC than the latter. Not much is reported on the slope for stage I crack growth and crack velocity for stage II, though such studies have been carried out on stainless steels [20]. Jones and Simonen [20] suggested that “m”, the slope for stage I crack growth, increases for alloys exhibiting higher KISCC and when the environment is more aggressive. Examination of Table 2 shows that between EV31A and AZ91E, the latter exhibits lower m (slope corresponding to stage I crack growth) and stage II crack growth rate than the
176
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
former. This is despite the fact that EV31A exhibits higher KISCC and produces less hydrogen than AZ91E in chloride containing environment (Fig. 3). An alloy should be more susceptible to hydrogen embrittlement and exhibit higher crack growth rate if it liberates more H2. The reasons for such a behavior are not quite clear, though the following explanations are plausible. Comparison of the microstructures of EV31A and AZ91E indicates that the latter possesses a complex microstructure (β phase and eutectic phases), which can act as crack arresters. Further, it seems these microstructures also cause localized attack, leading to possible crack blunting. This suggestion is made based on the fact that selective and localized attack has been found on the fracture feature of AZ91E alloy (Fig. 10), which is absent in EV31A alloy (Fig. 9). The effect of Al on the localized attack is worth discussing. It is known that in the presence of Al3+, the solution pH within cracks and crevices can decrease down to 4 because of hydrolysis as given by the following equation [21]: 3þ
Al
þ 3H2 O-AlðOHÞ3 þ 3Hþ
ð5Þ
Such reduction in pH can accelerate corrosion and possible crack blunting (Fig. 10b–Detail B). 5. Conclusion Stress corrosion crack growth behavior of a sand cast rare earth containing magnesium alloy, EV31A, in distilled water and 0.1 M NaCl solution saturated with Mg(OH)2 was studied using compact tension specimen under constant load condition for the first time. For comparison, a traditional Mg-alloy, AZ91E, was also investigated. The main conclusions of this study are as follows: (1) The KISCC in distilled water and Mg(OH)2 saturated solution of 0.1 M NaCl were determined for EV31A to be 10 and 8 MPa m1/2 and for AZ91E to be 8 and 6 MPa m1/2 respectively, indicating the superior SCC resistance of EV31A than AZ91E. However, AZ91E shows lower stage II crack growth rate than EV31A in both the test environments, which is attributed to a possible crack blunting and complex microstructures. (2) Slope, m, increases and KISCC decreases with the aggressiveness of the environment. (3) Despite higher H2 evolution tendency, AZ91E exhibits lower stage II crack growth rate than that of EV31A. Both the alloys exhibited increase in stage II crack growth rate due to chloride addition.
References [1] H.E. Friedrich, B.L. Mordike, Magnesium Technology, Metallurgy, Design Data, Applications, Springer Verlag, Berlin, Heidelberg, 2006.
[2] L.L. Rokhlin, Magnesium Alloys Containing Rare Earth Metals, Taylor and FrancisLondon, 2003. [3] J.F. Nie, X. Gao, S.M. Zhu, Scr. Mater. 53 (2005) 1049–1053. [4] C. Sanchez, C. Nussbaum, P. Azavant, H. Octor, Mater. Sci. Eng. A 221 (1996) 48–57. [5] P. Lyon, T. Wilks, I. Syed, in: N.R. Neelameggham, H.I. Kaplan, B.R. Powell (Eds.), Magnesium Technology 2005, Warrendale, PA, 2005, pp. 303–308. [6] G. Ben-Hamu, D. Eliezer, K.S. Shin, S. Cohen, J. Alloy Compds. 431 (2007) 269–276. [7] R.G. Song, C. Blawert, W. Dietzel, A. Atrens, Mater. Sci. Eng. A 399 (2005) 308–317. [8] H. Baker, M.M. Avedesian (Eds.), ASM Specialty Handbook: Magnesium and Magnesium Alloys, ASM International, USA, 1999. [9] M. Bobby Kannan, W. Dietzel, C. Blawert, A. Atrens, P. Lyon, Mater. Sci. Eng. A 480 (2008) 529–539. [10] N. Winzer, A. Atrens, G. Song, E. Ghali, W. Dietzel, K.U. Kainer, N. Hort, C. Blawert, Adv. Eng. Mater. 7 (2005) 659–693. [11] European Cooperation for Space Standardization, Material Selection for StressCorrosion Cracking, ECSS-Q-ST-70-36C, 6 March 2009. [12] B.S. Padekar, R.K.S. Raman, V.S. Raja, P. Lyon, Corros. Sci. 71 (2013) 1–9. [13] R.I. Stephens, C.D. Schrader, D.L. Goodenberger, K.B. Lease, V.V. Ogarevic, S.N. Perov, SAE, No. 930752, USA, 1993, pp. 843–857. [14] N. Winzer, A. Atrens, W. Dietzel, G. Song, K.U. Kainer, Adv. Eng. Mater. 10 (5) (2008) 453–458. [15] L. Choudhary, R.K.Singh Raman, Acta Biomater. 8 (2012) 916–923. [16] A. Moccari, C.R. Shastry, J. Mater Technol. (Zeitschrift für Werkstofftechnik) 10 (1979) 119–123. [17] K. Ebtehaj, D. Hardie, R.N. Parkins, Corros. Sci. 28 (1993) 811–829. [18] B.S. Padekar, V.S. Raja, R.K.Singh Raman, Eng. Fract. Mech. 102 (2013) 180–193. [19] M.O. Speidel, M.J. Blackburn, T.R. Beck, J.A. Feeney, Corrosion Fatigue: Chemistry, Mechanics and Microstructure, NACE-2, 1972, pp. 324–345. [20] R.H. Jones, E.P. Simonen, Mater. Sci. Eng. A160 (1993) 127–136. [21] M.O. Speidel, Metall. Trans. 6A (1975) 631–651. [22] W. Dietzel, P.B. Srinivasan, in: V.S. Raja, T. Shoji (Eds.), Stress Corrosion Cracking Theory and Practice, Woodhead Publishing, Oxford, UK, 2011, pp. 133–166. [23] P. Lyon, I. Syed, S. Heaney, Adv. Eng. Mater. 9 (2007) 793–798. [24] International Standard, Corrosion of Metals and Alloys—Stress Corrosion Testing, ISO 7539-6, 2006. [25] Standard Test Method for Linear-Elastic Plane-Strain Fracture Toughness KIC of Metallic Materials, ASTM E 399, 2009. [26] Standard Test Method for Determining Threshold Stress Intensity Factor for Environment-Assisted Cracking of Metallic Materials, ASTM E 1681, 2003. [27] Standard Test Methods for Tension Testing of Metallic Materials, ASTM E 8, 2011. [28] W. Dietzel, Guidelines for Fracture Mechanics SCC Testing, Characterization of Susceptibility of Metallic Materials to Environmentally Assisted Cracking, GKSS-Forschungszentrum Geesthacht GmbH, 1995. [29] Standard Test Method for Determining Average Grain Size, ASTM E 112, 2010. [30] Q. Li, Q. Wang, Y. Wang, X. Zeng, W. Ding, J. Alloy Compds. 427 (2007) 115–123. [31] R.K.Singh Raman, Metall. Mater. Trans. A 35 (2004) 2525–2531. [32] M. Liu, P.J. Uggowitzer, A.V. Nagasekhar, P. Schmutz, M. Easton, Corros. Sci. 51 (2009) 602–619. [33] Venugopal, R.D.Angal, V.S. Raja, Corrosion 52 (1996) 138–142. [34] G. Song, A. Atrens, D.H. StJohn, in: J. Hryn (Ed.), Magnesium Technology, TMS, New Orleans, PA, 2001, pp. 255–262. [35] Data sheet 455 and 456, Magnesium Elektron Ltd. UK, 〈http://www.magnesiu m-elektron.com/data/downloads/DS455.PDF〉. [36] B.S Padekar, V.S. Raja, R.K. Singh Raman, P. Lyon, NACE Corrosion Conference and Expo, 2012, pp. 1556–1559. [37] N. Winzer, A. Atrens, W. Dietzel, G. Song, K.U. Kainer, Mater. Sci. Eng. A 472 (2008) 97–106. [38] D.G. Chakrapani, E.N. Pugh, Metall. Trans. 7A (1976) 173–178. [39] E.I. Meletis, R.F. Hochman, Corrosion 40 (1) (1984) 39–45.
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Stress corrosion cracking behavior of magnesium alloys EV31A and AZ91E Bharat S. Padekar a,b,c, V.S. Raja b,n, R.K. Singh Raman c,d, P. Lyon e a
IITB-Monash Research Academy, India Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology Bombay, Mumbai 400076, India c Department of Mechanical and Aerospace Engineering, Monash University, Vic. 3800, Australia d Department of Chemical Engineering, Monash University, Vic. 3800, Australia e Magnesium Elektron Ltd, Manchester, United Kingdom b
art ic l e i nf o
a b s t r a c t
Article history: Received 11 March 2013 Received in revised form 26 June 2013 Accepted 29 June 2013 Available online 8 July 2013
The stress corrosion cracking of the sand cast magnesium alloys, Elektron 21 (ASTM-EV31A) and AZ91E, was studied using compact tension specimens in distilled water and in Mg(OH)2 saturated 0.1 M NaCl solution under constant load. EV31A showed a higher KISCC than AZ91E, but exhibited a higher stage II crack growth rate than the latter. Fractography showed transgranular cracking for AZ91E and EV31A in distilled water, whereas mixed intergranular and transgranular cracking was observed for EV31A in the chloride environment. The crack propagation involved hydrogen embrittlement. & 2013 Elsevier B.V. All rights reserved.
Keywords: Magnesium alloys Casting Fracture Hydrogen embrittlement
1. Introduction Magnesium alloys are attractive for light structural applications, such as automobile and aerospace, because their density is only 2/3rd of aluminum. A range of alloys, from magnesium– aluminum alloys (i.e., AZ91) to the high strength and high temperature magnesium–zirconium alloys (i.e., WE43), are commercially available. Rare-earth addition to magnesium alloys has a significant effect on their creep resistance, which is primarily attributed to the formation of rare-earth containing phases along the grain boundaries of such alloys [1–4]. Controlled addition of the individual rare-earth elements in magnesium alloys also influences their castability as well as age-hardening response [5]. Magnesium alloys can be divided into two broad categories [1,5]: (1) Zr-containing and (2) Zr-free alloys. Zr offers several advantages over magnesium alloys. It reacts with the impurity elements (that are extremely detrimental for corrosion resistance viz, Fe, Ni and Co), to form their intermetallics that are removed during melting, thus producing intrinsically “high purity” alloys with good corrosion resistance (note poor corrosion resistance is a major concern for Mg-alloys) [3]. Zr is also an effective grain refiner [1,3,6]. However, as Zr reacts with Al to form detrimental intermetallics, magnesium alloys would rarely contain both Al and Zr. Further, the Zr
n
Corresponding author. Tel.: +91 22 2576 7892; fax: +91 22 2572 3480. E-mail address: [email protected] (V.S. Raja).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.06.085
containing alloys often contain alloying elements like rare earths (RE), Ag, Th, etc. Zinc is added in most of the magnesium alloys in sufficient quantities to achieve precipitation hardening. In combination with Zn and Zr, addition of neodymium imparts hardness and strength upon artificial aging [5], as a result of precipitation both within the grains and at grain boundaries. Neodymium also improves castability of the magnesium alloys. Al containing magnesium alloys are susceptible to stress corrosion cracking (SCC) [2,7]. The literature on SCC of rare-earth containing magnesium alloys is very limited. Even the limited literature [2,8–10] reports contradictory views on SCC of such alloys. In a critical review of the SCC behavior of magnesium alloys Winzer et al. [10] have inferred that neodymium and zirconium have very little or no influence on SCC. In contrast, Rokhlin [2] reported that neodymium addition to Mg–Zn–Zr alloy increases the SCC resistance. Al-free magnesium–zinc alloys that are alloyed with either zirconium (i.e., ZK60) or rare earths (i.e., ZE10) have intermediate SCC resistance [8]. Kannan et al. [9] reported that rare-earth containing magnesium alloys such as ZE41, QE22 and EV31A possess superior SCC resistance than a RE free alloy AZ80. European Cooperation Space Standardization (ECSS) has also identified the RE-containing alloy EV31A T6 to have a considerably high resistance to SCC in comparison with the RE-free alloys [11]. It is noteworthy that the limited published literature on SCC behavior of RE containing Mg alloys is mainly concerned with studying smooth tensile specimens [9,12], which cannot address the potential role of surface defects. The fracture mechanics based
170
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
SCC tests that are designed specifically to account for such pre existing surface defects are in general limited [13–19], and such studies on rare-earth containing alloys [18] are even less. The fracture mechanics approach to SCC concerns the susceptibility of an existing flaw/crack to grow under the influence of corrosive environment and stress intensity at crack-tip. One of the common modes of quantifying this susceptibility is the crack growth rate (da/dt) vs. stress intensity (KI) curve [20,21]. Critical parameters in designing damage tolerant components susceptible to SCC are threshold stress intensity for SCC (KISCC), slope for stage I regime and stage II crack growth rate [20,22]. The slope (m) of curve for stage I crack growth is particularly important for life prediction but such detailed studies have not been carried out for Mg alloys. EV31A is a relatively new alloy developed for elevated temperature creep resistance and is suitable for investment casting and sand casting of relatively thicker sections (up to 75 mm) [23]. However, there is little reported on stress corrosion crack growth behavior of such alloy, and hence, the present study. This study also compares the stress corrosion crack growth behavior of EV31A with that of a relatively known alloy, AZ91E (sand cast and T6 heat treated). The employed testing procedures complied with the plane strain condition, which is a material property.
2. Investigation procedures 2.1. Materials Two magnesium alloys used in this study (EV31A and AZ91E) were received from Magnesium Electron UK as sand cast plates (250 150 22.5 mm) with respective T6 heat treatments. For EV31A, T6 heat treatment corresponded to 8 h hold at 520 1C, hot water quenching and then 16 h aging at 200 1C, whereas T6 for AZ91 it was 24 h hold at 413 1C, air cooling and then 16 h aging at 170 1C. The chemical composition in wt% of EV31A was 0.28 Zn–0.6Zr–2.9Nd–1.5Gd–0.14other RE (each o0.01%)– o0.001Ag– o0.0001Cu–0.003Fe– o0.0001Ni, and that of AZ91E was 8.5Al– 0.6Zn and 0.2Mn. Reflected light microscopy was used to observe the microstructure of the specimens developed upon standard metallography followed by etching with a solution of ethanol 100 ml, water 20 ml, acetic acid 7.5 ml and picric acid 7 g. Energy dispersive X-ray (EDX) analysis was carried out for broad chemical characterization of the secondary-phase particles. To observe preferential corrosion if any, due to microstructural features, abraded specimens were exposed to an aqueous solution of 2 wt% oxalic acid and examined under a scanning electron microscope (SEM) for localized corrosion.
2.2. Specimen design ISO 7539-6 [24], ASTM E399 [25] and ASTM E 1681[26] standards were followed to design compact tension (C(T)) specimens. Based on the available material size, specimen width (W) of 38.1 mm and specimen thickness (B) of 22 mm were designed. Straight-through notch was produced by EDM wire cut with 1 mm end slot of 0.16 mm radius. The notch was abraded manually using 1000 grit silicon carbide paper to remove EDM marks. The specimen sides were abraded till 5000 grit silicon carbide paper, with a final finish using a diamond paste of 0.25 mm grit. Marker lines to limit the fatigue pre-crack length were made on both the sides of specimens. Smooth tensile specimens of 6 mm diameter and 30 mm gage length were prepared as per ASTM E 8 [27] and were abraded in the axial direction to a final finish of 5000 grit size.
Fig. 1. Schematic of CLT test rig with CMOD measuring arrangement using dial indicator: (a) CT specimen, (b) clevis, (c) CMOD extension, (d) dial indicator, (e) load cell, (f) loading frame, (g) corrosive drip line, (h) loading spring, (i) banana clip, and (j) clock.
2.3. Pre-cracking The fatigue pre-cracking of C(T) specimens was carried out (till W/a 0.5) at the laboratory temperature with 4 Hz frequency and load ratio (R) as 0.1, using a load control sine wave on a servo hydraulic controlled Zwick/Roell UTM. Side grooves using an end mill cutter of 601 nose angle and end radius of 0.15 mm were milled to a depth of 5% of the specimen thickness ( 1.1 mm) per side [24,25] on both the sides of the specimens. Holes with M3 tapping were produced to fit the extension of PTFE for crack mouth opening displacement (CMOD) measurement using a dial test indicator (Fig. 1d). Deburring and cleaning were done before loading the specimen on a constant load test rig. 2.4. Constant load test setup A rig with horizontal loading axis was used so that drip of the corrosive solution could easily access the fatigue pre-crack and then drain down the specimen. A schematic of the test rig is shown in Fig. 1. Two identical rigs were used. Each rig is equipped with a load cell of 2000 kg capacity with a least count of 1 kg to measure the applied load. The load was applied through a compression spring having a spring index of about 400 N mm 1 such that the load did not decrease by greater than 1% of the applied load in the event of the change in load line deflection before resetting time (normally 12 h). The rotary motion of the pull rod was restricted by a locator to avoid twisting of specimen during loading and while adjusting the load whenever required. A clock having external battery connection with a banana clip was fixed to the end cups of the loading spring. When the specimen failed, the banana clip disconnected and the clock stopped, enabling determination of the exact time to fail. The pulling grips were provided with the alumina coated loading pins and wrapped with PTFE so as to facilitate the adjustment of the specimen to proper loading line and allow rotation of the specimen during testing [25]; also PTFE washers (0.5 mm thick) on both sides were provided to avoid friction and galvanic couple of specimen with the pulling grips. All the tests were done at the laboratory temperature (22 73 1C). The load was applied slowly so that the rate of the applied stress intensity (KI) will not exceed 100 MPa m1/2/min [26]. The drip of the corrosive solution was turned on when the applied load reached 100 N (10 kg), which provided some opening for the solution to reach the crack tip. The flow of solution was adjusted to 1.5 ml/min. Mg (OH)2 saturated solution of 0.1 M NaCl with initial pH as 10.5 and distilled water were employed as the test solutions. Each test started with a fresh solution, which was replaced with another fresh solution after every week of use. A schematic of the CMOD arrangement is shown in Fig. 1c–d. A mechanical dial indicator with a least count of 1 mm was used at a movement arm of 41 mm (above load line) to detect the crack growth. Deflection of the dial
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
increased as the crack grew under constant load test conditions. During the loading of C(T) specimen, the dial indicator of CMOD arrangement was set to 0 mm when there was no load. Load and CMOD were measured at every hour for the first 6 h, and subsequently, at least twice a day throughout the remaining test period. Tests were terminated after 1008 h if the specimen did not fail under the applied constant load.
2.5. Measurement of crack size Crack length and crack growth were measured on the fractured surface using a traveling microscope. Detailed procedures of measurement of crack size by the nine line method are mentioned in our recent publication [18]. The average crack length was obtained from the nine measurements using the formula [18,28] ( ) 8 1 a1 þ a9 a¼ þ ∑ ai ð1Þ 8 2 i¼2
2.6. Determination of KI In constant load test (CLT), the load was applied assuming an approximate crack size from the experience of specimens that were tested for determination of KIC. Parameters like B, thickness at the root of side groove (Bn), and W were measured before loading. After failure of the specimen or after the survived specimen was forced to fracture, the initial crack size (ai) and final crack size (af) were measured from the fractured surface. From the applied load and measured dimensions, KI was calculated using standard formula [24–26,28].
171
3. Results 3.1. Microstructure and corrosion resistance Microstructure of EV31A appears to have a greater density of fine second-phase particles within the grains as well as a few coarse precipitates at the grain boundaries (Fig. 2a). The average grain size of the alloy determined by the linear intercept method [29] was found to be 50 75 μm. Back scattered electron images of the etched microstructure reveal the fine second phase particles to have rod/plate like morphologies (Fig. 2a inset). EDX analysis revealed these precipitates to have Zn (1.54 at%), Zr (2.14at%), Nd (0.24 at%), Gd (0.43 at%) and Mg (94.64 at%). Here Zn and Zr contents are nearly in proportion to Zn2Zr3. Intergranular phases of Zn2Zr3 and Mg4Zn7 have been reported by Li et al. [30] for Mg– 5Zn–0.6Zr alloy. A more detailed analysis is made in our recent publication [12]. Microstructure of AZ91E alloy possesses α-phase matrix and second-phase precipitates, predominantly at the grain boundaries (Fig. 2b). The average grain size of the alloy was determined to be 320 715 μm. EDX analysis of the grain boundary precipitates, identified by SEM/BSE, was found to be rich in Mg and Al. The main precipitates in AZ91E alloy have been reported to be β-phase of composition Mg17Al12 [7,10]. SEM/BSE image (Fig. 2b) also reveals areas of eutectic (with lamellar morphology) occurring
2.7. Validity check Every C(T) specimen used in CLT was checked for validity of plane strain conditions to satisfy such parameters as B, a, and W–a, according to the criteria mentioned in the standards [24–26]. Also none of the nine dimensions measured for crack size differ by more than 0.05 B ( 1 mm) from the average value of crack, a [28].
2.8. Determination of crack growth rate Average crack growth rate (da/dt) of SCC was obtained from the difference of final and initial crack length (af ai) that was measured on the fractured surface and the time to fail or time of testing (for survived specimen) as [18,24] da SCC length ðmÞ ¼ dt Time of testing ðsÞ
ð2Þ
2.9. Fractography Stereo micrographs were obtained on fractured surfaces to observe broad features after subjecting them to cleaning in distilled water, drying, and finally ultrasonic cleaning in acetone and drying. However, for a detailed fractographic analysis, using SEM (Hitachi S-3400N), required regions were cut and chemical cleaning was carried out to remove surface oxides. Sections were cleaned in boiling solution of 20 wt% chromic acid in distilled water by immersing for 20 s and then rinsed in distilled water, dried, and ultrasonically cleaned with acetone before fractographic examination.
Fig. 2. SEM micrographs for (a) EV31A, showing coarse grain boundary precipitate, and precipitates in the interior of grain which are rod like and plate like (shown in the insets), and (b) AZ91E, showing heavy grain boundary precipitate (β), surrounded by eutectic, and α phase.
172
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
adjacent to the grain boundary precipitates. The etching in aqueous solution of 2 wt% oxalic acid produced galvanic attack adjacent to grain boundary precipitates which is consistent with the reported highly cathodic ( 300 mV) nature of β precipitate (Mg17Al12) as compared to α matrix [31]. Corrosion rate of magnesium alloys can be easily assessed by the rate of hydrogen evolution [32]. The hydrogen evolved during the immersion test was collected, following the procedure of Venugopal et al. [33] and Song et al. [34], and the results are presented in Fig. 3. The corrosion rates for the alloys EV31A and AZ91E were calculated to be 0.53 mg cm 2 day 1 (1.12 mm yr 1) and 2.16 mg cm 2 day 1 (4.54 mm yr 1) respectively in 0.1 M NaCl solution. AZ91E has exhibited higher corrosion rate than EV31A despite higher Al content of the former. The higher corrosion rate of AZ91E than EV31A can be attributed to the presence of cathodic phases namely Mg17Al12 in AZ91E. 3.2. Mechanical behavior
Fig. 4. Plot of CMOD vs. time to fail (arrows indicate specimens survived for 1008 h) for alloy AZ91E tested in 0.1 M NaCl solution saturated with Mg(OH)2.
Various mechanical properties obtained in the present study are summarized in Table 1. Tensile properties reported here are the average of triplicate test results obtained at 10 4 s 1 strain rate tested in air. The tensile properties such as ultimate tensile strength (UTS), yield strength (YS) (obtained at 0.2% offset strain), Young's modulus (E) and fracture toughness (KIC) values were in agreement with the material data sheet [35]. Fracture toughness tests provided a clear idea about the maximum load that may be employed in CLT. For the present geometry of C(T) specimens (i.e., for W and a), it was established that the specimens satisfied the plane strain criterion as long as the initial KI for the CLT specimens was below 14.8 MPa m1/2. A representative set of CMOD vs. time to fail plots of AZ91E tested in 0.1 M NaCl is presented in Fig. 4 for a maximum duration of 150 h (arrow indicates the specimens that survived in 1008 h). These plots suggest that there is no or very less incubation period to start SCC crack from fatigue pre-crack after attaining the applied
Fig. 5. Initial stress intensity vs. time to fail plots of EV31A and AZ91E in distilled water and Mg(OH)2 saturated 0.1 M NaCl solution, generated using CLT data (arrows indicate no failure).
load in CLT. Increasing CMOD indicates crack growth, and slope approaching infinity indicates fracture, i.e., when growing crack approached the critical size (acrit). Plots of initial KI vs. time to fail are presented in Fig. 5. When a specimen did not fail in 1008 h, the test was terminated. Such specimens are identified by arrows in Fig. 6. The initial KI approaches the threshold because of an exponential decay in the KI vs. time to fail relationship [13,15,16,18]. A generalized equation (Eq. (3)) that best fits the data of KI vs. time to fail in distilled water (Dw) and 0.1 M NaCl solution saturated with Mg (OH)2 for EV31A as well as AZ91E is given as follows: Fig. 3. Hydrogen evolution measurements for EV31A and AZ91E alloys in 0.1 M NaCl solution saturated with Mg(OH)2 with immersion time.
Table 1 Summary of the mechanical properties of the two alloys. Alloy
Elongation (El %)
YS (MPa)
UTS (MPa)
E (GPa)
KIC (MPa m1/2)
EV31A AZ91E
6 4.5
170 170
280 250
44 44
15 13.2
K I ¼ a exp
t f þc b
ð3Þ
where KI is the stress intensity (MPa m1/2), and tf is the time to fail (h). KISCC, a, b, and c are summarized in Table 2. KISCC (KI corresponding to time to fail tending to infinity) as calculated from Eq. (3) was determined to be 10 MPa m1/2 (in Dw) and 8 MPa m1/2 (in NaCl) for EV31A. At tf ¼0, KI is 15 MPa m1/2 which is the fracture toughness (KIC) of EV31A. Similarly, using Eq. (3), KISCC was determined for AZ91E to be 8 MPa m1/2 (in Dw)
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
173
and 6 MPa m1/2 (in NaCl), and KI (at tf ¼0) was found to be 13 MPa m1/2 which is the KIC of AZ91E. Plots of log crack growth rate vs. initial KI for the two alloys tested in distilled water and Mg(OH)2 saturated 0.1 M NaCl are presented in Fig. 6. Stage I data points of the alloys in the two environments fit to an empirical linear relation given by da ¼ mK I þ n log ð4Þ dt
Fig. 6. Plot of log crack growth rate vs. initial stress intensity factor for alloy EV31A and AZ91E in distilled water (Dw) and Mg(OH)2 saturated 0.1 M NaCl (0.1 M) solution using CLT.
Table 2 Summary of fracture mechanics data of two alloys obtained by fitting to the data of Fig. 5 using Eq. (3) and of Fig. 6 using Eq. (4). KISCC corresponds to crack growth rate ¼10 10 m s 1 obtained using Eq. (4). Alloy
EV31A AZ91E
Env.
Dw NaCl Dw NaCl
Eq. (3)
Eq. (4)
a
b
c
m
n
5 7 5 7
30 7 220 37
10 8 8 6
5.5 7 0.8 2
65 66 16.4 22
KISCC (MPa m1/2)
Stage II da/dt (m s 1)
10 8 8 6
7.05 10 9 7.25 10 8 2.69 10 9 3.98 10 8
Env, environment; Dw, distilled water; NaCl, 0.1 M NaCl saturated with Mg(OH)2.
where crack growth rate (da/dt) is in m s 1, KI in MPa m1/2, m is the slope and n is the intercept on the axis of log crack growth rate. These data are summarized in Table 2 for each alloy– environment combination. Table 2 also presents stage II crack growth rates. The KISCC values in the table correspond to a crack growth rate of 10 10 m s 1 [18,28]. These KISCC data (determined using Eq. (4)) are consistent with those determined from KI vs. time to fail plot (using Eq. (3)). KISCC data of EV31A alloy presented here are among the first such data which show higher KISCC (i.e., greater resistance to SCC crack propagation) than alloy AZ91E in both the test environments. 3.3. Fractography Representative stereo micrographs tested in each alloy–environment combination are shown in Fig. 7. Various regions, viz., machined notch region, fatigue pre-cracked region (that had typical parabolic crack front), the SCC region (demarcated by white lines), and mechanically failed region, are delineated on the fractographs. Fracture surface of C(T) specimen of EV31A tested in distilled water is intentionally presented in Fig. 7(a) to show the maximum extent of variation in the crack depth over the nine locations that was measured on fracture surface of specimen that met the validity requirements for variation in crack dimensions [28]. Side grooves are clearly visible on all the micrographs (Fig. 7). There is no visible evidence of deformation along the thickness (Fig. 7) which indicates that the specimens satisfied the plane strain condition.
Fig. 7. Stereo micrographs of the fractured surfaces of the two magnesium alloys in distilled water (Dw) and Mg(OH)2 saturated solution of 0.1 M NaCl (0.1 M): (a) EV31A—Dw, (b) EV31A—0.1 M, (c) AZ91E—Dw, and (d) AZ91E—0.1 M, showing different regions: (1) machined notch, (2) fatigue-pre-crack, (3) SCC region (demarcated by white line), and (4) mechanical failed region.
174
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
Fractographic features of the different broad regions of the fracture surfaces shown in Fig. 7 were observed at higher magnifications, as seen in Figs. 8–10. Fatigue pre-crack regions of both
the alloys have transgranular striations, as the general feature (Fig. 8a and c); however, there are areas of dissolution along grain boundary precipitates, particularly when tested in the 0.1 M NaCl
Fig. 8. Fractographs of EV31A (a) fatigue region, (b) mechanical failed region, and of AZ91E (c) fatigue region and (d) mechanical failed region. Note – the dissolution near heavy grain boundary precipitates is shown by arrows (in a and c).
Fig. 9. Fractographs of EV31A for SCC region (a) showing TGSCC in distilled water and parallel facets (Detail A), (b) predominantly TGSCC with isolated occurrence of IGSCC (shown by arrows) and parallel facets (Detail B) in Mg(OH)2 saturated 0.1 M NaCl.
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
175
Fig. 10. Fractographs of AZ91E for SCC region: (a) TGSCC region in distilled water reveals secondary cracks (shown by arrows) and parallel facets (shown in Detail A), and (b) TGSCC in 0.1 M NaCl saturated with Mg(OH)2 (showing secondary cracks marked by arrows) and parallel facets (Detail B). Fractograph shown in B reveals localized attack which could affect the otherwise sharp transgranular crack front.
solution (presumably that occurred during the long CLT in 1008 h). The mechanically failed regions have fine dimples (Fig. 8b and d), sign of ductile failure, similar to the fracture feature of the two alloys that failed under slow strain rate tensile (SSRT) tests in glycerol [36]. EV31A tested in distilled water showed mostly transgranular feature (Fig. 9a). The transgranular SCC (TGSCC) in distilled water propagated in a discontinuous manner having parallel facets (Fig. 9a—Detail A). The SCC region of EV31A tested in Mg(OH)2 saturated 0.1 M NaCl showed mainly TGSCC as shown in Fig. 9b and b—Detail B. There are a few evidences of dissolution near grain boundary precipitates (shown by arrows in Fig. 9b). Similar fractographic features were reported by Kannan et al. [9] for the failed SSRT specimens of EV31A that shows TGSCC when tested in distilled water, and mixed intergranular SCC (IGSCC) and TGSCC in 0.5 wt% NaCl. They attributed the IGSCC feature to the electrochemical dissolution along the continuous grain boundary precipitates in the rare-earth containing alloys (i.e. ZE41, QE22 and EV31A). On the other hand, the TGSCC feature is discontinuous and involves alternating mechanical and electrochemical processes. Similar to EV31A, the SCC region of AZ91E tested in distilled water showed mostly transgranular feature with parallel facets (Fig. 10a and Detail A). SCC region of specimen tested in Mg(OH)2 saturated 0.1 M NaCl also showed mostly the TGSCC feature (Fig. 10b and Detail B) with evidence of secondary cracks shown by arrows in Fig. 10a and b. TGSCC features having parallel facets and evidence of dissolution are observed at higher magnification (shown in Detail B). The fractographic features observed in the SCC region of C(T) specimens of alloy EV31A and AZ91 tested in distilled water and 0.1 M NaCl saturated with Mg(OH)2 have similar features to the author's earlier work on SSRT specimens tested in the same solution [12,36]. The transgranular features having flat parallel facets (Figs. 9 and 10) indicate discontinuous crack growth and the possible role of hydrogen in cracking.
Winzer et al. [37] reported TGSCC feature for AZ91 alloy tested in distilled water under constant extension rate test using cylindrical tensile specimens. Chakrapani and Pugh [38] suggested SCC of an Mg–Al alloy in NaCl–K2CrO4 solution to be a form of hydrogen embrittlement and the fracture to be cleavage like. For pure Mg in NaCl–K2CrO4 solution, Meletis et al. [39] have reported the cracking to be TGSCC, having a series of parallel crystallographic facets separated by steps and jogs, similar to the features observed in the present study.
4. Discussion EV31A and AZ91E alloys have shown the presence of stage I and stage II crack growth characteristics and the absence of stage III crack growth (Fig. 6), much the same as reported in the literature for magnesium alloys [10,17–19], and aluminum alloys [21]. The log (da/dt) vs. KI plots bring out the difference in crack growth tendency of the two alloys. Notably, EV31A exhibits higher crack growth rate than AZ91E in stage II, even though its KISCC is greater. Several factors seem to contribute to the SCC resistance of the alloys. The addition of chloride has caused a significant decrease in KISCC of both the alloys. Chloride ions increase the tendency for IGSCC as against TGSCC found in distilled water for EV31A alloy (Fig. 9). Between the two alloys, AZ91E liberates more hydrogen than EV31A (Fig. 3), which could be one of the reasons why the former exhibits lower KISCC than the latter. Not much is reported on the slope for stage I crack growth and crack velocity for stage II, though such studies have been carried out on stainless steels [20]. Jones and Simonen [20] suggested that “m”, the slope for stage I crack growth, increases for alloys exhibiting higher KISCC and when the environment is more aggressive. Examination of Table 2 shows that between EV31A and AZ91E, the latter exhibits lower m (slope corresponding to stage I crack growth) and stage II crack growth rate than the
176
B.S. Padekar et al. / Materials Science & Engineering A 583 (2013) 169–176
former. This is despite the fact that EV31A exhibits higher KISCC and produces less hydrogen than AZ91E in chloride containing environment (Fig. 3). An alloy should be more susceptible to hydrogen embrittlement and exhibit higher crack growth rate if it liberates more H2. The reasons for such a behavior are not quite clear, though the following explanations are plausible. Comparison of the microstructures of EV31A and AZ91E indicates that the latter possesses a complex microstructure (β phase and eutectic phases), which can act as crack arresters. Further, it seems these microstructures also cause localized attack, leading to possible crack blunting. This suggestion is made based on the fact that selective and localized attack has been found on the fracture feature of AZ91E alloy (Fig. 10), which is absent in EV31A alloy (Fig. 9). The effect of Al on the localized attack is worth discussing. It is known that in the presence of Al3+, the solution pH within cracks and crevices can decrease down to 4 because of hydrolysis as given by the following equation [21]: 3þ
Al
þ 3H2 O-AlðOHÞ3 þ 3Hþ
ð5Þ
Such reduction in pH can accelerate corrosion and possible crack blunting (Fig. 10b–Detail B). 5. Conclusion Stress corrosion crack growth behavior of a sand cast rare earth containing magnesium alloy, EV31A, in distilled water and 0.1 M NaCl solution saturated with Mg(OH)2 was studied using compact tension specimen under constant load condition for the first time. For comparison, a traditional Mg-alloy, AZ91E, was also investigated. The main conclusions of this study are as follows: (1) The KISCC in distilled water and Mg(OH)2 saturated solution of 0.1 M NaCl were determined for EV31A to be 10 and 8 MPa m1/2 and for AZ91E to be 8 and 6 MPa m1/2 respectively, indicating the superior SCC resistance of EV31A than AZ91E. However, AZ91E shows lower stage II crack growth rate than EV31A in both the test environments, which is attributed to a possible crack blunting and complex microstructures. (2) Slope, m, increases and KISCC decreases with the aggressiveness of the environment. (3) Despite higher H2 evolution tendency, AZ91E exhibits lower stage II crack growth rate than that of EV31A. Both the alloys exhibited increase in stage II crack growth rate due to chloride addition.
References [1] H.E. Friedrich, B.L. Mordike, Magnesium Technology, Metallurgy, Design Data, Applications, Springer Verlag, Berlin, Heidelberg, 2006.
[2] L.L. Rokhlin, Magnesium Alloys Containing Rare Earth Metals, Taylor and FrancisLondon, 2003. [3] J.F. Nie, X. Gao, S.M. Zhu, Scr. Mater. 53 (2005) 1049–1053. [4] C. Sanchez, C. Nussbaum, P. Azavant, H. Octor, Mater. Sci. Eng. A 221 (1996) 48–57. [5] P. Lyon, T. Wilks, I. Syed, in: N.R. Neelameggham, H.I. Kaplan, B.R. Powell (Eds.), Magnesium Technology 2005, Warrendale, PA, 2005, pp. 303–308. [6] G. Ben-Hamu, D. Eliezer, K.S. Shin, S. Cohen, J. Alloy Compds. 431 (2007) 269–276. [7] R.G. Song, C. Blawert, W. Dietzel, A. Atrens, Mater. Sci. Eng. A 399 (2005) 308–317. [8] H. Baker, M.M. Avedesian (Eds.), ASM Specialty Handbook: Magnesium and Magnesium Alloys, ASM International, USA, 1999. [9] M. Bobby Kannan, W. Dietzel, C. Blawert, A. Atrens, P. Lyon, Mater. Sci. Eng. A 480 (2008) 529–539. [10] N. Winzer, A. Atrens, G. Song, E. Ghali, W. Dietzel, K.U. Kainer, N. Hort, C. Blawert, Adv. Eng. Mater. 7 (2005) 659–693. [11] European Cooperation for Space Standardization, Material Selection for StressCorrosion Cracking, ECSS-Q-ST-70-36C, 6 March 2009. [12] B.S. Padekar, R.K.S. Raman, V.S. Raja, P. Lyon, Corros. Sci. 71 (2013) 1–9. [13] R.I. Stephens, C.D. Schrader, D.L. Goodenberger, K.B. Lease, V.V. Ogarevic, S.N. Perov, SAE, No. 930752, USA, 1993, pp. 843–857. [14] N. Winzer, A. Atrens, W. Dietzel, G. Song, K.U. Kainer, Adv. Eng. Mater. 10 (5) (2008) 453–458. [15] L. Choudhary, R.K.Singh Raman, Acta Biomater. 8 (2012) 916–923. [16] A. Moccari, C.R. Shastry, J. Mater Technol. (Zeitschrift für Werkstofftechnik) 10 (1979) 119–123. [17] K. Ebtehaj, D. Hardie, R.N. Parkins, Corros. Sci. 28 (1993) 811–829. [18] B.S. Padekar, V.S. Raja, R.K.Singh Raman, Eng. Fract. Mech. 102 (2013) 180–193. [19] M.O. Speidel, M.J. Blackburn, T.R. Beck, J.A. Feeney, Corrosion Fatigue: Chemistry, Mechanics and Microstructure, NACE-2, 1972, pp. 324–345. [20] R.H. Jones, E.P. Simonen, Mater. Sci. Eng. A160 (1993) 127–136. [21] M.O. Speidel, Metall. Trans. 6A (1975) 631–651. [22] W. Dietzel, P.B. Srinivasan, in: V.S. Raja, T. Shoji (Eds.), Stress Corrosion Cracking Theory and Practice, Woodhead Publishing, Oxford, UK, 2011, pp. 133–166. [23] P. Lyon, I. Syed, S. Heaney, Adv. Eng. Mater. 9 (2007) 793–798. [24] International Standard, Corrosion of Metals and Alloys—Stress Corrosion Testing, ISO 7539-6, 2006. [25] Standard Test Method for Linear-Elastic Plane-Strain Fracture Toughness KIC of Metallic Materials, ASTM E 399, 2009. [26] Standard Test Method for Determining Threshold Stress Intensity Factor for Environment-Assisted Cracking of Metallic Materials, ASTM E 1681, 2003. [27] Standard Test Methods for Tension Testing of Metallic Materials, ASTM E 8, 2011. [28] W. Dietzel, Guidelines for Fracture Mechanics SCC Testing, Characterization of Susceptibility of Metallic Materials to Environmentally Assisted Cracking, GKSS-Forschungszentrum Geesthacht GmbH, 1995. [29] Standard Test Method for Determining Average Grain Size, ASTM E 112, 2010. [30] Q. Li, Q. Wang, Y. Wang, X. Zeng, W. Ding, J. Alloy Compds. 427 (2007) 115–123. [31] R.K.Singh Raman, Metall. Mater. Trans. A 35 (2004) 2525–2531. [32] M. Liu, P.J. Uggowitzer, A.V. Nagasekhar, P. Schmutz, M. Easton, Corros. Sci. 51 (2009) 602–619. [33] Venugopal, R.D.Angal, V.S. Raja, Corrosion 52 (1996) 138–142. [34] G. Song, A. Atrens, D.H. StJohn, in: J. Hryn (Ed.), Magnesium Technology, TMS, New Orleans, PA, 2001, pp. 255–262. [35] Data sheet 455 and 456, Magnesium Elektron Ltd. UK, 〈http://www.magnesiu m-elektron.com/data/downloads/DS455.PDF〉. [36] B.S Padekar, V.S. Raja, R.K. Singh Raman, P. Lyon, NACE Corrosion Conference and Expo, 2012, pp. 1556–1559. [37] N. Winzer, A. Atrens, W. Dietzel, G. Song, K.U. Kainer, Mater. Sci. Eng. A 472 (2008) 97–106. [38] D.G. Chakrapani, E.N. Pugh, Metall. Trans. 7A (1976) 173–178. [39] E.I. Meletis, R.F. Hochman, Corrosion 40 (1) (1984) 39–45.