Corrosion Science 44 (2002) 729±749
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Corrosion behaviour of stressed magnesium alloys P.L. Bonora a, M. Andrei a,*, A. Eliezer b, E.M. Gutman b a
Department of Materials Engineering, Laboratory of Industrial Corrosion Control, University of Trento, via Messiano 77, I-38050 Trento, Italy b Department of Materials Engineering, Ben-Gurion University of the Negev, Beer-Sheva, Israel Received 3 July 2000; accepted 28 May 2001
Abstract Potentiodynamic polarisation and impedance measurements are used to examine the corrosion aspects of some Mg-based alloys, which were previously stressed in order to established the eect of mechanical deformation on surface electrochemical reactions. A ®rst approach was made for the unstressed alloys. The electrochemical tests were carried out in a sodium borate buer solution. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: Mechanochemical eect; Magnesium alloys; Potentiodynamic polarisation; Electrochemical impedance
1. Introduction The perception of magnesium as a rapidly corroding material has been a major obstacle to its growth in structural applications despite its other desirable physical properties. More importantly, the problem of stress corrosion is becoming a major one today in Mg alloys. New Mg alloys have been developed in recent years to meet the needs of structural applications. Many cast alloys were tested for stress corrosion cracking and their high sensitivity to stress corrosion was demonstrated [1,2]. Some magnesium and magnesium alloys applications, especially in the transportation industry require
*
Corresponding author. Tel.: +39-461-882403; fax: +39-461-881977. E-mail address:
[email protected] (M. Andrei).
0010-938X/02/$ - see front matter Ó 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 1 ) 0 0 1 0 1 - 9
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good fatigue properties [1]. Corrosion fatigue is the simultaneous action of corrosion and alternating stress on an alloy, leading to a lower fatigue limit than in the case of fatigue without any environmental impact. In general, the aim of every stress corrosion study is to develop a basis for quantitative prediction and qualitative diagnostic of the incidence of cracking and failure under impact of environments and stresses. A more realistic way is to separate stress corrosion with in¯uencing signi®cant phenomena and to study these phenomena with the hope to develop a more general mechanism in future. Such important phenomena are mechanochemical eects (MCEs) and creep at crack tip [3]. 2. Experimental method In this work, AM50 and AZ91D magnesium alloys were studied. The most common die casting alloy is AZ91D with 9% Al content. The combination of ¯uidity, strength and ductility of an alloy with 9% Al is undeniably one of the reasons for which AZ91D alloys usage is widespread. Nevertheless, AZ91D has some disadvantages for use in a die casting. Firstly, although the material has good shortterm strength at elevated temperatures, the creep resistance is poor. This has been attributed to the large amount of the presence of low melting eutectic (Mg17 Al12 ). Another disadvantage of higher Al content containing alloys is that ductility is usual sacri®ced to grain strength [4]. But Mg alloys having lower Al content, for example AM50 (5% Al) were found to be more ductile, especially during impact situation. However, many structural applications require an appreciable amount of energy absorption during their operation [3]. Thus, it is essential to study corrosion fatigue resistance of Mg alloys and to investigate the correlation of corrosion fatigue with the mechanochemical behaviour of the Mg alloys. The Mg alloys were submitted to the standard mechanical tests before the electrochemical ones. The behaviour of both stressed and non-stressed magnesium alloys was investigated by using potentiodynamic polarisation measurements (DC polarisation) and electrochemical impedance spectroscopy (EIS). All the electrochemical measurements were performed in a tetraborate buer solution that is a stable solution with a pH (9.7) at which Mg can cover itself with more or less protective oxide or hydroxide which checks the dissolution reaction [5]. A very used solution for the corrosion common studies of Mg alloys is NaCl Mg(OH)2 with a pH about 10.5. This solution does not have a stable pH value in the electrode surface and thus, the corrosion conditions are not stable. In our study the borate solution eliminate this unstable and the undesirable presence of anions Cl that even in small amounts, usually break down the protective ®lm on Mg [6,10]. The electrochemical testing was employed to study the main features of the processes taking place at the alloy/solution interface. The eect of dierent applied tensile stresses on the Mg-alloy corrosion resistance was studied. The corrosion resistance of Mg alloys was pointed out by EIS measurements performed during the free immersion time and under polarisation and the eect of the dierent tensile stresses was studied. The evolution of the electrode/electrolyte interface at dierent
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immersion times was also studied. The corrosion rate has been calculated by the analysis of the potentiodynamic polarisation that was carried out by the linear polarisation method. The ¯at specimens of magnesium alloys (AM50 and AZ91D) used for this study were prepared by the casting procedure (tcast 660°C and 650°C, respectively). Chemical composition (wt.%) and mechanical properties are given in Tables 1 and 2. The electrochemical tests using both DC (potentiodynamic polarisation) and AC techniques (EIS) were carried out in an aerated 0.05 M sodium tetraborate solution (pH 9:7). For all measurements a three electrode electrochemical cell was used, with an Ag/AgCl as reference electrode and a platinum counter electrode. The working electrode was prepared from the Mg-alloy samples after the application of the tensile strength. Mg-alloy samples were embedded in an acrylic resin to provide electrical isolation of the sample surface. The samples were air dried at room temperature. The potentiodynamic curves were obtained using a PAR 273 potentiostat, with a voltage scan rate of 0.2 mV/s. The impedance measurements were carried out using a Solartron 1250 frequency response analyser coupled with the potentiostat. All the experiments were controlled by a PC, which was also used for the acquisition, storage and plotting of data. The scanned frequency ranged from 6 mHz to 100 kHz and the perturbation amplitude was of 5 mV (it was observed that a variation of the amplitude did not change the frequency response of the electrode/electrolyte interface). The impedance measurements were performed at open circuit potential (EOC ) and also under potentiostatic conditions. Table 1 Chemical analysis after die casting in wt.% Mg alloy
Al
Mn
Zn
Si
Cu
Fe
Ni
Be
AM50 AZ91D
5.1 8.4
0.57 0.17
0 0.85
0.013 0.01
0.0007 0.0008
0.0074 0.0013
0.0006 0.0007
0.0013 0.0003
Table 2 Mechanical properties of samples Mg alloy
TYSa (MPa)
UTS (MPa)
AM50 AM50 AM50 AM50 AM50 AM50 AM50 AZ91D AZ91D AZ91D AZ91D
Without 120.3 129.9 131.2 129.9 135.6 118.7 Without 166.4 166.7 179.2
Loading 125.4 147.6 152.1 151.7 166.0 160.4 Loading 174.2 180.4 193.7
a
TYS is the tensile yield strength.
Elongation (%) 0.30 1.94 2.13 2.26 3.56 4.39 0.29 0.45 0.71
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All the measurements of this study were taken after stabilisation of the open circuit potential, i.e. 20 min after immersion of the electrode in the solution. All the potential values are reported with respect to the reference electrode (EAg=AgCl 0:210 V versus ENH).
3. Results and discussion 3.1. Potentiodynamic polarisation measurements The potentiodynamic curves in 0.05 M sodium tetraborate solution (pH 9:7) of unstressed Mg alloys and pure Mg are shown in Fig. 1. As a ®rst remark we note that the shape of the polarisation curves is almost the same for the two alloys. In addition, a current plateau begins at 1:2 and 1:3 V for AM50 and AZ91D, respectively, the current values being quite similar and relatively high. We can assume that the alloy samples present a ``pseudo-passivation'', which diers with respect to pure magnesium. We also note that by alloying the corrosion potential values become much nobler and the anodic current densities reduce. Figs. 2 and 3 show the potentiodynamic curves obtained for both Mg alloys as a function of the applied stress. The eect of mechanical deformation on the corrosion behaviour can be seen. The corrosion parameters estimate the MCE, which means the increasing of the anodic current densities and the disennobeling of the potential values caused by the deformation. The corrosion rate (icorr ) was determined for small deviations from the corrosion potential. The obtained values are plotted as a function of strain in the Fig. 4. The graphs in Fig. 4 also contain the corrosion potential values. First of all, we remark that both corrosion current and corrosion potential depend strongly on the applied stress. So, we can note that icorr increases with the elongation up to a value of 1.94% and 0.29% for AM50 and AZ91D, respectively. After that, the corrosion current is
Fig. 1. Potentiodynamic curves of pure Mg and of AM50 and AZ91D alloys at v 0:2 mV/s.
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Fig. 2. Potentiodynamic curves of stressed and unstressed AM50 alloy at v 0:2 mV/s.
Fig. 3. Potentiodynamic curves of stressed and unstressed AZ91D alloy at v 0:2 mV/s.
reduced. Thus, the corrosion rate passes over a maximum in relation to the increase in plastic deformation, according to the general theory [2]. For the same values of the strain for that a maximum value of the corrosion current was obtained, the corrosion potential reaches the most negative values. A signi®cant MCE for both corrosion rate and corrosion potential manifests itself in the magnesium alloys under study. It can be also seen that a new slower increase of icorr is produced at 3.56% and after 0.45% of residual strain for AM50 and AZ91D, respectively. Experimental con®rmation of the correlation between the mechanochemical behaviour and the strain hardening stages (intensive strain hardening and dynamic recovery) could be see in Fig. 5, which regroups the stress±strain curves and the corrosion rate variations for both magnesium alloys. The observed correlation could be explained by the change of dislocation substructure during the plastic deformation as was already done in the case of other materials alloys [7]. Under intense strain
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Fig. 4. Corrosion current and corrosion potential for AM50 (a) and AZ91D (b) magnesium alloys as a function of strain.
hardening, stresses grow and planar dislocation pile-ups appear. This leads to a sharp increase in the MCE. At the ®nal stage of the dynamic recovery, planar pileups are destroyed due to dislocation cross-slip and partial annihilation. This leads to a decrease in the MCE value, which, thus, should pass over a maximum in the process of plastic deformation. The MCE should grow intensively under plastic deformation at the stage of strain hardening. In particular, the acceleration of metal anodic dissolution is caused by a local reduction of equilibrium potential in the vicinity of dislocations. This eect is much lower at the ®nal recovery stage. At this stage, strain hardening is suppressed because of the development of dislocation cross-slip processes. In a previous stress corrosion study [1] a signi®cant mechanoelectrochemical eect for corrosion rate, anodic current under constant potential and corrosion potential for the dierent Mg alloys in alkaline media (pH 10:5) was already observed. It was also noted that for the AM20 and AM50 alloys the mechanical eect show a maximum when the strain hardening stage transfers to the dynamic recovery stage. It is also important that the features of mechanical behaviour are the same in both active and pseudo-passive states, independently of surface ®lm existence. Thus, the analysis of the potentiodynamic curves was also carried out at the cross-sections of
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Fig. 5. Anodic current (under constant potential) in the active state and stress curves as a function of strain for AM50 (a) and AZ91D (b) magnesium alloys.
curves at dierent selected potential levels to ®nd the corresponding current densities. The potential levels were chosen in both active and ``pseudo-passive'' ranges. The obtained values are plotted in Fig. 6. For a better illustration of the MCE, Fig. 6 also regroups the potential values under constant current densities. All the values are plot as a function of strain. Therefore, it can be seen that the anodic current increment and the decrease of potential manifest itself in both active and passive states in the case of AM50 alloy. For the AZ91D magnesium alloy, a slight decrease in anodic current and a slow increment of potential are observed in both active and passive states at the beginning of the applied stress, contrary to the corrosion current and potential. The two alloys under study contain the intermetallic compound Mg17 Al12 . A low content of Al leads to a reduced amount of phase Mg17 Al12 [4]. It was observed that the b-phase is inert to the chloride solution in comparison to the magnesium matrix and acts as a corrosion barrier depending on the manner that the phase Mg17 Al12 is distributed in the alloy matrix [8]. So, the b-phase present in the cast alloy is generally more resistant than the surrounding matrix alloy. Corrosion resistance of the b-phase is related to its passive behaviour within a much wider pH range than its pure components.
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Fig. 6. Anodic current densities (under constant potential) and potential values (under constant current) in both active and pseudo-passivation states for AM50 (a) and AZ91D (b) magnesium alloys, as a function of strain.
Corrosion rate measurements of binary alloys showed that Al was the only element which caused a decrease in the corrosion rate of Mg; the corrosion rate decreased with increasing Al content [9]. Indeed, in our study realised in a free chloride media it was seen that in the nonstressed conditions, the corrosion rate of the alloy with higher Al content is lower than that for the alloy with less Al. This is largely due to the presence and distribution of the b-phase (Fig. 7a) that better protects in the AZ91D alloy with a higher content of Al and so with an elevated amount of b-phase. For AM50 with low Al content, the fraction b-phase is small and is only formed as discrete islands in the structure (Fig. 7b). But hard secondary phase promotes strain hardening and thus, increases chemical potential of atoms, i.e. they create the necessary conditions for mechanochemical dissolution [2]. Consequently it is expected that most resistance to developing MCE will be obtained for the alloys with highest Al contents. Indeed, deep pitting corrosion occurs and the lifetime of the alloy decreases as the amount of b-phase increases [10]. Our potentiodynamic study demonstrated that the AZ91D had high corrosion rate in deformed state that AM50 under stress. In addition, some recent studies
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Fig. 7. Typical microstructures of die casting AZ91D (a) and AM50 (b) alloys by optical microscope.
proved that the highest sensitivity to corrosion creep and corrosion fatigue [11] in a corrosive environment is observed in the alloy with highest Al content. It was also shown that the borate anions act as a corrosion inhibitors at the ®rst stage of creep [10]. 3.2. Electrochemical impedance spectroscopy measurements A corroding metal is, as a ®rst approximation, modelled as a simple electrochemical system consisting of a double-layer capacitance, a solution resistance and a charge transfer resistance. A system like this can be studied by using an AC signal, which can provide more information than a DC polarisation. Thus, applying a 5 mV sinusoidal potential through a potentiostatic circuit, the potential±current response plots provide the impedance values.
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Fig. 8. Nyquist plots at free immersion potential and under anodic and cathodic polarisation for AM50 (a) and AZ91D (b) magnesium alloys.
The impedance diagrams are recorded at the initial time (t 0 h) immediately after the stabilisation of the steady-state potential. In the present study, EIS measurements were performed at open circuit and also under polarisation. Thus, the typical Nyquist impedance plots showed in Fig. 8 are obtained for unstressed Mg alloys in the tetraborate solution at free immersion potential and under anodic and cathodic polarisation.
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The Nyquist plots of both magnesium alloys at open circuit exhibit two capacitive loops, one for high and intermediate frequencies and the other, the small one, for low frequencies. The ®rst capacitive loop is attributed at the charge transfer process. Thus, for the frequencies higher than 1 Hz, a resistor Rp and a capacitor Cdl in parallel can model the electrode/electrolyte interface. A partial data ®tting made with the Boukamp circuit equivalent software [12] for the charge transfer process produced the Rp (polarisation resistance) and Cdl (double-layer capacitance) values. The Rp of the charge transfer process is 207.7 and 374 X cm2 for unstressed AM50 and AZ91D alloys, respectively. The obtained capacitance values are 22.6 and 68 lF cm 2 , for AM50 and AZ91D, respectively and they are associated at a doublelayer capacity, Cdl , (50 lF cm 2 ). The slightly lower value of Cdl for the AM50 alloy implies the formation of a thick, protective ®lm on the electrode surface; the much lower Cdl values being already reported for other Mg-based alloys [13]. The second small capacitive loop is generally attributed at the masse transfer in the solid phase [14], which consists of the oxide/hydroxide layers. The EIS spectra (Fig. 8) obtained under anodic polarisation inside the potential range of the MgO formation exhibit one capacitive loop followed by a linear part for both magnesium alloys. As a ®rst remark we note the increase of the Rp , which is signi®cant in the case of AZ91D alloy. The increment of Rp suggests the layer growing on the electrode surface. The linear part of the Nyquist diagrams suggests the diusion process. The impedance data obtained under anodic polarisation was simulated with the Boukamp equivalent circuit software. Thus the values of Rp are 236:8 X cm2 for AM50 and 596:4 X cm2 for AZ91D, whereas the Cdl values are 63.84 and 127.31 lF cm 2 for AM50 and AZ91D, respectively. The equivalent circuit consists of a resistor (Rp ) in series with a constant phase element (CPE), the two being connected with a capacitor (Cdl ) in parallel. The CPE can be assumed to be a Warburg diusion according to the n values close to 0.5. Thus, under anodic polarisation, the corrosion process is controlled by the mass transfer of the corrosion products through the oxide layers [15]. The Nyquist plots for both Mg alloys obtained under cathodic polarisation present one loop capacitive (Fig. 8), which is attributed at water reduction. The EIS data under polarisation show that both magnesium alloys have quite similar behaviour in free immersion conditions and under anodic polarisation. Moreover, for the AZ91D alloy, the charge transfer parameters obtained under anodic polarisation in the active state indicate that the formation and growth of corrosion layer products is more signi®cant than in the case of AM50 alloy. The EIS measurements are further performed at the free immersion potential. For the stressed alloys, the Nyquist diagrams (Fig. 9) obtained at the potential of open circuit present a capacitive loop at high and intermediate frequencies as in the case of unstressed magnesium alloys. In addition, at low frequencies a small capacitive loop was always observed for all the samples, this loop being more or less reproducible. For a good clarity of the pictures in Fig. 9, the small capacitive loop was not plotted. The EIS data for the ®rst capacitive loop associated at the transfer charge process can be ®tted with the Boukamp circuit equivalent as in the case of
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Fig. 9. Nyquist plots at free immersion potential for stressed and unstressed AM50 (a) and AZ91D (b) magnesium alloys.
unstressed alloys, the equivalent circuit consisting of a Rp and a Cdl in parallel. The obtained 1=Rp (which is proportional to the corrosion rate) and Cdl values are plotted as a function of strain (Figs. 10 and 11), for both AM50 and AZ91D alloys. Moreover, Fig. 10 regroups the corrosion current values determined on the potentiodynamic polarisation curves (DC measurements). It can be seen that the maximal 1=Rp value is obtained for a strain of 1.94% and of 0.29% for AM50 alloy and
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Fig. 10. Corrosion rate (1=Rp ) from EIS measurements and corrosion current from potentiodynamic polarisation as a function of strain for AM50 (a) and AZ91D (b) magnesium alloys.
AZ91D, respectively. Thus the 1=Rp values obtained on the EIS measurements pass over the maximum as the amount of increase in the plastic deformation, in the same way as the corrosion current determined on the DC polarisation. Moreover, the Cdl values plotted as a function of strain (Fig. 11) also pass over a maximum. Fig. 11 also plots the stress±strain curves for both magnesium alloys, in order to emphasise the correlation between the strain hardening stages and the electrochemical parameters. The strain values for the maximum of Cdl are of 1.94% and 0.29% for AM50 and AZ91D, respectively. The corrosion rate, proportionally with 1=Rp , and
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Fig. 11. Double-layer capacity from EIS measurements as a function of the strain and the stress±strain curves for AM50 (a) and AZ91D (b) magnesium alloys.
also the Cdl values manifest a signi®cant MCE. These results show that the EIS measurement could be a rapid and ecient test to highlight the MCE. The electrode/electrolyte interface behaviour as a function of immersion time was also studied. The Nyquist and Bode diagrams obtained at dierent immersion times are plotted in Figs. 12, 13, 15 and 16 for unstressed alloys and also for dierent levels of stress for magnesium alloys. For both magnesium alloys (stressed and unstressed), the capacitive loop attributed to the charge transfer process generally increases with time, whereas the capacitive tail of the low frequencies does not show a uniform evolution. For unstressed AM50 alloy (Fig. 12), for a relatively short immersion time (1 and 3 h), the electrode/electrolyte interface does not undergo important modi®cations. For a time of over 24 h the polarisation resistance increases and so the surface protection is increased because of the formation of corrosion products layers. In
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Fig. 12. Nyquist (a,b) and Bode-phase (c) diagrams for the unstressed AM50 alloy at dierent immersion times.
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Fig. 13. Nyquist (a,b) and Bode-phase (c) diagrams for the AM50 alloy at strain of 2.26% at dierent immersion times.
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Fig. 14. Charge transfer resistance (Rp ) and double-layer capacity (Cdl ) as a function of immersion time for AM50 alloy at strain of (a) 0%; (b) 1.94%; (c) 2.13%; (d) 2.26%; (e) 3.56% and (f) 4.39%.
addition, the double-layer capacity increases as can be seen in Fig. 14(a), which plots the Rp and Cdl values obtained with the Boukamp ®tting equivalent circuit. For a time of under 96 h the corrosion process can be ®tted with only one time constant, which corresponds at the charge transfer. For t P 96 h the Bode-phase diagram presents two time constants and the equivalent circuit for the process ®tting changes. The same behaviour was practically observed for all the stressed AM50-alloy samples that can be seen, for example, in Fig. 13. This behaviour suggests that after an immersion time, which diers from a sample to another (but falls between 24 and 96 h), the corrosion mechanism changes; the charge transfer occurs through the corrosion products layers, which can also be suggested by the drastic decrease in the Cdl values. The equivalent circuit, which ®ts the impedance results, consists of two parallel RC circuits connected in series. The ®rst equivalent circuit corresponds to the oxide/hydroxide magnesium ®lm, which becomes quite protective so that the
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Fig. 15. Nyquist (a) and Bode-phase (b) diagrams for unstressed AZ91D alloy at dierent immersion times.
charge transfer process ®tted with the second equivalent circuit reaches a high polarisation resistance (Fig. 14). For the AZ91D alloy (Figs. 15 and 16) the evolution of the electrolyte/electrode interface is not similar. No change of corrosion mechanism is produced even for longer immersion times. The values of Rp and Cdl of the corrosion process obtained with a Boukamp partial ®tting are plotted in Fig. 17 for the AZ91D-alloy samples under study. Generally, both corrosion parameters increase with time, which suggests the formation of corrosion products layers. A suitable equivalent circuit relevant to the lower frequencies is hardly found due to the irregular behaviour of the Mg-alloy samples.
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(a)
(b)
Fig. 16. Nyquist (a) and Bode-phase (b) diagrams for the AZ91D alloy at strain of 0.29% at dierent immersion times.
4. Conclusion The plastic deformation eects on both potentiodyanmic polarisation curves and impedance diagrams have been studied in the case of two magnesium-based alloys (AM50 and AZ91D, respectively). For both magnesium alloys, it was obtained that the anodic current density determined on the potentiodynamic curves passes over a maximum as a function of the level of the plastic deformation (as predicted in the theory). It was also shown that
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Fig. 17. Charge transfer resistance (Rp ) and double-layer capacity (Cdl ) as a function of immersion time for AZ91D alloy at strain of (a) 0%; (b) 0.29%; (c) 0.45% and (d) 0.71%.
the disennobeling of potential manifests itself in both magnesium alloys. The correlation between the mechanochemical behaviour and strain hardening stages (intensive strain hardening and dynamic recovery) was also related to the known change of dislocation substructure during the plastic deformation. It was also shown that the AZ91D had a high corrosion rate in deformed state that AM50 under stress, while in the non-loading state, the corrosion rate was found to be higher for the AM50 alloy. This behaviour con®rms the MCE theory and also the alloys behaviour at creep. In addition, the inverse of resistance polarisation
1=Rp and the double-layer capacitance (Cdl ) determined on the EIS spectra pass also over the maximum as the amount of increase of the plastic deformation for both Mg alloys. The obtained results indicate that the EIS measurement could be a suitable and ecient experimental test in order to evidence the MCE. Thus experimental con®rmation of the correlation between the MCE and the strain hardening stages by both electrochemical techniques (DC polarisation and EIS, respectively) was provided. Only DC polarisation was generally used in order to de®ne the mechanochemical properties of dierent alloys that will be employed under stress corrosion conditions.
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[3] A. Eliezer, E.M. Gutman, E. Abramov, Ya. Unigosky, G. Agiv, E. Aghion, in: E. Aghion, D. Eliezer (Eds.), Magnesium 2000, Proceedings of the Second Israeli International Conference on Magnesium Science and Technology, 2000, p. 356. [4] C.F. Baker, New alloys and Shot Delivery System Developments, in: B.L. Mordike, F. Hehmann (Eds.), Magnesium Alloys and their Applications, Informationsgesellschaft, 1992, pp. 77±87. [5] M. Pourbaix, Atlas of Electrochemical Equilibria in Aqueous Solutions, Cebelcor, Brussesls, 1974, p. 141. [6] Metal Handbook, ninth ed., vol. 13, ASM, Metal Parks, Ohio, 1979, p. 742. [7] E.M. Gutman, G. Solovio, D. Eliezer, Corros. Sci. 38 (7) (1996) 1141. [8] O. Lunder, J.E. Lein, T.Kr. Aune, K. Nisancioglu, Corrosion 45 (1989) 741. [9] G.L. Makar, J. Kruger, J. Electrochem. Soc. 137 (2) (1990) 414. [10] E.M. Gutman, A. Eliezer, Ya. Unigovski, E. Abramov, Mechanoelectrochemical behavior and creep corrosion of magnesium alloys, Material Science Engineering A, vol. A302, 2001, p. 63. [11] A. Eliezer, E.M. Gutman, E. Abramov, Y. Unigovski, E. Aghion, Corrosion fatigue and corrosion creep of magnesium alloys, Proceedings of the International Congress: Magnesium 2000, Magnesium Alloys and their Applications, 27±28 September, 2000, Munich, Germany, p. 498. [12] Boukamp, Equivalent Circuit Software, Users Manual, The Netherlands University of Twente, 1988, pp. 6±26. [13] R. Udhayan, D.P. Prokash Bhatt, J. Power Sources 63 (1996) 103. [14] S. Turgoose, R.A. Cottis, The impedance response of ®lm-covered metals, in: Scully, Silverman, Kendig (Eds.), Electrochemical Impedance: Analysis and Interpretation, 1993, pp. 173±191. [15] C.H. (Raymond) Tsai, Analysis of EIS data for common corrosion processes, in: Scully, Silverman, Kendig (Eds.), Electrochemical Impedance: Analysis and Interpretation, 1993, pp. 37±52.
