Influence of the β-phase morphology on the corrosion of the Mg alloy AZ91

Corrosion Science 50 (2008) 1939–1953

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Influence of the b-phase morphology on the corrosion of the Mg alloy AZ91 Ming-Chun Zhao a,b,c, Ming Liu a,b, Guangling Song a,d, Andrej Atrens a,b,* a

Division of Materials, The University of Queensland, St. Lucia, Brisbane, Qld 4072, Australia Swiss Federal Laboratories for Materials Science and Technology, EMPA, Dept 136, Überlandstrasse 129, CH-8600 Dubendorf, Switzerland School of Material Science and Engineering, Central South University, Changsha 410083, China d CAST Cooperative Research Centre, Materials Engineering, University of Queensland, Brisbane, Qld 4072, Australia b c

a r t i c l e

i n f o

Article history: Received 20 December 2007 Accepted 11 April 2008 Available online 26 April 2008 Keywords: Magnesium Microstructure Weight loss Microgalvanic corrosion Hydrogen evolution

a b s t r a c t The influence of the microstructure, particularly the morphology of the b-phase, on the corrosion of Mg alloys has been studied using AZ91 as a model Mg alloy. The corrosion behaviour was characterized for five different types of microstructure produced by heat treatment of as-cast AZ91. The influence of microstructure can be understood from the interaction of the following three factors: (i) the surface films can be more or less effective in hindering corrosion and more or less effective in controlling the form of corrosion as uniform corrosion or localised corrosion, (ii) the second phase (the b-phase in AZ91) can cause micro-galvanic acceleration of corrosion and (iii) the second phase can act as a corrosion barrier and hinder corrosion propagation in the matrix, if the second phase is in the form of a continuous network. It is expected that these factors are important for all multi-phase Mg alloys because all known second phases have corrosion potentials more positive than that of the a-phase. A particular example of the corrosion barrier effect is provided by the fine (a + b) lamellar micro-constituent; when a b-phase plate nucleates this micro-constituent, the b-phase plate acts as a corrosion barrier. In contrast, nano-sized b precipitates, produced by aging, caused micro-galvanic corrosion acceleration of the adjacent a-phase. However, it is an important finding that the corrosion rate of the a-phase was decreased by the aging treatments that caused the precipitation of the nano-sized b particles. Ó 2008 Elsevier Ltd. All rights reserved.

1. Introduction Mg alloys are of significant interest to the automobile and aerospace industries due to their low densities and adequate strength/ weight ratios [1–3]. A significant limitation, however, is their corrosion performance [4–11]. It is important to understand the factors that influence their corrosion and to understand the conditions necessary to achieve adequate corrosion performance. Their corrosion behaviour has been investigated [12–22] over the past decade in order to facilitate their use in structural applications. Mg alloys are often multi-phase and their corrosion performance is influenced by their microstructure, in particular the amount and distribution of the different phases. AZ91 is one of the most popular of the cast magnesium alloys, with nominal composition Mg–9 wt% Al–1 wt% Zn. The AZ91 ascast microstructure has typically a primary a-phase matrix and a divorced eutectic distributed along the a-phase grain boundaries [6,14,16,23–27]. The divorced eutectic typically consists of large bphase particles and the eutectic a-phase. The eutectic a-phase is

* Corresponding author. Address: Division of Materials, The University of Queensland, St. Lucia, Brisbane, Qld 4072, Australia. Tel.: +61 733653748; fax: +61 733653888. E-mail address: [email protected] (A. Atrens). 0010-938X/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2008.04.010

super-saturated with Al and can transform, by discontinuous precipitation of the b-phase during cooling from the eutectic temperature, to form a fine lamellar arrangement of a + b [24–27]. There have been several studies [14,16,21,28] of the role of microstructure on the corrosion of AZ91. The corrosion behaviour of the a-phase and the b-phase are the foundations on which to build an understanding of the influence of the microstructure of multi-phase alloys like AZ91. Song et al [13,14] have shown that the free corrosion potential of the b-phase (1.3 V) is 0.3 V more positive than the free corrosion potential of the a-phase (1.6 V) in sodium chloride solutions. The a-phase corrodes due to its very negative free corrosion potential and there is the tendency for the corrosion rate of the a-phase to be accelerated by micro-galvanic coupling between the a-phase and the b-phase. The b-phase, in contrast, has a relatively lower corrosion rate, is a more efficient site for the cathodic reaction and may act as a barrier against corrosion propagation [6,13,14,29]. Thus, the b-phase has two different influences on the corrosion behaviour: the bphase can act as a galvanic cathode to accelerate corrosion and the b-phase can act as a corrosion barrier to hinder corrosion. The micro-galvanic corrosion acceleration is dependent on the anode (a)/cathode (b) area ratio whereas the b-phase acts to hinder corrosion if it is finely divided and continuous. The corrosion behaviour of AZ91 can be changed by changes in the amount

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and distribution of the b-phase. Moreover, there is this tendency for micro-galvanic corrosion in all multi-phase Mg alloys because all known second phases have corrosion potentials more positive than that of the a-phase [5]. The microstructure of AZ91, like other cast Mg alloys, is determined by the casting method, the rate of solidification and subsequent heat treatments used to improve mechanical properties. The distribution, configuration and size of the b-phase can be changed, which may result in different corrosion behaviour. Our prior work [21] showed that the second b-phase in an as-cast AZ91 ingot was in the form of large b particles and within a fine lamellar (a + b) micro-constituent. Homogenisation annealing heat treatment (HA) at 380 °C and 410 °C can cause the dissolution of the fine lamellar (a + b) micro-constituent without significantly altering the coarse b particles or can also cause the progressive dissolution of the coarse b particles. Solution heat treatment can dissolve the b-phase and the produce a microstructure consisting largely of the a-phase. The a-phase is super-saturated with Al for AZ91 in the solution heat treated condition and in the homogenisation annealed condition. Subsequent ageing at a lower temperature, e.g. 200 °C, precipitates the Al as fine b precipitates [15,30] in the a-phase, thereby decreasing the Al concentration in the a-phase to a low level (2% is the equilibrium Al concentration according to the Al–Mg phase diagram [31]). It is important to understand the corrosion behaviour of these different microstructures. The present work studied the influence of microstructure on corrosion using AZ91 as a model Mg alloy with the aim to build on the foundations of the prior studies [14,16,21,28]. Five different types of microstructures were produced as follows: (i) the as-cast microstructure of the AZ91 ingot, (ii) the as-cast microstructure as modified by homogenization annealing heat treatments at 380 °C and 410 °C, (iii) solution heat treatment to produce a largely a-phase microstructure, (iv) the solution heat treated, largely a-phase microstructure, aged to precipitate fine b precipitates and (v) homogenization annealed conditions aged to precipitate fine b precipitates. The aim was to carry out a comprehensive evaluation of the corrosion behaviour of all these microstructures using relatively short immersion tests of 48 h duration during which the rate of corrosion could be continuously monitored and the area corroding could be characterised at the completion of the 48 h immersion. Immersion tests of longer duration, 96 h, were carried out to determine if any new phenomena occurred over the longer time span. These longer term test took more effort and consequently they were carried out for a smaller pallet of alloys, the most important ones were chosen to check the trends that were becoming evident from the shorter term tests of 48 h duration. Detailed observations were carried out on the micro-corrosion morphology of a small selected number of samples. The corrosion of pure Mg has been included as a reference standard. Pure Mg has a homogenous single-phase microstructure consisting only of the a-phase. This comprehensive investigation into the influence of the microstructure on the corrosion behaviour of AZ91 has been carried out to provide a better understanding of the factors that influence corrosion of two-phase Mg alloys like AZ91.

2. Experimental procedure 2.1. Materials The AZ91D 0.14 wt% Mn, <0.002 wt% Cu Mg had the 0.008 wt% Mn,

had the composition of 8.26 wt% Al, 0.69 wt% Zn, <0.002 wt% Fe, 0.002 wt% Ni, <0.001 wt% Cr, <0.002 wt% Zr and balance Mg. The reference pure composition of 99.94 wt% Mg, 0.0137 wt% Si, 0.0045 wt% Fe, 0.00092 wt% Ni and 0.0066 wt% Al.

These two Mg alloys are high purity because they both contain low levels of the impurity elements Fe, Ni and Cu, so the results of the present research should be directly comparable with the prior studies [5,6,13–16]. The AZ91 specimens were cut from an as-cast ingot and were subjected to heat treatment as follows. Homogenization annealing (HA) heat treatments of as-cast specimens were carried out at 380 °C and 410 °C for 1–25 h followed in each case by air cooling (AC); the resulting specimens are designed as HA3805, HA3810, HA4105, HA4110 and HA4125, where the first two numbers indicate the temperature of the HA heat treatment (380 °C and 410 °C) and the second two numbers designate the time in hours of the HA heat treatment. The solid solution (SS) heat treatment consisted of heat-treating as-cast specimens at 410 °C for 100 h followed by water quenching (WQ), designed as 4100S (or SS condition). Solid solution treatment plus ageing (SA) consisted of the solid solution heat treatment followed by an ageing heat treatment at 200 °C for 5, 10, 24 and 48 h; the full designation was 4100S205A, 4100S210A, 4100S224A and 4100S248A, respectively; the short designation was SA205, SA210, SA224 and SA248. HA and aging was carried out to produce 4125HA225A (HA for 25 h at 410 °C, AC, aging for 25 h at 200 °C), 3801S205A (HA for 1 h at 380 °C, WC, aging for 5 h at 200 °C) and 3810S205A (HA for 10 h at 380 °C, WC, aging for 5 h at 200 °C). The microstructure was examined by optical microscopy and scanning electron microscopy (SEM) after metallographic preparation by mechanical grinding successively to 1200 grit SiC paper, polishing successively to 0.5 lm diamond, washing with distilled water, drying with warm flowing air and etching in 3% nital. X-ray diffraction (XRD) using CuKa radiation was used to characterize the phases present in these samples. 2.2. Corrosion evaluation The corrosion behaviour was evaluated using separate immersion tests of duration of either 48 h or 96 h, at room temperature, in 1 N NaCl aqueous solution, which was made with analytical grade reagent and distilled water. The corrosion rate was evaluated by measuring (i) the evolved hydrogen during corrosion in the 1 N NaCl solution and (ii) the weight lost by the specimen. The AZ91 specimens were encapsulated in epoxy resin so that a surface, with dimension 18 mm  27 mm, was exposed to the solution. The working surface was mechanically ground to 1200 grit SiC paper, washed with distilled water, dried with warm flowing air, dried in a desiccator for 1 to 2 days and weighed, to give the specimen weight before exposure, Wb. The specimen was horizontally immersed in 1500 ml of test solution and the hydrogen evolved during the corrosion experiment was collected in a burette above the corroding sample. The evolved hydrogen is a direct measure of the corrosion rate [4–6,32–33], as, in the overall magnesium corrosion reaction

Mg + Hþ + H2 O = Mg2þ + OH + H2

ð1Þ

one molecule of hydrogen is evolved for each atom of corroded magnesium. After the immersion test, the corrosion products were removed by immersion at room temperature for 5–10 min in a chromic acid cleaning solution, the sample was washed with distilled water, dried with warm flowing air, dried in a desiccator for 1–2 days and weighted to determine the sample weight after the immersion test, Wa. The weight loss data is presented as (weight loss = Wb  Wa [mg])/(specimen area [cm])/(exposure time [d]). The chromic acid cleaning solution composition was (200 g CrO3 + 10 g AgNO3)/L; previous work has shown that this chemical cleaning solution causes almost no weight loss for noncorroded AZ alloy specimens [16]. Separate blank experiments

M.-C. Zhao et al. / Corrosion Science 50 (2008) 1939–1953

showed that there was negligible weight change of the encapsulating epoxy resin by similar exposures to the 1 N NaCl solution or to the acid cleaning solution. The hydrogen evolution rate can be related to the weight loss rate, and the weight loss rate can be related to the corrosion rate, using the following conversions [5,6,20,34– 36]:

Weight loss rate ½mg=cm2 =d ¼ 1:085 2

ðHydrogen evolution rate ½ml=cm =dÞ;

ð2Þ 2

Corrosion rate ½mm=y ¼ 2:10 ðweight loss rate ½mg=cm =dÞ: ð3Þ The corrosion macro-morphology was examined using optical microscopy after the completion of each immersion test; this examination of the corrosion macro-morphology included recording the macro-morphology using macro-photography. Corrosion micro-morphology for the as-cast condition and representative HA conditions was examined using optical microscopy and using another series of metallographic prepared specimens, which were removed from the 1 N NaCl solution after various immersion times. After removal from the test solution, these specimens were rinsed with distilled water, dried using flowing air and observed using optical microscopy. 3. Results 3.1. Microstructure-as-cast The as-cast AZ91 microstructure had some micro-porosity as illustrated in Fig. 1a for the polished, un-etched surface. The dark spots in the low-magnification optical micrograph correspond to micro-pores. The occurrence of micro-porosity in as-cast AZ91 is consistent with [14]. It is possible to discern the microstructure as the a-Mg matrix with the second b-phase, because the difference in the hardness between the a-phase and the b-phase causes a difference in polishing rates. The detailed microstructure for as-cast AZ91, as revealed by etching, is presented in the SEM micrographs (Fig. 1a and b). The microstructure, as illustrated at low magnification in Fig. 1a, contained some dark regions in the vicinity of grain boundaries. These regions are identified as the eutectic a-phase with an Al content higher than the primary a-phase [14]; the dark appearance is attributed to different etching by the nital. The microstructure also contained large eutectic b-phase particles and an (a + b) microconstituent in a fine lamellar arrangement as illustrated at higher magnification in Fig. 1b. The fine lamellar (a + b) micro-constituent

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has the same outline as the eutectic a-phase and is the result of the discontinuous precipitation of the b-phase. These fine lamellar (a+b) micro-constituents and associated large b particles were interconnected and formed an interconnected network throughout the microstructure. In addition, the remainder of the microstructure contained large eutectic b-phase particles surrounded by the fine lamellar (a + b) micro-constituent as isolated entities in the a-phase matrix. The microstructure provides clues to the mechanism of formation of the fine lamellar (a + b) micro-constituent in the as-cast AZ91 microstructure. The fine lamellar (a + b) micro-constituent had a morphology similar to the typical lamellar pearlite colony in steel and hence the formation mechanism may be similar [37– 39]. There were the cases where it appeared that there was a b plate at the edge of the fine lamellar (a + b) micro-constituent and it appeared that this b plate had nucleated the fine lamellar (a + b) micro-constituent in a mechanism that was similar to the nucleation of lamellar pearlite in steels as illustrated in Fig. 2 [37–39]. The formation of the fine lamellar (a + b) micro-constituent occurs after solidification is complete and it occurs by a solidstate transformation of an a-Mg phase super-saturated in Al. During the transformation, Al is rejected from the a-Mg. If the Al diffusion is not sufficiently fast, Al accumulates at the interface and a b-plate forms when the Al content reaches a critical value. Therefore, the rejection of Al from the super-saturated a-Mg in the ascast AZ91 can lead to the formation of an interface b plate that can nucleate the fine lamellar (a + b) micro-constituent. After nucleation, the fine lamellar (a + b) micro-constituent grows due to the slow diffusion of Al in the Al-rich matrix. This proposed formation mechanism for the fine lamellar (a + b) micro-constituent is also consistent with instances [16] where the fine lamellar (a + b) micro-constituent forms when die-cast AZ91 is aged at a low temperature, e.g. 200 °C. In Mg die-casting, molten metal is rapidly injected into a steel die, solidification is relatively rapid and the formation of the fine lamellar (a+b) micro-constituent is prevented. During ageing of die-cast AZ91 at a low temperature, the fine lamellar (a + b) micro-constituent nucleates from large b particles. Therefore, it is reasonable that the fine lamellar (a + b) micro-constituent nucleates from a large b particle and grows away from the b particle as shown in Fig. 2 for the nucleation and growth of pearlite in steels. Fig. 2 presents the mechanism of transformation of austentite (c) to pearlite (the eutectoid microconstituent containing ferrite (a) and cementite in a lamellar arrangement). The pearlite is nucleated by a cementite plate at the nucleating grain boundary [37–39].

Fig. 1. The microstructure of as-cast AZ91.

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homogenization annealing temperature and/or a longer time produced a more homogeneous distribution of the b-phase throughout the microstructure. 3.3. Microstructure-solid solution (SS) The microstructure in the solid solution (SS) condition was homogenous, largely single a-phase and the b-phase was almost completely dissolved as illustrated in Fig. 4a. 3.4. Microstructure-solution heat treated and aged (SA)

Fig. 2. Pearlite nucleation at a cementite plate during the transformation of austenite (c) to ferrite (a) in steel. Pearlite is a lamellar arrangement of cementite (carbide, Fe3C, grey) and ferrite (a, white). As the austenite (c) transforms to ferrite (a), carbon is rejected from the ferrite and consequently the carbon concentration builds up at the a/c interface until a carbide forms at the interface. Subsequently pearlite grows from this carbide [37–39].

Furthermore, in the as-cast AZ91 microstructure, there were different orientations for the packets of the fine lamellar (a + b) micro-constituent, as shown in Fig. 1b, due to multiple nuclei. 3.2. Microstructure-homogenization anneal (HA) Homogenization annealing (HA) changed the as-cast AZ91 microstructure. The volume fraction of the b-phase decreased and its distribution became more homogeneous with an increase of the homogenization annealing temperature and/or homogenization annealing time. Fig. 3a shows that, after HA for 5 h at 380 °C (HA3805), most of the fine lamellar (a + b) micro-constituent had dissolved but the large b particles were largely unchanged. The microstructure consisted of the a-Mg matrix and the isolated large b particles associated with a small amount of the fine (a + b) lamellar micro-constituent remaining at the a-Mg grain boundaries. Fig. 3b shows that, after HA for 10 h at 380 °C (HA3810), the fine lamellar (a + b) micro-constituent had completely dissolved with little change of the large b particles; the microstructure consisted of the a-Mg matrix and isolated large b particles. Fig. 3c shows that, after HA for 5 h at 410 °C (HA4105), the fine lamellar (a + b) micro-constituent had completely dissolved and the large b particles had partly dissolved, showing a microstructure similar to that of HA3810, but the large b particle in HA 4105 were smaller in size due to part dissolution. Fig. 3d shows that, after HA for 10 h at 410 °C (HA4110), the fine (a + b) lamellar micro-constituent had completely dissolved and most of the b particles had dissolved, leaving relatively few fine b particles homogeneously distributed throughout the matrix. Fig. 3e shows that, after HA for 25 h at 410 °C (HA4125) the microstructure was similar to that of HA4110, the fine (a + b) lamellar micro-constituent had mostly dissolved and most of the b particles had dissolved, leaving relatively few fine b particles homogeneously distributed throughout the matrix and a trace along the a-phase grain boundaries of the fine (a + b) in the lamellar arrangement. HA3805 and HA3810 caused small changes to the large b particles and mainly caused the dissolution of the lamellar (a + b) micro-constituent. In contrast, HA4105 and HA4110 caused some dissolution of the large b particles and dissolution of the fine lamellar (a + b) micro-constituent. This indicates that a higher

The solid solution and aged microstructures, Fig. 4b–e, consisted of the a-Mg matrix with numerous nano-sized b precipitates, which were different in morphology to the b-phase particles in the as-cast and the HA conditions. The b precipitates in the solution treated and aged (SA) conditions where needle like, had an average size of 100 nm in width and had a relatively homogenous distribution throughout the microstructure. Furthermore, the volume fraction of b precipitates increased noticeably during aging and their size grew somewhat with an increase of the ageing time as shown in Fig. 4b–e. Nevertheless, all the b precipitates where needle like with most of the needles oriented in the same direction in each grain examined. 3.5. Microstructure-homogenization anneal and aged (SA) The 4125HA225A microstructure (Fig. 4f, HA for 25 h at 410 °C, AC, aging for 25 h at 200 °C) was somewhat similar to SA224 (Fig. 4d). The 4125HA225 microstructure consisted of the a-Mg matrix with numerous needle-shaped b precipitates, size of 100 nm in width, but in 4125HA225 the needles were oriented in a number of distinct directions compared with SA224 wherein the needles were largely oriented in the same direction. For the 3801S205A microstructure, Fig. 4g, the HA for 1 h at 380 °C did not cause a large change to the as-cast microstructure; the microstructure contained large eutectic b particles and the fine lamellar (a + b) micro-constituent similar to the as-cast structure. There were, in addition, numerous needle-shaped b precipitates, size of 100 nm in width, with an appearance of being oriented in a number of distinct directions. These were precipitated from the super-saturated a-phase solid solution during the aging treatment (5 h at 200 °C). For the 3810S205A microstructure, Fig. 4h, it would be expected that the HA for 10 h at 380 °C would have dissolved most of the fine lamellar (a + b) micro-constituent and would have dissolved some of the large b particles; however the microstructure, Fig. 4h did consist of the a-Mg matrix, large b particles and a significant amount of the fine lamellar (a + b) micro-constituent; the fine lamellar (a + b) micro-constituent might have reformed during the air cooling from the HA treatment temperature or during the aging treatment. In addition, there were, numerous fine interconnected b precipitates, size of 100 nm in width, with some appearance of being oriented. These were precipitated from the super-saturated a-phase solid solution during the aging treatment (5 h at 200 °C). 3.6. XRD The XRD spectra of specimens in the as-cast condition, representative HA conditions and a representative SA condition, Fig. 5, indicated that the second b-phase particles were Mg17Al12 because all of peaks in the XRD spectra corresponded to either Mg or Mg17Al12, although the second b-phase was reported to sometimes be present as Mg4Al3 [40–41]. This indicates that AZ91 had a matrix of a-Mg grains with the second b-phase consisting of the inter-

M.-C. Zhao et al. / Corrosion Science 50 (2008) 1939–1953

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Fig. 3. The microstructure after homogenization annealing (HA).

metallic Mg17Al12 for the as-cast condition and for the various heat treatments. Fig. 5 also indicates that the relative peak heights for the peaks associated with a-Mg were different for the different heat treated conditions. The largest difference was that between the XRD spectrum for the as-cast condition and the other spectra. This difference is associated with the changes in the microstructure. In the as-cast condition, there was three forms of a-Mg: (i) primary a, (ii) eutectic a-phase and (iii) a-phase in the fine a + b lamellar arrangement. In contrast, after heat treatment, there was largely only the primary a-phase left, as the other micro-constituents had dissolved during the heat treatment.

3.7. Hydrogen evolution Fig. 6 presents the corrosion rate as measured by hydrogen evolution. The measured values for pure Mg were in good agreement with prior measurements [5,6,14] as were the measurements for as-cast AZ91D [5,6,14,15]. This provides confidence in the validity of the data. Most specimens exhibited an increase in hydrogen evolution rate with increasing immersion time. In contrast, the hydrogen evolution volume for pure Mg, also depicted in Fig. 6, increased linearly with exposure time. The increasing hydrogen evolution meant that the evolved volumes during the 96 h immersions

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Fig. 4. The microstructures of SS and aged conditions.

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Fig. 5. XRD patterns of specimens in the as-cast condition, representative HA conditions and a representative SA condition. The peaks marked with a diamond are from Mg17Al12, all other peaks are from a-Mg.

were significantly larger for the 46 h immersion; compare Fig. 6a with b. The data of hydrogen evolution volumes are presented in two groups. Group A includes the as-cast condition, the HA conditions and the solid solution condition. Group B includes the solid solution condition, the aged conditions and pure Mg. Group A had higher hydrogen volumes. For the 48 h immersion tests for Group A, the highest hydrogen evolution volume was for the AZ91D specimen HA for 10 h at 380 °C and the lowest hydrogen evolution volume was for the solid solution condition. The hydrogen evolution volume of all the group A conditions in Fig. 6a can be ranked as a decreasing series: HA for 10 h at 380 °C > HA for 5 h at 380 °C > HA for 5 h at 410 °C > HA for 10 h at 410 °C  HA for 25 h at 410 °C > the as-cast condition > the solid solution condition. For the 48 h immersion test for Group B, the highest hydrogen evolution volume was for the specimen SA for 48 h at 200 °C and the lowest hydrogen evolution volume was for pure Mg. The hydrogen evolution volume for all the group B conditions in Fig. 6a can be ranked as a decreasing series: SA for 48 h at 200 °C > the solid solution condition > SA for 24 h at 200 °C > HA for 25 h at 410 °C + aged for 25 h at 200 °C  SA for 10 h at 200 °C > SA for 5 h at 200 °C > HA for 1 h at 380 °C + aged for 5 h at 200 °C  HA for 10 h at 380 °C + aged for 5 h at 200 °C > pure Mg. Similar results were obtained for 96 h immersion test for Groups A and B, that is, the hydrogen evolution volume of all the group A conditions in Fig. 6b can be ranked as a decreasing series: HA for 10 h at 380 °C > HA for 5 h at 380 °C > the solid solution condition > the as-cast condition. The hydrogen evolution volume of all the group B conditions in Fig. 6b can be ranked as a decreasing

series: the solid solution condition > SA for 48 h at 200 °C > SA for 24 h at 200 °C > SA for 10 h at 200 °C > SA for 5 h at 200 °C > pure Mg. However, there an exception for the solid solution condition: (i) the hydrogen evolution volume of the solid solution condition was a little lower than the as-cast condition until about 60 h immersion and thereafter there was a reverse tendency, and (ii) the hydrogen evolution volume of the solid solution condition was a little lower than the SA for 48 h at 200 °C until about 80 h immersion and thereafter there is a reverse tendency. 3.8. Weight loss data Fig. 7 presents weight loss measurements. Fig. 7a presents the weight loss data for 48 h immersion and Fig. 7b presents the data for 96 h immersion. The corrosion rates were significantly higher in the 96 h immersions; this is consistent with the accelerating corrosion rate as shown in Fig. 6. Furthermore, the trends revealed in the hydrogen evolution data of Fig. 6 are reproduced in the weight loss measurements presented in Fig. 7. It appeared that the fine lamellar a + b micro-constituent in ascast AZ91 was somewhat beneficial to corrosion. This is proposed because the as-cast AZ91, with a microstructure containing the fine lamellar a + b micro-constituent, had a lower hydrogen evolution rate than all the HA conditions without the lamellar a + b microconstituent in the microstructure, Figs. 6 and 7. Also the HA3805 condition (HA for 5 h at 380 °C) had a lower hydrogen evolution rate than the HA3810 condition (HA for 10 h at 380 °C); for these two conditions, the only microstructure difference was that HA3805 condition had a small amount of the fine lamellar (a + b)

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M.-C. Zhao et al. / Corrosion Science 50 (2008) 1939–1953 16 as-cast 3805-H 3810-H 4105-H 4110-H 4100-S 4125-H

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Fig. 6. Hydrogen evolution for the various microstructures during immersion in 1 N NaCl for the stated period.

micro-constituent in the a-Mg grain boundary whereas the HA3810 condition did not. The weight loss rate is directly related to the hydrogen evolution rate as shown in Fig. 8. This direct correlation is expected from the overall magnesium corrosion reaction (1) that one molecule of hydrogen is evolved for each atom of corroded magnesium and from prior research [4–6,8,32,33,36]. Fig. 8 provides a cross plot of the independent measurements of weight loss rate and hydrogen evolution rate, for the 48 h immersion tests and for the 96 immersion tests. The line in Fig. 8 is the plot of Eq. (2). Fig. 6 shows that (i) the weight loss rate data is consistent with the hydrogen evolution rate data and (ii) that the relationship between these independent, experimental measurements follows the theoretical expectation from the corrosion reaction (1) as embodied in Eq. (2). This provides confidence in the experimental measurements, including confidence (i) that the acid cleaning procedure has removed all the corrosion products, (ii) that the acid cleaning proce-

dure has removed negligible Mg metal and (iii) that there was negligible weight change of the encapsulating epoxy resin due to exposure to the 1 N NaCl solution or to the acid cleaning solution. 3.9. Macroscopic corrosion morphology The corrosion of AZ91 initiated as localized corrosion at some sites on the surface and subsequently expanded over surface. The advance of the corrosion over the surface was different for different conditions. At the end of the 48 or 96 h immersion time, the active corrosion area was covered by a thick layer of corrosion products and was different for the various conditions as presented by the macro-photographs (27 mm  18 mm) in Figs. 9 and 10. In most cases, the area corroding after 96 h immersion was larger than after 48 h immersion. In marked contrast to the local corrosion of AZ91, the corrosion for pure Mg immersed in 1 N NaCl was uniform corrosion; there

M.-C. Zhao et al. / Corrosion Science 50 (2008) 1939–1953 8 HA3810

48 h immersion

7

2

Weight loss rate (mg/cm day)

HA3805

6 5 HA4105

4 HA4110 HA4125

3

SA248

as-cast SS

2 SA210

SA224

4125HA225

SA205

1

3801S205 3810S205

pure Mg

0 1

2

Weight loss rate (mg/cm day)

20

2

3

4HA3810 5

6

7

8

9

10

11

13 14 15 9612h immersion

15 HA3805 SS

10

SA248 as-cast SA224

5 SA205

SA210

pure Mg

0 1

2

3

4

5

6

7

8

9

Fig. 7. Weight loss data for the various microstructures during immersion in 1 N NaCl for the stated period.

48 h 96 h Y=1.085X

2

Weight loss rate (mg/cm day)

20

15

10

5

0 0

10

20

30

1947

set of samples in the as-cast and HA conditions exposed to the test solution. After a few minutes immersion in the test solution, some small hydrogen bubbles were evolved from all over the whole working surface. After a longer immersion time (1 h), several sites suffered corrosion in the form of localized corrosion. The corrosion in these local areas enlarged in area slowly and a large number of hydrogen bubbles were evolved from these local areas of corrosion. Fig. 11a presents the micro-corrosion morphology for the as-cast condition after immersion in the test solution for 16 h; the samples had the same surface preparation as the samples for the hydrogen gas collection. Optical microscope examination led to the conclusion that the fine lamellar a + b micro-constituent and the coarse b particles provided some action as corrosion barriers; the adjacent a-Mg matrix had undergone the majority of the corrosion, with some corrosion in the fine lamellar a + b micro-constituent. The same surface preparation was used as was used in the corrosion exposures with hydrogen collection, with the necessary consequence that the quality of the optical micrographs could not be expected to equal that of a metallographic prepared sample. To obtain good quality micrographs the corrosion exposures were repeated with samples that were metallographic polished but were not etched; ground successively to 1,200 grit SiC paper, polished successively to 0.5 lm diamond, washed with distilled water and dried with warm flowing air. The corrosion morphologies for these metallographic polished samples were essentially the same as for the samples subjected to only grinding to 1200 SiC paper. Optical microscope examination of the metallographic polished samples exposed in 1 N NaCl is presented in Fig. 11b. Fig. 11a and b show the following three different corrosion cases for the as-cast condition: (i) corrosion of the adjacent a-Mg with no significant corrosion of the fine lamellar a + b micro-constituent and the coarse b particles (ii) corrosion of the a-Mg up to the fine lamellar a + b micro-constituent, with the appearance of the corrosion of the a-Mg stopped at the interface of the fine lamellar a + b micro-constituent and (iii) corrosion proceeding into the fine lamellar a + b microconstituent, with both phases having been corroded to a similar extent. In marked contrast, the HA conditions had no continuous micro-constituents and had only isolated b particles as shown in Fig. 3. Fig. 12 presents optical microscope photos of the metallographic polished samples exposed to 1 N NaCl for (HA for 5 h at 380 °C) and (HA for 10 h at 410 °C). Some isolated b particles were in the forefront of the corrosion advance. However, the distance between the b particles in these two microstructures was large, and the corrosion advance was not at all blocked by the b-phase particles, although the b particles were themselves not significantly corroded.

4. Discussion

2

H2 evolution volume, ml/cm day Fig. 8. Cross plot of the independent measurements of weight loss rate and hydrogen evolution rate, for 48 h and 96 h immersion tests. The line is a plot of Eq (2).

were no preferential sites for corrosion. At the end of the 48 and 96 h immersion periods, the whole exposed surface of pure Mg was homogenously covered by a layer of corrosion products. The fact that the whole surface of pure Mg was corroding at the end of both the 48 and 96 h exposure periods is consistent with the linear rate of hydrogen evolution presented in Fig. 6. 3.10. Micro-corrosion morphology The influence of the microstructure on the corrosion behaviour was studied by observing the micro-corrosion morphologies of a

4.1. Macroscopic corrosion morphology The corrosion of AZ91 in the various heat treated conditions was localised corrosion, initiated at some sites on the surface and subsequently expanded over surface. The advance of the corrosion over the surface was different for different conditions, as presented by the macro-photographs (27 mm  18 mm) in Figs. 9 and 10. In most cases, the area corroding after 96 h immersion was larger than after 48 h immersion. This provides part of the explanation for (i) the increasing rate of evolved hydrogen with immersion time (Fig. 6) (ii) the greater evolved hydrogen volume after 96 h immersion (Fig. 6) and the greater weight loss after 96 h immersion (Fig. 7). The fact that there was still un-corroded surface after 96 h immersion indicates (i) that the air formed film provides a certain resistance to corrosion initiation, (ii) that the air formed

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Fig. 9. Macro-appearance (27 mm  18 mm) of the corrosion morphology after 48 h immersion for different microstructures.

film provides a certain resistance to corrosion propagation and (iii) that steady state corrosion conditions had not been reached. In marked contrast to the local corrosion of AZ91, the corrosion for pure Mg immersed in 1 N NaCl was uniform corrosion and there was a linear rate of hydrogen evolution, Fig. 6. 4.2. Corrosion mechanism – homogenization anneal Figs. 9 and 10 indicate that the corroded area was approximately the same for all the homogenization annealed conditions after both 48 h and 96 h immersion, although the corroded area after 96 h immersion was greater than after 48 h immersion. As it is possible to exclude the influence of the surface area, surface state and the initiation of corrosion, it is possible to consider the influence of the microstructure on the propagation of corrosion as presented in Figs. 6 and 7. AZ91 is a two-phase alloy, in which different micro-constituents may each form a micro-cell. The corrosion behaviour and morphology of different microstructure conditions can be explained on the basis of the microstructure characteristics and the corresponding corrosion performance of the a-Mg matrix, the fine lamellar a+b micro-constituent, coarse b particles and fine b particles. When the samples are immersed in the test solution, hydrogen evolution preferentially starts at some sites and then the corrosion attack advances. The b-phase can have two roles in corrosion: (i) as a corrosion barrier and (ii) as a galvanic cathode. The b-phase itself is more stable in the test solution and it has a lower corrosion rate. However, the cathodic hydrogen evolution on the b-phase surface is

much faster than that on the surface of the a-phase and thus the bphase is a more effective cathode [14]. A competition between micro-galvanic corrosion and the corrosion barrier effect can explain the corrosion [14]. If the b-phase is present as a small fraction, the b-phase serves mainly as a galvanic cathode and accelerates the overall corrosion of the a matrix. If b-phase fraction is high, the bphase may act mainly as a barrier against the corrosion of the a matrix. In the as-cast condition, Fig. 11 shows that the a-Mg matrix adjacent to coarse b particles with lack of the surrounding lamellae (a + b) eutectic has undergone the majority of the corrosion, with there being less corrosion of the b particles and the fine lamellar a + b micro-constituent. The corrosion is initiated and gradually advanced by the dissolution of the a-Mg matrix adjacent to the region of the continuous lamellar micro-constituent plus coarse b particles. When the propagation of the corrosion attack reaches the region of the continuous (a + b) micro-constituent plus coarse b particles, the corrosion is retarded to a certain extent as shown in Fig. 11. In this case, even though the a-Mg matrix surrounding the (a + b) plus coarse b particles is corroded, some should be under the continuous lamellae (a + b) plus coarse b particles as the a-Mg matrix in the top layer has corroded. In fact, the corrosion damage as above-mentioned is quite common for AZ91D in a corrosive environment. There is evidence for some b particles being undermined from corroding areas of AZ91D under an atmospheric corrosion [28]. Section 3.1 evaluated the formation of the fine lamellar (a + b) micro-constituent. This evaluation led to the realisation that a

M.-C. Zhao et al. / Corrosion Science 50 (2008) 1939–1953

1949

Fig. 9 (continued)

plate of the b-phase could be the cause of the nucleation of the fine (a + b) lamellae micro-constituent. Such b-phase plates could be the mechanism by which corrosion is stopped by the fine (a + b) lamellae micro-constituent. It is proposed that, when a b-phase plate nucleates this micro-constituent, the b-phase plate acts as a corrosion barrier In the 380 °C HA conditions (5 h and 10 h), isolated b particles are distributed throughout the microstructure, the distance between the isolated b particles is large and the corrosion attack is not effectively blocked by this type of isolated b particles. However, the corrosion attack of the a-phase is accelerated by the micro-galvanic coupling with the b-phase. This means that the a-Mg matrix and the isolated coarse b particles in the 380 °C HA conditions (5 h and 10 h) undergo micro-galvanic corrosion in addition to the spontaneous dissolution of the a-Mg in the aqueous solution (owing to its very negative open circuit potential). The HA condition at 10 h for 380 °C shows marked micro-galvanic corrosion due to complete elimination of lamellar eutectic (a + b) in the microstructure leaving only the isolated b particles. On the other hand, the predominant factors determining the rate of galvanic corrosion include the anode-to-cathode area ratios. The HA treatment for 5 h at 410 °C had caused some dissolution of the large b particles and the size of these b particles was smaller compared to the HA treatment for 10 h at 380 °C, therefore, the galvanic corrosion effect was decreased. After the HA treatment for 10 h at 410 °C and the HA treatment for 25 h at 410 °C, the

microstructure contained only some fine isolated b particles, such that the smaller cathode area caused less galvanic corrosion to accelerate the spontaneous corrosion of the a-Mg. Similar corrosions rates are derived by the similar microstructures for the HA treatment for 10 h at 410 °C and the HA treatment for 25 h at 410 °C. The above analysis indicates that the b-phase acts mainly as a galvanic cathode for the various HA conditions during the initial period of corrosion, as the b fraction is low. During corrosion, the a grains dissolve preferentially whereas most of the b-phase particles remain except for isolated b-phase particles which are undermined and which fall out. Hence the b fraction on the sample surface cannot remain constant during corrosion. It changes from the initial b surface fraction characteristic of the sample surface to a steady state surface fraction determined by the preferential corrosion of the a grains and the undermining of the b-phase particles. As a consequence there is the phenomenon that the hydrogen evolution volume for the as-cast condition with continuous (a + b) plus b-phase in the microstructure is slightly higher than that for the solid solution condition with homogenous single a phase in the microstructure in the beginning after 48 h immersion, but become slightly lower with the advance of the corrosion attack at 96 h immersion, as shown in Fig. 6 (compare ‘‘as-cast” and ‘‘4100-S”); this is attributed to an increased surface coverage by the b-phase for the as-cast condition and a higher barrier effect by the b-phase in the area that is corroding. This barrier effect

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Fig. 10. Macro-appearance (27 mm  18 mm) of the corrosion morphology after 96 h immersion for different microstructures.

Fig. 11. Micro-corrosion morphologies after immersion for as-cast condition.

M.-C. Zhao et al. / Corrosion Science 50 (2008) 1939–1953

1951

Fig. 12. Typical corrosion morphology for HA3805 and HA4110 conditions for metallographic polished specimens; arrows indicate isolated fine b particles.

explains the change in relative corrosion rates of the as-cast and SS conditions; the increased area of corrosion explains why their increased absolute corrosion rates after 96 h immersion are both higher than after 48 h immersion. There was a similar effect between 4100-S and 4100-248 samples, Fig. 6, and a similar explanation is proposed. 4.3. Corrosion mechanism – solid solution Many studies [12,16,29] have reported that there is a decrease in corrosion rate for Mg–Al alloys with an increasing Al content in the range 2–11% Al in a-Mg. These studies are consistent. However, comparison of the corrosion rates for Mg–Al alloys from the work of Song et al [16] indicate that the corrosion rates for the Mg–Al alloys are all greater than the corrosion rate for pure Mg as measured by Song et al [13,14]. The data in the present work are consistent with these prior measurements [12–14,16,29]. Figs. 6 and 7 shows that the solid solution condition of AZ91 had a corrosion rate higher than that of pure Mg; this comparison is particularly emphasized by Fig. 6. The solid solution condition of AZ91, with a homogenous single phase, has a higher Al content than pure Mg with a homogenous single phase and had a higher corrosion rate. This implies that the corrosion behaviour of pure Mg is characteristically different from that of AZ91. Figs. 9 and 10 shows that the corrosion was uniform for pure Mg immersed in 1 N NaCl and that there were no preferential sites for corrosion, i.e. the whole exposed surface of pure Mg was homogenously covered by a layer of corrosion film. This was markedly different from solution treated AZ91 which showed local corrosion. This may explain the observation that the solid solution condition of AZ91 has a corrosion rate higher than that of pure Mg. The steady state surface film on pure Mg does not hinder further corrosion as the corrosion rate is constant with exposure time, Fig. 6 whereas the prior work indicates that the film on Mg–Al alloys is partially protective, because the corrosion rate decreases with increasing Al content for Mg-Al alloys heat treated to be homogeneous single a-phase alloys. However, the film on the solid solution treated AZ91 (4100-S) allowed localised corrosion, and the corrosion rate increased with immersion time, Fig. 6. 4.4. Corrosion mechanism – aged conditions The corrosion of the aged samples SA205, SA210, SA224 and SA248 is consistent with the increased micro-galvanic corrosion acceleration of an a-phase that has an intrinsic corrosion rate significantly lower than that of the SS condition and approaching that of pure Mg. The aged samples SA205, SA210, SA224 and SA248 were aged for 5, 10, 24 and 48 h at 200 °C. Increased aging time increased the amount of the b-phase, Fig. 4b–e which would provide

increased micro-galvanic corrosion acceleration. Furthermore, for these samples, there was an additional effect due to the fact that the area corroded increased in the sequence SA205, SA210, SA224 and SA248 for the 48 h immersions, Figs. 9 and 10 shows a similar tendency after 96 h immersion. The increased measured weight loss in the sequence SA205, SA210, SA224 and SA248 can be attributed to a combination of these two effects: (i) the increased micro-galvanic corrosion directly increased the corrosion rate and (ii) the greater surface area corroding means that the sample has a greater weight loss. However, it is not entirely clear why the a-phase would have a corrosion rate significantly lower than that of the SS condition and approaching that of pure Mg, particularly as aging would be expected to decrease the Al content of the a-phase to 2% if equilibrium is approached as predicted by the phase diagram [31]. The microstructure for 4125HA225A is very similar to that of SA224, so it is to be expected that they have similar corrosion rates. Similarly, the microstructures of 3801S205 and 3810S205 are similar as are their corrosion rates. Their low corrosion rates are attributed to the fact that their microstructures contain a continuous network of the b-phase, which would be expected to act as a barrier to corrosion propagation. 4.5. Longer immersion tests The longer-term immersion tests for 96 h, Figs. 6, 7 and 10, exhibited the same trends and corrosion mechanism as shown by the shorter 48 h immersion tests. This provided confidence that the deductions were indeed valid that were being drawn from the short tests. 4.6. Corrosion and mechanical properties A homogenization annealing (HA) heat treatment was proposed in our prior work [21] for mechanical property enhancement for AZ91; HA for 10 h at 410 °C causes an improvement in hardness, ultimate tensile strength and ductility without loss of corrosion properties, i.e. corrosion properties are similar to those of the ascast condition. The improvement of the mechanical properties is due to the absence of a continuous easy crack path which is present in the as-cast microstructure due to the interconnected network throughout the microstructure of the fine lamellar (a + b) micro-constituent and associated large b particles. In the as-cast microstructure, the interconnected network throughout the microstructure of the fine lamellar (a + b) micro-constituent and associated large b particles provides a benefit for corrosion by acting as a corrosion barrier. However, this same microstructure causes inferior mechanical properties due to the continuous easy crack path through this network.

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In contrast, the peak strength condition (HA for 10 h at 410 °C) has isolated fine b particles homogeneously distributed through the a-phase matrix. There is no degradation of the mechanical properties by a continuous easy crack path, and consequently the mechanical properties are superior to those of the as-cast condition. As presented above, the isolated fine b particles do not lead to an obvious loss of corrosion because they act as a small cathode connected to a large anode (a-Mg matrix) so that the peak strength condition (HA for 10 h at 410 °C) has negligible micro-galvanic corrosion. There have been a number of investigations [16,42–49] on the influence of aging on the mechanical properties of AZ91. These have investigated the classic two step SA heat treatment. The two heat treatment steps of SA are (i) first a solution heat treatment at a higher temperature followed by quick cooling and (ii) second an ageing heat treatment at a lower temperature. These studies show that the optimum combination of strength and ductility are obtained for aging 5–10 h at 200 °C. Longer aging times leads to a decrease in mechanical strength. Fig. 7 show that the corrosion rate also increases with longer aging time. Our prior work proposed [21], for mechanical property enhancement, the homogenization annealing heat treatment: HA for 10 h at 410 °C. Fig. 7 indicates that the corrosion rate for this proposed HA heat treatment was greater than for SA205 and SA210, i.e. solution heat treatment at 410 °C for 100 h, water quench and an ageing heat treatment at 200 °C for 5 and 10 h. However, HA only needs reheating up to a temperature for several hours followed by air cooling rather than the two heat treatment steps of SA, which has (i) first a solution heat treatment at a higher temperature followed by quick cooling and (ii) second an ageing heat treatment at a lower temperature. As a consequence, HA has potentially a cost lower than the SA treatment.

5. Conclusions 1. The fine lamellar (a + b) micro-constituent, in as-cast AZ91, may be nucleated by a b-phase plate. This b-phase plate may be effective as a corrosion barrier to subsequent corrosion. 2. Most AZ91 specimens exhibited an increase in corrosion rate with increasing immersion time for immersion times of up to 96 h in 1 N NaCl. Part of the reason for the increasing corrosion rate is the increase in area corroded with increasing immersion time. This means that long term immersion tests lasting two weeks to a month may be needed to measure steady state corrosion rates. 3. Relatively short immersion tests like 48 h are valuable as they reveal the same trends as longer immersion tests 96 h, and most probably also correlate with steady state corrosion conditions. 4. The weight loss rate is directly related to the hydrogen evolution rate by the relationship: weight loss rate [mg/cm2/d] = 1.085 (Hydrogen evolution rate [ml/cm2/d]), as expected from the overall magnesium corrosion reaction that one molecule of hydrogen is evolved for each atom of corroded magnesium. 5. There was considerable difference in corrosion rates and areas corroded for the various microstructures. 6. AZ91SS (i.e. AZ91 subjected to a solid solution heat treatment of 100 h at 410 °C and water quenched to have a microstructure largely single a-phase) had a corrosion rate higher than that of pure Mg despite the fact that the corrosion covered part of the surface for AZ91SS whereas corrosion was uniform for pure Mg and covered the whole surface of the pure Mg. Thus the higher corrosion rate for AZ91SS is not an effect of surface area corroding. An impurity explanation is ruled out by the fact that both AZ91 and pure Mg were high purity.

7. The microstructural influence of corrosion can be understood from the interaction of the following three factors: (i) the surface film formed on the surface of the a-phase can be more or less effective in hindering corrosion and more or less effective in controlling the form of corrosion as uniform corrosion or localised corrosion, (ii) the second phase (the b-phase in AZ91) can cause micro-galvanic acceleration of corrosion of the matrix a-phase and (iii) the second phase (the b-phase in AZ91) can act as a corrosion barrier and hinder the corrosion propagation in the matrix a-phase. 8. It is expected that the same factors are important for all multiphase Mg alloys because all known second phases have corrosion potentials more positive than that of the a-phase.

Acknowledgements This work was supported by the ARC Center of Excellence, Design of Light Alloys. CAST CRC was established under, and is supported in part by, the Australian Government’s Cooperative Research Centres scheme. DH StJohn is thanked for valuable discussion on the formation of the fine lamellar (a + b) microconstituent. References [1] K.Y. Sohn, J.A. Yurko, F.C. Chen, J.W. Jones, J.E. Allison, in: Proc. Auto Alloys II, San Antonio, TX, The Material, Metals and Minerals Society, 1988, p. P81. [2] R.S. Busk, Magnesium Products Design, 499, Marcel Dekker, New York, 1987. [3] B.B. Clow, Advanced Materials and Processes 150 (1996) 33. [4] G. Song, A. Atrens, Advanced Engineering Materials 1 (1999) 11–33. [5] G.L. Song, A. Atrens, Advanced Engineering Materials 5 (2003) 837. [6] G. Song, Advanced Engineering Materials 7 (2005) 563. [7] N. Winzer, A. Atrens, G. Song, E. Ghali, W. Dietzel, K.U. Kainer, N. Hort, C. Blawert, Advanced Engineering Materials 7 (2005) 659–693. [8] G. Song, A. Atrens, Advanced Engineering Materials 9 (2007) 177–183. [9] J.X. Jia, G.L. Song, A. Atrens, Corrosion Science 48 (2006) 2133–2153. [10] N. Winzer, A. Atrens, W. Dietzel, G. Song, K.U. Kainer, Materials Science and Engineering A 472 (2008) 97–106. [11] N. Winzer, A. Atrens, W. Dietzel, V.S. Raja, G. Song, K.U. Kainer, Materials Science and Engineering A (2008), doi:10.1016/j.msea.2007.11.064. [12] R. Ambat, N.N. Aung, W. Zhou, Corrosion Science 42 (2000) 1433. [13] G.L. Song, A. Atrens, X.L. Wu, B. Zhang, Corrosion Science 40 (1998) 1769. [14] G.L. Song, A. Atrens, M. Dargusch, Corrosion Science 41 (1999) 249. [15] R.K. Singh Raman, Metallurgical and Materials Transactions A 35A (2004) 2525. [16] G. Song, A.L. Bowles, D.H. St. John, Materials Science and Engineering A 366 (2004) 74. [17] W. Zhang, S. Jin, E. Ghali, R. Trmblay, M. Shehata, E. Es-Sadiqi, Advanced Engineering Materials 8 (2006) 973. [18] N. Winzer, A. Atrens, W. Dietzel, G. Song, K.U. Kainer, Materials Science and Engineering A 466 (2007) 18. [19] G. Ben-Hamu, D. Eliezer, K.S. Shin, Materials Science and Engineering A 447 (2007) 35. [20] M.C. Zhao, M. Liu, G. Song, A. Atrens, Advanced Engineering Materials 10 (2008) 104–111. [21] M.C. Zhao, M. Liu, G. Song, A. Atrens, Advanced Engineering Materials 10 (2008) 93–103. [22] M . Liu, D. Qiu, M.C. Zhao, G. Song, A. Atrens, Scripta Materialia 58 (2008) 421– 424. [23] C. Suman, SAE Transactions 99 (5) (1990) 849. [24] A.K. Dahle, Y.C. Lee, M.D. Nave, P.L. Schaffer, D.H. St. John, Journal of Light Metals 1 (2001) 61. [25] C.J. Bettles, Materials Science and Engineering A 348 (2003) 280. [26] M.X. Zhang, P.M. Kelly, Scripta Materialia 48 (2003) 647. [27] A. Srinivasan, J. Swaminathan, U.T.S. Pillai, K. Guguloth, B.C. Pai, Materials Science and Engineering A (2007), doi:10.1016/j.msea.2007.09.059. [28] .Y. Wan, J. Tan, G. Song, C. Yan, Metallurgical and materials Transactions 37A (7) (2006) 2313–2316. [29] O. Lunder, J.E. Lein, T.K. Aune, K. Nisancioglu, Corrosion 45 (1989) 741–748. [30] S. Prem Kumar, S. Kumaran, T. Srinivasa Rao, Materials Science and Technology 20 (2004) 835. [31] ASM International, ASM Handbook Online, 2003. [32] G.L. Song, A. Atrens, D.H. St. John, J. Nairn, Y. Lang, Corrosion Science 39 (1997) 855–875. [33] G.L. Song, A. Atrens, X. Wu, B. Zhang, Corrosion Science 40 (1998) 1769– 1791.

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