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Stress corrosion cracking behavior of the wrought magnesium alloy AZ31 under controlled cathodic potentials
Materials Science and Engineering A 531 (2012) 171–177
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Stress corrosion cracking behavior of the wrought magnesium alloy AZ31 under controlled cathodic potentials Yoshihiko Uematsu a,∗ , Toshifumi Kakiuchi a , Masaki Nakajima b a b
Department of Mechanical and Systems Engineering, Gifu University, 1-1 Yanagido, Gifu 501-1193, Japan Toyota National College of Technology, 2-1 Eisei-cho, Toyota, Aichi 471-8525, Japan
a r t i c l e
i n f o
Article history: Received 23 May 2011 Received in revised form 14 October 2011 Accepted 18 October 2011 Available online 24 October 2011 Keywords: Magnesium alloy Stress corrosion cracking Hydrogen embrittlement Anodic dissolution Crack propagation rate Stress intensity factor
a b s t r a c t Stress corrosion cracking (SCC) tests were performed using compact tension (CT) specimens in a NaCl solution under hydrogen-charged conditions to investigate the SCC behavior of the wrought magnesium alloy AZ31. The effect of NaCl concentration was estimated by conducting the SCC tests under a constant cathodic potential of −1.4 V, with NaCl concentrations of 0.5, 3.0 and 8.0%. The crack propagation rate (da/dt) became accelerated with increasing concentration due to the enhanced anodic dissolution by chloride ions. Subsequently, SCC tests were performed using a 3%NaCl solution with cathodic potentials at −1.4, −2.5 and −3.0 V to investigate the effect of cathodic potentials on SCC behavior. According to the Pourbaix diagram, the cathodic potential of −3.0 V corresponds to the immunity region, whereas −1.4 V is within the corrosion region and −2.5 V is the boundary between the two regions. In the immunity region, the SCC due to hydrogen embrittlement occurred when the KIscc value was higher than that of the SCC dominated by anodic dissolution. The KIscc value decreased and the crack propagation rate increased with decreasing cathodic potential in the immunity region at which SCC occurred due to hydrogen embrittlement. At the cathodic potential of −3.0 V, the da/dt was insensitive to the stress intensity factor because highly charged hydrogen gave rise to brittle fracture. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Magnesium (Mg) alloys have excellent properties, such as low weight, high specific strength and stiffness, machinability and recyclability. Because of these advantages, Mg alloys are very attractive structural materials that can be used in a wide variety of applications; in particular, they can be utilized as components for air planes or transportation vehicles in which saving weight is extremely important. The major drawback of this alloy is poor corrosion resistance compared with other light metals, such as aluminum and titanium alloys. Hence, it is very important to understand the corrosion properties of Mg alloys to be able to use these alloys as structural components, and thus, many studies have been performed on corrosion [1–4], corrosion fatigue [5–7] and stress corrosion cracking (SCC) [8–25]. Song and Atrens [3,4] reviewed the corrosion behavior of Mg alloys in detail and proposed a partially protective surface film model, a mono-valent magnesium ion model, a particle undermining model and a magnesium hydride model for the hydrogen evolution reaction (HER) during corrosion. All models suggest that hydrogen generation is strongly related
∗ Corresponding author. Tel.: +81 58 293 2501; fax: +81 58 293 2491. E-mail address: [email protected] (Y. Uematsu). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.10.052
to the corrosion of Mg alloys. From an engineering viewpoint, the SCC behavior of Mg alloy has been widely investigated because SCC causes the unexpected sudden failure of structures. In addition, hydrogen embrittlement plays an important role due to the presence of HER during corrosion. Winzer et al. [8] surveyed the SCC behavior of Mg alloys, and demonstrated the effect of alloy composition, the manufacturing process and heat treatment on the SCC susceptibility of materials. They also postulated that hydrogen should be involved in the mechanism of SCC because Mg had a negative difference effect (NDE) [3,4,8]. Thus, it is important to evaluate the effect of hydrogen on SCC behavior. Most of the investigations of SCC in Mg alloys are based on slow strain rate (SSR) testing in which the SCC susceptibility of materials could be easily evaluated. For example, Kannan et al. [15] reported that rare-earth elements could improve SCC resistance relative to the conventional structural Mg alloy AZ80. Ben-Hamu et al. [18] showed the beneficial effect of Si incorporation on the SCC resistance. From the viewpoint of designing mechanical components, quantitative analysis of crack propagation rate is inevitable. Marrow et al. [12] investigated the crack growth path using fourpoint bend specimens with small holes as crack initiators, and Winzer et al. [17,20] attempted to characterize the SCC susceptibility by threshold stress, which was determined using a DC potential drop method during the SSR test. However, to estimate
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Fig. 1. Microstructure of material.
Fig. 2. Compact tension (CT) specimen configuration.
2.2. Stress corrosion cracking tests crack propagation rates based on fracture mechanics parameters, i.e., stress intensity factor (K), the SCC tests should be performed using specimens with long cracks, such as compact tension (CT) and center-cracked panel specimens. Furthermore, the SCC could be attributed to the combined effects of anodic dissolution and hydrogen embrittlement at the crack tip. As mentioned above, hydrogen embrittlement plays an important role in the SCC of Mg alloys [8–11]. For example, Song et al. [13] demonstrated that the preexposure of Mg alloy to a NaCl solution tends to decrease tensile strength and elongation due to hydrogen embrittlement. However, it is difficult to investigate the effects of anodic dissolution and hydrogen embrittlement independently. Kannan et al. [14] conducted SSR tests under a continuously hydrogen-charged condition to enhance the effect of hydrogen on SCC behavior and found a significant decrease in tensile strength, whereas the crack propagation rate was not evaluated. This study focuses on the SCC behavior of the wrought magnesium alloy AZ31 and provides a qualitative estimation of the crack propagation rate based on fracture mechanics. The SCC tests were performed using CT specimens with long cracks that were placed in a NaCl solution under hydrogen-charged conditions. By controlling both the NaCl concentration and the cathodic potential, the effects of anodic dissolution and hydrogen embrittlement on SCC behavior were decoupled in this study.
Prior to the SCC tests, a pre-crack with a length of 2 mm was introduced from an artificial notch root of the CT specimen by a fatigue test at a frequency of 10 Hz, a stress ratio of 0.1 and an initial K value of 3 MPam1/2 . The SCC test specimens with fatigue pre-crack were attached to a creep testing machine equipped with a NaCl solution tank, as shown in Fig. 3. The crack length was monitored by a crack gauge (KV-25B: KYOWA) bonded to the segment of the CT specimen where the crack might grow. The Ag/AgCl and Pt electrodes were used as reference and counter electrodes, respectively. A cathodic potential of −1.4 V was applied using a potentiostat, and the SCC tests were performed in a NaCl solution with concentrations of 0.5%, 3% and 8%, to investigate the SCC behavior that is dominated by the anodic dissolution at the crack tip. Notably, the cathodic potential of −1.4 V corresponds to the corrosion region according to the Pourbaix diagram [26], as shown in Fig. 4. Subsequently, the SCC tests were conducted using a 3% NaCl solution at the cathodic potentials of −2.5 and −3.0 V to investigate the SCC behavior that is controlled by hydrogen embrittlement. The cathodic potential of −3.0 V is in the immunity region, while the potential at −2.5 V is at the boundary between the corrosion and immunity regions based on the Pourbaix diagram. It should be noted that the potential is established on the surface of the specimen and not at the crack tip where the SCC occurs. Thus, it is more important to understand the actual potential at the crack tip for a more precise prediction of the crack propagation rates, although the measurement of the potential at the crack tip is quite difficult.
2. Experimental details 2.1. Material and specimen configuration The material used in this study is the wrought magnesium alloy AZ31. The microstructure of AZ31 is shown in Fig. 1. The chemical composition (mass%) is as follows: Al: 2.7, Zn: 0.79, Mn: 0.44, Fe: 0.0012, Ni: 0.0009, Cu: 0.0011, Si: 0.004, Ca: 0.001, Sn < 0.001, and Mg: bal. The mechanical properties of this material are summarized in Table 1. CT specimens with a width of 50.8 mm and a thickness of 6 mm were used for the SCC tests, as shown in Fig. 2, using a CT specimen configuration that complies with the ASTM E647 standard. The specimens were machined from the as-received AZ31 plates with a thickness of 6 mm in the L–T orientation, where L is the longitudinal direction, or the extrusion direction, and T is the transverse direction. The surface of the specimen was not modified, with an average roughness of 3.2 m.
Fig. 3. SCC test configuration.
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Table 1 Mechanical properties of material. 0.2% proof stress 0.2 (MPa)
Tensile strength B (MPa)
Elongation ı (%)
Vickers hardness HV
Elastic modulus E (GPa)
Grain size d (m)
170
248
17
54
46
29
0.03
Crack length a (m)
The SCC test specimens were maintained in the NaCl solution at a given cathodic potential for 24 h with no load to charge the hydrogen. The initial value of K was approximately 7 MPam1/2 , and the load was gradually increased every 24 h until the crack began to propagate. Consequently, it took 3–7 days before the SCC began, depending on the testing environments. Thus, the precharging time was assumed to be long enough for the hydrogen to diffuse to the crack tip, where the triaxial stress is dominant. Once the crack began to propagate, the load was sustained, and the crack propagation was monitored by a crack gauge. The sustained load during crack propagation leads to an increasing K value in the SCC test. Therefore, the threshold value of KIscc is defined as the K value calculated from the load step just before the crack propagation occurs. In this case, the increment in the K value with the increasing load was approximately 0.4–0.5 MPam1/2 for each step.
Magnesium Alloy AZ31 Cathodic potential −1.4V
0.02
0.01
0
0
20000 Time (s)
3. Results 3.1. Effect of the NaCl concentration
as at 8%, resulted in much thicker corrosion products. The corrosion products seem to be thicker at lower K values due to the lower crack growth rate, although it is difficult to distinguish between the corrosion products formed during the SCC or afterwards. The EDX analysis of the corrosion products on the fractured surface of the 3.0% NaCl solution was performed. The data demonstrate that the corrosion products are Mg(OH)2 because the quantitative analysis revealed an atomic ratio of oxygen to magnesium of 2:1. Some part of the fractured surfaces were not completely covered by corrosion
10
10
a 2
Mg2+ Corrosion
E(V)
Mg(OH)2 Passivation Test conditions
-2
Crack growth rate da/dt (m/s)
0 -2 -4 -6
-1
10
10
10
-6
Mg Immunity 0
−3
−4
Magnesium alloy AZ31 Cathodic potential −1.4V 0.5%NaCl 3.0%NaCl 8.0%NaCl
−5
−6
−7
0
3
-3
40000
Fig. 5. Relationship between crack length and testing time at a cathodic potential of −1.4 V.
The change in the crack length as a function of testing time is illustrated in Fig. 5. It should be noted that the crack propagated in a stable manner under the sustained loading condition regardless of the NaCl concentration. Fig. 6 shows the relationship between the crack propagation rate (da/dt) and the stress intensity factor (K), where da/dt was calculated from Fig. 5. The critical values of K for crack propagation, i.e., the KIscc , were measured to be 7.0, 8.6 and 7.1 MPam1/2 for the NaCl concentrations of 0.5, 3 and 8%, respectively, which indicates that the critical values are insensitive to the NaCl concentration. However, the crack propagation rate increases with increasing NaCl concentration. Fig. 7 reveals the SEM micrographs of typical fractured surfaces at different K levels in 0.5%, 3.0% and 8.0% NaCl solutions, respectively. Fractured surfaces are covered by corrosion products that are irrelevant to the K values, where the high NaCl concentration, such
0
0.5%NaCl 3.0%NaCl 8.0%NaCl
1 10
−8
3
7
14
PH Fig. 4. A typical Pourbaix diagram for magnesium–water system at 25 ◦ C.
5
8
10
20
30
1/2
Stress intensity factor K (MPam ) Fig. 6. Relationship between crack propagation rate, da/dt, and stress intensity factor, K, at a cathodic potential of −1.4 V.
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Fig. 7. SEM micrographs of fractured surfaces at a cathodic potential of −1.4 V: (a) 0.5%NaCL, K = 8.2 (MPam1/2 ), (b) 0.5%NaCL, K = 18.4 (MPam1/2 ), (c) 3.0%NaCl, K = 8.1 (MPam1/2 ), (d) 3.0%NaCl, K = 18.2 (MPam1/2 ), (e) 8.0%NaCl, K = 7.9 (MPam1/2 ), and (f) 8.0%NaCl, K = 14.8 (MPam1/2 ). The crack growth direction is from the left to the right.
products at the lowest NaCl concentration of 0.5%, revealing that the characteristic fractured surfaces are covered with packets in fine steps, where the packet size appears to coincide with the grain size. This appearance is similar to the transgranular, quasi-cleavage fracture surface morphology observed under environmentally induced cracking conditions in the literature [9,10,14,24,25].
Fig. 8 illustrates the change in the crack length as a function of the testing time at which a stable crack propagation occurred under the cathodic potentials of −2.5 and −3.0 V, which correspond to the boundary and the immunity region, respectively. The relationships between da/dt and K under different cathodic potentials are shown in Fig. 9. The KIscc , were measured to be 8.6, 19 and 14 MPam1/2 for the cathodic potentials of −1.4, −2.5 and −3.0 V, respectively, which indicate that the critical values are much lower in the corrosion region. The critical value at −3.0 V was lower than the value observed at −2.5 V. The crack propagation rates are much faster in the corrosion region (−1.4 V) than in the immunity region
Magnesium Alloy AZ31 3%NaCl solution −1.4V: 2.1kN (sustained load) −2.5V: 5.5kN −3.0V: 4.1kN Crack length a (m)
3.2. Effect of the cathodic potential
0.02
0.01
0 0
2500
5000
Time (s) Fig. 8. Relationship between crack length and testing time in 3.0%NaCl solution.
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Crack growth rate da/dt (m/s)
Magnesium alloy AZ31 3%NaCl −1.4V −2.5V −3.0V
10
10
−5
−6
7
10
30
20 1/2
Stress intensity factor K (MPam ) Fig. 9. Relationship between crack propagation rate, da/dt, and stress intensity factor, K, in 3.0%NaCl solution.
(−3.0 V) and are found to be the slowest at the boundary between the regions (−2.5 V). Furthermore, it seems that the crack propagation rates are insensitive to the value of the stress intensity factor at −3.0 V. At −2.5 V, KIscc is the highest (19 MPam1/2 ), and the crack growth rate was faster than 2 × 10−5 m/s when the K value became greater than 23 MPam1/2 . Thus, the SCC test was terminated at a relatively early stage of crack propagation as shown in Fig. 8. Fig. 10 reveals the SEM micrographs of typical fractured surfaces at the cathodic potentials of −2.5 and −3.0 V, respectively. Fractured surfaces were covered by corrosion products at −1.4 V in 3% NaCl solution, as shown in Fig. 7(c) and (d). In contrast, no corrosion products are recognized on the fractured surfaces at −2.5 and −3.0 V. A plausible explanation for these observations is that the SCC due to anodic dissolution occurs at the cathodic potential of −1.4 V, whereas hydrogen embrittlement dominates the SCC without anodic dissolution at −3.0 V. This explanation is in line with the Pourbaix diagram, where the −3.0 V corresponds to the immunity region. At the cathodic potential of −2.5 V, the fractured surfaces appear the same as those at −3.0 V, again without any corrosion products. In addition, the Pourbaix diagram implies that the anodic dissolution could take place at −2.5 V but at a lower concentration than 10−4 mol/L of Mg2+ . Thus, based on the fractographic analysis and the Pourbaix diagram showing a small amount of anodic dissolution, we postulate that the hydrogen embrittlement is the controlling factor of the SCC at −2.5 V. 4. Discussion 4.1. SCC behavior in the corrosion and immunity regions As shown in Figs. 6 and 7, da/dt increased and the corrosion products on the fractured surfaces became thicker with increasing NaCl concentration under a cathodic potential of −1.4 V, where the anodic dissolution occurs at the crack tip. It is well known that chloride ions significantly aggravate corrosion conditions of most metals. Consequently, the high NaCl concentration, i.e., the high
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chloride ion concentration, accelerated the crack propagation rate because its presence facilitated anodic dissolution. In the 3% NaCl solution, the KIscc value was much lower in the corrosion region (−1.4 V) than in the immunity region (−3.0 V) or the boundary region (−2.5 V), as shown in Fig. 9. At the cathodic potentials of −2.5 and −3.0 V, the anodic dissolution occurs to a very low degree, which means that the SCC mainly occurs due to hydrogen embrittlement. Subsequently, it could be concluded that the SCC by anodic dissolution exhibits a much lower KIscc than the SCC by hydrogen embrittlement. In the corrosion region, the SCC due to hydrogen embrittlement could be superimposed with the SCC by anodic dissolution. However, the total K values where SCC occurred at −1.4 V in a 3% NaCl solution (K = 8.6–15 MPam1/2 ) were lower than the KIscc value at −3.0 V in a 3% NaCl solution (KIscc = 14 MPam1/2 ), indicating that the anodic dissolution was the controlling factor for the SCC at −1.4 V in a 3% NaCl solution. Thus, the KIscc value in the corrosion region in Fig. 6 represents that of the SCC by anodic dissolution. In the corrosion region, the interaction between the anodic dissolution and hydrogen embrittlement could be one of the reasons for the accelerated growth rate under the more concentrated NaCl solution. However, much higher KIscc values at −2.5 V and −3.0 V imply that the effect of interaction would be small, and thus, we conclude that the SCC in the corrosion region was mainly due to the anodic dissolution at the crack tip and is accelerated by chloride ions. The fractured surface in the corrosion region without any corrosion products at −1.4 V in a 0.5% NaCl solution (Fig. 7(b)) showed fine steps in grains similar to those in the immunity region. As shown in Fig. 10, the space of each step was much finer in the fractured surfaces than in the immunity region. These very fine steps have a characteristic SCC appearance due to hydrogen embrittlement. The fatigue-fractured surfaces of the same alloy, AZ31, are shown in Fig. 11 for comparison. In this case, the fatigue crack propagation test was performed using the same specimen configuration, as shown in Fig. 2 under dry air conditions. It is clear that the fractured surfaces are much flatter than those of the SCC, and the characteristic fine steps in Fig. 10 are not observed. Further fractographic analyses of fatigue-fractured surfaces of this alloy under humid conditions are presented in Ref. [27].
4.2. Effect of cathodic potential on SCC behavior in the immunity region At the cathodic potentials of −2.5 V and −3.0 V, the SCC was dominated by hydrogen embrittlement, whereas the KIscc was lower at −3.0 V than at −2.5 V, as shown in Fig. 9. According to the Pourbaix diagram, the generation of H2 should be dependent on the cathodic potentials, where lower cathodic potential leads to accelerated hydrogen charging. Consequently, lower KIscc at −3.0 V could be attributed to a larger amount of charged hydrogen. The SCC fractured surfaces are similar at −2.5 and −3.0 V, as shown in Fig. 10, which implies that the mechanism of fracture is independent of the cathodic potential, where SCC occurs due to hydrogen embrittlement. Therefore, the higher crack propagation rates at −3.0 V relative to those at −2.5 V could also be attributed to the larger amount of charged hydrogen, as discussed above. At the cathodic potential of −3.0 V, the crack propagation rates seem to be insensitive to the stress intensity factor. At −3.0 V, the hydrogen charging occurred at a greater intensity and resulted in the SCC in which the hydrogen embrittlement was the main factor. Thus, the brittle fracture of the hydrogen embrittlement-type SCC strongly depends on the amount of charged hydrogen; consequently, the effect of the stress intensity factor on the crack propagation rate had decreased in a gradual manner. Although the KIscc value was 14 MPam1/2 at −3.0 V, it should be noted that KIscc could be further
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Fig. 10. SEM micrographs of fractured surfaces in 3.0%NaCl solution: (a) −2.5 V, K = 19.4 (MPam1/2 ), (b) −2.5 V, K = 20.9 (MPam1/2 ), (c) −3.0 V, K = 14.4 (MPam1/2 ), and (d) −3.0 V, K = 17.6 (MPam1/2 ). The crack growth direction is from the left to the right.
Fig. 11. SEM micrograph showing fatigue-fractured surfaces of AZ31 in dry air at R = 0.05: (a) K = 2.5 (MPam1/2 ) and (b) K = 9.7 (MPam1/2 ). The crack growth direction is from the left to the right.
decreased by a higher amount of charged hydrogen at a lower cathodic potential than −3.0 V. Several hydrogen embrittlement mechanisms have been proposed for ferrous materials in which hydrogen effects, such as slip localization [28,29], lattice hardening [30], softening [31] and hydrogen–dislocation interaction [32] may occur. For example, Murakami and Matsuoka [33] conducted fatigue crack propagation tests using hydrogen-charged Cr–Mo steel and proposed that the hydrogen concentration at the crack tip led to a hydrogeninduced local slip deformation instead of the lattice decohesion. It should be noted that the hydrogen decohesion hypothesis is one of the major models explaining hydrogen embrittlement. In our case, however, fractographic analysis revealed a brittle appearance, especially in the immunity region. For nonferrous metals, Birnbaum and coworkers had reported that the stress-induced hydride formation could occur in ␣-titanium [34] and niobium [35] alloys.
It is believed that hydrides precipitate in the stress singularity filed near the advancing crack tip, which results in brittle fracture of hydrides. Because magnesium forms hydrides, stress-induced hydride formation could be a reasonable mechanism of the hydrogen embrittlement-type SCC; thus the detection of hydrides around the crack tip is currently being examined. Although neither the hydrogen content of the specimen nor the hydrogen distribution around the crack tip was evaluated in this study, it would be reasonable to consider the basic hydrogen decohesion hypothesis and stress-induced hydride formation to be applicable to the SCC of Mg alloy in the immunity region. 5. Conclusions SCC tests of wrought Mg alloy (AZ31) were performed under controlled NaCl concentrations and cathodic potentials. Based on
Y. Uematsu et al. / Materials Science and Engineering A 531 (2012) 171–177
the experimental observations, the following conclusions can be made: 1. At a constant cathodic potential of −1.4 V, which corresponds to the corrosion region according to the Pourbaix diagram, the SCC was dominated by anodic dissolution at the crack tip and the crack propagation rate (da/dt) was accelerated with increasing NaCl concentration. 2. At a constant NaCl concentration of 3%, the SCC due to hydrogen embrittlement occurred in the immunity region (−3.0 V) and at the boundary between the two regions (−2.5 V), at which no corrosion products were recognized on the fractured surfaces. 3. The KIscc value for the SCC by anodic dissolution was much lower than that of the SCC due to hydrogen embrittlement. For the SCC dominated by hydrogen embrittlement, the KIscc value became lower and the crack propagation rate became higher with decreasing cathodic potential. These observations could be attributed to a higher amount of charged hydrogen at a lower cathodic potential. 4. The crack propagation rate was insensitive to the stress intensity factor at −3.0 V due to the brittle nature of the SCC that was induced by hydrogen embrittlement. References [1] B.A. Shaw, ASM Handbook 13A (2003) 692–696. [2] R.B. Alvarez, H.J. Martin, M.F. Horstemeyer, M.Q. Chandler, N. Williams, P.T. Wang, A. Ruiz, Corros. Sci. 52 (2010) 1635–1648. [3] G. Song, A. Atrens, Adv. Eng. Mater. 1 (1999) 11–33. [4] G. Song, A. Atrens, Adv. Eng. Mater. 5 (2003) 837. [5] Y. Uematsu, K. Tokaji, T. Ohashi, Strength Mater. 40 (2008) 130–133. [6] Z.Y. Nan, S. Ishihara, T. Goshima, Int. J. Fatigue (2008) 1181–1188.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Stress corrosion cracking behavior of the wrought magnesium alloy AZ31 under controlled cathodic potentials Yoshihiko Uematsu a,∗ , Toshifumi Kakiuchi a , Masaki Nakajima b a b
Department of Mechanical and Systems Engineering, Gifu University, 1-1 Yanagido, Gifu 501-1193, Japan Toyota National College of Technology, 2-1 Eisei-cho, Toyota, Aichi 471-8525, Japan
a r t i c l e
i n f o
Article history: Received 23 May 2011 Received in revised form 14 October 2011 Accepted 18 October 2011 Available online 24 October 2011 Keywords: Magnesium alloy Stress corrosion cracking Hydrogen embrittlement Anodic dissolution Crack propagation rate Stress intensity factor
a b s t r a c t Stress corrosion cracking (SCC) tests were performed using compact tension (CT) specimens in a NaCl solution under hydrogen-charged conditions to investigate the SCC behavior of the wrought magnesium alloy AZ31. The effect of NaCl concentration was estimated by conducting the SCC tests under a constant cathodic potential of −1.4 V, with NaCl concentrations of 0.5, 3.0 and 8.0%. The crack propagation rate (da/dt) became accelerated with increasing concentration due to the enhanced anodic dissolution by chloride ions. Subsequently, SCC tests were performed using a 3%NaCl solution with cathodic potentials at −1.4, −2.5 and −3.0 V to investigate the effect of cathodic potentials on SCC behavior. According to the Pourbaix diagram, the cathodic potential of −3.0 V corresponds to the immunity region, whereas −1.4 V is within the corrosion region and −2.5 V is the boundary between the two regions. In the immunity region, the SCC due to hydrogen embrittlement occurred when the KIscc value was higher than that of the SCC dominated by anodic dissolution. The KIscc value decreased and the crack propagation rate increased with decreasing cathodic potential in the immunity region at which SCC occurred due to hydrogen embrittlement. At the cathodic potential of −3.0 V, the da/dt was insensitive to the stress intensity factor because highly charged hydrogen gave rise to brittle fracture. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Magnesium (Mg) alloys have excellent properties, such as low weight, high specific strength and stiffness, machinability and recyclability. Because of these advantages, Mg alloys are very attractive structural materials that can be used in a wide variety of applications; in particular, they can be utilized as components for air planes or transportation vehicles in which saving weight is extremely important. The major drawback of this alloy is poor corrosion resistance compared with other light metals, such as aluminum and titanium alloys. Hence, it is very important to understand the corrosion properties of Mg alloys to be able to use these alloys as structural components, and thus, many studies have been performed on corrosion [1–4], corrosion fatigue [5–7] and stress corrosion cracking (SCC) [8–25]. Song and Atrens [3,4] reviewed the corrosion behavior of Mg alloys in detail and proposed a partially protective surface film model, a mono-valent magnesium ion model, a particle undermining model and a magnesium hydride model for the hydrogen evolution reaction (HER) during corrosion. All models suggest that hydrogen generation is strongly related
∗ Corresponding author. Tel.: +81 58 293 2501; fax: +81 58 293 2491. E-mail address: [email protected] (Y. Uematsu). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.10.052
to the corrosion of Mg alloys. From an engineering viewpoint, the SCC behavior of Mg alloy has been widely investigated because SCC causes the unexpected sudden failure of structures. In addition, hydrogen embrittlement plays an important role due to the presence of HER during corrosion. Winzer et al. [8] surveyed the SCC behavior of Mg alloys, and demonstrated the effect of alloy composition, the manufacturing process and heat treatment on the SCC susceptibility of materials. They also postulated that hydrogen should be involved in the mechanism of SCC because Mg had a negative difference effect (NDE) [3,4,8]. Thus, it is important to evaluate the effect of hydrogen on SCC behavior. Most of the investigations of SCC in Mg alloys are based on slow strain rate (SSR) testing in which the SCC susceptibility of materials could be easily evaluated. For example, Kannan et al. [15] reported that rare-earth elements could improve SCC resistance relative to the conventional structural Mg alloy AZ80. Ben-Hamu et al. [18] showed the beneficial effect of Si incorporation on the SCC resistance. From the viewpoint of designing mechanical components, quantitative analysis of crack propagation rate is inevitable. Marrow et al. [12] investigated the crack growth path using fourpoint bend specimens with small holes as crack initiators, and Winzer et al. [17,20] attempted to characterize the SCC susceptibility by threshold stress, which was determined using a DC potential drop method during the SSR test. However, to estimate
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Fig. 1. Microstructure of material.
Fig. 2. Compact tension (CT) specimen configuration.
2.2. Stress corrosion cracking tests crack propagation rates based on fracture mechanics parameters, i.e., stress intensity factor (K), the SCC tests should be performed using specimens with long cracks, such as compact tension (CT) and center-cracked panel specimens. Furthermore, the SCC could be attributed to the combined effects of anodic dissolution and hydrogen embrittlement at the crack tip. As mentioned above, hydrogen embrittlement plays an important role in the SCC of Mg alloys [8–11]. For example, Song et al. [13] demonstrated that the preexposure of Mg alloy to a NaCl solution tends to decrease tensile strength and elongation due to hydrogen embrittlement. However, it is difficult to investigate the effects of anodic dissolution and hydrogen embrittlement independently. Kannan et al. [14] conducted SSR tests under a continuously hydrogen-charged condition to enhance the effect of hydrogen on SCC behavior and found a significant decrease in tensile strength, whereas the crack propagation rate was not evaluated. This study focuses on the SCC behavior of the wrought magnesium alloy AZ31 and provides a qualitative estimation of the crack propagation rate based on fracture mechanics. The SCC tests were performed using CT specimens with long cracks that were placed in a NaCl solution under hydrogen-charged conditions. By controlling both the NaCl concentration and the cathodic potential, the effects of anodic dissolution and hydrogen embrittlement on SCC behavior were decoupled in this study.
Prior to the SCC tests, a pre-crack with a length of 2 mm was introduced from an artificial notch root of the CT specimen by a fatigue test at a frequency of 10 Hz, a stress ratio of 0.1 and an initial K value of 3 MPam1/2 . The SCC test specimens with fatigue pre-crack were attached to a creep testing machine equipped with a NaCl solution tank, as shown in Fig. 3. The crack length was monitored by a crack gauge (KV-25B: KYOWA) bonded to the segment of the CT specimen where the crack might grow. The Ag/AgCl and Pt electrodes were used as reference and counter electrodes, respectively. A cathodic potential of −1.4 V was applied using a potentiostat, and the SCC tests were performed in a NaCl solution with concentrations of 0.5%, 3% and 8%, to investigate the SCC behavior that is dominated by the anodic dissolution at the crack tip. Notably, the cathodic potential of −1.4 V corresponds to the corrosion region according to the Pourbaix diagram [26], as shown in Fig. 4. Subsequently, the SCC tests were conducted using a 3% NaCl solution at the cathodic potentials of −2.5 and −3.0 V to investigate the SCC behavior that is controlled by hydrogen embrittlement. The cathodic potential of −3.0 V is in the immunity region, while the potential at −2.5 V is at the boundary between the corrosion and immunity regions based on the Pourbaix diagram. It should be noted that the potential is established on the surface of the specimen and not at the crack tip where the SCC occurs. Thus, it is more important to understand the actual potential at the crack tip for a more precise prediction of the crack propagation rates, although the measurement of the potential at the crack tip is quite difficult.
2. Experimental details 2.1. Material and specimen configuration The material used in this study is the wrought magnesium alloy AZ31. The microstructure of AZ31 is shown in Fig. 1. The chemical composition (mass%) is as follows: Al: 2.7, Zn: 0.79, Mn: 0.44, Fe: 0.0012, Ni: 0.0009, Cu: 0.0011, Si: 0.004, Ca: 0.001, Sn < 0.001, and Mg: bal. The mechanical properties of this material are summarized in Table 1. CT specimens with a width of 50.8 mm and a thickness of 6 mm were used for the SCC tests, as shown in Fig. 2, using a CT specimen configuration that complies with the ASTM E647 standard. The specimens were machined from the as-received AZ31 plates with a thickness of 6 mm in the L–T orientation, where L is the longitudinal direction, or the extrusion direction, and T is the transverse direction. The surface of the specimen was not modified, with an average roughness of 3.2 m.
Fig. 3. SCC test configuration.
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Table 1 Mechanical properties of material. 0.2% proof stress 0.2 (MPa)
Tensile strength B (MPa)
Elongation ı (%)
Vickers hardness HV
Elastic modulus E (GPa)
Grain size d (m)
170
248
17
54
46
29
0.03
Crack length a (m)
The SCC test specimens were maintained in the NaCl solution at a given cathodic potential for 24 h with no load to charge the hydrogen. The initial value of K was approximately 7 MPam1/2 , and the load was gradually increased every 24 h until the crack began to propagate. Consequently, it took 3–7 days before the SCC began, depending on the testing environments. Thus, the precharging time was assumed to be long enough for the hydrogen to diffuse to the crack tip, where the triaxial stress is dominant. Once the crack began to propagate, the load was sustained, and the crack propagation was monitored by a crack gauge. The sustained load during crack propagation leads to an increasing K value in the SCC test. Therefore, the threshold value of KIscc is defined as the K value calculated from the load step just before the crack propagation occurs. In this case, the increment in the K value with the increasing load was approximately 0.4–0.5 MPam1/2 for each step.
Magnesium Alloy AZ31 Cathodic potential −1.4V
0.02
0.01
0
0
20000 Time (s)
3. Results 3.1. Effect of the NaCl concentration
as at 8%, resulted in much thicker corrosion products. The corrosion products seem to be thicker at lower K values due to the lower crack growth rate, although it is difficult to distinguish between the corrosion products formed during the SCC or afterwards. The EDX analysis of the corrosion products on the fractured surface of the 3.0% NaCl solution was performed. The data demonstrate that the corrosion products are Mg(OH)2 because the quantitative analysis revealed an atomic ratio of oxygen to magnesium of 2:1. Some part of the fractured surfaces were not completely covered by corrosion
10
10
a 2
Mg2+ Corrosion
E(V)
Mg(OH)2 Passivation Test conditions
-2
Crack growth rate da/dt (m/s)
0 -2 -4 -6
-1
10
10
10
-6
Mg Immunity 0
−3
−4
Magnesium alloy AZ31 Cathodic potential −1.4V 0.5%NaCl 3.0%NaCl 8.0%NaCl
−5
−6
−7
0
3
-3
40000
Fig. 5. Relationship between crack length and testing time at a cathodic potential of −1.4 V.
The change in the crack length as a function of testing time is illustrated in Fig. 5. It should be noted that the crack propagated in a stable manner under the sustained loading condition regardless of the NaCl concentration. Fig. 6 shows the relationship between the crack propagation rate (da/dt) and the stress intensity factor (K), where da/dt was calculated from Fig. 5. The critical values of K for crack propagation, i.e., the KIscc , were measured to be 7.0, 8.6 and 7.1 MPam1/2 for the NaCl concentrations of 0.5, 3 and 8%, respectively, which indicates that the critical values are insensitive to the NaCl concentration. However, the crack propagation rate increases with increasing NaCl concentration. Fig. 7 reveals the SEM micrographs of typical fractured surfaces at different K levels in 0.5%, 3.0% and 8.0% NaCl solutions, respectively. Fractured surfaces are covered by corrosion products that are irrelevant to the K values, where the high NaCl concentration, such
0
0.5%NaCl 3.0%NaCl 8.0%NaCl
1 10
−8
3
7
14
PH Fig. 4. A typical Pourbaix diagram for magnesium–water system at 25 ◦ C.
5
8
10
20
30
1/2
Stress intensity factor K (MPam ) Fig. 6. Relationship between crack propagation rate, da/dt, and stress intensity factor, K, at a cathodic potential of −1.4 V.
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Fig. 7. SEM micrographs of fractured surfaces at a cathodic potential of −1.4 V: (a) 0.5%NaCL, K = 8.2 (MPam1/2 ), (b) 0.5%NaCL, K = 18.4 (MPam1/2 ), (c) 3.0%NaCl, K = 8.1 (MPam1/2 ), (d) 3.0%NaCl, K = 18.2 (MPam1/2 ), (e) 8.0%NaCl, K = 7.9 (MPam1/2 ), and (f) 8.0%NaCl, K = 14.8 (MPam1/2 ). The crack growth direction is from the left to the right.
products at the lowest NaCl concentration of 0.5%, revealing that the characteristic fractured surfaces are covered with packets in fine steps, where the packet size appears to coincide with the grain size. This appearance is similar to the transgranular, quasi-cleavage fracture surface morphology observed under environmentally induced cracking conditions in the literature [9,10,14,24,25].
Fig. 8 illustrates the change in the crack length as a function of the testing time at which a stable crack propagation occurred under the cathodic potentials of −2.5 and −3.0 V, which correspond to the boundary and the immunity region, respectively. The relationships between da/dt and K under different cathodic potentials are shown in Fig. 9. The KIscc , were measured to be 8.6, 19 and 14 MPam1/2 for the cathodic potentials of −1.4, −2.5 and −3.0 V, respectively, which indicate that the critical values are much lower in the corrosion region. The critical value at −3.0 V was lower than the value observed at −2.5 V. The crack propagation rates are much faster in the corrosion region (−1.4 V) than in the immunity region
Magnesium Alloy AZ31 3%NaCl solution −1.4V: 2.1kN (sustained load) −2.5V: 5.5kN −3.0V: 4.1kN Crack length a (m)
3.2. Effect of the cathodic potential
0.02
0.01
0 0
2500
5000
Time (s) Fig. 8. Relationship between crack length and testing time in 3.0%NaCl solution.
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Crack growth rate da/dt (m/s)
Magnesium alloy AZ31 3%NaCl −1.4V −2.5V −3.0V
10
10
−5
−6
7
10
30
20 1/2
Stress intensity factor K (MPam ) Fig. 9. Relationship between crack propagation rate, da/dt, and stress intensity factor, K, in 3.0%NaCl solution.
(−3.0 V) and are found to be the slowest at the boundary between the regions (−2.5 V). Furthermore, it seems that the crack propagation rates are insensitive to the value of the stress intensity factor at −3.0 V. At −2.5 V, KIscc is the highest (19 MPam1/2 ), and the crack growth rate was faster than 2 × 10−5 m/s when the K value became greater than 23 MPam1/2 . Thus, the SCC test was terminated at a relatively early stage of crack propagation as shown in Fig. 8. Fig. 10 reveals the SEM micrographs of typical fractured surfaces at the cathodic potentials of −2.5 and −3.0 V, respectively. Fractured surfaces were covered by corrosion products at −1.4 V in 3% NaCl solution, as shown in Fig. 7(c) and (d). In contrast, no corrosion products are recognized on the fractured surfaces at −2.5 and −3.0 V. A plausible explanation for these observations is that the SCC due to anodic dissolution occurs at the cathodic potential of −1.4 V, whereas hydrogen embrittlement dominates the SCC without anodic dissolution at −3.0 V. This explanation is in line with the Pourbaix diagram, where the −3.0 V corresponds to the immunity region. At the cathodic potential of −2.5 V, the fractured surfaces appear the same as those at −3.0 V, again without any corrosion products. In addition, the Pourbaix diagram implies that the anodic dissolution could take place at −2.5 V but at a lower concentration than 10−4 mol/L of Mg2+ . Thus, based on the fractographic analysis and the Pourbaix diagram showing a small amount of anodic dissolution, we postulate that the hydrogen embrittlement is the controlling factor of the SCC at −2.5 V. 4. Discussion 4.1. SCC behavior in the corrosion and immunity regions As shown in Figs. 6 and 7, da/dt increased and the corrosion products on the fractured surfaces became thicker with increasing NaCl concentration under a cathodic potential of −1.4 V, where the anodic dissolution occurs at the crack tip. It is well known that chloride ions significantly aggravate corrosion conditions of most metals. Consequently, the high NaCl concentration, i.e., the high
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chloride ion concentration, accelerated the crack propagation rate because its presence facilitated anodic dissolution. In the 3% NaCl solution, the KIscc value was much lower in the corrosion region (−1.4 V) than in the immunity region (−3.0 V) or the boundary region (−2.5 V), as shown in Fig. 9. At the cathodic potentials of −2.5 and −3.0 V, the anodic dissolution occurs to a very low degree, which means that the SCC mainly occurs due to hydrogen embrittlement. Subsequently, it could be concluded that the SCC by anodic dissolution exhibits a much lower KIscc than the SCC by hydrogen embrittlement. In the corrosion region, the SCC due to hydrogen embrittlement could be superimposed with the SCC by anodic dissolution. However, the total K values where SCC occurred at −1.4 V in a 3% NaCl solution (K = 8.6–15 MPam1/2 ) were lower than the KIscc value at −3.0 V in a 3% NaCl solution (KIscc = 14 MPam1/2 ), indicating that the anodic dissolution was the controlling factor for the SCC at −1.4 V in a 3% NaCl solution. Thus, the KIscc value in the corrosion region in Fig. 6 represents that of the SCC by anodic dissolution. In the corrosion region, the interaction between the anodic dissolution and hydrogen embrittlement could be one of the reasons for the accelerated growth rate under the more concentrated NaCl solution. However, much higher KIscc values at −2.5 V and −3.0 V imply that the effect of interaction would be small, and thus, we conclude that the SCC in the corrosion region was mainly due to the anodic dissolution at the crack tip and is accelerated by chloride ions. The fractured surface in the corrosion region without any corrosion products at −1.4 V in a 0.5% NaCl solution (Fig. 7(b)) showed fine steps in grains similar to those in the immunity region. As shown in Fig. 10, the space of each step was much finer in the fractured surfaces than in the immunity region. These very fine steps have a characteristic SCC appearance due to hydrogen embrittlement. The fatigue-fractured surfaces of the same alloy, AZ31, are shown in Fig. 11 for comparison. In this case, the fatigue crack propagation test was performed using the same specimen configuration, as shown in Fig. 2 under dry air conditions. It is clear that the fractured surfaces are much flatter than those of the SCC, and the characteristic fine steps in Fig. 10 are not observed. Further fractographic analyses of fatigue-fractured surfaces of this alloy under humid conditions are presented in Ref. [27].
4.2. Effect of cathodic potential on SCC behavior in the immunity region At the cathodic potentials of −2.5 V and −3.0 V, the SCC was dominated by hydrogen embrittlement, whereas the KIscc was lower at −3.0 V than at −2.5 V, as shown in Fig. 9. According to the Pourbaix diagram, the generation of H2 should be dependent on the cathodic potentials, where lower cathodic potential leads to accelerated hydrogen charging. Consequently, lower KIscc at −3.0 V could be attributed to a larger amount of charged hydrogen. The SCC fractured surfaces are similar at −2.5 and −3.0 V, as shown in Fig. 10, which implies that the mechanism of fracture is independent of the cathodic potential, where SCC occurs due to hydrogen embrittlement. Therefore, the higher crack propagation rates at −3.0 V relative to those at −2.5 V could also be attributed to the larger amount of charged hydrogen, as discussed above. At the cathodic potential of −3.0 V, the crack propagation rates seem to be insensitive to the stress intensity factor. At −3.0 V, the hydrogen charging occurred at a greater intensity and resulted in the SCC in which the hydrogen embrittlement was the main factor. Thus, the brittle fracture of the hydrogen embrittlement-type SCC strongly depends on the amount of charged hydrogen; consequently, the effect of the stress intensity factor on the crack propagation rate had decreased in a gradual manner. Although the KIscc value was 14 MPam1/2 at −3.0 V, it should be noted that KIscc could be further
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Fig. 10. SEM micrographs of fractured surfaces in 3.0%NaCl solution: (a) −2.5 V, K = 19.4 (MPam1/2 ), (b) −2.5 V, K = 20.9 (MPam1/2 ), (c) −3.0 V, K = 14.4 (MPam1/2 ), and (d) −3.0 V, K = 17.6 (MPam1/2 ). The crack growth direction is from the left to the right.
Fig. 11. SEM micrograph showing fatigue-fractured surfaces of AZ31 in dry air at R = 0.05: (a) K = 2.5 (MPam1/2 ) and (b) K = 9.7 (MPam1/2 ). The crack growth direction is from the left to the right.
decreased by a higher amount of charged hydrogen at a lower cathodic potential than −3.0 V. Several hydrogen embrittlement mechanisms have been proposed for ferrous materials in which hydrogen effects, such as slip localization [28,29], lattice hardening [30], softening [31] and hydrogen–dislocation interaction [32] may occur. For example, Murakami and Matsuoka [33] conducted fatigue crack propagation tests using hydrogen-charged Cr–Mo steel and proposed that the hydrogen concentration at the crack tip led to a hydrogeninduced local slip deformation instead of the lattice decohesion. It should be noted that the hydrogen decohesion hypothesis is one of the major models explaining hydrogen embrittlement. In our case, however, fractographic analysis revealed a brittle appearance, especially in the immunity region. For nonferrous metals, Birnbaum and coworkers had reported that the stress-induced hydride formation could occur in ␣-titanium [34] and niobium [35] alloys.
It is believed that hydrides precipitate in the stress singularity filed near the advancing crack tip, which results in brittle fracture of hydrides. Because magnesium forms hydrides, stress-induced hydride formation could be a reasonable mechanism of the hydrogen embrittlement-type SCC; thus the detection of hydrides around the crack tip is currently being examined. Although neither the hydrogen content of the specimen nor the hydrogen distribution around the crack tip was evaluated in this study, it would be reasonable to consider the basic hydrogen decohesion hypothesis and stress-induced hydride formation to be applicable to the SCC of Mg alloy in the immunity region. 5. Conclusions SCC tests of wrought Mg alloy (AZ31) were performed under controlled NaCl concentrations and cathodic potentials. Based on
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the experimental observations, the following conclusions can be made: 1. At a constant cathodic potential of −1.4 V, which corresponds to the corrosion region according to the Pourbaix diagram, the SCC was dominated by anodic dissolution at the crack tip and the crack propagation rate (da/dt) was accelerated with increasing NaCl concentration. 2. At a constant NaCl concentration of 3%, the SCC due to hydrogen embrittlement occurred in the immunity region (−3.0 V) and at the boundary between the two regions (−2.5 V), at which no corrosion products were recognized on the fractured surfaces. 3. The KIscc value for the SCC by anodic dissolution was much lower than that of the SCC due to hydrogen embrittlement. For the SCC dominated by hydrogen embrittlement, the KIscc value became lower and the crack propagation rate became higher with decreasing cathodic potential. These observations could be attributed to a higher amount of charged hydrogen at a lower cathodic potential. 4. The crack propagation rate was insensitive to the stress intensity factor at −3.0 V due to the brittle nature of the SCC that was induced by hydrogen embrittlement. References [1] B.A. Shaw, ASM Handbook 13A (2003) 692–696. [2] R.B. Alvarez, H.J. Martin, M.F. Horstemeyer, M.Q. Chandler, N. Williams, P.T. Wang, A. Ruiz, Corros. Sci. 52 (2010) 1635–1648. [3] G. Song, A. Atrens, Adv. Eng. Mater. 1 (1999) 11–33. [4] G. Song, A. Atrens, Adv. Eng. Mater. 5 (2003) 837. [5] Y. Uematsu, K. Tokaji, T. Ohashi, Strength Mater. 40 (2008) 130–133. [6] Z.Y. Nan, S. Ishihara, T. Goshima, Int. J. Fatigue (2008) 1181–1188.
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