See discussions, stats, and author profiles for this publication at: https://www.researchgate.net/publication/223194981 Corrosion Behaviour of Magnesium/Aluminium Alloys in 3.5wt.% NaCl Article in Corrosion Science · March 2008 DOI: 10.1016/j.corsci.2007.11.005 CITATIONS READS 567 2,371 6 authors, including: Maria Concepción Merino Raul Arrabal Complutense University of Madrid Complutense University of Madrid 93 PUBLICATIONS 4,077 CITATIONS 141 PUBLICATIONS 7,510 CITATIONS SEE PROFILE SEE PROFILE Fernando Viejo Endzhe Matykina Industrial University of Santander Complutense University of Madrid 90 PUBLICATIONS 3,395 CITATIONS 131 PUBLICATIONS 6,534 CITATIONS SEE PROFILE Some of the authors of this publication are also working on these related projects: Corrosion Protection of Light Alloys View project SSF-Magnesium View project All content following this page was uploaded by Fernando Viejo on 21 November 2018. The user has requested enhancement of the downloaded file. SEE PROFILE Available online at www.sciencedirect.com Corrosion Science 50 (2008) 823–834 www.elsevier.com/locate/corsci Corrosion behaviour of magnesium/aluminium alloys in 3.5 wt.% NaCl A. Pardo a,*, M.C. Merino a, A.E. Coy a, R. Arrabal b, F. Viejo b, E. Matykina b b a Departamento de Ciencia de Materiales, Facultad de Quı́micas, Universidad Complutense, 28040 Madrid, Spain Corrosion and Protection Centre, School of Materials, The University of Manchester, P.O. Box 88, Sackville Street, Manchester M60 1QD, United Kingdom Received 13 July 2007; accepted 7 November 2007 Available online 24 November 2007 Abstract Corrosion behaviour of commercial magnesium/aluminium alloys (AZ31, AZ80 and AZ91D) was investigated by electrochemical and gravimetric tests in 3.5 wt.% NaCl at 25 °C. Corrosion products were analysed by scanning electron microscopy, energy dispersive X-ray analysis and low-angle X-ray diffraction. Corrosion damage was mainly caused by formation of a Mg(OH)2 corrosion layer. AZ80 and AZ91D alloys revealed the highest corrosion resistance. The relatively fine b-phase (Mg17Al12) network and the aluminium enrichment produced on the corroded surface were the key factors limiting progression of the corrosion attack. Preferential attack was located at the matrix/b-phase and matrix/MnAl intermetallic compounds interfaces. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: A. Magnesium; B. Polarization; B. Weight loss; C. Corrosion 1. Introduction AZ (Mg–Al–Zn) system, containing 2–10% Al with minor additions of Zn and Mn, is the most widely used among Mg–Al alloys. They are characterised by low cost of production and also by relatively good corrosion resistance and satisfactory mechanical properties from 95 to 120 °C [1]. However, corrosion performance of magnesium alloys has been a major obstacle to their growth in structural applications, despite their high stiffness/weight ratio, ease of machinability, high damping capacity, castability, weldability and recyclability [2]. These properties make them suitable for aerospace and automotive applications, where light metals are mandatory in order to reduce weight and greenhouse gas emissions. Over the past 30 years, alloy design development, new surface treatments and improved * Corresponding author. Tel.: +34 1 3944348; fax: +34 1 3944357. E-mail address: [email protected] (A. Pardo). 0010-938X/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2007.11.005 knowledge of corrosion mechanisms has lead to an increase of the real and potential applications of magnesium alloys [3]. Corrosion resistance of Mg alloys depends on many factors: (i) environment, (ii) alloy composition and microstructure, and (iii) properties of the film developed in the medium to which they are exposed. Concerning the environment, corrosion resistance of magnesium alloys in chloride containing solutions greatly depends on pH and Cl concentration, with no significant influence of oxygen concentration [4,5]. In general, magnesium and its alloys dissolve at a very low rate in alkaline or poorly buffered sodium chloride solutions, where the pH can increase, due to the formation of a partially protective Mg(OH)2 layer [4]. On the other hand, chloride ions promote rapid attack in neutral aqueous solutions and even higher in acidic solutions [4,6]. It is also common to find higher corrosion rates with increasing Cl ion concentration at all pH levels [6]. The corrosion of magnesium alloys in non-oxidizing neutral or alkaline chloride solutions at free corrosion potential typically initiates as irregular pits, which spread 824 A. Pardo et al. / Corrosion Science 50 (2008) 823–834 laterally and cover the whole surface [7,8]. However, the mechanism is different [9,10] from the auto-catalytic pitting experienced by stainless steels [11], since there does not seem to be much tendency for deep pitting, possibly as a result of pH increase and magnesium hydroxide film formation. However, this is not always true, since there is a significant influence of the microstructure on the corrosion mechanism, especially in two-phase magnesium alloys. On the topic of alloy composition and microstructure it is known that alloying elements not only modify the mechanical properties of magnesium, but also impart a significant impact on the corrosion behaviour. Alloying elements can form secondary particles, which are noble with respect to the magnesium matrix, thereby facilitating corrosion, or enrich the corrosion product thereby possibly inhibiting the corrosion rate [12]. In general, it is reported that increasing Al concentrations in Mg–Al alloys have a beneficial effect on the corro- sion behaviour in chloride media [13–15], but the specific mechanism and influence of Al is still not well understood. For instance, Lunder [16] found that increasing concentrations of 2–8 wt.% Al in die cast AS, AM and AE alloys decrease the corrosion rate in 5% NaCl, however, the decrease in corrosion rate appears to be related to a decrease in the impurity level with increasing Al content (it is known that Al reduces the iron tolerance limit from 170 wt.-ppm to 20 wt.-ppm [17]). More recent data compared the corrosion rates of high purity alloys in 3% NaCl [18]. Results showed that HP–Mg5Al alloy had a corrosion rate significantly higher than that of HP–Mg due to the micro-galvanic corrosion acceleration of the corrosion of the a-phase by the adjacent b-phase. On the other hand, the high purity two phase commercial Mg alloys (MEZ, AM60 and AZ91D) had similar corrosion rates to that of HP–Mg, despite the fact that these commercial alloys each had a two phase microstructure which would cause a Table 1 Nominal composition of the materials tested Material Chemical composition (wt.%) Al Zn Mn Si Cu Fe Ni Ca Zr Others Mg (99%) AZ31 AZ80 AZ91D 0.006 3.1 8.2 8.8 0.014 0.73 0.46 0.68 0.03 0.25 0.13 0.30 0.019 0.02 0.01 0.01 0.001 <0.001 <0.001 <0.001 0.004 0.005 0.004 0.004 <0.001 <0.001 <0.01 <0.001 <0.30 <0.30 <0.30 <0.008 Fig. 1. SEM micrographs of: (a) unalloyed Mg; (b) AZ31 alloy; (c) AZ80 alloy and (d) AZ91D alloy. A. Pardo et al. / Corrosion Science 50 (2008) 823–834 2. Experimental procedure Fig. 2. Mass loss versus time for the materials immersed in 3.5 wt.% NaCl solution. Table 2 Kinetic laws of the materials immersed in 3.5 wt.% NaCl solution Material Kinetic law: y = b t; [y (mg/cm2), t (h)] Mg (99%) AZ31 AZ80 AZ91D y = 13.03t y = 2.31 101t y = 3.10 103t y = 3.70 103t a (r2 = 0.99) (r2 = 0.99) (r2 = 0.98) (r2 = 0.98) 0 6 t 6 16 0 6 t 6 240 0 6 t 6 240 0 6 t 6 240 -1.1 AZ31 -1.2 E SSE (V) -1.3 -1.4 -1.5 -1.6 -1.7 -1.8 -7 10 -6 10 -5 10 -4 10 -4 10 10 -3 10 -2 -3 10 2 I (A/cm ) b -1.2 AZ91D -1.3 -1.4 ESSE (V) galvanic acceleration of the corrosion of the a-phase [19,20]. Therefore, apart from impurities and Al concentration it is also necessary to study the effect of phase distribution. Normally, in Mg–Al alloys the aluminium is partly in solid solution and partly precipitated in the form of Mg17Al12 along grain boundaries as a continuous phase as well as part of a lamellar structure. It is known that Mg17Al12 exhibits a passive behaviour over a wider pH range than either of its component aluminium and magnesium [21], and it was found that the distribution of the Mg17Al12 phase determines the corrosion resistance of the Mg–Al alloys [22]. Song et al. [17] suggested that the bphase mainly served as a galvanic cathode and accelerated the corrosion process of the a-matrix if the volume fraction of b-phase was small; however for a high volume fraction, the b-phase might act as an anodic barrier to inhibit the overall corrosion of the alloy. It has also been reported that the ratio of the b-phase to the surrounding Al-rich-a or eutectic-a, which can be about 12 wt.% compared to 1.5 wt.% in the grain centre [23], plays an important role in the formation of galvanic cells. Thus, Raman [24] found that an increase in the relative size of the b-phase at the expense of the Al-rich-a area, increasing the cathode to anode area ratio, results in an increase in the localized corrosion. However, the existence of a surrounding Al-rich-a area with independent electrochemical identity appears to be greatly dependant on the thermal history and composition of the alloy. There are also contradictory results depending on aluminium concentration in the a-phase. For instance, Lunder et al. [21] proposed that aluminium accelerate anodic dissolution below 8% whereas it decreases corrosion above 10%. However, Song [25,26] found that an increase of Al in the a-phase improves the corrosion resistance. Moreover, results obtained for rapid solidified alloys with different additions of Al and no precipitation of b-phase revealed a positive effect of Al on corrosion resistance [27]. In summary, it seems that Al improves the corrosion resistance of Mg–Al alloys in chloride environments but its effect is strongly influenced by the b-phase morphology and distribution. The aim of this work was to study the corrosion behaviour of three commercial magnesium/aluminium alloys in marine environment. The effect of immersion time and Al concentration on corrosion resistance was monitored by electrochemical and gravimetric tests, scanning and electron microscopy (SEM), energy dispersive X-ray analysis (EDX) and low-angle X-ray diffraction (XRD). 825 -1.5 -1.6 -1.7 2.1. Test materials Chemical compositions of the tested magnesium alloys, namely AZ31, AZ80 and AZ91D, are listed in Table 1. Unalloyed Mg was used as the reference material. Pure Mg and AZ31 alloy were fabricated in wrought condition, -1.8 -7 10 -6 10 -5 10 10 -2 2 I (A /cm ) Fig. 3. Double cyclic polarization curves after the immersion in 3.5 wt.% NaCl solution during 1 h for the alloys: (a) AZ31 and (b) AZ91D. 826 A. Pardo et al. / Corrosion Science 50 (2008) 823–834 whereas AZ80 (chill casting) and AZ91D alloys were manufactured by ingot casting process. All the materials were supplied by Magnesium Elektron Ltd. cases, the tests were performed in duplicate to guarantee the reliability of the results. 2.3. Gravimetric measurements 2.2. Preparation and surface characterization Gravimetric measurements were performed using rectangular (Mg and AZ31) and cylindrical (AZ80 and AZ91D) specimens of working area 16 cm2 immersed in 3.5 wt.% NaCl solution at room temperature (pH 5.6, 25 °C). Prior to the tests, specimens were measured and weighed. Once the test was finished for each immersion time, the samples were extracted, rinsed with isopropyl alcohol, dried in hot air, and then weighed again in order to calculate the mass gain per unit surface area. Corrosion rate was calculated from the mass losses per unit of surface area, calculated from the expression (Mi Mf)/A, where Mi is the initial mass, Mf the final mass and A the exposed surface area. In order to obtain further information, pH changes of the test solution during the gravimetric experiments were recorded with a standard pH-meter. For the same reason, hydrogen evolution was measured using a simple procedure described elsewhere [9]. For metallographic characterization, samples were wet ground through successive grades of silicon carbide abrasive papers from P120 to P2000 followed by diamond finishing to 0.1 lm. Two etching reagents were used: (a) Nital, 5 ml HNO3 + 95 ml ethanol, to reveal the constituents and general microstructure of Mg, AZ80 and AZ91D materials and (b) Vilella reagent, 0.6 g picric acid + 10 ml ethanol + 90 ml H2O, to reveal grain boundaries of AZ31 alloy. The constituents were examined by SEM using a JEOL JSM-6400 microscope equipped with Oxford Link EDX microanalysis hardware. For low-angle XRD studies a Philips X́Pert diffractometer Ka Cu = 1.54056 Å) was used. Prior to corrosion tests, specimens were wet ground to grade P1200 grit, followed by rinsing with isopropyl alcohol in an ultrasonic bath and drying in warm air. In all -1.0 -1.0 Mg ESSE (V) -1.1 -1.2 -1.2 -1.3 -1.3 -1.4 -1.4 -1.5 -1.5 -1.6 -1.6 -1.7 -1.7 -1.8 -8 10 -7 10 AZ31 -6 10 -5 -4 10 10 2 I (A /cm ) -3 10 -2 10 -1.8 -8 10 ESSE (V) -5 10 -4 10 -4 10 10 -3 10 -2 -3 10 -1.0 AZ91 1h 1d 3d 7d -1.1 -1.2 -1.3 -1.3 -1.4 -1.4 -1.5 -1.5 -1.6 -1.6 -1.7 -1.7 -1.8 -8 10 -6 10 I (A /cm ) AZ80 -1.2 -7 10 2 -1.0 -1.1 1h 1d 3d 7d -1.1 1h 1d -7 10 -6 10 -5 -4 10 10 2 I (A /cm ) -3 10 -2 10 -1.8 -8 10 1h 1d 3d 7d -7 10 -6 10 -5 10 10 2 I (A /cm ) Fig. 4. Influence of the immersion time on the anodic behaviour for the materials tested in 3.5 wt.% NaCl solution. -2 A. Pardo et al. / Corrosion Science 50 (2008) 823–834 2.4. Electrochemical measurements DC electrochemical measurements were performed using specimens with 2 cm2 of working area exposed to 3.5 wt.% NaCl at room temperature (25 °C). An AUTOLAB model PGSTAT 30 potentiostat connected to a three-electrode cell was used for the electrochemical measurements; the working electrode was the test material, whereas the counter and reference electrodes were graphite and Ag/AgCl (SSE), respectively. Solution concentration inside the reference electrode compartment was KCl 3 M, with a potential of 0.197 V with respect to hydrogen. Anodic, cyclic and double cyclic polarization measurements were carried out at a scan rate of 0.3 mV/s, from 100 mV to +400 mV with respect to the corrosion potential (Ecorr). For cyclic polarization, the scan direction was reversed when the samples reached the anodic corrosion current of 5 mA and potential was scanned back to the starting potential. 2.5. Characterization of corrosion products After the tests, magnesium alloys were observed by SEM in order to study the morphology and evolution of corrosion products formed on the material surface. Also, the composition of the corrosion layer was analysed by low-angle XRD. ited linear kinetics of mass loss associated with magnesium matrix dissolution. Unalloyed Mg presented the highest mass loss and it was completely dissolved after 16 h of immersion with a mass loss of 196 mg/cm2. On the other hand, the addition of aluminium increased notably the corrosion resistance. Thus, 3 wt.% Al in AZ31 alloy reduced the mass loss to 55.2 mg/cm2 after 10 days of immersion, and 8–9 wt.% Al in AZ80 and AZ91D alloys diminished the mass loss to 0.75 and 1.05 mg/cm2 at the end of the test, respectively. Table 2 shows the kinetic laws, calculated from the experimental data, corresponding to the growth of corrosion products layer generated during the gravimetric test. In each case, data have been approximated by a linear equation y = b t, where ‘‘y” coordinate represents the mass gain in units of mg/cm2 and ‘‘t” is the immersion time in hours. Aluminium reduced the magnesium reactivity, i.e. 3 wt.% Al reduced corrosion rate by a factor of 56 and 8– 9 wt.% Al did it by a factor of 300. In summary, Mg and AZ31 alloy manifested very low corrosion resistance in seawater, whereas AZ80 and AZ91D alloys presented a moderate resistance after 10 days of immersion. a -1.0 3.1. Microstructural characterization -1.2 Ecorr (VSSE) 3. Results -1.1 Mg AZ31 AZ80 AZ91 -1.3 -1.4 -1.5 -1.6 0 b Fig. 2 shows the mass loss versus time of tested materials immersed in 3.5 wt.% NaCl solution. All materials exhib- 2 3 4 5 6 7 5 10 Mg AZ31 AZ80 AZ91 4 10 2 3 10 2 10 1 10 0 3.2. Gravimetric results 1 Time (d) Rp (Ω·cm ) Fig. 1 shows the SEM microstructure of tested materials. Unalloyed Mg only reveals equiaxial grains with average size of 45 lm, whereas the presence of aluminium (3.1 wt.%) and manganese (0.25 wt.%) in AZ31 alloy favours the formation of intermetallic phases, mainly in form of MnAl2 inclusions (Figs. 1a, b). On the other hand, AZ80 and AZ91D ingot casting alloys disclose two different types of solidification microstructures. AZ80 alloy shows a biphasic microstructure with a a-Mg solid solution and a discontinuous precipitation in lamellar form of bphase (Mg17Al12), which starts at the grain boundaries from solid solution (Fig. 1c). AZ91D alloy reveals a-Mg primary dendrites and eutectic a-Mg/Mg17Al12 in the interdendritic region, which appears completely in divorced form with respect to solid solution (Fig. 1d). Indeed, the rapid solidification obtained by chill casting process used for AZ80 alloy promotes a refinement of the microstructure compare to AZ91D alloy. Finally, MnAl2 intermetallic inclusions are also found for AZ80 and AZ91D alloys. 827 1 2 3 4 5 6 7 Time (d) Fig. 5. Influence of the immersion time for materials tested in 3.5 wt.% NaCl solution on: (a) polarization resistance (Rp) and (b) corrosion potential (Ecorr). 828 A. Pardo et al. / Corrosion Science 50 (2008) 823–834 Fig. 6. SEM micrographs of alloys immersed in 3.5 wt.% NaCl solution for 10 days: (a) AZ31; (b) AZ80 and (c) AZ91D. Fig. 7. SEM micrographs of alloys immersed in 3.5 wt.% NaCl solution for 2 h: (a) AZ31; (b) AZ80 and (c) AZ91D. A. Pardo et al. / Corrosion Science 50 (2008) 823–834 3.3. Electrochemical results In order to obtain further information, the gravimetric study was supplemented with potentiodynamic polarization measurements for different times of immersion in 3.5 wt.% NaCl solution. Double cycle polarization curves for AZ31 and AZ91D alloys after immersion for 1 h reveal a different behaviour compared to materials with a stable passive layer (Fig. 3). The onset of pitting is not visible in the forward scan of both cycles, since pitting potential (Epit) is very close to Ecorr and, consequently, it should be expected that these alloys suffer pitting attack immediately after their immersion in the aggressive media at the open circuit potential. However, the current densities of the cathodic branch and, therefore, the growth of corrosion products on the material surface are quite high, suggesting general corrosion attack as the main mechanism of degradation. Fig. 4 discloses the evolution of the electrochemical behaviour of each material with increasing immersion time. The high dissolution rate observed for unalloyed magnesium limits this study, since it was dissolved after 1 day of immersion according to gravimetric results. Nevertheless, longer immersion times shifted its curves to greater current densities, possibly associated with the formation of a low-protective oxide layer, which does not impede the corrosion attack progression. AZ80 and AZ91 alloys revealed a slight shift of Ecorr towards more noble values, and corrosion current densities (icorr) were approximately 829 one order of magnitude lower. AZ31 presented the most significant change on the Ecorr values, possibly due to the high reactivity of this alloy which promotes formation of a thick corrosion layer, limiting the attack progression and increasing Ecorr values (Fig. 5a). Concerning polarization resistance (Rp) values, they remained low after 7 days of immersion (between 102 and 104 X cm2) (Fig. 5b). However, an increase of Rp values and, therefore, a better corrosion behaviour, it is observed for Mg and AZ31 materials, possibly related to the nucleation and growth of a thick and semi-protective layer of corrosion products. 3.4. Characterization of corrosion products The SEM micrographs of the surface of the tested materials after 10 days of immersion in 3.5 wt.% NaCl solution are shown in Fig. 6. Unalloyed magnesium is not presented due to its complete dissolution after 24 h of exposure. Although AZ31 alloy revealed higher corrosion resistance than unalloyed magnesium, its surface was completely covered by a thick and uneven film of corrosion products. An increase of aluminium content up to 8–9 wt.% reduced the reactivity of the magnesium alloys and, thereby, AZ80 and AZ91D alloys presented a lower degree of corrosion. Initial stages of corrosion for AZ31, AZ80 and AZ91D alloys immersed in 3.5 wt.% NaCl revealed localised corrosion around MnAl2 inclusions and b-phase interfaces, which form a galvanic couple with the surrounding Mg matrix (Fig. 7). Polarization tests did not show a clearly Fig. 8. Cross-section BSE image and corresponding X-ray maps of Mg, Al and O for the AZ31 alloy immersed in 3.5 wt.% NaCl for 10 days. 830 A. Pardo et al. / Corrosion Science 50 (2008) 823–834 defined pitting potential, so the corrosion process should be associated with a localised attack, but not to a typical pitting attack reported for materials with a stable passive layer, such as aluminium alloys, stainless steels, etc. X-ray map analysis of the cross section of AZ31 alloy after immersion for 10 days in the corrosive medium confirmed the formation of a thick corrosion layer with an irregular thickness between 200 and 400 lm and mainly constituted by magnesium oxides and/or hydroxides (Fig. 8). AZ80 alloy revealed a discontinuous and cracked corrosion film of lower thickness than AZ31 alloy (<200 lm) (Fig. 9). Furthermore, aluminium enrichment was found at the metal/corrosion layer interface, possibly produced by preferential Mg dissolution during the immersion test. The formation of Al oxide species is likely to be one of the main reasons for the improved corrosion resis- tance exhibited by this alloy. Additionally, it can be observed that the front of the corrosion attack was stopped when it reached the lamellar aggregate of b-phase (Mg17Al12). Thus, the higher corrosion resistance of AZ80 alloy may be associated with the presence of aluminium and its direct or indirect intervention in corrosion mechanism, consisting of two steps: (a) formation of a semi-protective Al-rich oxide layer and (b) reduction of the corrosion progression near to the lamellar b-phase aggregate. Finally, Figs. 10 and 11 show the study of the corrosion layer formed for AZ91D alloy after immersion for 10 days. Like for the AZ80 alloy, corrosion layer was irregular, cracked, and 200 lm thick (Fig. 10). However, the positive effect of Al on the corrosion performance cannot be explained in the same way as above, since aluminium Fig. 9. Cross-section BSE image, profile line and corresponding X-ray maps of Mg, Al and O for the AZ80 alloy immersed in 3.5 wt.% NaCl for 10 days. A. Pardo et al. / Corrosion Science 50 (2008) 823–834 831 rosion mechanism of magnesium needs further investigation, it is generally reported that the following anodic and cathodic reactions take place in marine environments: MgðsÞ ! Mg2þ ðaqÞ þ 2e ðanodic reactionÞ ð1Þ 2H2 O þ 2e ! H2 þ 2OH ðaqÞ ðcathodic reactionÞ ð2Þ 2+ Firstly, magnesium dissolves and Mg (aq) cations are produced (Eq. (1)), possibly through intermediate steps involving monovalent magnesium ion [14,30]. Secondly, magnesium dissolution is accompanied by hydrogen evolution (Eq. (2)), since magnesium in neutral and low pH aqueous solutions is well below the region of water stability. Finally, pH raises along with the cathodic reaction due to the formation of OH–, which favours the formation of Mg(OH)2(s) according to Pourbaix diagram (Fig. 13a) [31]. Thus, the overall reaction could be expressed as Mg2þ ðaqÞ þ 2OH ðaqÞ ! MgðOHÞ2 ðsÞðcorrosion productÞ ð3Þ Fig. 10. Cross-section BSE image and corresponding X-ray map of Al for the AZ91D alloy immersed in 3.5 wt.% NaCl for 10 days. enriched eutectic aggregate solidifies at the interdendritic spaces resulting in an a-primary phase with lower Al content than the predicted by the equilibrium phase diagram. Therefore, (a) there was not an appreciable Al enrichment observed on the surface after the Mg dissolution during the corrosion attack. Although, as mentioned above, (b) the presence of a relatively fine b-phase network, which for this particular alloy is extended all along the surface, partially impeded the corrosion attack (Fig. 11). Low-angle XRD study (incident angle of 1°) of the corrosion layer produced after immersion in 3.5 wt.% NaCl for 10 days revealed brucite (Mg(OH)2) as the main corrosion product, and its peaks exhibited higher intensity for unalloyed Mg and AZ31 alloy due to the formation of a thicker corrosion layer during the severe attack that both materials suffered (Fig. 12). Unlike other works, evidences of magnesite (MgCO3) formation, due to CO2 presence in the atmosphere, were not found [28]. 4. Discussion When unalloyed magnesium is exposed to the atmosphere or aqueous solutions, a grey oxide (mainly magnesium hydroxide, brucite) forms on its surface, which is stable for the basic range of pH values [29]. Nevertheless, in presence of chloride anions, this surface film breaks down and magnesium appears unprotected. Although cor- Since pH changes and H2 evolution can be easily measured, it is feasible to study the corrosion evolution of magnesium with these parameters [19]. Fig. 13b shows the variation of pH for all tested materials immersed in 3.5 wt.% NaCl (pH 5.6). For commercially pure Mg, test solution revealed the highest pH raise, pH >10 after immersion for 1 h. On the other hand, even though aluminium reduces the magnesium reactivity, pH values were registered between 8 and 9.5 after 1 h, the pH increasing up to 10 for longer immersion times. Therefore, these technique may be convenient and reliable for measurement of initial corrosion of magnesium and its alloys, however it becomes gradually insensitive for long times, since the solution is saturated with Mg(OH)2, and concentrations of Mg2+ and OH cannot further increase while corrosion is still proceeding [19]. Concerning H2 evolution Fig. 13c compares mass loss obtained by gravimetric measurements with H2 evolution, assuming that the dissolution of one atom of magnesium generates one molecule of hydrogen (Eqs. (1) and (2)). Results confirmed the higher corrosion resistance of AZ80 and AZ91D alloys and the high dissolution rate of pure Mg in the test solution. In all cases, corrosion rate estimated by H2 evolution was lower than the one obtained by gravimetric tests, possibly due to small particles falling out of the surface after having been undermined by the corrosion process [9]. Therefore, although pH and H2 evolution provide information regarding the corrosion of magnesium alloys, it is still suggested to perform gravimetric measurements, since they are the most basic and reliable method. In general, it is reported that the presence of aluminium in single-phase Mg–Al alloys has a beneficial effect on the corrosion behaviour in chloride media [13–15]. In the present work, this behaviour has been clearly revealed for the AZ31 alloy, which exhibited higher corrosion resistance than commercially pure Mg. However, aluminium influence in two-phase Mg–Al alloys is still not well understood 832 A. Pardo et al. / Corrosion Science 50 (2008) 823–834 Fig. 11. Cross-section BSE image of detail of the corrosion layer, profile line and corresponding X-ray maps of Mg, Al and O for the AZ91D alloy immersed in 3.5 wt.% NaCl for 10 days. Fig. 12. Low-angle XRD (incidence angle: 1°) of tested materials immersed in 3.5 wt.% NaCl solution for 10 days. since depends on many factors, such as impurities, b-phase morphology and distribution, aluminium content and size of the a-Mg primary dendrites and the eutectic-a, aluminium enrichment on the corrosion product layer [12], etc. Magnesium and magnesium alloys present an active position in both the electromotive force (EMF) and the galvanic series for seawater. As a result, galvanic interactions between magnesium and other metals are a serious concern. Thus, the corrosion resistance of Mg–Al alloys will depend on the presence of alloy elements acting as active cathode, such as manganese, iron, etc. The most detrimental potential cathodes in Mg–Al alloys are iron-rich and MnAl second phase particles [22]. The former were not found in any of the studied alloys, whereas MnAl2 intermetallic compounds were identified as preferential sites for localised corrosion in AZ31, AZ80 and AZ91D alloys. A. Pardo et al. / Corrosion Science 50 (2008) 823–834 833 a b 12 Mg AZ31 AZ80 AZ91 pH 11 10 9 8 0 1 2 3 4 5 6 7 Time (d) c Fig. 13. (a) Pourbaix Diagram for magnesium; (b) pH variation and (c) mass loss and H2 evolution rates for all materials tested in 3.5 wt.% NaCl for 10 days. Previous investigations reported that distribution of the b-phase (Mg17Al12) determines the corrosion resistance of the Mg–Al alloys [17,19–22,24,32]. According to this, for a low volume fraction, the b-phase serves as a galvanic cathode and accelerates the corrosion process of the aphase, whereas large volume fractions act as an anodic barrier and the overall corrosion diminishes [26,33]. Moreover, it has also been reported that the ratio of the b-phase to the surrounding Al-rich-a or eutectic-a plays an important role in the formation of galvanic cells. Apart from the microstructure, the composition of the corrosion layer may improve the corrosion resistance. For instance, Nordlien et al. [34] observed that alumina components forms a continuous skeletal structure in an oxide layer of the magnesium alloy, which passivating properties are much better than Mg(OH)2 and MgO layers. Therefore, keeping in mind these observations and the results obtained in the present work, the following considerations about the corrosion mechanism of studied Mg–Al alloys can be summarized: (i) an increase of the aluminium concentration in the nominal composition of the alloys reduced the activity of the pure Mg in 3.5 wt.% NaCl, with AZ80 and AZ91D showing the highest corrosion resistance; (ii) an independent electrochemical Al-rich-a area surrounding the b-phase was not observed by metallographic characterization; (iii) AZ91D alloy exhibits a slightly worse corrosion behaviour than AZ80 alloy, although the aluminium content in its nominal composition is slightly higher and both alloys reveal the presence of relatively fine b-phase network segregated in the matrix alloy, which acts as an effective barrier against progression of corrosion attack. As a result, the main reason for this different corrosion behaviour observed for the AZ80 and AZ91D alloys could be explained in terms of aluminium enrichment on the metallic surface during the magnesium matrix dissolution. According to this, aluminium content for both alloys in the a-Mg primary dendrites was estimated by EDX. Chill casting employed for AZ80 alloy promotes high aluminium contents in the solid solution (13.3 at.%), meanwhile ingot casting, with lower solidification rate, promotes a coarse dendritic microstructure for the AZ91D alloy and less aluminium content in the solid solution (8.4 at.%). Hence, AZ80 alloy revealed aluminium enrichment on the metallic surface after immersion in 834 A. Pardo et al. / Corrosion Science 50 (2008) 823–834 3.5 wt.% NaCl for 10 days, whereas this aluminium-rich oxide layer was not observed for the AZ91D alloy (Figs. 9 and 11). 5. Conclusions 1. Corrosion attack of all tested magnesium alloys occurs at the a-magnesium matrix/Al–Mn and Mg17Al12 intermetallic compounds interfaces, by means of the formation of galvanic couples. Later, the nucleation and growth of an irregular and less protective corrosion layer consisted mainly of Mg(OH)2 is produced from a-Mg matrix. 2. An increase of the aluminium concentration in the nominal composition of the alloys reduced the activity of the pure Mg in 3.5 wt.% NaCl. However, the AZ31 alloy (3.1 wt.% Al) still presented high corrosion rates. 3. The principal cause of higher corrosion resistance of AZ80 alloy is associated with a dual mechanism. Firstly, the magnesium matrix dissolution, during the corrosion attack, favours aluminium enrichment on the metallic surface and allows the formation of a semi-protective Al-rich oxide layer which improves the corrosion resistance of the alloy. Additionally, the lamellar aggregate network of b-phase acts as a barrier to the progression of the corrosion attack. 4. AZ91D presented similar corrosion behaviour as AZ80 alloy, but the different solidification microstructure changes the mechanism of the corrosion attack for this alloy. In this case, the corrosion resistance is exclusively attributed to the presence of a network of eutectic aggregate with higher Al content, which limits the advance of the corrosion attack. Acknowledgements The authors wish to thank the MCYT for the financial support given to this work (Project MAT2006-13179C02-01-02). References [1] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37. [2] E. Ghali, Magnesium and magnesium alloys, in: R.W. Revie (Ed.), Uhlig’s Corrosion Handbook, John Wiley & Sons, New York, 2000, p. 793. [3] D. Eliezer, E. Aghion, F.H. Froes, Adv. Perform. Mater. 5 (1998) 201. View publication stats [4] W.S. Loose, Corrosion and Protection of Magnesium, in: L.M. Pidgeon, J.C. 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