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Corrosion Science 50 (2008) 823–834

Corrosion behaviour of magnesium/aluminium alloysin 3.5 wt.% NaCl

A. Pardo a,*, M.C. Merino a, A.E. Coy a, R. Arrabal b, F. Viejo b, E. Matykina b

a Departamento de Ciencia de Materiales, Facultad de Quımicas, Universidad Complutense, 28040 Madrid, Spainb 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 2007Available online 24 November 2007

Abstract

Corrosion behaviour of commercial magnesium/aluminium alloys (AZ31, AZ80 and AZ91D) was investigated by electrochemical andgravimetric tests in 3.5 wt.% NaCl at 25 �C. Corrosion products were analysed by scanning electron microscopy, energy dispersive X-rayanalysis and low-angle X-ray diffraction. Corrosion damage was mainly caused by formation of a Mg(OH)2 corrosion layer. AZ80 andAZ91D alloys revealed the highest corrosion resistance. The relatively fine b-phase (Mg17Al12) network and the aluminium enrichmentproduced on the corroded surface were the key factors limiting progression of the corrosion attack. Preferential attack was located at thematrix/b-phase and matrix/MnAl intermetallic compounds interfaces.� 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 withminor additions of Zn and Mn, is the most widely usedamong Mg–Al alloys. They are characterised by low costof production and also by relatively good corrosion resis-tance and satisfactory mechanical properties from 95 to120 �C [1].

However, corrosion performance of magnesium alloyshas been a major obstacle to their growth in structuralapplications, despite their high stiffness/weight ratio, easeof machinability, high damping capacity, castability, wel-dability and recyclability [2]. These properties make themsuitable for aerospace and automotive applications, wherelight metals are mandatory in order to reduce weight andgreenhouse gas emissions. Over the past 30 years, alloydesign development, new surface treatments and improved

0010-938X/$ - see front matter � 2007 Elsevier Ltd. All rights reserved.

doi:10.1016/j.corsci.2007.11.005

* Corresponding author. Tel.: +34 1 3944348; fax: +34 1 3944357.E-mail address: anpardo@quim.ucm.es (A. Pardo).

knowledge of corrosion mechanisms has lead to an increaseof the real and potential applications of magnesium alloys[3].

Corrosion resistance of Mg alloys depends on manyfactors: (i) environment, (ii) alloy composition and micro-structure, and (iii) properties of the film developed in themedium to which they are exposed. Concerning theenvironment, corrosion resistance of magnesium alloys inchloride containing solutions greatly depends on pH andCl� concentration, with no significant influence of oxygenconcentration [4,5]. In general, magnesium and its alloysdissolve at a very low rate in alkaline or poorly bufferedsodium 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 rapidattack in neutral aqueous solutions and even higher inacidic solutions [4,6]. It is also common to find higher cor-rosion rates with increasing Cl� ion concentration at all pHlevels [6]. The corrosion of magnesium alloys in non-oxidiz-ing neutral or alkaline chloride solutions at free corrosionpotential 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, themechanism is different [9,10] from the auto-catalytic pittingexperienced by stainless steels [11], since there does notseem to be much tendency for deep pitting, possibly as aresult of pH increase and magnesium hydroxide film for-mation. However, this is not always true, since there is asignificant influence of the microstructure on the corrosionmechanism, especially in two-phase magnesium alloys.

On the topic of alloy composition and microstructure itis known that alloying elements not only modify themechanical properties of magnesium, but also impart asignificant impact on the corrosion behaviour. Alloying ele-ments can form secondary particles, which are noble withrespect to the magnesium matrix, thereby facilitatingcorrosion, or enrich the corrosion product thereby possiblyinhibiting the corrosion rate [12].

In general, it is reported that increasing Al concentra-tions in Mg–Al alloys have a beneficial effect on the corro-

Table 1Nominal composition of the materials tested

Material Chemical composition (wt.%)

Al Zn Mn Si Cu

Mg (99%) 0.006 0.014 0.03 0.019 0.00AZ31 3.1 0.73 0.25 0.02 <0.00AZ80 8.2 0.46 0.13 0.01 <0.00AZ91D 8.8 0.68 0.30 0.01 <0.00

Fig. 1. SEM micrographs of: (a) unalloyed Mg; (b) A

sion behaviour in chloride media [13–15], but the specificmechanism and influence of Al is still not well understood.For instance, Lunder [16] found that increasing concentra-tions of 2–8 wt.% Al in die cast AS, AM and AE alloysdecrease the corrosion rate in 5% NaCl, however, thedecrease in corrosion rate appears to be related toa decrease in the impurity level with increasing Al content(it is known that Al reduces the iron tolerance limit from170 wt.-ppm to 20 wt.-ppm [17]). More recent data com-pared the corrosion rates of high purity alloys in 3% NaCl[18]. Results showed that HP–Mg5Al alloy had a corrosionrate significantly higher than that of HP–Mg due to themicro-galvanic corrosion acceleration of the corrosion ofthe 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 ofHP–Mg, despite the fact that these commercial alloys eachhad a two phase microstructure which would cause a

Fe Ni Ca Zr Others

1 0.004 <0.0011 0.005 <0.001 <0.01 <0.001 <0.301 0.004 <0.301 0.004 <0.008 <0.30

Z31 alloy; (c) AZ80 alloy and (d) AZ91D alloy.

Fig. 2. Mass loss versus time for the materials immersed in 3.5 wt.% NaClsolution.

Table 2Kinetic laws of the materials immersed in 3.5 wt.% NaCl solution

Material Kinetic law: y = b � t; [y (mg/cm2), t (h)]

Mg (99%) y = 13.03t 0 6 t 6 16 (r2 = 0.99)AZ31 y = 2.31 � 10�1t 0 6 t 6 240 (r2 = 0.99)AZ80 y = 3.10 � 10�3t 0 6 t 6 240 (r2 = 0.98)AZ91D y = 3.70 � 10�3t 0 6 t 6 240 (r2 = 0.98)

10-7 10-6 10-5 10-4 10-3 10-2-1.8

-1.7

-1.6

-1.5

-1.4

-1.3

-1.2

-1.1

ESS

E(V

)

I (A/cm2)

AZ31

10-7 10-6 10-5 10-4 10-3 10-2-1.8

-1.7

-1.6

-1.5

-1.4

-1.3

-1.2

ESS

E(V

)

I (A /cm2)

AZ91D

a

b

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.

A. Pardo et al. / Corrosion Science 50 (2008) 823–834 825

galvanic acceleration of the corrosion of the a-phase[19,20]. Therefore, apart from impurities and Al concentra-tion it is also necessary to study the effect of phasedistribution.

Normally, in Mg–Al alloys the aluminium is partly insolid solution and partly precipitated in the form ofMg17Al12 along grain boundaries as a continuous phaseas well as part of a lamellar structure. It is known thatMg17Al12 exhibits a passive behaviour over a wider pHrange than either of its component aluminium and magne-sium [21], and it was found that the distribution of theMg17Al12 phase determines the corrosion resistance of theMg–Al alloys [22]. Song et al. [17] suggested that the b-phase mainly served as a galvanic cathode and acceleratedthe corrosion process of the a-matrix if the volume fractionof b-phase was small; however for a high volume fraction,the b-phase might act as an anodic barrier to inhibit theoverall corrosion of the alloy. It has also been reported thatthe ratio of the b-phase to the surrounding Al-rich-a oreutectic-a, which can be about 12 wt.% compared to1.5 wt.% in the grain centre [23], plays an important rolein the formation of galvanic cells. Thus, Raman [24] foundthat an increase in the relative size of the b-phase at theexpense of the Al-rich-a area, increasing the cathode toanode area ratio, results in an increase in the localized cor-rosion. However, the existence of a surrounding Al-rich-aarea with independent electrochemical identity appears tobe greatly dependant on the thermal history and composi-tion of the alloy. There are also contradictory resultsdepending on aluminium concentration in the a-phase.For instance, Lunder et al. [21] proposed that aluminiumaccelerate anodic dissolution below 8% whereas itdecreases corrosion above 10%. However, Song [25,26]found that an increase of Al in the a-phase improves thecorrosion resistance. Moreover, results obtained for rapidsolidified alloys with different additions of Al and no pre-cipitation of b-phase revealed a positive effect of Al on cor-rosion resistance [27]. In summary, it seems that Alimproves the corrosion resistance of Mg–Al alloys in chlo-ride environments but its effect is strongly influenced by theb-phase morphology and distribution.

The aim of this work was to study the corrosion behav-iour of three commercial magnesium/aluminium alloys inmarine environment. The effect of immersion time and Alconcentration on corrosion resistance was monitored byelectrochemical and gravimetric tests, scanning and elec-tron microscopy (SEM), energy dispersive X-ray analysis(EDX) and low-angle X-ray diffraction (XRD).

2. Experimental procedure

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. PureMg and AZ31 alloy were fabricated in wrought condition,

826 A. Pardo et al. / Corrosion Science 50 (2008) 823–834

whereas AZ80 (chill casting) and AZ91D alloys weremanufactured by ingot casting process. All the materialswere supplied by Magnesium Elektron Ltd.

2.2. Preparation and surface characterization

For metallographic characterization, samples were wetground through successive grades of silicon carbide abra-sive papers from P120 to P2000 followed by diamond fin-ishing to 0.1 lm. Two etching reagents were used: (a)Nital, 5 ml HNO3 + 95 ml ethanol, to reveal the constitu-ents and general microstructure of Mg, AZ80 andAZ91D materials and (b) Vilella reagent, 0.6 g picricacid + 10 ml ethanol + 90 ml H2O, to reveal grain bound-aries of AZ31 alloy. The constituents were examined bySEM using a JEOL JSM-6400 microscope equipped withOxford Link EDX microanalysis hardware. For low-angleXRD studies a Philips XPert diffractometer Ka Cu =1.54056 A) was used.

Prior to corrosion tests, specimens were wet ground tograde P1200 grit, followed by rinsing with isopropyl alco-hol in an ultrasonic bath and drying in warm air. In all

10-8 10-7 10-6 10-5 10-4 10-3 10-2-1.8

-1.7

-1.6

-1.5

-1.4

-1.3

-1.2

-1.1

-1.0

10-8 10-7 10-6 10-5 10-4 10-3 10-2-1.8

-1.7

-1.6

-1.5

-1.4

-1.3

-1.2

-1.1

-1.0

ESS

E (

V)

I (A /cm2)

1 h 1 d

Mg

ESS

E(V

)

I (A /cm2)

1 h 1 d 3 d 7 d

AZ80

Fig. 4. Influence of the immersion time on the anodic behav

cases, the tests were performed in duplicate to guaranteethe reliability of the results.

2.3. Gravimetric measurements

Gravimetric measurements were performed using rect-angular (Mg and AZ31) and cylindrical (AZ80 andAZ91D) specimens of working area �16 cm2 immersed in3.5 wt.% NaCl solution at room temperature (pH 5.6,25 �C). Prior to the tests, specimens were measured andweighed. Once the test was finished for each immersiontime, the samples were extracted, rinsed with isopropylalcohol, dried in hot air, and then weighed again in orderto calculate the mass gain per unit surface area. Corrosionrate was calculated from the mass losses per unit of surfacearea, calculated from the expression (Mi �Mf)/A, whereMi is the initial mass, Mf the final mass and A the exposedsurface area. In order to obtain further information, pHchanges of the test solution during the gravimetric experi-ments were recorded with a standard pH-meter. For thesame reason, hydrogen evolution was measured using asimple procedure described elsewhere [9].

10-8 10-7 10-6 10-5 10-4 10-3 10-2-1.8

-1.7

-1.6

-1.5

-1.4

-1.3

-1.2

-1.1

-1.0

10-8 10-7 10-6 10-5 10-4 10-3 10-2-1.8

-1.7

-1.6

-1.5

-1.4

-1.3

-1.2

-1.1

-1.0

I (A /cm2)

1 h 1 d 3 d 7 d

AZ91

I (A /cm2)

1 h 1 d 3 d 7 d

AZ31

iour for the materials tested in 3.5 wt.% NaCl solution.

A. Pardo et al. / Corrosion Science 50 (2008) 823–834 827

2.4. Electrochemical measurements

DC electrochemical measurements were performedusing specimens with 2 cm2 of working area exposed to3.5 wt.% NaCl at room temperature (25 �C). An AUTO-LAB model PGSTAT 30 potentiostat connected to athree-electrode cell was used for the electrochemical mea-surements; the working electrode was the test material,whereas the counter and reference electrodes were graphiteand Ag/AgCl (SSE), respectively. Solution concentrationinside the reference electrode compartment was KCl 3 M,with a potential of 0.197 V with respect to hydrogen. Ano-dic, cyclic and double cyclic polarization measurementswere carried out at a scan rate of 0.3 mV/s, from�100 mV to +400 mV with respect to the corrosion poten-tial (Ecorr). For cyclic polarization, the scan direction wasreversed when the samples reached the anodic corrosioncurrent of 5 mA and potential was scanned back to thestarting potential.

2.5. Characterization of corrosion products

After the tests, magnesium alloys were observed bySEM in order to study the morphology and evolution ofcorrosion products formed on the material surface. Also,the composition of the corrosion layer was analysed bylow-angle XRD.

0 1 2 3 4 5 6 7-1.6

-1.5

-1.4

-1.3

-1.2

-1.1

-1.0

Eco

rr (

VSS

E)

Time (d)

Mg AZ31 AZ80 AZ91

1

102

103

104

105

Rp

(Ω·c

m2 )

MgAZ31AZ80AZ91

a

b

3. Results

3.1. Microstructural characterization

Fig. 1 shows the SEM microstructure of tested materi-als. Unalloyed Mg only reveals equiaxial grains with aver-age size of 45 lm, whereas the presence of aluminium(3.1 wt.%) and manganese (0.25 wt.%) in AZ31 alloyfavours the formation of intermetallic phases, mainly inform of MnAl2 inclusions (Figs. 1a, b). On the other hand,AZ80 and AZ91D ingot casting alloys disclose two differ-ent types of solidification microstructures. AZ80 alloyshows a biphasic microstructure with a a-Mg solid solutionand a discontinuous precipitation in lamellar form of b-phase (Mg17Al12), which starts at the grain boundariesfrom solid solution (Fig. 1c). AZ91D alloy reveals a-Mgprimary dendrites and eutectic a-Mg/Mg17Al12 in the inter-dendritic region, which appears completely in divorcedform with respect to solid solution (Fig. 1d). Indeed, therapid solidification obtained by chill casting process usedfor AZ80 alloy promotes a refinement of the microstruc-ture compare to AZ91D alloy. Finally, MnAl2 intermetallicinclusions are also found for AZ80 and AZ91D alloys.

0 1 2 3 4 5 6 710

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) corrosionpotential (Ecorr).

3.2. Gravimetric results

Fig. 2 shows the mass loss versus time of tested materialsimmersed in 3.5 wt.% NaCl solution. All materials exhib-

ited linear kinetics of mass loss associated with magnesiummatrix dissolution. Unalloyed Mg presented the highestmass loss and it was completely dissolved after 16 h ofimmersion with a mass loss of 196 mg/cm2. On the otherhand, the addition of aluminium increased notably the cor-rosion resistance. Thus, 3 wt.% Al in AZ31 alloy reducedthe mass loss to 55.2 mg/cm2 after 10 days of immersion,and 8–9 wt.% Al in AZ80 and AZ91D alloys diminishedthe 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 theexperimental data, corresponding to the growth of corro-sion products layer generated during the gravimetric test.In each case, data have been approximated by a linearequation y = b � t, where ‘‘y” coordinate represents themass gain in units of mg/cm2 and ‘‘t” is the immersion timein 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 andAZ31 alloy manifested very low corrosion resistance in sea-water, whereas AZ80 and AZ91D alloys presented a mod-erate resistance after 10 days of immersion.

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.

828 A. Pardo et al. / Corrosion Science 50 (2008) 823–834

A. Pardo et al. / Corrosion Science 50 (2008) 823–834 829

3.3. Electrochemical results

In order to obtain further information, the gravimetricstudy was supplemented with potentiodynamic polariza-tion measurements for different times of immersion in3.5 wt.% NaCl solution. Double cycle polarization curvesfor AZ31 and AZ91D alloys after immersion for 1 h reveala different behaviour compared to materials with a stablepassive layer (Fig. 3). The onset of pitting is not visible inthe forward scan of both cycles, since pitting potential(Epit) is very close to Ecorr and, consequently, it should beexpected that these alloys suffer pitting attack immediatelyafter their immersion in the aggressive media at the opencircuit potential. However, the current densities of thecathodic branch and, therefore, the growth of corrosionproducts on the material surface are quite high, suggestinggeneral corrosion attack as the main mechanism ofdegradation.

Fig. 4 discloses the evolution of the electrochemicalbehaviour of each material with increasing immersion time.The high dissolution rate observed for unalloyed magne-sium limits this study, since it was dissolved after 1 dayof immersion according to gravimetric results. Neverthe-less, longer immersion times shifted its curves to greatercurrent densities, possibly associated with the formationof a low-protective oxide layer, which does not impedethe corrosion attack progression. AZ80 and AZ91 alloysrevealed a slight shift of Ecorr towards more noble values,and corrosion current densities (icorr) were approximately

Fig. 8. Cross-section BSE image and corresponding X-ray maps of Mg, A

one order of magnitude lower. AZ31 presented the mostsignificant change on the Ecorr values, possibly due to thehigh reactivity of this alloy which promotes formation ofa thick corrosion layer, limiting the attack progressionand increasing Ecorr values (Fig. 5a). Concerning polariza-tion resistance (Rp) values, they remained low after 7 daysof immersion (between 102 and 104 X cm2) (Fig. 5b). How-ever, an increase of Rp values and, therefore, a better cor-rosion behaviour, it is observed for Mg and AZ31materials, possibly related to the nucleation and growthof a thick and semi-protective layer of corrosion products.

3.4. Characterization of corrosion products

The SEM micrographs of the surface of the tested mate-rials after 10 days of immersion in 3.5 wt.% NaCl solutionare shown in Fig. 6. Unalloyed magnesium is not presenteddue to its complete dissolution after 24 h of exposure.Although AZ31 alloy revealed higher corrosion resistancethan unalloyed magnesium, its surface was completely cov-ered by a thick and uneven film of corrosion products. Anincrease of aluminium content up to 8–9 wt.% reduced thereactivity of the magnesium alloys and, thereby, AZ80 andAZ91D alloys presented a lower degree of corrosion.

Initial stages of corrosion for AZ31, AZ80 and AZ91Dalloys immersed in 3.5 wt.% NaCl revealed localised corro-sion around MnAl2 inclusions and b-phase interfaces,which form a galvanic couple with the surrounding Mgmatrix (Fig. 7). Polarization tests did not show a clearly

l 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 shouldbe associated with a localised attack, but not to a typicalpitting attack reported for materials with a stable passivelayer, such as aluminium alloys, stainless steels, etc.

X-ray map analysis of the cross section of AZ31 alloyafter immersion for 10 days in the corrosive medium con-firmed the formation of a thick corrosion layer with anirregular thickness between 200 and 400 lm and mainlyconstituted by magnesium oxides and/or hydroxides(Fig. 8). AZ80 alloy revealed a discontinuous and crackedcorrosion film of lower thickness than AZ31 alloy(<200 lm) (Fig. 9). Furthermore, aluminium enrichmentwas found at the metal/corrosion layer interface, possiblyproduced by preferential Mg dissolution during the immer-sion test. The formation of Al oxide species is likely to beone of the main reasons for the improved corrosion resis-

Fig. 9. Cross-section BSE image, profile line and corresponding X-ray maps of

tance exhibited by this alloy. Additionally, it can beobserved that the front of the corrosion attack was stoppedwhen it reached the lamellar aggregate of b-phase(Mg17Al12). Thus, the higher corrosion resistance ofAZ80 alloy may be associated with the presence of alumin-ium and its direct or indirect intervention in corrosionmechanism, consisting of two steps: (a) formation of asemi-protective Al-rich oxide layer and (b) reduction ofthe corrosion progression near to the lamellar b-phaseaggregate.

Finally, Figs. 10 and 11 show the study of the corrosionlayer 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 posi-tive effect of Al on the corrosion performance cannot beexplained in the same way as above, since aluminium

Mg, Al and O for the AZ80 alloy immersed in 3.5 wt.% NaCl for 10 days.

Fig. 10. Cross-section BSE image and corresponding X-ray map of Al forthe AZ91D alloy immersed in 3.5 wt.% NaCl for 10 days.

A. Pardo et al. / Corrosion Science 50 (2008) 823–834 831

enriched eutectic aggregate solidifies at the interdendriticspaces resulting in an a-primary phase with lower Al con-tent than the predicted by the equilibrium phase diagram.Therefore, (a) there was not an appreciable Al enrichmentobserved on the surface after the Mg dissolution during thecorrosion attack. Although, as mentioned above, (b) thepresence of a relatively fine b-phase network, which for thisparticular alloy is extended all along the surface, partiallyimpeded the corrosion attack (Fig. 11).

Low-angle XRD study (incident angle of 1�) of the cor-rosion layer produced after immersion in 3.5 wt.% NaClfor 10 days revealed brucite (Mg(OH)2) as the main corro-sion product, and its peaks exhibited higher intensity forunalloyed Mg and AZ31 alloy due to the formation of athicker corrosion layer during the severe attack that bothmaterials suffered (Fig. 12). Unlike other works, evidencesof magnesite (MgCO3) formation, due to CO2 presence inthe atmosphere, were not found [28].

4. Discussion

When unalloyed magnesium is exposed to the atmo-sphere or aqueous solutions, a grey oxide (mainly magne-sium hydroxide, brucite) forms on its surface, which isstable for the basic range of pH values [29]. Nevertheless,in presence of chloride anions, this surface film breaksdown and magnesium appears unprotected. Although cor-

rosion mechanism of magnesium needs further investiga-tion, it is generally reported that the following anodicand cathodic reactions take place in marine environments:

MgðsÞ !Mg2þðaqÞ þ 2e� ðanodic reactionÞ ð1Þ2H2Oþ 2e� ! H2 þ 2OH�ðaqÞ ðcathodic reactionÞ ð2Þ

Firstly, magnesium dissolves and Mg2+(aq) cations areproduced (Eq. (1)), possibly through intermediate stepsinvolving monovalent magnesium ion [14,30]. Secondly,magnesium dissolution is accompanied by hydrogen evolu-tion (Eq. (2)), since magnesium in neutral and low pHaqueous solutions is well below the region of water stabil-ity. Finally, pH raises along with the cathodic reaction dueto the formation of OH–, which favours the formation ofMg(OH)2(s) according to Pourbaix diagram (Fig. 13a)[31]. Thus, the overall reaction could be expressed as

Mg2þðaqÞ þ 2OH�ðaqÞ !MgðOHÞ2ðsÞðcorrosion productÞð3Þ

Since pH changes and H2 evolution can be easily measured,it is feasible to study the corrosion evolution of magnesiumwith these parameters [19]. Fig. 13b shows the variation ofpH for all tested materials immersed in 3.5 wt.% NaCl (pH5.6). For commercially pure Mg, test solution revealed thehighest pH raise, pH >10 after immersion for 1 h. On theother hand, even though aluminium reduces the magne-sium reactivity, pH values were registered between 8 and9.5 after 1 h, the pH increasing up to 10 for longer immer-sion times. Therefore, these technique may be convenientand reliable for measurement of initial corrosion of magne-sium and its alloys, however it becomes gradually insensi-tive for long times, since the solution is saturated withMg(OH)2, and concentrations of Mg2+ and OH� cannotfurther increase while corrosion is still proceeding [19].

Concerning H2 evolution Fig. 13c compares mass lossobtained by gravimetric measurements with H2 evolution,assuming that the dissolution of one atom of magnesiumgenerates one molecule of hydrogen (Eqs. (1) and (2)).Results confirmed the higher corrosion resistance ofAZ80 and AZ91D alloys and the high dissolution rate ofpure Mg in the test solution. In all cases, corrosion rateestimated by H2 evolution was lower than the one obtainedby gravimetric tests, possibly due to small particles fallingout of the surface after having been undermined by the cor-rosion process [9]. Therefore, although pH and H2 evolu-tion provide information regarding the corrosion ofmagnesium alloys, it is still suggested to perform gravimet-ric measurements, since they are the most basic and reliablemethod.

In general, it is reported that the presence of aluminiumin single-phase Mg–Al alloys has a beneficial effect on thecorrosion behaviour in chloride media [13–15]. In the pres-ent work, this behaviour has been clearly revealed for theAZ31 alloy, which exhibited higher corrosion resistancethan commercially pure Mg. However, aluminium influ-ence in two-phase Mg–Al alloys is still not well understood

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 alloyimmersed in 3.5 wt.% NaCl for 10 days.

Fig. 12. Low-angle XRD (incidence angle: 1�) of tested materialsimmersed in 3.5 wt.% NaCl solution for 10 days.

832 A. Pardo et al. / Corrosion Science 50 (2008) 823–834

since depends on many factors, such as impurities, b-phasemorphology and distribution, aluminium content and sizeof the a-Mg primary dendrites and the eutectic-a, alumin-ium enrichment on the corrosion product layer [12], etc.

Magnesium and magnesium alloys present an activeposition in both the electromotive force (EMF) and the gal-vanic series for seawater. As a result, galvanic interactionsbetween magnesium and other metals are a serious con-cern. Thus, the corrosion resistance of Mg–Al alloys willdepend on the presence of alloy elements acting as activecathode, such as manganese, iron, etc. The most detrimen-tal potential cathodes in Mg–Al alloys are iron-rich andMnAl second phase particles [22]. The former were notfound in any of the studied alloys, whereas MnAl2 interme-tallic compounds were identified as preferential sites forlocalised corrosion in AZ31, AZ80 and AZ91D alloys.

0 1 2 3 4 5 6 78

9

10

11

12

pH

Time (d)

Mg AZ31 AZ80 AZ91

a

b

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 10days.

A. Pardo et al. / Corrosion Science 50 (2008) 823–834 833

Previous investigations reported that distribution of theb-phase (Mg17Al12) determines the corrosion resistance ofthe Mg–Al alloys [17,19–22,24,32]. According to this, fora low volume fraction, the b-phase serves as a galvaniccathode and accelerates the corrosion process of the a-phase, whereas large volume fractions act as an anodic bar-rier and the overall corrosion diminishes [26,33]. Moreover,it has also been reported that the ratio of the b-phase to thesurrounding Al-rich-a or eutectic-a plays an important rolein the formation of galvanic cells. Apart from the micro-structure, the composition of the corrosion layer mayimprove the corrosion resistance. For instance, Nordlienet al. [34] observed that alumina components forms a con-tinuous skeletal structure in an oxide layer of the magne-sium alloy, which passivating properties are much betterthan Mg(OH)2 and MgO layers.

Therefore, keeping in mind these observations and theresults obtained in the present work, the following consid-erations about the corrosion mechanism of studied Mg–Alalloys can be summarized: (i) an increase of the aluminiumconcentration in the nominal composition of the alloysreduced the activity of the pure Mg in 3.5 wt.% NaCl, with

AZ80 and AZ91D showing the highest corrosion resis-tance; (ii) an independent electrochemical Al-rich-a areasurrounding the b-phase was not observed by metallo-graphic characterization; (iii) AZ91D alloy exhibits aslightly worse corrosion behaviour than AZ80 alloy,although the aluminium content in its nominal composi-tion is slightly higher and both alloys reveal the presenceof relatively fine b-phase network segregated in the matrixalloy, which acts as an effective barrier against progressionof corrosion attack. As a result, the main reason for thisdifferent corrosion behaviour observed for the AZ80 andAZ91D alloys could be explained in terms of aluminiumenrichment on the metallic surface during the magnesiummatrix dissolution. According to this, aluminium contentfor both alloys in the a-Mg primary dendrites was esti-mated by EDX. Chill casting employed for AZ80 alloy pro-motes high aluminium contents in the solid solution(13.3 at.%), meanwhile ingot casting, with lower solidifica-tion rate, promotes a coarse dendritic microstructure forthe AZ91D alloy and less aluminium content in the solidsolution (8.4 at.%). Hence, AZ80 alloy revealed aluminiumenrichment 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-richoxide layer was not observed for the AZ91D alloy (Figs.9 and 11).

5. Conclusions

1. Corrosion attack of all tested magnesium alloys occursat the a-magnesium matrix/Al–Mn and Mg17Al12 inter-metallic compounds interfaces, by means of the forma-tion of galvanic couples. Later, the nucleation andgrowth of an irregular and less protective corrosionlayer consisted mainly of Mg(OH)2 is produced froma-Mg matrix.

2. An increase of the aluminium concentration in thenominal composition of the alloys reduced the activityof the pure Mg in 3.5 wt.% NaCl. However, the AZ31alloy (3.1 wt.% Al) still presented high corrosion rates.

3. The principal cause of higher corrosion resistance ofAZ80 alloy is associated with a dual mechanism. Firstly,the magnesium matrix dissolution, during the corrosionattack, favours aluminium enrichment on the metallicsurface and allows the formation of a semi-protectiveAl-rich oxide layer which improves the corrosion resis-tance of the alloy. Additionally, the lamellar aggregatenetwork of b-phase acts as a barrier to the progressionof the corrosion attack.

4. AZ91D presented similar corrosion behaviour as AZ80alloy, but the different solidification microstructurechanges the mechanism of the corrosion attack for thisalloy. In this case, the corrosion resistance is exclusivelyattributed to the presence of a network of eutectic aggre-gate with higher Al content, which limits the advance ofthe corrosion attack.

Acknowledgements

The authors wish to thank the MCYT for the financialsupport given to this work (Project MAT2006-13179-C02-01-02).

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