Corrosion behaviour of Mg/Al alloys in high humidity atmospheres

326 DOI: 10.1002/maco.200905538 Materials and Corrosion 2011, 62, No. 4

Corrosion behaviour of Mg/Al alloys in high humidityatmospheres

R. Arrabal*, A. Pardo, M. C. Merino, S. Merino, M. Mohedano andP. Casajus

The influence of relative humidity (80–90–98% RH) and temperature (25 and

50 8C) on the corrosion behaviour of AZ31, AZ80 and AZ91D magnesium alloys

was evaluated using gravimetric measurements. The results were compared

with the data obtained for the same alloys immersed in Madrid tap water. The

corrosion rates of AZ alloys increased with the RH and temperature and were

influenced by the aluminium content and alloy microstructure for RH values

above 90%.The initiation of corrosionwas localised around the Al–Mn inclusions

in the AZ31 alloy and at the centre of the a-Mg phase in the AZ80 and AZ91D

alloys. The b-Mg17Al12 phase acted as a barrier against corrosion.

1 Introduction

The atmospheric corrosion of magnesium alloys is receiving

more attention in the recent years because of their potential for

lightweight transport applications [1]. Magnesium alloys contain-

ing aluminium and zinc (AZ series) are amongst the most

corrosion resistant magnesium alloys [2]. The atmospheric

corrosion behaviour of these alloys is greatly influenced by the

relative humidity (RH), temperature, presence of chloride ions,

carbon dioxide (CO2) levels and alloy composition and micro-

structure.

Studies by Le Bozec et al. [3] illustrated that the corrosion rate

of AZ91D and AM50magnesium alloys exposed to high humidity

atmospheres was strongly connected to the amount of sodium

chloride (NaCl) on the surface, regardless of the composition of

the alloy and surface state. A linear dependence between the

corrosion rate and the amount of NaCl was noticed, which was

explained by an increase of electrolyte conductivity. Increasing the

temperature usually increases the corrosion rate [3, 4] and

atmospheric CO2 is known to react with magnesium corrosion

products to form several magnesium carbonates, which provide

some surface protection [5].

In general, immersion tests of AZ alloys in chloride aqueous

solutions show better corrosion resistance with increasing

aluminium content [6, 7]. However, the effect of aluminium

greatly depends on the microstructural features of the alloy,

R. Arrabal, A. Pardo, M. C. Merino, M. Mohedano, P. Casajus

Facultad de Quımicas, Departamento de Ciencia de Materiales,

Universidad Complutense, 28040 Madrid (Spain)

E-mail: [email protected]

S. Merino

Departamento de Tecnologıa Industrial, Universidad Alfonso X El Sabio,

Villanueva de la Canada, 28691 Madrid (Spain)

� 2011 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

especially on the distribution and morphology of the b-phase [8].

Therefore, considering the previous and that the atmospheric

corrosion behaviour of magnesium alloys is different from that

observed in aqueous solutions [9], the present study focuses on

the influence of aluminium content, alloy microstructure, RH

and temperature on the corrosion behaviour of Mg–Al alloys

exposed to high humidity environments.

2 Experimental

2.1 Test materials

Chemical compositions of tested Mg–Al alloys (AZ31, AZ80 and

AZ91D) are listed in Table 1. Commercially pure or ‘low-purity’

Mg was used as the reference material. Mg and AZ31 materials

were fabricated in wrought condition and supplied in plates of

3mm thickness. AZ80 and AZ91D alloys weremanufactured by a

casting process in the form of billets of 300 and 250mm in

diameter. All the materials were supplied byMagnesium Elektron

Ltd.

2.2 Materials preparation and characterisation

For metallographic characterisation, samples were wet ground

through successive grades of SiC abrasive papers from P120 to

P2000, followed by finishing with 0.1mm diamond paste. Two

etching reagents were used: (a) 5% Nital (5mL HNO3þ 95mL

ethanol) during 30 s, to reveal the constituents and general

microstructure of Mg, AZ80 and AZ91D materials and (b) Vilella

reagent (0.6 g picric acidþ 10mL ethanolþ 90mL H2O) during

10 s, to reveal grain boundaries of AZ31 alloy.

For characterisation of Al–Mn particles imbedded in the

studied Mg–Al alloys, carbon extraction replicas were prepared

and examined in a JEOL 2000–FX transmission electron

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Materials and Corrosion 2011, 62, No. 4 Corrosion behaviour of Mg/Al alloys 327

Table 1. Nominal composition of test materials

Material Elements (wt%)

Al Zn Mn Si Cu Fe Ni Ca Zr Others

Mg (99.9%) 0.006 0.014 0.03 0.019 0.001 0.004 <0.001

AZ31 3.1 0.73 0.25 0.02 <0.001 0.005 <0.001 <0.01 <0.001 <0.30

AZ80 8.2 0.46 0.13 0.01 <0.001 0.004 <0.30

AZ91D 8.8 0.68 0.30 0.01 <0.001 0.004 <0.008 <0.30

microscope (TEM) operating at 200 kV and equipped with Oxford

link EDX microanalysis hardware.

Surface potential maps of polished specimens were obtained

with a MultiMode scanning probe microscope (Veeco–Digital

Instruments) working in tapping mode. A silicon tip with a

platinum coating of 20 nm thickness was used for simultaneous

acquisition of topographic and surface potential images. The

AFM tip was calibrated by performing a potential ramp on a

reference sample consisting of aluminium coated with a thin gold

layer. The tip to sample distance was kept constant at 100 nm

using a two-pass technique, where the height data is recorded in

tappingmode during the first pass and, in the second pass, the tip

lifts above the surface to an adjustable lift height and scans the

same line while following the height profile recorded in the first

pass. All measurements were made at room temperature with a

RH in the range of 40–65%.

2.3 Gravimetric measurements

Prior to corrosion tests, specimens with a working area of

�15 cm2 were ground to a P1200 SiC finish, degreased in

isopropyl alcohol in an ultrasonic bath and dried in warm air.

High RH tests consisted of 24 h cycles performed in a

saturated water vapour at various RH values (80–90–98% RH)

and temperatures (25–50� 1 8C) during 672 h (28 days) simu-

lated by a humidity condensation cabinet CCK 300 (Dycometal).

Temperature and humidity were verified using digitals thermo-

meter and hygrometer. At the end of the tests, the specimens were

Figure 1. Optical micrographs of test materials: (a) Mg; (b) AZ31;(c) AZ80 and (d) AZ91D

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rinsed with deionised water and dried in warm air. The

specimens were weighed before and after the tests using a

Sartorius BP 211D scale with an accuracy of 0.00001 g. Mass

changes per unit of surface area were calculated according to the

expression (Mf �Mi)/A, whereMf is the final mass,Mi the initial

mass and A the exposed surface area.

A comparative study was conducted with the same

magnesium materials immersed in Madrid tap water (pH 8;

100–200mS/cm conductivity at 20 8C; 5.8–34.8mg/L Ca; 1.3–

4.3mg/L Mg; 4–20mg/L Na; zero free chlorine and 1–1.5mg/L

combined chlorine for disinfection). Duplicated tests were

conducted for each specimen to ensure reproducible test results.

2.4 Characterisation of corrosion products

Tested specimens were examined by scanning electronmicroscopy

(SEM) using a JEOL JSM–6400 microscope equipped with Oxford

Link energy-dispersive X-ray (EDX) microanalysis hardware in

Figure 2. Bright field TEM micrographs and electron diffraction

patterns of some of the Al–Mn inclusions in the (a) AZ80 and (b) AZ91Dalloys


328 Arrabal, Pardo, Merino, Merino, Mohedano, and Casajus Materials and Corrosion 2011, 62, No. 4

order to study themorphology and evolution of corrosion products

formed on the surfaces of the alloys. The composition of the

corrosion products was determined by low angle X-ray diffraction

(XRD) using a Philips X’Pert diffractometer (CuKa¼ 1.54056 A).

3 Results

3.1 Microstructural characterisation

3.1.1 SEM/TEM

Plan view optical micrographs of studied materials are presented

in Fig. 1. Commercially pure Mg and the AZ31 alloy consist

of equiaxial grains with dimensions in the range of 40–50 and 5–

100mm, respectively (Fig. 1a and b). The AZ80 and AZ91D

casting alloys show two different solidification microstructures.

The AZ80 alloy reveals a biphasic microstructure with a-Mg

grains and grains with precipitation of a fine lamellar aggregate,

Figure 3. Topographic images, surface potential maps and potential prof


a-MgRb-Mg17Al12 (Fig. 1c). The AZ91D alloy consists of a-Mg

primary dendrites and a partially divorced eutectic a-Mg/b-

Mg17Al12, with the b-phase in the form of coarse particles at the

interdendritic regions (Fig. 1d). The presence of manganese in

the AZ alloys favoured the formation of Al–Mn intermetallic

inclusions with various stoichiometries (Al2Mn, Al3Mn,

Al4Mn, Al11Mn4, Al8Mn5, Al19Mn4 and Al10Mn3) according to

TEM characterisation (Fig. 2). The size of these particles varied

from 1 to 15mm.

3.1.2 Surface potential maps

Topographic images and surface potential maps of the Al–Mn

inclusions and b-phase indicated that both constituents are

cathodic with respect to the magnesium matrix (Fig. 3). The

potential profile across these phases revealed variations above

200mVand below 100mV for the Al–Mn inclusions and b-phase,

respectively. These values suggested a stronger galvanic couple at

iles of (a) AZ31 and (b) AZ91D alloys

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-0.10

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Mas

s gai

n (m

g/cm

2 )

Mg AZ31 AZ80 AZ91D

CC 80% RH 25ºC CC 80% RH 50 ºC

Mas

s gai

n (m

g/cm

2 ) CC 90% RH 50 ºCCC 90% RH 25 ºC

Madrid water 22 ºC

Mas

s gai

n (m

g/cm

2 )

CC 98% RH 25 ºC

Time (d)

CC 98% RH 50 ºC

Mas

s gai

n (m

g/cm

2 )

Time (d)

Figure 4. Mass gain versus time of test materials exposed in humid environments and Madrid water

the a-Mg/Al–Mn interface than that formed between the a-Mg

and b-Mg17Al12 phases. In the case of the Al–Mn inclusions, the

potential variations fitted with previous results [10], however, the

potential difference between the a-Mg and b-Mg17Al12 phases

was smaller than that observed in the die cast AZ91D alloy [10].

This suggested that galvanic corrosion phenomena are less likely

to occur at the a-Mg/b-Mg17Al12 interface when the alloy is in the

as-cast condition.

3.2 Gravimetric results

Figure 4 discloses the mass gain versus time of test materials

exposed to high humidity environments and Madrid water for

28 days. At 25 8Cand 80%RH, all thematerials revealedmass loss

during the first 4 days of exposure due to magnesium dissolution,

followed by a plateau until the end of the test that indicated

similar rates for magnesium dissolution and formation of

corrosion products. Under these conditions, neither the

alloy composition nor the microstructure constituents had a

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significant effect on the corrosion rate and all the materials

presented a very similar behaviour. At 90–98% RH and 25 8C, allthe materials presented smaller mass loss at the initial stages

compared with the results observed at 80% RH and exhibited

mass gain for exposure times over 4 days due to accumulation of

corrosion products on the surface. For the specimens exposed at

50 8C, the initial mass loss was always observed and themass gain

after 4 days increased with the RH. The effect of alloy

composition/microstructure was only noticeable for RH values

starting at 90%. Thus, at 25–50 8C and 90–98% RH, the alloys

with higher aluminium content (AZ80 and AZ91D) revealed the

lowest mass gain values.

In Madrid tap water at 22� 1 8C for 28 days, Mg, AZ80 and

AZ91D materials presented mass loss during the entire test,

indicating faster magnesium dissolution than formation of

corrosion products. The AZ31 alloy showed mass gain, possibly

as a consequence of a more uniform and stable corrosion layer.

The water pH for each specimen at the end of the test is presented

in Table 2. In aqueous solutions the pH increases along with the



Table 2. Water pH before and after immersion tests in Madrid water at

22 8C for 28 days

pH Mg AZ31 AZ80 AZ91D

Madrid water

8 10.1 9.9 9.5 9.4

process of magnesium corrosion due to the following cathodic

reaction (1):

2H2Oþ 2e� ! H2 þ 2OH� ðaqÞ ðcathodic reactionÞ (1)

Thus, according to the measured pH values (Table 2), the

intensity of corrosion in the immersion tests decreased with

increasing aluminium content in the bulk composition of the

alloy.

Table 3 shows the kinetic laws calculated from the

gravimetric data using a linear expression (y¼ bt; y: mass gain;

t: time) for exposure times where the mass changes followed a

linear trend. The kinetic laws for the specimens exposed to 80–

90% RH/25–50 8C were of the same order of magnitude. The

most noticeable change was observed at 98% RH, especially at

Figure 5. Surface appearance of test materials after exposure to humid

environments and Madrid water for 28 days

Table 3. Kinetic laws of test materials exposed to humid environments

and Madrid water

Relativehumidity

T(8C)

Material Kinetic law: y¼ bt;[y (mg/cm2), t(d)]

r2

80% 25 Mg (99.9%) y¼ 1.0� 10�4t 7� t� 28 0.98

AZ31 y¼ 0.9� 10�4t 7� t� 28 0.98

AZ80 y¼ 1.0� 10�4t 7� t� 28 0.98

AZ91D y¼ 0.6� 10�4t 7� t� 28 0.98

50 Mg (99.9%) y¼�5.3� 10�4t 7� t� 28 0.94

AZ31 y¼ 5.4� 10�4t 7� t� 28 0.98

AZ80 y¼ 1.4� 10�4t 7� t� 28 0.98

AZ91D y¼ 1.9� 10�4t 7� t� 28 0.97

90% 25 Mg (99.9%) y¼�1.6� 10�4t 7� t� 28 0.99

AZ31 y¼�1.4� 10�4t 7� t� 28 0.94

AZ80 y¼�2.4� 10�4t 7� t� 28 0.96

AZ91D y¼�0.19� 10�4t 7� t� 28 0.96

50 Mg (99.9%) y¼�1.2� 10�3t 7� t� 28 0.99

AZ31 y¼�5.7� 10�4t 7� t� 28 0.98

AZ80 y¼�3.2� 10�4t 7� t� 28 0.96

AZ91D y¼�2.1� 10�4t 7� t� 28 0.97

98% 25 Mg (99.9%) y¼ 5.1� 10�3t 0� t� 28 0.91

AZ31 y¼ 3.3� 10�3t 0� t� 28 0.91

AZ80 y¼ 2.6� 10�3t 0� t� 28 0.99

AZ91D y¼ 1.7� 10�3t 0� t� 28 0.93

50 Mg (99.9%) y¼ 1.6� 10�2t 7� t� 28 0.99

AZ31 y¼ 1.1� 10�2t 7� t� 28 0.92

AZ80 y¼ 0.8� 10�2t 7� t� 28 0.92

AZ91D y¼ 0.8� 10�2t 7� t� 28 0.92

Madrid

water

Mg (99.9%) y¼�1.4� 10�2t 0� t� 28 0.94

AZ31 y¼ 0.2� 10�2t 4� t� 8 0.96

AZ80 y¼�0.7� 10�2t 0� t� 8 0.95

AZ91D y¼�1.0� 10�2t 0� t� 8 0.94


50 8C, where thematerials revealed kinetic constants�2 orders of

magnitude higher than those observed at 80 and 90% RH. The

kinetic constants calculated for the specimens immersed in

Madrid water were of the same order of magnitude that those

estimated in the corrosion test at 98% RH and 50 8C.

3.3 Morphology and characterisation of corrosion

products

3.3.1 Surface appearance

Figure 5 shows the surface appearances of test materials after

exposure to different humidity environments at 25 and 50 8C and

Madrid tap water for 28 days. In general, surface degradation at

80–90% RH/25 8C was negligible, however at 98% RH, the

specimens became duller and signs of surface corrosion were

readily visible. At 50 8C and 98% RH, the degradation of the

specimens was more severe and their surface was almost entirely

covered by corrosion products after 28 days of exposure (Fig. 5).

The surface appearance of the specimens immersed in water

for 28 days revealed higher corrosion attack than in high humidity

atmospheres. Clear signs of localised corrosion were observed for

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Figure 6. SEMmicrographs of test alloys after exposure to 98%RH–50 8C for (a–c) 4 days and (d–f) 28 days: (a and d) AZ31; (b and e) AZ80 and (c and

f) AZ91D

the AZ80 and AZ91D alloys, whereas Mg and the AZ31 alloy

revealed a more uniform corrosion attack (Fig. 5).

3.3.2 SEM

Figure 6a–c shows the scanning electron micrographs of the

initial stages of corrosion of AZ31, AZ80 and AZ91D alloys after

exposure to 98% RH and 50 8C. For the AZ31 alloy, the corrosionattack commenced at the a-Mg/Al–Mn interfaces as it was

predicted by the surface potential maps. For the AZ80 alloy,

localised corrosion was observed in the a-Mg grains and it often

started around the Al–Mn inclusions. The grains with the aRb

lamellar aggregate did not show significant signs of corrosion.

The early stages of corrosion of the AZ91D alloy revealed

preferential corrosion in the centre of the a-Mg dendrites without

the Al–Mn inclusions having a noticeable effect.

After 28 days at 98% RH and 50 8C the cross-sectional

scanning electron micrographs of the test specimens (Fig. 6d–f)

revealed thin layers of corrosion products. The thickness of the

corrosion layer formed on the AZ31 alloy varied from 1 to 5mm

and it covered its entire surface. For the AZ80 and AZ91D alloys,

the corrosion products were mainly accumulated on the a-Mg

phases and their thicknesses were below 2mm.

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The cross-sections of the test materials after immersion in

Madrid water for 28 days (Fig. 7) revealed higher severe corrosion

attack compared with the specimens exposed to high humidity

atmospheres. Thus, the AZ31 alloy revealed general corrosion

with formation of a corrosion layer of 20mm thickness (Fig. 7a

and d) and the AZ80 and AZ91D alloys developed localised

corrosion with a maximum pit depth of 100 and 350mm,

respectively (Figs. 7b and 10c). Figure 7e and f shows that the

corrosion attack progressed through the a-Mg grains for the

AZ80 and AZ91D alloys and that the b-phase remained

uncorroded, acting as a barrier against corrosion.

3.3.3 XRD

Figure 8 discloses the low angle XRD (incident angle 18) study ofinvestigated alloys after the corrosion tests for 28 days. The XRD

peaks corresponding to corrosion products (Mg(OH)2and Mg5(CO3)4(OH)2 � 4H2O) presented higher intensity for

the specimens immersed in Madrid water compared with those

exposed to high humidity atmospheres. Peaks of CaCO3 were also

detected. This could be due to the high Ca concentration of

Madrid water (5.8–34.8mg/L Ca). The presence of metallic

magnesium in all the XRD patterns was caused by the effect of the



Figure 7. BSEmicrographs of the cross-sections of tested alloys after immersion inMadrid water at 22 8C for 28 days. (a and d) AZ31; (b and e) AZ80and (c and f) AZ91D

substrates and the low X-ray absorption of the thin layers of

corrosion products.

4 Discussion

The AZ31, AZ80 and AZ91D magnesium alloys presented an

excellent corrosion performance without visible signs of surface

degradation in atmospheres with an RH of 80%. The increase of

either the temperature or RH leads to degradation of the surface

appearance with formation of corrosion products mainly

consisting of Mg(OH)2 and Mg5(CO3)4(OH)2 � 4H2O. These

corrosion products are commonly found during atmospheric

corrosion of magnesium alloys [11]. The increasing corrosion

attack could be explained by an acceleration of the kinetics of the

electrochemical corrosion reactions with the temperature [3] and

by the increased thickness of the surface electrolyte layer with

the increase of RH [10]. Nevertheless, at 98% RH and 50 8C,corrosion rates of the test materials were still below those

observed during immersion in tap water at 22 8C. This higher

corrosion rate was possibly associated with the higher con-

ductivity of tap water compared to that of the surface electrolyte

layer in high humidity environments and/or the presence of a


diffusion layer continuously renewed during the immersion tests

that delays the increase of pH and the subsequent formation of

stable corrosion products.

The effect of the alloy composition and/or microstructure

was only noticeable for atmospheres with RH values above 90%,

where the alloys with higher aluminium content revealed smaller

surface degradation. The beneficial effect of aluminium is

normally attributed to several factors [12–14]. The findings of the

work suggested that this improvement was associated with the

higher aluminium content in the a-Mg phase of the AZ alloys,

compared with commercially pure Mg, and with the b-phase

acting as physical barrier against corrosion attack in the case of

the AZ80 and AZ91D magnesium alloys.

Surface characterisation of the AZ31 and AZ80 alloys after

exposure for short times in a high humidity atmosphere (98%

RH, 50 8C) revealed localised corrosion attack around the Al–Mn

intermetallic compounds, which revealed much higher surface

potential than the a-Mg phase. However, the small cathodic area

corresponding to the Al–Mn intermetallic compounds was not

decisive during the corrosion attack for increased times of

exposure. Thus, after 28 days, the AZ31 alloy presented general

corrosion with formation of a relatively thick layer of corrosion

products and the AZ80 alloy revealed deep pits associated with

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Figure 8. Low angle XRD patterns (incident angle 18) of test materials

after exposure to 98% RH–50 8C and Madrid water for 28 days

Figure 9.Metallographic cross-section in BSE mode of the AZ91D alloy

after exposure to 98% RH–50 8C for 28 days

Figure 10. SEM detail of the b-phase in the AZ91D alloy

Table 4. EDX analysis from Fig. 10

Spectrum Elements (wt%)

Mg Al Si Fe

1 63.98 35.93 0.05 0.05

2 87.62 12.26 0.07 0

3 91.24 8.64 0.08 0

4 93.51 6.37 0.10 0

5 94.90 5.03 0.07 0

preferential corrosion of a-Mg grains. For the AZ91D alloy,

corrosion started and progressed through the centre of the

primary a-Mg grains, whereas the b-phase/a-Mg interfaces did

not reveal signs of galvanic coupling as predicted by the surface

potential maps. In the case of the Al–Mn particles, galvanic

corrosion only occurred when these particles were situated at the

centre of the a-Mg grains, but not when they were near the b-

phase [15] (Fig. 9). These features were considered to be related to

the distribution of aluminium in the alloy. As illustrated in Fig. 10

and Table 4, aluminium content increased towards the grain

boundaries as a result of coring during solidification. Therefore,

the reduced aluminium content (5wt%) made the centre of a-Mg

grains more active, especially when Al–Mn particles were located

in these areas.

5 Conclusions

1. M

ww

g–Al alloys revealed good corrosion resistance in environ-

ments with RH up to 90% at room temperature. Higher

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temperature and RH lead to increased formation of corrosion

products consisting of brucite (Mg(OH)2) and hydromagnesite

(Mg5(CO3)4(OH)2 � 4H2O). Corrosion resistance of test mate-

rials exposed to 98% RH and 50 8C, was slightly higher than

that observed during immersion in tap water at 22 8C.
2. T he initiation of corrosion attack was localised at the Al–Mn/
a-Mg interfaces in the AZ31 and AZ80 alloys. Further

corrosion resulted in a uniform corrosion layer for the AZ31



�

alloy and deep pits for the AZ80 alloy due to dissolution of

a-Mg grains. For the AZ91D alloy, corrosion attack progressed

through the centre of a-Mg grains without considerable

galvanic corrosion at the a-Mg/b-Mg17Al12 interfaces. Intense

galvanic couples around the Al–Mn inclusions were only

found in areas with low aluminium content.

3. T
he AZ80 and AZ91D alloys revealed the highest corrosion
resistance of all tested materials due to higher aluminium

content in the composition of the alloy and to the b-Mg17Al12phase acting as physical barrier against corrosion.

Acknowledgements: The authors wish to thank MCYT (Project

MAT2006-13179-C02-01-02) and Comunidad deMadrid (ESTRU-

MAT_CM/MAT/77) for the financial support given to this work.

6 References

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[2] A. Froats, T. K. Aune, D. Hawke, W. Unsworth, J. E. Hillis,Corrosion of Magnesium and Magnesium alloys, Metal Hand-

2011 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

book, 9th edn., Vol. 13, ASM International, Materials Park,Ohio, USA, 1987.

[3] N. Le Bozec, M. Jonsson, D. Thierry, Corrosion 2004, 60, 356.[4] P. V. Strekalov, Prot. Met. 1993, 29, 673.[5] M. Jonsson, D. Persson, D. Thierry, Corros. Sci. 2007, 49,

1540.[6] A. Pardo, M. C. Merino, A. E. Coy, R. Arrabal, F. Viejo, E.

Matykina, Corros. Sci. 2008, 50, 823.[7] A. Pardo, M. C. Merino, A. E. Coy, F. Viejo, R. Arrabal, S.

Feliu, Jr., Electrochim. Acta 2008, 53, 7890.[8] D. Eliezer, P. Uzan, E. Aghion, J. Mater. Sci. 2003, 419–422,

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Easton, G. Song, A. Atrens, Corros. Sci. 2009, 51, 602.[13] R. Ambat, N. N. Aung, W. Zhou, Corros. Sci. 2000, 42, 1433.[14] S. J. Splinter, N. S. McIntyre, Surf. Sci. 1994, 314, 157.[15] M. Jonsson, D. Persson, C. Leygraf, Corros. Sci. 2008, 50,

1406.

(Received: September 29, 2009)

(Accepted: October 26, 2009)

W5538

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Corrosion behaviour of Mg/Al alloys in high humidity atmospheres

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