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Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y x x x ( 2 0 1 7 ) 1 e1 3
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Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions L.F. Zhou a,b, Z.Y. Liu a,b,*, W. Wu a,b, X.G. Li a,b,c, C.W. Du a,b,c, B. Jiang d a
Corrosion and Protection Center, University of Science and Technology Beijing, Beijing 100083, PR China Key Laboratory for Corrosion and Protection (MOE), Beijing 100083, PR China c State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China d College of Materials Science and Engineering, Chongqing University, Chongqing 400030, China b
article info
abstract
Article history:
The stress corrosion cracking (SCC) behavior of ZK60 magnesium alloy was investigated
Received 30 June 2017
under different conditions, i.e. thin electrolyte layer (TEL) and solution, by slow strain rate
Received in revised form
tensile tests, electrochemical techniques, Auger electron spectroscopy, scanning electron
22 August 2017
microscopy coupled with electron backscattered diffraction, and time of flight secondary
Accepted 23 August 2017
ion mass spectrometry. Results indicated that the ZK60 magnesium alloy in solution ex-
Available online xxx
hibits a higher SCC susceptibility with a combined SCC mechanism of weaker anodic dissolution (AD) and stronger hydrogen embrittlement (HE) compared to under TEL.
Keywords: Magnesium alloys
Moreover, the HE mechanism under various conditions was discussed. © 2017 Published by Elsevier Ltd on behalf of Hydrogen Energy Publications LLC.
Thin electrolyte layer Stress corrosion cracking Anodic dissolution Hydrogen embrittlement
Introduction Among structural materials such as steel, titanium, and aluminum alloys, etc., magnesium alloys exhibit the lowest density (r z 1.74 g$cm3), making them attractive for automotive and aerospace industries because of the significance of weight [1e3]. In addition, magnesium alloys also possess several other advantages, such as excellent castability, electromagnetic interference shielding, non-magnetic property and recyclability. Despite these intriguing properties, their use has historically been restricted considerably because of their chemical reactivity, making them less resistant to corrosion
[4e9], corrosion fatigue [10,11], and stress corrosion cracking (SCC) [12e25]. SCC is typically related to mechanisms associated with anodic dissolution (AD) [26e29] or hydrogen embrittlement (HE) [30e38]. Several studies [17,18,24] have reported that HE is the typical mechanism of SCC in magnesium alloys. However, some others [13,14,25] have concluded that SCC in magnesium alloys is caused by the combination of AD and HE. Thus, the exact mechanism of SCC in magnesium alloys remains unclear. Moreover, several mechanisms may simultaneously occur depending on the material and other variables (e.g., strain rate and ambient conditions). However, the two above-mentioned
* Corresponding author. Corrosion and Protection Center, University of Science and Technology Beijing, Beijing 100083, PR China. E-mail address: [email protected] (Z.Y. Liu). http://dx.doi.org/10.1016/j.ijhydene.2017.08.161 0360-3199/© 2017 Published by Elsevier Ltd on behalf of Hydrogen Energy Publications LLC. Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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mechanisms that play a dominant role in the SCC of magnesium alloys are still controversial. Magnesium alloys most likely suffer from atmospheric corrosion, which is an electrochemical process occurring on a metal surface covered with a thin electrolyte layer (TEL) [39]. Among all corrosion modes, the SCC behavior of metals under TEL is significantly different from that in a bulk solution [40e42]. The corrosion of magnesium alloys is thought to involve the cathodic reaction of oxygen reduction, and this cathodic process is significantly enhanced for samples tested under TEL. However, oxygen reduction is not as important as hydrogen evolution during the corrosion of alloys in solution. Thus, the associated mechanisms under TEL are possibly related to the combined effect of AD and HE, while HE is believed to be the most common mechanism for the SCC of magnesium alloys in solution. To identify the dominant mechanisms related to the SCC of magnesium alloys, and to support the increased use of magnesium alloys for stressed components in aerospace and automobile industry in the marine atmosphere, the SCC mechanisms of magnesium alloys under TEL and in solution must be systematically investigated. Hence, with the limited knowledge about the SCC behavior of the ZK60 magnesium alloy, slow strain rate tensile (SSRT) tests are conducted using the ZK60 magnesium alloy under TEL and in solution. In addition, the surface appearance and fractography of all samples were observed by scanning electron microscopy (SEM); the surface film was examined using high-resolution Auger electron spectroscopy (AES); the cracking mode was characterized by electron backscattered diffraction (EBSD); and the distribution of elements in the vicinity of cracks was examined by time-of-flight secondary ion mass spectroscopy (ToF-SIMS). The results of this study may provide a comprehensive understanding on the dominant SCC mechanisms of magnesium alloys under different conditions.
Experimental Materials and environment Commercial ZK60 Magnesium alloy with a nominal composition of Mge6Zne0.5Zr (wt%) was used. Inductively coupled plasmaeatomic emission spectroscopy analysis of the asreceived alloy revealed a Zr content of 0.34 wt%. This value is close to the low limit for alloys with this denomination. Table 1 summarizes the chemical composition. The experiment device shown in Fig. 1 comprised a test medium, maintained at ambient temperature (approximately 25 C), of a modified simulated TEL atmosphere, which facilitated easy oxygen transfer. The solution in the atomizer was 0.1%NaCl þ 0.05%Na2SO4 þ 0.05%CaCl2 with a pH of
approximately 5.1 related to CO2 dissolution. By contrast, the test medium in the solution was the same as that in the atomizer. However, the transfer of oxygen in solution was weaker; hence, the HE effect is possibly stronger than that under TEL. Samples for stereoscopic microscope observations were mechanically polished to 1 mm using ethanol-based diamond suspensions followed by chemical etching using a solution of 5 g picric acid þ 5 g acetic acid þ 20 mL distilled water þ 100 mL anhydrous ethanol. Similarly, the samples for EBSD characterization were mechanically polished and electropolished using 10% perchloric acid in ethanol at 0 C and at a polishing voltage of 15 V for 30 s. The EBSD samples were observed in the vicinity of the cracks. Automatic EBSD scans were recorded using TSL data acquisition software with a step size of 1 mm. The EBSD data were analyzed using TSL OIM software.
SSRT tests SSRT tests were carried out to investigate the SCC behavior of ZK60 alloy under different conditions. A flat-plate tensile specimen was prepared based on GB/T 15970 [43]. Prior to each SCC test, the tensile samples were prepared from electrodischarge machining. The samples were successively ground using silicon carbide papers up to 2000 grit, ultrasonically cleaned in absolute ethanol for 10 min, and dried under a flow of cool air. SSRT tests were carried out on a WDML-30kN Materials Test System with a strain rate of 1 106 s1 (Fig. 1a). Before each test, the specimen was maintained under 100% RH for 1 h to ensure the formation of an electrolyte film on the surface. Moisture was continuously pumped during the entire experiment. After failure, one piece of the fracture was analyzed for characterization, and the other was applied for the measurement of the hydrogen concentration. First, the percentage of elongation and area reduction of each specimen were calculated after the removal of corrosion products. In addition, fracture morphologies were observed by SEM. Each test was reproduced at least three times to ensure the reliability of experimental data, and the standard deviation of the data was measured. Next, the cross-sectional morphologies of the fracture were observed to investigate the localized corrosion severity and cracking mode. Moreover, the hydrogen concentration of the tested samples was measured using a hydrogen analyzer (G4 Phoenix DH).
Electrochemical measurements Fig. 1b shows the measurement system comprising the typical three-electrode system. The working electrode is ZK60 specimen with a working area of 1 cm2, the counter electrode is a Pt bulk circling around the working electrode, and the
Table 1 e Chemical composition of the ZK60 magnesium alloy. Alloy
ZK60
Chemical composition (mass%) Al
Zn
Mn
Fe
Si
Cu
Ni
Zr
Mg
e
5.12
0.02
0.0024
0.003
0.0019
0.0011
0.34
Bal.
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Fig. 1 e Schematic of experimental devices for (a) SSRT and (b) electrochemical measurements.
reference electrode is a saturated Ag/AgCl electrode [44,45]. After a pumping time of 20 min for electrochemical impedance spectroscopy (EIS) measurement, the scanning frequency ranged from 100 kHz to 10 mHz, with a perturbing AC amplitude of 10 mV. The polarization curve was obtained on an electrochemical working station (PARSTAT2273, Princeton). Potentiodynamic polarization curves were recorded at a potential sweep rate of 0.5 mV/s, which ranged from 1750 mV to 1250 mV (vs. Ag/AgCl). For better reproducibility, polarization curves and EIS measurements were repeated more than three times. All tests were conducted at ambient temperature.
AES and TOF-SIMS measurements AES experiments were conducted on a UHV scanning Auger nanoprobe (PHI 700, ULVAC-PHI). The Ar GCIB-sputtered sample was rapidly transferred into the AES chamber using a vacuum transfer vessel capable of maintaining a low vacuum level and minimizing sample contamination during transfer. AES measurements were carried out at a beam voltage and current of 5 kV and 10 nA, respectively, with a SiO2 sputter rate of 17 nm/min. TOF-SIMS was employed for local chemical analysis around the cracks using an ION-TOF GmbH TOF-SIMS 5-100 analyzer. A highly focused primary Csþ ion
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beam in a rectangle of 50 mm was scanned across the cracks. Secondary ions, resulting from abrasion, were analyzed at each point to reconstruct a micrograph. Different signals were collected during the analysis. The high-resolution mode was utilized to avoid interferences between the different signals.
Table 2 e Hydrogen concentration of ZK60 samples after SSRT under different conditions. Test conditions
Air
TEL
Solution
CH(ppm)
1.31
12.57
17.89
300
Initial microstructure
250
Fig. 2 shows the initial microstructure and texture of ZK60 magnesium alloy. The material exhibited a microstructure with a grain size of 5e200 mm (Fig. 2a), which was also observed in the EBSD spectrum in Fig. 2b. Moreover, it was obviously seen that a thimbleful of twin boundaries subsist in the grain.
200
Hydrogen concentration and SCC behavior of the ZK60 magnesium alloy
Stress(MPa)
Results
150
100
0
After SSRT failure, the concentrations of hydrogen in the samples were measured (Table 2). The concentrations of hydrogen in the samples sharply increased in a sequence of in air (1.31 ppm), under TEL (12.57 ppm), and in solution (17.89 ppm). Fig. 3 shows the stressestrain curves of ZK60 tested under different conditions. Table 3 summarizes the change and decrease in terms of strength and plasticity, respectively, for the comparison of their mechanical properties. The ZK60 magnesium alloys exhibited different values of s0.2 and ultimate tensile strength (UTS) under various environments (Table 3). High values of s0.2 and UTS were observed for samples tested in air, while lower values were observed for samples tested under TEL and in solution. The concentrations of hydrogen in the samples tested under TEL and in solution were different because of different experimental environments, thereby leading to different mechanical properties of the ZK60 magnesium alloys. Compared to samples exposed to air, those exposed to a corrosive medium exhibited a considerable reduction in the s0.2 and UTS values. Samples tested in air, under TEL, and in solution exhibited a considerable difference in their tensile properties. The plasticity of samples tested under TEL was slightly lower than that of the samples tested in solution (Table 3). Moreover, the
in air under TEL in solution
50
0
5
10 15 Strain(%)
20
25
Fig. 3 e Stress versus strain curves for the ZK60 magnesium alloy at a strain rate of 1 £ 10¡6 s¡1 under different conditions.
Table 3 e Summary of the mechanical properties of the ZK60 magnesium alloy samples tested under various conditions. Alloy Conditions s0.2 (MPa) UTS (MPa) d(%) j(%) Decrease of s0.2 (%) Decrease of UTS (%) Decrease of d (%) Decrease of j (%)
ZK60 Air
TEL
Solution
240.6 ± 7.3 267.2 ± 7.2 21.4 ± 1.5 20.7 ± 1.6 e e e e
192.3 ± 5.4 205.7 ± 5.6 8.0 ± 0.6 8.3 ± 0.6 20.1 23.0 62.6 60
e 179.0 ± 4.3 3.7 ± 0.3 5.2 ± 0.4 e 33.0 82.7 94.7
tensile properties of the samples tested under TEL and in solution significantly decreased in comparison with those in air. Moreover, the values for the elongation-to-failure (d) and area reduction (j) tested in solution were lower than that of the
Fig. 2 e Initial microstructure of ZK60 magnesium alloy: (a) metallograph, (b) inverse pole figure. Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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samples tested under TEL. In addition, the following two equations, Id ¼ ð1 dmedium Þ=dair 100%
(1)
Ij ¼ ð1 jmedium Þ=jair 100%
(2)
were applied to calculate the loss of d and j. Herein, dair and jair represent the d and j values in air, dmedium and jmedium represent the d and j values under TEL and in solution, respectively. To quantify the SCC susceptibility of the samples tested under different conditions, the SCC susceptibility index (Id and Ij) was calculated by the decline in elongation and area reduction. Id and Ij are high for a system with high SCC susceptibility, whereas Id and Ij tend to be close to 0 for low susceptibility. The analysis above indicates that to a certain extent, the SCC susceptibility of samples tested in solution, is greater than that of samples tested under TEL.
Electrochemical characterization Fig. 4 shows the potentiodynamic polarization curves and EIS curves of the ZK60 magnesium alloy under different conditions. The corrosion potentials (Ecorr) and corrosion current
-1.1
(a)
-1.2
E(VAg/AgCl)
-1.3
-1.5 -1.6 -1.7 -1.8 -1.9 1E-8
2500
1E-7
-Zim(Ω.cm2)
1E-6 1E-5 i (A/cm2)
1E-4
1E-3
(b) under TEL in solution
2000
CPE R
1500
CPE + R R +
L
1000 500 0 0
500
densities (Icorr) under the two conditions were obtained (Fig. 4a) by the Tafel fitting method [46]. Table 4 shows the obtained results. From the thermodynamic perspective, the corrosion potential of ZK60 in solution shifted to the noble direction compared to the samples tested under TEL, possibly related to the formation of a thicker film formed on the matrix; this point is discussed later with the AES results. However, in terms of dynamic perspective, the anodic Tafel slope (ba) and cathodic Tafel slope (bc) of the samples tested under TEL were slightly greater than those of the samples in solution, indicative of various electrochemical reactions. Moreover, the cathodic process of magnesium alloys tested under TEL was dominated by hydrogen evolution [40], and the presence of oxygen inhibited hydrogen evolution. Three loops were observed in the EIS curvesdone high-frequency capacitive loop, followed an intermediate frequency capacitive loop, and a final low-frequency inductive loopdrespectively. The equivalent circuit in Fig. 4b was used for EIS modeling, where Rs represents the solution resistance, (CPE)dl represents the double layer capacitance, Rt represents the charge-transfer þ resistance, (CPE)þ Mg represents the capacitance of the Mg reaction on the breaking area of the partial protective film [47], þ Rþ Mg represents the resistance of the Mg reaction, and L represents the inductance of the partial protective film [47]. Moreover, the EIS results revealed that the impedance of the samples tested under TEL slightly decreases. This result is in agreement with the low corrosion current density value mentioned above.
Cross-sectional and secondary cracking observations
under TEL in solution
-1.4
5
1000
1500
Zre(Ω.cm2)
2000
2500
Fig. 4 e (a) Potentiodynamic polarization curves and (b) electrochemical impedance spectroscopy curves of the ZK60 magnesium alloy under different conditions.
To investigate the cracking mechanism, the cross-sectional and secondary cracks of the samples subjected to tensile tests under different conditions were observed (Fig. 5). Samples tested under TEL exhibited a slight localized corrosion on the surface (Fig. 5a). Meanwhile, several secondary cracks initiated at the localized corrosion areas (Fig. 5a). On the contrary, the corrosion attack of samples tested in solution suffered relatively greater attack with deep corrosion ditches on the surface (Fig. 5b). Moreover, the density of secondary cracks of the samples tested in solution was considerably less than those of the samples tested under TEL (Fig. 5b). Based on their cross-sectional morphologies, an interesting phenomenon was observed for the samples tested under TEL: the severity of localized corrosion areas was slight; however, a number of secondary cracks distributed at or nearby the localized corrosion areas were easily observed (Fig. 5c). For the samples in solution, the localized corrosion areas were relatively more severe; however, fewer secondary cracks were observed (Fig. 5d). To further disclose the cracking mode and initiation location of cracks, EBSD analyses of the two sample surfaces were carried out (Fig. 6). Fig. 6a and c shows the image qualityinverse pole figure (IQ-IPF) maps of two samples. The colors of grains in the maps corresponded to the crystallographic axes of grains in the inverse pole figure (Fig. 2b). Cracks were observed under two conditions, with propagation in the transgranular direction. To confirm the stress concentration in the grains at the tips of cracks, the kernel average misorientation (KAM) values were measured by SEM-EBSD (Fig. 6b
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Table 4 e Electrochemical parameters fitted from the potentiodynamic polarization curves of the ZK60 magnesium alloy under different conditions. Conditions TEL Solution
Ecorr (VAg/AgCl) 1.605 1.547
Icorr (A/cm2) 6
4.724 10 9.652 106
and d). A high KAM value typically reflects the distribution of the geometry necessary dislocation (GND) related to the localized deformation [48] that contributes to the film rupture. No considerable difference in the KAM values between the cracking paths and their surrounding area was observed, and
ba (mVAg/AgCl/decade)
bc (mVAg/AgCl/decade)
311.5 156.2
389.9 181.7
both cracks propagated along the high KAM value region (as highlighted by the arrows). As the pileup of dislocations promoted the local anodic dissolution and hydrogen transfer and aggregation in front of SCC areas [49], EBSD results revealed that the SCC of the ZK60 magnesium alloy is possibly
Fig. 5 e Microstructural analysis of the failed samples tested under different conditions: (a,b) secondary electron images of the sample cross-sectional surfaces under TEL and in solution, respectively; (c,d) secondary electron images of the secondary cracks of samples under TEL and in solution, respectively; (e,f) are high-magnification images of the rectangle area of (c,d), respectively. Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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Fig. 6 e EBSD IQ-IPF maps (a,c) and IP-KAM maps (b,d) recorded near the cracks in Fig. 5e and f under TEL (a,b) and in solution (c, d).
controlled by both AD and HE which we would further explain it in discussion.
Fractography Fig. 7 shows the SEM morphologies of the fracture surface under different conditions. Marginal necking was observed for all samples (Fig. 7a, c, and e). The fracture surface of ZK60 tested in air showed features consistent with quasi-cleavage and plastic dimples (Fig. 7b). For the samples tested under TEL, the primary part of the fracture surface was smooth, which showed typical transgranular stress corrosion cracking features (Fig. 7d) [15]. On one hand, the fracture surface showed considerable dissolution, microvoids, and minor cracks, indicating that vapor easily penetrates in the matrix, leading to a complex phenomenon. On the other hand, the transgranular SCC (TGSCC) feature was discontinuous, involving alternating mechanical and electrochemical processes [13]. Meanwhile, for the samples tested in solution, the TGSCC in solution propagated in a discontinuous manner with parallel facets; similarly, evidence of dissolution, microvoids, and minor cracks was observed (Fig. 7f). However, the severity of localized corrosion and the number of minor cracks and microvoids were greater than those of the samples tested under TEL. Transgranular features characterized by flat parallel facets (Fig. 7d and f) indicated the growth of discontinuous cracks and
the possible role of hydrogen in cracking. Padekar et al. [13] have mostly observed TGSCC for EV31A and AZ91E in distilled water. They have proposed that this feature is related to the possible role of hydrogen in cracking. Chakrapani and Pugh [31,32] have inferred that TGSCC occurs by a brittle cleavage mechanism involving H, resulting in stepped, faceted interlocking fracture surfaces similar to the features observed herein.
Local chemical analysis AES compositional depth profiling (Fig. 8) was carried out to examine the films formed on the surfaces of the failed tensile samples tested under different conditions. Analyses were performed on the uncorroded area distributed throughout the surface. In all cases, we reached to the alloy substrate, which exhibited a weak O signal and strong Mg signal (Fig. 8). C caused by surface contamination characterized in the AES spectra was not included in the AES compositional depth profiles. The surface films formed on the tensile samples tested under TEL (~24 nm) and in solution (~55 nm) were thicker than that in air (~19 nm). Interestingly, compared with the surface films for the tensile samples tested under TEL, that for the samples tested in solution were thicker. Irrespective of the test conditions, the top part of the film comprised hydromagnesite [6,7]. The O/Mg ratio gradually decreased in the deeper part of the film, indicating the presence of Mg(OH)2 and MgO. Notably, the lack of straightforward evidence for Mg(OH)2 and MgO in the film in Fig. 8 does not rule out the
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Fig. 7 e Fractography of samples under different conditions: (a,b) in air; (c,d) under TEL; (e,f) in solution. presence of Mg(OH)2 and MgO layer at the film/metal interface as reported by several studies [50,51]. For evaluating the local chemical composition along and near the cracks and to detect possibly absorbed hydrogen or hydrides, ToF-SIMS ion images of H and O around the cracks of the ZK60 samples developed during SSRT under various conditions were recorded. Fig. 9 shows the SIMS maps of the hydrogen (Fig. 9a and c) and oxygen signals (Fig. 9b and d) respectively, for two cracks grown in a ZK60 sample subjected to SSRT under different conditions. The hydrogen and oxygen signals were considerably different under both conditions. Moreover, most of the hydrogen signals in the cracks can be assumed to correspond to the adsorbed hydrogen or hydrides, i.e., MgH2. Meanwhile, oxygen signals were scarcely detected in the cracks. For the crack developed under TEL (Fig. 9a and b), several areas showed high, superimposed signals for hydrogen and oxygen,
related to the formation of Mg(OH)2. However, contrary to the sample tested under TEL, other hydrogen signals were observed in addition to the hydrogen signals observed for the cracks for the sample tested in solution (Fig. 9c and d). Moreover, more hydrogen and oxygen signals in highly superimposed areas were observed compared to the sample tested under TEL, which is also related to the generated Mg(OH)2.
Discussion Effect of the environment on the SCC susceptibility and mechanical properties As mentioned above, the ZK60 magnesium alloy exhibited higher SCC susceptibility in bulk solution than that observed for under TEL, indicating that oxygen reduction is not as
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100
Concentration(at.%)
80
60
Mg
O
air TEL solution
40
20
0
0
20
40
60 Depth(nm)
80
100
120
Fig. 8 e AES depth profiles (area ¼ 2 £ 2 mm2) of the uncorroded area of ZK60 magnesium alloy surface tested under different conditions.
important as hydrogen evolution during SCC. The conditions for the sample in solution were clearly more favorable for hydrogen evolution than those for the sample tested under TEL, which is clearly more beneficial for oxygen reduction. The environment is thought to considerably affect SCC
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susceptibility. Based on Equations (1) and (2), the calculated Id and Ij values for the samples tested under TEL were 0.63 and 0.6, suggesting the samples tested under TEL were moderately susceptible to SCC. However, the calculated Id and Ij for the samples tested in solution were 0.83 and 0.95, suggesting that the SCC susceptibility is greater in solution than that under TEL. Moreover, the mechanical properties were also considerably affected by the environment. The s0.2 and UTS values were considerably less than those observed in air for all samples exposed to corrosive medium, possibly related to the HE [18,33,52]. In addition, the higher SCC susceptibility and the decrease in the mechanical properties of samples tested in solution are related to the weaker oxygen reduction and stronger hydrogen evolution compared to the samples tested under TEL.
Relationship between the electrochemical properties and SCC behavior The conventional electrochemical measurement result indicated that the Ecorr of the samples tested under TEL was relatively more negative compared to that of the samples in solution (Fig. 4a). The Tafel fitting and EIS results indicated that the samples tested in solution exhibited a considerably higher Icorr and slightly lower impedance than those of
Fig. 9 e ToF-SIMS ion images of (a, c) H and (b, d) O recorded on the cracks of the ZK60 samples after SSRT under various conditions, (a, b) under TEL and (c, d) in solution. Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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samples tested under TEL, respectively (Fig. 4). The lower Ecorr of the samples tested under TEL is related to the thinner film formed on the surface of the matrix; thus, the corrosion under TEL is prone to occur. However, lower Icorr and higher impedance of the samples tested under TEL are related to two main reasons. On the one hand, oxygen reduction is not as important as hydrogen evolution during corrosion; on the other hand, the formation or deposition of Mg(OH)2 is easier for samples under TEL, which may deter corrosion. These results are in agreement with those obtained from AES and secondary cracking observations. Hence, on first blush, the AD of samples occurred, resulting in localized corrosion and stress concentration in this area, which could easily become SCC nucleation sites. These sites in turn can serve as pathways for hydrogen to diffuse into the matrix [33,38].
AD role in SCC mechanism According to several books and studies [51e55], possible reactions involving anodic and cathodic reactions are observed between the samples under TEL and in solution; in this study, these reactions are listed as follows: Anodic reaction Mg/Mgþ þ e
(3)
Mgþ /Mg2þ þ e
(4)
2Mgþ /Mg þ Mg2þ /Mg$Mg2þ
(5)
Mg$Mg2þ þ 2Hþ /2Mg2þ þ H2
(6)
Cathodic reaction: TEL : 2H2 O þ 2e /H2 þ 2OH
(7)
O2 þ 2H2 O þ 4e /4OH
(8)
Solution : 2Hþ þ 2e /H2
(9)
Effect of HE on the SCC mechanism
Salt formation: Mg2þ þ 2OH /MgðOHÞ2
diffusion in solution adjacent to the surface is relatively fast. Hence, the presence of the film accompanying a rapid anodic dissolution process in solution significantly accelerates the dissolution kinetics of the film, leading to a considerably high corrosion rate of ZK60. Although the dissolved Mg2þ can react with OH and form Mg(OH)2 deposits on the surface, the loosely deposited Mg(OH)2 may deposit onto the uncorroded area other than the broken areas. The evolved hydrogen from the corroding areas can stir the deposited Mg(OH)2 and prevent the corroding areas from being fully covered by the deposited Mg(OH)2. This observation is in agreement with the AES profiles. The transgranular stress corrosion cracking is related to the electrochemical corrosion or mechanical rupture of protective films, permitting the evolved hydrogen to enter the alloy matrix and eventually causing embrittlement and premature fracture [58]. Therefore, the nature and stability of the oxide or hydroxide films that develop on magnesium alloys is an important parameter for their SCC. The samples tested under TEL and in solution distinctly formed different oxide or hydroxide scales. Moreover, the samples tested in solution developed a more robust and sturdy layer; hence, the films are less susceptible to mechanical rupture, and the electrolyte can only infiltrate into the matrix to create a less number of cracks compared to the samples under TEL (Fig. 5c and d). However, once the films were undermined by the applied strain rate and chloride ions, the detrimental ions can easily penetrate into the matrix to cause localized corrosion [36,52]. Therefore, the AD mechanisms may be greater, and the HE mechanisms may be weaker for samples tested under TEL compared to the samples tested in solution. Moreover, the severity of localized disruption also increased the possibility for hydrogen to enter the Mg matrix, which explains the increase of SCC susceptibility in solution.
(10)
Based on the reactions above, the bulk solution is acidic with a pH of approximately 5.1; however, it is proverbial that the corrosion of Mg and its alloys results in the increase of solution pH, related to the surface alkalization effect [56,57]. Herein, the acidity of the electrolyte on the specimen surface can increase to attain a high pH value, and the corrosion continues to proceed at this higher pH value. At this time, the deposition of Mg(OH)2 (reaction (10)) dominates at this pH level. Whereas for the samples tested under TEL, only a small number of aqueous drops or a thin electrolyte film remains on the surface of the tested samples, while the samples tested in solution are all immersed. Hence, the corrosion of the sample under TEL (Fig. 5a) is slower than that of the sample in solution (Fig. 5b) because of the strong surface alkalization effect. In addition, the hydrolysis of Mg2þ in solution will momentarily decrease the pH of the surface, contributing to the aggravation of the reaction (9). In nature, a dynamic process of OH
To further examine the effect of HE on the SCC behavior of ZK60, the hydrogen and oxygen signals in the vicinity of cracks were examined (Fig. 9). The “H” signal was completely detected in the cracks, but no “O” signal was detected. The surface film on Mg is well known to not be thick or compact, thereby offering only a limited level of protection. In the early stage of SSRT, because of the pseudo-protection of the film and mechanical rupture, an electrochemical cell occurs between active area of anode and cathode which is capable of preventing film repair and sustaining enhanced anodic dissolution. Moreover, hydrogen can easily diffuse into the bulk of the material and be absorbed at the crack tips or to form hydrides. H diffusion has been proposed to be critical for several SCC mechanisms for magnesium alloys [36,37]. Moreover, the transgranular cracks intruded into the grains and a high density of dislocations was formed near the intruded cracks because the KAM value was proportional to the microstrain caused by crystal defects, such as dislocations (Fig. 6b and d). The grains exhibited a higher KAM compared to the grain boundaries possibly because of the difference in the mobility of dislocations between the grains and grain boundaries. Namely, the grains with a high initial density of dislocations could possibly undergo local plasticity because of
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Fig. 10 e Schematic of the HE effect for the samples tested under different conditions: (a) under TEL and (b) in solution. the high mobility of dislocations enhanced by H [49], resulting in the propagation of crack into the interior of the grains. With continuous SSRT, the pileup of dislocations can further promote the local anodic dissolution and hydrogen transfer and aggregation in front of SCC areas [27e29,49]. Thus, TGSCC extends towards higher KAM value in both conditions. To clearly demonstrate the effect of HE on the samples tested under two conditions, the overall schematic illustration is shown in Fig. 10. Generally, localized corrosion areas can assist SCC by providing film-free surfaces for hydrogen evolution and easy ingress of H [40,59,60]. For the samples tested under TEL, more film-free areas and localized areas were observed, related to a thinner protective film, in comparison to the samples tested in solution. Thus, hydrogen atoms preferentially penetrate and accumulate in the film-free areas and localized corrosion areas (Fig. 10a). Moreover, with the increase in the immersion time, the film-free areas suffered from localized corrosion, and the localized corrosion areas could further expand. Thus, a high H concentration gradient of hydrogen atoms is observed between the localized corrosion areas and the surrounding Mg matrix [18,34]. H is assumed to form brittle hydrides or produce decohesion in brittle fracture, and the low surface energy plane (i.e., {0001}, {10-10}, {10-11}, {31-40}, and {2-203}) provides preferential pathways for the diffusion of hydrogen in Mg alloys [33,38,61]. Thus, the evolved hydrogen atoms can easily penetrate the Mg matrix through those low surface energy planes underneath the localized corrosion areas. Then, the diffused hydrogen atoms will gradually conglomerate at these planes. With continuous SSRT, the gradually conglomerated hydrogen atoms decreased the atomic bonding at these cleavage facets (Fig. 7d) [35]. With the application of stress, these places where hydrogen atoms conglomerate will be preferentially act as the crack initiation sites (Fig. 10a). Thus, the cracking mechanism for the samples tested under TEL is the hydrogen enhanced de-cohesion (HEDE) mechanism (Fig. 7d). On the contrary, for the samples tested in solution, the localized corrosion was remarkably aggravated, and the concentration of hydrogen atoms in the localized corrosion areas was relatively higher (Fig. 10b). Similarly, the diffusion of hydrogen atoms preferentially occurred along low-surface energy planes (Fig. 10b). However, the localized microvoid coalescence observed was
more severe (Fig. 7f) than that observed under TEL, related to the hydrogen enhanced local plasticity (HELP) mechanism. Moreover, the stepped, faceted interlocking fracture surfaces observed in (Fig. 7f) were related to hydride formation, i.e., the delayed hydride cracking (DHC) mechanism. As a result, the HE mechanism for the samples tested in solution was HELP and DHC. Moreover, besides the oxygen signal, hydrogen signals for the cracks were detected, related to the facile diffusion of hydrogen in the magnesium alloys. Thus, assemblages in the vicinity of cracks were possibly absorbed hydrogen or hydrides other than hydroxide. Based on the above analysis, the present model for HE mechanism is different, mainly related to the role of hydrogen under different conditions.
Conclusions The SCC behavior of the ZK60 magnesium alloy was characterized by AES, EBSD, and ToF-SIMS measurement under various conditions for the first time to the best of our knowledge. The main conclusions of this study were listed as follows: (1) The ZK60 magnesium alloy was less susceptible to SCC under TEL compared to that in solution. The difference result in the more intensive anodic dissolution of samples under TEL, proved by electrochemical characterization, which contributed to lower hydrogen concentration than that of samples in solution. (2) The AES profiles and cross-sectional observations further indicated that the ZK60 magnesium alloy exhibits more severe localized corrosion and a lesser number of cracks in solution compared to those observed under TEL because of the stronger hydrogen evolution and thicker film formed on the matrix, respectively. (3) From EBSD maps and fractography, the cracking mode under both conditions is dominated by TGSCC features. The observation of SIMS maps further revealed that the TGSCC features under both conditions are related to different HE mechanisms.
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Acknowledgements This study is supported by the National Environmental Corrosion Platform (NECP), the National Basic Research Program of China (973 Program project, No. 2014CB643300), and the National Natural Science Foundation of China (No. 51471034).
references
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Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions L.F. Zhou a,b, Z.Y. Liu a,b,*, W. Wu a,b, X.G. Li a,b,c, C.W. Du a,b,c, B. Jiang d a
Corrosion and Protection Center, University of Science and Technology Beijing, Beijing 100083, PR China Key Laboratory for Corrosion and Protection (MOE), Beijing 100083, PR China c State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China d College of Materials Science and Engineering, Chongqing University, Chongqing 400030, China b
article info
abstract
Article history:
The stress corrosion cracking (SCC) behavior of ZK60 magnesium alloy was investigated
Received 30 June 2017
under different conditions, i.e. thin electrolyte layer (TEL) and solution, by slow strain rate
Received in revised form
tensile tests, electrochemical techniques, Auger electron spectroscopy, scanning electron
22 August 2017
microscopy coupled with electron backscattered diffraction, and time of flight secondary
Accepted 23 August 2017
ion mass spectrometry. Results indicated that the ZK60 magnesium alloy in solution ex-
Available online xxx
hibits a higher SCC susceptibility with a combined SCC mechanism of weaker anodic dissolution (AD) and stronger hydrogen embrittlement (HE) compared to under TEL.
Keywords: Magnesium alloys
Moreover, the HE mechanism under various conditions was discussed. © 2017 Published by Elsevier Ltd on behalf of Hydrogen Energy Publications LLC.
Thin electrolyte layer Stress corrosion cracking Anodic dissolution Hydrogen embrittlement
Introduction Among structural materials such as steel, titanium, and aluminum alloys, etc., magnesium alloys exhibit the lowest density (r z 1.74 g$cm3), making them attractive for automotive and aerospace industries because of the significance of weight [1e3]. In addition, magnesium alloys also possess several other advantages, such as excellent castability, electromagnetic interference shielding, non-magnetic property and recyclability. Despite these intriguing properties, their use has historically been restricted considerably because of their chemical reactivity, making them less resistant to corrosion
[4e9], corrosion fatigue [10,11], and stress corrosion cracking (SCC) [12e25]. SCC is typically related to mechanisms associated with anodic dissolution (AD) [26e29] or hydrogen embrittlement (HE) [30e38]. Several studies [17,18,24] have reported that HE is the typical mechanism of SCC in magnesium alloys. However, some others [13,14,25] have concluded that SCC in magnesium alloys is caused by the combination of AD and HE. Thus, the exact mechanism of SCC in magnesium alloys remains unclear. Moreover, several mechanisms may simultaneously occur depending on the material and other variables (e.g., strain rate and ambient conditions). However, the two above-mentioned
* Corresponding author. Corrosion and Protection Center, University of Science and Technology Beijing, Beijing 100083, PR China. E-mail address: [email protected] (Z.Y. Liu). http://dx.doi.org/10.1016/j.ijhydene.2017.08.161 0360-3199/© 2017 Published by Elsevier Ltd on behalf of Hydrogen Energy Publications LLC. Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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mechanisms that play a dominant role in the SCC of magnesium alloys are still controversial. Magnesium alloys most likely suffer from atmospheric corrosion, which is an electrochemical process occurring on a metal surface covered with a thin electrolyte layer (TEL) [39]. Among all corrosion modes, the SCC behavior of metals under TEL is significantly different from that in a bulk solution [40e42]. The corrosion of magnesium alloys is thought to involve the cathodic reaction of oxygen reduction, and this cathodic process is significantly enhanced for samples tested under TEL. However, oxygen reduction is not as important as hydrogen evolution during the corrosion of alloys in solution. Thus, the associated mechanisms under TEL are possibly related to the combined effect of AD and HE, while HE is believed to be the most common mechanism for the SCC of magnesium alloys in solution. To identify the dominant mechanisms related to the SCC of magnesium alloys, and to support the increased use of magnesium alloys for stressed components in aerospace and automobile industry in the marine atmosphere, the SCC mechanisms of magnesium alloys under TEL and in solution must be systematically investigated. Hence, with the limited knowledge about the SCC behavior of the ZK60 magnesium alloy, slow strain rate tensile (SSRT) tests are conducted using the ZK60 magnesium alloy under TEL and in solution. In addition, the surface appearance and fractography of all samples were observed by scanning electron microscopy (SEM); the surface film was examined using high-resolution Auger electron spectroscopy (AES); the cracking mode was characterized by electron backscattered diffraction (EBSD); and the distribution of elements in the vicinity of cracks was examined by time-of-flight secondary ion mass spectroscopy (ToF-SIMS). The results of this study may provide a comprehensive understanding on the dominant SCC mechanisms of magnesium alloys under different conditions.
Experimental Materials and environment Commercial ZK60 Magnesium alloy with a nominal composition of Mge6Zne0.5Zr (wt%) was used. Inductively coupled plasmaeatomic emission spectroscopy analysis of the asreceived alloy revealed a Zr content of 0.34 wt%. This value is close to the low limit for alloys with this denomination. Table 1 summarizes the chemical composition. The experiment device shown in Fig. 1 comprised a test medium, maintained at ambient temperature (approximately 25 C), of a modified simulated TEL atmosphere, which facilitated easy oxygen transfer. The solution in the atomizer was 0.1%NaCl þ 0.05%Na2SO4 þ 0.05%CaCl2 with a pH of
approximately 5.1 related to CO2 dissolution. By contrast, the test medium in the solution was the same as that in the atomizer. However, the transfer of oxygen in solution was weaker; hence, the HE effect is possibly stronger than that under TEL. Samples for stereoscopic microscope observations were mechanically polished to 1 mm using ethanol-based diamond suspensions followed by chemical etching using a solution of 5 g picric acid þ 5 g acetic acid þ 20 mL distilled water þ 100 mL anhydrous ethanol. Similarly, the samples for EBSD characterization were mechanically polished and electropolished using 10% perchloric acid in ethanol at 0 C and at a polishing voltage of 15 V for 30 s. The EBSD samples were observed in the vicinity of the cracks. Automatic EBSD scans were recorded using TSL data acquisition software with a step size of 1 mm. The EBSD data were analyzed using TSL OIM software.
SSRT tests SSRT tests were carried out to investigate the SCC behavior of ZK60 alloy under different conditions. A flat-plate tensile specimen was prepared based on GB/T 15970 [43]. Prior to each SCC test, the tensile samples were prepared from electrodischarge machining. The samples were successively ground using silicon carbide papers up to 2000 grit, ultrasonically cleaned in absolute ethanol for 10 min, and dried under a flow of cool air. SSRT tests were carried out on a WDML-30kN Materials Test System with a strain rate of 1 106 s1 (Fig. 1a). Before each test, the specimen was maintained under 100% RH for 1 h to ensure the formation of an electrolyte film on the surface. Moisture was continuously pumped during the entire experiment. After failure, one piece of the fracture was analyzed for characterization, and the other was applied for the measurement of the hydrogen concentration. First, the percentage of elongation and area reduction of each specimen were calculated after the removal of corrosion products. In addition, fracture morphologies were observed by SEM. Each test was reproduced at least three times to ensure the reliability of experimental data, and the standard deviation of the data was measured. Next, the cross-sectional morphologies of the fracture were observed to investigate the localized corrosion severity and cracking mode. Moreover, the hydrogen concentration of the tested samples was measured using a hydrogen analyzer (G4 Phoenix DH).
Electrochemical measurements Fig. 1b shows the measurement system comprising the typical three-electrode system. The working electrode is ZK60 specimen with a working area of 1 cm2, the counter electrode is a Pt bulk circling around the working electrode, and the
Table 1 e Chemical composition of the ZK60 magnesium alloy. Alloy
ZK60
Chemical composition (mass%) Al
Zn
Mn
Fe
Si
Cu
Ni
Zr
Mg
e
5.12
0.02
0.0024
0.003
0.0019
0.0011
0.34
Bal.
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Fig. 1 e Schematic of experimental devices for (a) SSRT and (b) electrochemical measurements.
reference electrode is a saturated Ag/AgCl electrode [44,45]. After a pumping time of 20 min for electrochemical impedance spectroscopy (EIS) measurement, the scanning frequency ranged from 100 kHz to 10 mHz, with a perturbing AC amplitude of 10 mV. The polarization curve was obtained on an electrochemical working station (PARSTAT2273, Princeton). Potentiodynamic polarization curves were recorded at a potential sweep rate of 0.5 mV/s, which ranged from 1750 mV to 1250 mV (vs. Ag/AgCl). For better reproducibility, polarization curves and EIS measurements were repeated more than three times. All tests were conducted at ambient temperature.
AES and TOF-SIMS measurements AES experiments were conducted on a UHV scanning Auger nanoprobe (PHI 700, ULVAC-PHI). The Ar GCIB-sputtered sample was rapidly transferred into the AES chamber using a vacuum transfer vessel capable of maintaining a low vacuum level and minimizing sample contamination during transfer. AES measurements were carried out at a beam voltage and current of 5 kV and 10 nA, respectively, with a SiO2 sputter rate of 17 nm/min. TOF-SIMS was employed for local chemical analysis around the cracks using an ION-TOF GmbH TOF-SIMS 5-100 analyzer. A highly focused primary Csþ ion
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beam in a rectangle of 50 mm was scanned across the cracks. Secondary ions, resulting from abrasion, were analyzed at each point to reconstruct a micrograph. Different signals were collected during the analysis. The high-resolution mode was utilized to avoid interferences between the different signals.
Table 2 e Hydrogen concentration of ZK60 samples after SSRT under different conditions. Test conditions
Air
TEL
Solution
CH(ppm)
1.31
12.57
17.89
300
Initial microstructure
250
Fig. 2 shows the initial microstructure and texture of ZK60 magnesium alloy. The material exhibited a microstructure with a grain size of 5e200 mm (Fig. 2a), which was also observed in the EBSD spectrum in Fig. 2b. Moreover, it was obviously seen that a thimbleful of twin boundaries subsist in the grain.
200
Hydrogen concentration and SCC behavior of the ZK60 magnesium alloy
Stress(MPa)
Results
150
100
0
After SSRT failure, the concentrations of hydrogen in the samples were measured (Table 2). The concentrations of hydrogen in the samples sharply increased in a sequence of in air (1.31 ppm), under TEL (12.57 ppm), and in solution (17.89 ppm). Fig. 3 shows the stressestrain curves of ZK60 tested under different conditions. Table 3 summarizes the change and decrease in terms of strength and plasticity, respectively, for the comparison of their mechanical properties. The ZK60 magnesium alloys exhibited different values of s0.2 and ultimate tensile strength (UTS) under various environments (Table 3). High values of s0.2 and UTS were observed for samples tested in air, while lower values were observed for samples tested under TEL and in solution. The concentrations of hydrogen in the samples tested under TEL and in solution were different because of different experimental environments, thereby leading to different mechanical properties of the ZK60 magnesium alloys. Compared to samples exposed to air, those exposed to a corrosive medium exhibited a considerable reduction in the s0.2 and UTS values. Samples tested in air, under TEL, and in solution exhibited a considerable difference in their tensile properties. The plasticity of samples tested under TEL was slightly lower than that of the samples tested in solution (Table 3). Moreover, the
in air under TEL in solution
50
0
5
10 15 Strain(%)
20
25
Fig. 3 e Stress versus strain curves for the ZK60 magnesium alloy at a strain rate of 1 £ 10¡6 s¡1 under different conditions.
Table 3 e Summary of the mechanical properties of the ZK60 magnesium alloy samples tested under various conditions. Alloy Conditions s0.2 (MPa) UTS (MPa) d(%) j(%) Decrease of s0.2 (%) Decrease of UTS (%) Decrease of d (%) Decrease of j (%)
ZK60 Air
TEL
Solution
240.6 ± 7.3 267.2 ± 7.2 21.4 ± 1.5 20.7 ± 1.6 e e e e
192.3 ± 5.4 205.7 ± 5.6 8.0 ± 0.6 8.3 ± 0.6 20.1 23.0 62.6 60
e 179.0 ± 4.3 3.7 ± 0.3 5.2 ± 0.4 e 33.0 82.7 94.7
tensile properties of the samples tested under TEL and in solution significantly decreased in comparison with those in air. Moreover, the values for the elongation-to-failure (d) and area reduction (j) tested in solution were lower than that of the
Fig. 2 e Initial microstructure of ZK60 magnesium alloy: (a) metallograph, (b) inverse pole figure. Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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samples tested under TEL. In addition, the following two equations, Id ¼ ð1 dmedium Þ=dair 100%
(1)
Ij ¼ ð1 jmedium Þ=jair 100%
(2)
were applied to calculate the loss of d and j. Herein, dair and jair represent the d and j values in air, dmedium and jmedium represent the d and j values under TEL and in solution, respectively. To quantify the SCC susceptibility of the samples tested under different conditions, the SCC susceptibility index (Id and Ij) was calculated by the decline in elongation and area reduction. Id and Ij are high for a system with high SCC susceptibility, whereas Id and Ij tend to be close to 0 for low susceptibility. The analysis above indicates that to a certain extent, the SCC susceptibility of samples tested in solution, is greater than that of samples tested under TEL.
Electrochemical characterization Fig. 4 shows the potentiodynamic polarization curves and EIS curves of the ZK60 magnesium alloy under different conditions. The corrosion potentials (Ecorr) and corrosion current
-1.1
(a)
-1.2
E(VAg/AgCl)
-1.3
-1.5 -1.6 -1.7 -1.8 -1.9 1E-8
2500
1E-7
-Zim(Ω.cm2)
1E-6 1E-5 i (A/cm2)
1E-4
1E-3
(b) under TEL in solution
2000
CPE R
1500
CPE + R R +
L
1000 500 0 0
500
densities (Icorr) under the two conditions were obtained (Fig. 4a) by the Tafel fitting method [46]. Table 4 shows the obtained results. From the thermodynamic perspective, the corrosion potential of ZK60 in solution shifted to the noble direction compared to the samples tested under TEL, possibly related to the formation of a thicker film formed on the matrix; this point is discussed later with the AES results. However, in terms of dynamic perspective, the anodic Tafel slope (ba) and cathodic Tafel slope (bc) of the samples tested under TEL were slightly greater than those of the samples in solution, indicative of various electrochemical reactions. Moreover, the cathodic process of magnesium alloys tested under TEL was dominated by hydrogen evolution [40], and the presence of oxygen inhibited hydrogen evolution. Three loops were observed in the EIS curvesdone high-frequency capacitive loop, followed an intermediate frequency capacitive loop, and a final low-frequency inductive loopdrespectively. The equivalent circuit in Fig. 4b was used for EIS modeling, where Rs represents the solution resistance, (CPE)dl represents the double layer capacitance, Rt represents the charge-transfer þ resistance, (CPE)þ Mg represents the capacitance of the Mg reaction on the breaking area of the partial protective film [47], þ Rþ Mg represents the resistance of the Mg reaction, and L represents the inductance of the partial protective film [47]. Moreover, the EIS results revealed that the impedance of the samples tested under TEL slightly decreases. This result is in agreement with the low corrosion current density value mentioned above.
Cross-sectional and secondary cracking observations
under TEL in solution
-1.4
5
1000
1500
Zre(Ω.cm2)
2000
2500
Fig. 4 e (a) Potentiodynamic polarization curves and (b) electrochemical impedance spectroscopy curves of the ZK60 magnesium alloy under different conditions.
To investigate the cracking mechanism, the cross-sectional and secondary cracks of the samples subjected to tensile tests under different conditions were observed (Fig. 5). Samples tested under TEL exhibited a slight localized corrosion on the surface (Fig. 5a). Meanwhile, several secondary cracks initiated at the localized corrosion areas (Fig. 5a). On the contrary, the corrosion attack of samples tested in solution suffered relatively greater attack with deep corrosion ditches on the surface (Fig. 5b). Moreover, the density of secondary cracks of the samples tested in solution was considerably less than those of the samples tested under TEL (Fig. 5b). Based on their cross-sectional morphologies, an interesting phenomenon was observed for the samples tested under TEL: the severity of localized corrosion areas was slight; however, a number of secondary cracks distributed at or nearby the localized corrosion areas were easily observed (Fig. 5c). For the samples in solution, the localized corrosion areas were relatively more severe; however, fewer secondary cracks were observed (Fig. 5d). To further disclose the cracking mode and initiation location of cracks, EBSD analyses of the two sample surfaces were carried out (Fig. 6). Fig. 6a and c shows the image qualityinverse pole figure (IQ-IPF) maps of two samples. The colors of grains in the maps corresponded to the crystallographic axes of grains in the inverse pole figure (Fig. 2b). Cracks were observed under two conditions, with propagation in the transgranular direction. To confirm the stress concentration in the grains at the tips of cracks, the kernel average misorientation (KAM) values were measured by SEM-EBSD (Fig. 6b
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Table 4 e Electrochemical parameters fitted from the potentiodynamic polarization curves of the ZK60 magnesium alloy under different conditions. Conditions TEL Solution
Ecorr (VAg/AgCl) 1.605 1.547
Icorr (A/cm2) 6
4.724 10 9.652 106
and d). A high KAM value typically reflects the distribution of the geometry necessary dislocation (GND) related to the localized deformation [48] that contributes to the film rupture. No considerable difference in the KAM values between the cracking paths and their surrounding area was observed, and
ba (mVAg/AgCl/decade)
bc (mVAg/AgCl/decade)
311.5 156.2
389.9 181.7
both cracks propagated along the high KAM value region (as highlighted by the arrows). As the pileup of dislocations promoted the local anodic dissolution and hydrogen transfer and aggregation in front of SCC areas [49], EBSD results revealed that the SCC of the ZK60 magnesium alloy is possibly
Fig. 5 e Microstructural analysis of the failed samples tested under different conditions: (a,b) secondary electron images of the sample cross-sectional surfaces under TEL and in solution, respectively; (c,d) secondary electron images of the secondary cracks of samples under TEL and in solution, respectively; (e,f) are high-magnification images of the rectangle area of (c,d), respectively. Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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Fig. 6 e EBSD IQ-IPF maps (a,c) and IP-KAM maps (b,d) recorded near the cracks in Fig. 5e and f under TEL (a,b) and in solution (c, d).
controlled by both AD and HE which we would further explain it in discussion.
Fractography Fig. 7 shows the SEM morphologies of the fracture surface under different conditions. Marginal necking was observed for all samples (Fig. 7a, c, and e). The fracture surface of ZK60 tested in air showed features consistent with quasi-cleavage and plastic dimples (Fig. 7b). For the samples tested under TEL, the primary part of the fracture surface was smooth, which showed typical transgranular stress corrosion cracking features (Fig. 7d) [15]. On one hand, the fracture surface showed considerable dissolution, microvoids, and minor cracks, indicating that vapor easily penetrates in the matrix, leading to a complex phenomenon. On the other hand, the transgranular SCC (TGSCC) feature was discontinuous, involving alternating mechanical and electrochemical processes [13]. Meanwhile, for the samples tested in solution, the TGSCC in solution propagated in a discontinuous manner with parallel facets; similarly, evidence of dissolution, microvoids, and minor cracks was observed (Fig. 7f). However, the severity of localized corrosion and the number of minor cracks and microvoids were greater than those of the samples tested under TEL. Transgranular features characterized by flat parallel facets (Fig. 7d and f) indicated the growth of discontinuous cracks and
the possible role of hydrogen in cracking. Padekar et al. [13] have mostly observed TGSCC for EV31A and AZ91E in distilled water. They have proposed that this feature is related to the possible role of hydrogen in cracking. Chakrapani and Pugh [31,32] have inferred that TGSCC occurs by a brittle cleavage mechanism involving H, resulting in stepped, faceted interlocking fracture surfaces similar to the features observed herein.
Local chemical analysis AES compositional depth profiling (Fig. 8) was carried out to examine the films formed on the surfaces of the failed tensile samples tested under different conditions. Analyses were performed on the uncorroded area distributed throughout the surface. In all cases, we reached to the alloy substrate, which exhibited a weak O signal and strong Mg signal (Fig. 8). C caused by surface contamination characterized in the AES spectra was not included in the AES compositional depth profiles. The surface films formed on the tensile samples tested under TEL (~24 nm) and in solution (~55 nm) were thicker than that in air (~19 nm). Interestingly, compared with the surface films for the tensile samples tested under TEL, that for the samples tested in solution were thicker. Irrespective of the test conditions, the top part of the film comprised hydromagnesite [6,7]. The O/Mg ratio gradually decreased in the deeper part of the film, indicating the presence of Mg(OH)2 and MgO. Notably, the lack of straightforward evidence for Mg(OH)2 and MgO in the film in Fig. 8 does not rule out the
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Fig. 7 e Fractography of samples under different conditions: (a,b) in air; (c,d) under TEL; (e,f) in solution. presence of Mg(OH)2 and MgO layer at the film/metal interface as reported by several studies [50,51]. For evaluating the local chemical composition along and near the cracks and to detect possibly absorbed hydrogen or hydrides, ToF-SIMS ion images of H and O around the cracks of the ZK60 samples developed during SSRT under various conditions were recorded. Fig. 9 shows the SIMS maps of the hydrogen (Fig. 9a and c) and oxygen signals (Fig. 9b and d) respectively, for two cracks grown in a ZK60 sample subjected to SSRT under different conditions. The hydrogen and oxygen signals were considerably different under both conditions. Moreover, most of the hydrogen signals in the cracks can be assumed to correspond to the adsorbed hydrogen or hydrides, i.e., MgH2. Meanwhile, oxygen signals were scarcely detected in the cracks. For the crack developed under TEL (Fig. 9a and b), several areas showed high, superimposed signals for hydrogen and oxygen,
related to the formation of Mg(OH)2. However, contrary to the sample tested under TEL, other hydrogen signals were observed in addition to the hydrogen signals observed for the cracks for the sample tested in solution (Fig. 9c and d). Moreover, more hydrogen and oxygen signals in highly superimposed areas were observed compared to the sample tested under TEL, which is also related to the generated Mg(OH)2.
Discussion Effect of the environment on the SCC susceptibility and mechanical properties As mentioned above, the ZK60 magnesium alloy exhibited higher SCC susceptibility in bulk solution than that observed for under TEL, indicating that oxygen reduction is not as
Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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100
Concentration(at.%)
80
60
Mg
O
air TEL solution
40
20
0
0
20
40
60 Depth(nm)
80
100
120
Fig. 8 e AES depth profiles (area ¼ 2 £ 2 mm2) of the uncorroded area of ZK60 magnesium alloy surface tested under different conditions.
important as hydrogen evolution during SCC. The conditions for the sample in solution were clearly more favorable for hydrogen evolution than those for the sample tested under TEL, which is clearly more beneficial for oxygen reduction. The environment is thought to considerably affect SCC
9
susceptibility. Based on Equations (1) and (2), the calculated Id and Ij values for the samples tested under TEL were 0.63 and 0.6, suggesting the samples tested under TEL were moderately susceptible to SCC. However, the calculated Id and Ij for the samples tested in solution were 0.83 and 0.95, suggesting that the SCC susceptibility is greater in solution than that under TEL. Moreover, the mechanical properties were also considerably affected by the environment. The s0.2 and UTS values were considerably less than those observed in air for all samples exposed to corrosive medium, possibly related to the HE [18,33,52]. In addition, the higher SCC susceptibility and the decrease in the mechanical properties of samples tested in solution are related to the weaker oxygen reduction and stronger hydrogen evolution compared to the samples tested under TEL.
Relationship between the electrochemical properties and SCC behavior The conventional electrochemical measurement result indicated that the Ecorr of the samples tested under TEL was relatively more negative compared to that of the samples in solution (Fig. 4a). The Tafel fitting and EIS results indicated that the samples tested in solution exhibited a considerably higher Icorr and slightly lower impedance than those of
Fig. 9 e ToF-SIMS ion images of (a, c) H and (b, d) O recorded on the cracks of the ZK60 samples after SSRT under various conditions, (a, b) under TEL and (c, d) in solution. Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161
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samples tested under TEL, respectively (Fig. 4). The lower Ecorr of the samples tested under TEL is related to the thinner film formed on the surface of the matrix; thus, the corrosion under TEL is prone to occur. However, lower Icorr and higher impedance of the samples tested under TEL are related to two main reasons. On the one hand, oxygen reduction is not as important as hydrogen evolution during corrosion; on the other hand, the formation or deposition of Mg(OH)2 is easier for samples under TEL, which may deter corrosion. These results are in agreement with those obtained from AES and secondary cracking observations. Hence, on first blush, the AD of samples occurred, resulting in localized corrosion and stress concentration in this area, which could easily become SCC nucleation sites. These sites in turn can serve as pathways for hydrogen to diffuse into the matrix [33,38].
AD role in SCC mechanism According to several books and studies [51e55], possible reactions involving anodic and cathodic reactions are observed between the samples under TEL and in solution; in this study, these reactions are listed as follows: Anodic reaction Mg/Mgþ þ e
(3)
Mgþ /Mg2þ þ e
(4)
2Mgþ /Mg þ Mg2þ /Mg$Mg2þ
(5)
Mg$Mg2þ þ 2Hþ /2Mg2þ þ H2
(6)
Cathodic reaction: TEL : 2H2 O þ 2e /H2 þ 2OH
(7)
O2 þ 2H2 O þ 4e /4OH
(8)
Solution : 2Hþ þ 2e /H2
(9)
Effect of HE on the SCC mechanism
Salt formation: Mg2þ þ 2OH /MgðOHÞ2
diffusion in solution adjacent to the surface is relatively fast. Hence, the presence of the film accompanying a rapid anodic dissolution process in solution significantly accelerates the dissolution kinetics of the film, leading to a considerably high corrosion rate of ZK60. Although the dissolved Mg2þ can react with OH and form Mg(OH)2 deposits on the surface, the loosely deposited Mg(OH)2 may deposit onto the uncorroded area other than the broken areas. The evolved hydrogen from the corroding areas can stir the deposited Mg(OH)2 and prevent the corroding areas from being fully covered by the deposited Mg(OH)2. This observation is in agreement with the AES profiles. The transgranular stress corrosion cracking is related to the electrochemical corrosion or mechanical rupture of protective films, permitting the evolved hydrogen to enter the alloy matrix and eventually causing embrittlement and premature fracture [58]. Therefore, the nature and stability of the oxide or hydroxide films that develop on magnesium alloys is an important parameter for their SCC. The samples tested under TEL and in solution distinctly formed different oxide or hydroxide scales. Moreover, the samples tested in solution developed a more robust and sturdy layer; hence, the films are less susceptible to mechanical rupture, and the electrolyte can only infiltrate into the matrix to create a less number of cracks compared to the samples under TEL (Fig. 5c and d). However, once the films were undermined by the applied strain rate and chloride ions, the detrimental ions can easily penetrate into the matrix to cause localized corrosion [36,52]. Therefore, the AD mechanisms may be greater, and the HE mechanisms may be weaker for samples tested under TEL compared to the samples tested in solution. Moreover, the severity of localized disruption also increased the possibility for hydrogen to enter the Mg matrix, which explains the increase of SCC susceptibility in solution.
(10)
Based on the reactions above, the bulk solution is acidic with a pH of approximately 5.1; however, it is proverbial that the corrosion of Mg and its alloys results in the increase of solution pH, related to the surface alkalization effect [56,57]. Herein, the acidity of the electrolyte on the specimen surface can increase to attain a high pH value, and the corrosion continues to proceed at this higher pH value. At this time, the deposition of Mg(OH)2 (reaction (10)) dominates at this pH level. Whereas for the samples tested under TEL, only a small number of aqueous drops or a thin electrolyte film remains on the surface of the tested samples, while the samples tested in solution are all immersed. Hence, the corrosion of the sample under TEL (Fig. 5a) is slower than that of the sample in solution (Fig. 5b) because of the strong surface alkalization effect. In addition, the hydrolysis of Mg2þ in solution will momentarily decrease the pH of the surface, contributing to the aggravation of the reaction (9). In nature, a dynamic process of OH
To further examine the effect of HE on the SCC behavior of ZK60, the hydrogen and oxygen signals in the vicinity of cracks were examined (Fig. 9). The “H” signal was completely detected in the cracks, but no “O” signal was detected. The surface film on Mg is well known to not be thick or compact, thereby offering only a limited level of protection. In the early stage of SSRT, because of the pseudo-protection of the film and mechanical rupture, an electrochemical cell occurs between active area of anode and cathode which is capable of preventing film repair and sustaining enhanced anodic dissolution. Moreover, hydrogen can easily diffuse into the bulk of the material and be absorbed at the crack tips or to form hydrides. H diffusion has been proposed to be critical for several SCC mechanisms for magnesium alloys [36,37]. Moreover, the transgranular cracks intruded into the grains and a high density of dislocations was formed near the intruded cracks because the KAM value was proportional to the microstrain caused by crystal defects, such as dislocations (Fig. 6b and d). The grains exhibited a higher KAM compared to the grain boundaries possibly because of the difference in the mobility of dislocations between the grains and grain boundaries. Namely, the grains with a high initial density of dislocations could possibly undergo local plasticity because of
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Fig. 10 e Schematic of the HE effect for the samples tested under different conditions: (a) under TEL and (b) in solution. the high mobility of dislocations enhanced by H [49], resulting in the propagation of crack into the interior of the grains. With continuous SSRT, the pileup of dislocations can further promote the local anodic dissolution and hydrogen transfer and aggregation in front of SCC areas [27e29,49]. Thus, TGSCC extends towards higher KAM value in both conditions. To clearly demonstrate the effect of HE on the samples tested under two conditions, the overall schematic illustration is shown in Fig. 10. Generally, localized corrosion areas can assist SCC by providing film-free surfaces for hydrogen evolution and easy ingress of H [40,59,60]. For the samples tested under TEL, more film-free areas and localized areas were observed, related to a thinner protective film, in comparison to the samples tested in solution. Thus, hydrogen atoms preferentially penetrate and accumulate in the film-free areas and localized corrosion areas (Fig. 10a). Moreover, with the increase in the immersion time, the film-free areas suffered from localized corrosion, and the localized corrosion areas could further expand. Thus, a high H concentration gradient of hydrogen atoms is observed between the localized corrosion areas and the surrounding Mg matrix [18,34]. H is assumed to form brittle hydrides or produce decohesion in brittle fracture, and the low surface energy plane (i.e., {0001}, {10-10}, {10-11}, {31-40}, and {2-203}) provides preferential pathways for the diffusion of hydrogen in Mg alloys [33,38,61]. Thus, the evolved hydrogen atoms can easily penetrate the Mg matrix through those low surface energy planes underneath the localized corrosion areas. Then, the diffused hydrogen atoms will gradually conglomerate at these planes. With continuous SSRT, the gradually conglomerated hydrogen atoms decreased the atomic bonding at these cleavage facets (Fig. 7d) [35]. With the application of stress, these places where hydrogen atoms conglomerate will be preferentially act as the crack initiation sites (Fig. 10a). Thus, the cracking mechanism for the samples tested under TEL is the hydrogen enhanced de-cohesion (HEDE) mechanism (Fig. 7d). On the contrary, for the samples tested in solution, the localized corrosion was remarkably aggravated, and the concentration of hydrogen atoms in the localized corrosion areas was relatively higher (Fig. 10b). Similarly, the diffusion of hydrogen atoms preferentially occurred along low-surface energy planes (Fig. 10b). However, the localized microvoid coalescence observed was
more severe (Fig. 7f) than that observed under TEL, related to the hydrogen enhanced local plasticity (HELP) mechanism. Moreover, the stepped, faceted interlocking fracture surfaces observed in (Fig. 7f) were related to hydride formation, i.e., the delayed hydride cracking (DHC) mechanism. As a result, the HE mechanism for the samples tested in solution was HELP and DHC. Moreover, besides the oxygen signal, hydrogen signals for the cracks were detected, related to the facile diffusion of hydrogen in the magnesium alloys. Thus, assemblages in the vicinity of cracks were possibly absorbed hydrogen or hydrides other than hydroxide. Based on the above analysis, the present model for HE mechanism is different, mainly related to the role of hydrogen under different conditions.
Conclusions The SCC behavior of the ZK60 magnesium alloy was characterized by AES, EBSD, and ToF-SIMS measurement under various conditions for the first time to the best of our knowledge. The main conclusions of this study were listed as follows: (1) The ZK60 magnesium alloy was less susceptible to SCC under TEL compared to that in solution. The difference result in the more intensive anodic dissolution of samples under TEL, proved by electrochemical characterization, which contributed to lower hydrogen concentration than that of samples in solution. (2) The AES profiles and cross-sectional observations further indicated that the ZK60 magnesium alloy exhibits more severe localized corrosion and a lesser number of cracks in solution compared to those observed under TEL because of the stronger hydrogen evolution and thicker film formed on the matrix, respectively. (3) From EBSD maps and fractography, the cracking mode under both conditions is dominated by TGSCC features. The observation of SIMS maps further revealed that the TGSCC features under both conditions are related to different HE mechanisms.
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Acknowledgements This study is supported by the National Environmental Corrosion Platform (NECP), the National Basic Research Program of China (973 Program project, No. 2014CB643300), and the National Natural Science Foundation of China (No. 51471034).
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Please cite this article in press as: Zhou LF, et al., Stress corrosion cracking behavior of ZK60 magnesium alloy under different conditions, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/j.ijhydene.2017.08.161