Effect of ball milling on the corrosion resistance of magnesium in aqueous media

Electrochimica Acta 49 (2004) 2461–2470

Effect of ball milling on the corrosion resistance of magnesium in aqueous media Marie-Hélène Grosjean a , Moussa Zidoune a , Lionel Roué a,∗ , Jacques Huot b,1 , Robert Schulz b,1 a

INRS-Énergie, Matériaux et Télécommunications, 1650 Boulevard Lionel-Boulet, Varennes, Que, Canada J3X 1S2 b HERA Hydrogen Storage Systems Inc., 577 rue Le Breton, Longueuil, Que., Canada J4G 1R9 Received 19 September 2003; received in revised form 27 January 2004; accepted 1 February 2004

Abstract The influence of the high-energy ball milling on the corrosion behavior of magnesium in aqueous media has been investigated through electrochemical experiments complemented by morphological, structural, chemical and surface analyses. The milling duration was varied from 0 to 40 h. Polarization curves show that the milling procedure improves the magnesium corrosion resistance in passive conditions (KOH solution) and in more active corrosion conditions (borate solution). This is illustrated by the corrosion potential which becomes nobler with milling. The variation of the polarization resistance and related corrosion current with milling time is also an indication of the improvement of the Mg corrosion resistance due to the milling. Moreover, the passive current is significantly lower for milled Mg. The Mg crystallite size is reduced from >100 to 34 nm after 10 h of milling and does not decrease significantly with further milling. The iron contamination of the Mg powder due to the erosion of the milling tools is very low (0.09 wt.% after 40 h of milling). In contrast, a significant oxygen contamination occurs during milling (2.6 wt.% after 40 h of milling). XPS and AES data indicate MgO enrichment in the bulk of the milled Mg without significant MgO increase at the powder surface. The corrosion improvement was attributed to the increase through the milling process of the density of surface defects and grain boundaries susceptible to increase the number of nucleation sites for Mg hydroxylation in aqueous media, leading to the rapid formation of a dense and protective Mg(OH)2 layer. © 2004 Elsevier Ltd. All rights reserved. Keywords: Magnesium; Nanocrystalline structure; Ball milling; Corrosion behavior; Passive film

1. Introduction Due to its low density and its low price, magnesium is a very attractive element for light alloys in the automotive and aerospace industries [1]. In addition, magnesium is able to absorb large quantities of hydrogen which makes it a particularly interesting element for the elaboration of Mg-based metal hydrides as hydrogen source for fuel cells [2] and as negative electrode for Ni–MH batteries [3]. However, application of Mg and its alloys is limited because of their high sensitivity to oxidation and their relatively low corrosion resistance. For example, the low cycle lifetime of ∗ Corresponding author. Tel.: +1-450-929-8185; fax: +1-450-929-8102. E-mail address: [email protected] (L. Rou´e). 1 Present address: Hydro-Qu´ ebec Research Institute, 1800 Boulevard Lionel-Boulet, Varennes, Que., Canada J3X 1S1.

0013-4686/$ – see front matter © 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2004.02.001

Mg-based electrodes in Ni–MH batteries (typically few cycles compared to several hundreds of cycles for commercial LaNi5 -based electrodes) is associated with the irreversible oxidation of the Mg-based alloy by the KOH electrolyte leading to the formation of a Mg(OH)2 layer on the surface. This consumes active material, affects the charge transfer across the alloy/electrolyte interface and acts as a barrier for hydrogen diffusion into and from the alloy [4]. Similarly, the lifetime of a hydrogen or heat generator based on the reaction of Mg with water is greatly reduced due to passivation of the magnesium by an unreactive MgO/Mg(OH)2 layer that interrupts the hydrolysis reaction [5]. The influence of environmental, structural and chemical factors on the corrosion behavior of magnesium and magnesium alloys has been largely investigated as indicated in several reviews [6–9]. The corrosion of magnesium in aqueous solutions proceeds by the reduction of water to produce magnesium hydroxide and hydrogen gas. TEM

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investigation on pure Mg indicated that the film formed in water has a three-layered structure, consisting of a hydrated inner layer, an apparently dense intermediate layer and an outer layer with platelet-like morphology [10]. The low corrosion resistance of Mg and Mg-based alloys in aqueous media is mainly related to (i) the high dissolution tendency of magnesium at pH less than about 11; (ii) the permeability and the low stability of the magnesium hydroxide film and (iii) the sensitivity of magnesium to impurities and alloying elements that promote internal galvanic attacks. Significant improvements in the corrosion resistance of Mg have been achieved by reduction of the metal impurity content or alloying with elements such as aluminium [11] or rare earth elements [12]. Further corrosion resistance improvements have been obtained with magnesium alloys elaborated by non-equilibrium methods such as rapid solidification processing [13–15], magnetron sputtering [16] or mechanical alloying [17]. These synthesis techniques induce several microstructural modifications (e.g. increased microstructure homogeneity, extended solid solubility of the alloying elements, formation of new corrosion resistance phase and metallic glass alloys) and thus, structural and chemical effects cannot be easily separated to determinate the origin of the higher corrosion resistance of these non-equilibrium compounds. The general conclusion is however that a highly corrosion resistance Mg-alloy is ideally a single-phase, chemically homogeneous system with passivating alloying elements in sufficient concentration [15]. In order to focus on the influence of the structure on the corrosion resistance of magnesium, we propose in this study, to compare the corrosion behavior of pure Mg in polycrystalline and nanocrystalline states. Nanocrystalline Mg was synthesized by ball milling. In this technique, repeated mechanical deformations caused by ball to powder collisions introduce strain in the powder and, as a result, the crystals fracture into smaller pieces down to the nanometer range [18]. Electrochemical investigations on polycrystalline and nanocrystalline magnesium were performed in aqueous alkaline media. The electrochemical tests were complemented by morphological, structural, chemical and surface analyses in order to correlate the Mg microstructure to its corrosion behavior.

2. Experimental 2.1. Sample preparation Polycrystalline magnesium consisted of pure elemental magnesium powder (99.8 wt.% purity, −325 mesh, Alfa Aesar). Nanocrystalline magnesium was prepared by high energy ball milling of polycrystalline Mg powder. The ball milling was performed with a SPEX 8000 ball mill. The vial and the balls were made in stainless steel. The vial has an internal diameter of 38.1 mm and a length of 47.6 mm, corresponding to a capacity of about 55 ml. 3 g of Mg

powder were introduced into the vial with two 11.1 mm and one 14.3 mm diameter steel balls, corresponding to a ball-to-powder mass ratio of 8:1. The vials were sealed with an O-ring under argon atmosphere in a glove box. The milling duration was varied from 30 min to 40 h. 2.2. Sample analyses The oxygen and nitrogen contents of the different powders were measured by inert gas fusion technique with a TC-600 oxygen/nitrogen detector from LECO. The detection limit of the apparatus is 10 ppm and the accuracy of the analysis is 5% of the measured value. The iron content in the powder was determined by neutron activation analysis. The detection limit of the analysis is 50 ppm and its accuracy is 5% of the measured value. X-ray diffraction (XRD) was performed on a Bruker AXSD8 Siemens diffractometer with Cu K␣ radiation. Scanning electron microscopy (SEM) observations were made using a Jeol JSM-6300F microscope. The Auger electron spectroscopy (AES) depth profile analysis was performed on a PHI 660 from Perkin Elmer with a sputtering rate of ∼15 nm min−1 . X-ray photoelectron spectroscopy (XPS) measurements were performed with a VG Escalab 220I-XL equipped with an Al K␣ monochromatic source under ultra-high vacuum conditions (10−8 Pa). The C 1s peak of the adsorbed carbon at 284.6 eV was used as an internal reference to calibrate the peak position. 2.3. Electrochemical measurements The working electrode was a pellet made by cold pressing 500 mg of Mg powder in a 16 mm diameter stainless steel cylindrical die. A load of 10,000 kg cm−2 was applied during 10 min leading to a nearly fully dense pellet without apparent porosity (confirmed by SEM observations). The pellets were polished with an abrasive solution (methanol and alumina powder) over a polishing cloth, starting with alumina grains size of 3, 1 ␮m and finally 0.3 ␮m. Then, they were rinsed with methanol, dried in air and stored under an inert atmosphere of argon before use. This procedure permits to obtain electrodes with a similar effective surface independently of the particle size distribution of the powder. The polished pellets were mounted on a glass tube with an epoxy resin. The electrical contact was made by sticking a Cu wire to the back of the pellet with a silver epoxy resin. Only one face of the pellet was in contact with the electrolyte. The geometric area of the working electrode exposed to the electrolyte is 2 cm2 . Electrochemical measurements were performed at room temperature (23±1 ◦ C) using a potentiostat/galvanostat/FRA VoltaLab 40, type PGZ301 (Radiometer Analytical) in 1 M KOH solution (pH = 14) and in borate buffer solution (0.3 M Na2 B4 O7 /H3 BO3 , pH = 8.4) using a three electrodes cell. The counter electrode was a Pt wire placed in a separate compartment. The reference electrode was either a Hg/HgO (KOH 1 M) electrode for tests in the KOH solution

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or a saturated calomel electrode for tests in borate solution. A Luggin capillary was used to reduce the Ohmic drop. All measured potential values were referred to the saturated calomel electrode (SCE). Before each experiment, oxygen was removed from the solution by bubbling with nitrogen and a nitrogen flux was maintained over the electrolyte during the experiment. After 1 h in open circuit conditions, polarization curves started at −0.3 V with respect to the open circuit potential and progressed in the positive direction at 0.5 mV s−1 . For each curve, the reproducibility of the measurement was confirmed. The polarization curves were corrected for the uncompensated resistance determined by ac impedance spectroscopy. The corrosion parameters were calculated according to the first and the second Stern methods [19,20].

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3. Results and discussion 3.1. Powder morphology SEM observations of the Mg powders indicate that milling affects the morphology of the powders as illustrated in Fig. 1, which shows typical micrographs of unmilled and 40 h milled Mg powders. The milled powders consist of irregular particles with a grain size distribution varying from few ␮m to several tens of ␮m. The large particles are due to the well-known cold welding of the particles occurring concomitantly to the repeated fracturing of the powder during milling. The mean diameter of the particles was estimated from SEM observations and plotted versus milling time in Fig. 2. The curve indicates a decrease of the powder

Fig. 1. Scanning electron micrographs of (a) unmilled and (b) 40 h milled Mg powders.

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M.-H. Grosjean et al. / Electrochimica Acta 49 (2004) 2461–2470 Table 1 Crystallites size and strain (estimated from a Rietveld refinement analysis of the X-ray patterns) for Mg powders as a function of the milling time Milling time (h)

Crystallite size (nm) Strain (%)

Fig. 2. Mean particle size (estimated from SEM observations) of the Mg powder vs. the milling time.

size with increasing milling time. After 40 h of milling, a steady state is nearly reached with a mean particle size equals to 20 ␮m compared to 40 ␮m for the unmilled powder. By assuming that particles have a spherical shape, the estimated specific surface area of the 40 h milled powder is ca. 0.17 m2 g−1 compared to 0.085 m2 g−1 for unmilled Mg powder. 3.2. Powder structure Fig. 3 shows the X-ray diffraction patterns of Mg powder as a function of the milling time. The position of the diffraction peaks remains constant indicating that the lattice parameters are unchanged by the ball milling, i.e. solid solution resulting from possible contamination with milling tools does not occur. On the other hand, extra peaks attributed to

0

0.5

3

10

40

>100

63

42

34

31

0.000

0.000

0.018

0.036

0.056

a MgO phase are discernable and their intensity increases slightly with prolonged milling. Referring to the relative intensity of the diffraction peaks, the texture of the Mg powders is unmodified by the milling. However, an exception is observed for the Mg powder milled 30 min which shows an increase of the relative intensity of the Mg(0 0 2) peak at 34.4◦ . Its relative intensity is around 70 compared to 30 for the other samples. The increase of the Mg(0 0 2) peak intensity was also observed with unmilled Mg in pellet form (i.e. pressed powder). These observations indicate that a preferential orientation along the c-axis is induced during the initial stage of milling and by cold compaction of the unmilled powder. It may be connected to the fact that energy required to carry out a deformation by slip perpendicular to the c-axis is very weak or at least, much weaker than in another direction. When low deformation energy is involved such as at the beginning of milling or during cold pressing of the initial powder, dislocations generate mainly deformations by slip perpendicular to the c-axis and a texture (0 0 2) appears. During prolonged milling (t ≥ 3 h), large energy is accumulated into the powder and dislocations occur in all directions, i.e. deformations are random and the texture disappears. A Rietveld refinement analysis of the X-ray patterns was performed to determine the crystallite sizes and the internal strain and the results are presented in Table 1. The crystallite

Fig. 3. X-ray diffraction patterns of Mg powder as a function of the milling time.

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size is reduced to 34 nm after 10 h of milling and does not decrease significantly with further milling. This is consistent with results found by Hwang et al. [21] indicating a stabilization of the Mg crystallite size to about 40 nm after 6–7 h of milling. The relatively high limit value of the Mg crystallite size compared to other ball-milled metals can be related to its low melting temperature which facilitates the recovery process rather than the plastic deformation of the material during the milling process [21,22]. The decrease of the crystallite size is accompanied by the increase of the internal strain associated to the presence of lattice distortion at the grain boundaries and the accentuation of the dislocation density with prolonged milling. However, one must note that the maximum strain measured after 40 h of milling is low (0.056%), confirming the high recovery rate of magnesium during milling. 3.3. Powder contamination Fig. 4 shows the variation of the concentration of contaminants (O, N, and Fe) into the Mg powder as a function of the milling time. The iron contamination of the Mg powder due to the erosion of the vial and the balls increases linearly with milling time but remains very low (ca. 0.09 wt.% after 40 h milling). It is much lower than normally found in most of the powders milled with steel grinding medium (typically, 1–4 wt.% Fe [18]). This can be explained by the low strength/hardness of the Mg powder compared to the steel milling tools. Although the container is carefully sealed with an O-ring under an argon atmosphere, a contamination during milling cannot be avoided completely. Based on the nitrogen content, air contamination appears very limited (N content is inferior to 0.2 wt.% after 40 h of milling). However, the oxygen contamination is much higher than expected from the N content due to the much larger affinity of Mg to oxygen. Moreover, O contamination does not increase linearly with milling time as observed for the N contamination but

Fig. 4. Variation of the concentration of contaminants (O, N, and Fe) into the Mg powder as a function of the milling time.

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rather increases abruptly in a first step and then saturates around 2.6 wt.% after 10 h of milling. Assuming that oxygen reacts with Mg to form a MgO phase, the oxygen content corresponds to an equivalent proportion of MgO varying from 1.0 wt.% for unmilled powder to a maximum value of 6.5 wt.% for 10 h and 40 h milled powders. The O contamination may be related to the oxidation of the powder surface with ambient atmosphere once the sample has been taken out of the vial. However, based on the previous estimation of the surface area increase by milling (see SEM observations), it cannot explain the totality of the O contamination. So, it appears clearly that a large part of the Mg oxidation occurs during milling. In other words, the variation of the O content with milling duration in Fig. 4 reflects the increasing sensitivity of magnesium to oxidation through the fracturing process which provides fresh and highly reactive surface to be oxidized before subsequent cold welding event. This implies that oxygen is incorporated in the bulk of the material as confirmed by AES depth profile analysis (see below). The source of oxygen during milling can be residual/leaking air present in the container and more largely, oxides present at the surface of the milling tools. 3.4. Electrochemical behavior The variation of the open circuit potential value (EOCP ) with immersion time in 1 M KOH solution for unmilled and 10 h milled magnesium electrodes is given in Fig. 5. The change in EOCP (equivalent to the free corrosion potential) is assumed to be related to the growing of the Mg(OH)2 passive layer onto the electrode. For unmilled Mg electrode, EOCP increases sharply during the first 5 min and shifts slowly for further immersion time to reach a value of −1.41 V versus SCE after 1 h in KOH solution. For milled Mg electrode, the open circuit potential starts at a more positive value, increases more drastically during the first 10 min and reaches progressively −1.30 V versus SCE after 1 h of immersion. The larger EOCP variation for milled Mg may indicate a

Fig. 5. Variation of the open circuit potential with immersion time in 1 M KOH solution for unmilled and 10 h milled magnesium electrodes.

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Fig. 6. Polarization curves of unmilled and 10 h milled magnesium electrodes in 1 M KOH solution. Scan rate = 0.5 mV s−1 .

Fig. 7. Corrosion current density (icorr ) of Mg electrodes in KOH solution as a function of milling time. icorr was calculated using the first Stern method.

greater reactivity with the electrolyte to form the Mg(OH)2 passive layer. The significantly nobler EOCP value for milled Mg in the initial state as well as in the near steady state illustrate the positive effect of the ball milling process on the corrosion resistance of Mg. Fig. 6 shows typical polarization curves of unmilled and 10 h milled magnesium electrodes in 1 M KOH solution. The curves have a similar shape with no active–passive transition observed on the anodic part of the polarization curves, confirming the passive behavior of magnesium in 1 M KOH solution independently of its microstructure. However, a significant shift in the positive direction of the corrosion potential is observed with milled Mg, confirming the previous OCP measurements. Moreover, current densities in the passive region are much lower than those measured for unmilled Mg, indicating the formation of a more protective surface film on milled Mg electrode. Finally, the anodic current observed from about 0.5 V versus SCE, which is related to the oxygen evolution reaction with possible concomitant breakdown of the passive film, starts at a lower potential and increases more drastically for milled Mg electrode. The corrosion parameters deduced from polarization measurements in 1 M KOH solution are summarized in Table 2 as a function of the milling time. It shows that the corrosion potential (Ecorr ) becomes more positive with milling,

the corrosion current density (icorr ) decreases and the polarization resistance (Rp ) increases when milling is prolonged. Furthermore, the passive current density (ipass ) measured at −0.2 V versus SCE decreases substantially with increasing milling time. These data confirm the improvement of the corrosion resistance of magnesium with milling. It is more evident in Fig. 7 where icorr is plotted versus the milling duration. The corrosion resistance improvement is already observable after 30 min of milling and becomes more apparent with increasing milling time. No further improvement is observed after 10 h of milling, indicating that a stabilized state is reached. On the other hand, the apparent breakdown potential (Eb ) becomes less positive with milling, suggesting an easier rupture of the passive film. However, the shift of Eb may also be related to an improvement of the oxygen evolution reaction and may reflect the better electronic transfer through the passive film due to its small thickness which facilitates electron tunnelling. In order to confirm the positive effect of milling on the Mg corrosion resistance, polarization experiments were performed in borate solution (pH = 8.4), i.e. in active corrosion conditions. As shown in Fig. 8, the anodic polarization curve for unmilled magnesium electrode presents a current peak characteristic of an active–passive transition around

Table 2 Corrosion parameters of Mg electrodes in 1 M KOH as a function of milling time Milling time (h)

Ecorr (V vs. SCE)

icorr (␮A cm−2 )a

Rp (k cm2 )b

i corr ; (␮A cm−2 )c

ipass (␮A cm−2 )d

Eb (V)

0 0.5 3 10 40

−1.40 −1.35 −1.29 −1.28 −1.30

23 14 7 3 4

2.5 3.9 7.1 13.2 9.7

20.8 13.4 7.3 3.9 5.4

63 25 16 12 12

0.56 0.53 0.45 0.46 0.47

a b c d

Estimated by extrapolation of the Tafel slopes (first Stern method). Estimated from the tangent slope of the parabola at the zero-current potential (second Stern method). Calculated from Rp values using the Stern–Geary relationship with the Tafel coefficients bc = 120 mV dec−1 and ba = ∞. Measured at −0.2 V vs. SCE.

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Fig. 8. Polarization curves in borate buffer solution (pH = 8.4) for unmilled and 10 h milled magnesium electrodes. Scan rate = 0.5 mV s−1 .

−1.4 V versus SCE. A pseudo-passive region appears from ca. −0.6 V versus SCE with relatively high passive current density. The polarization curve of 10 h milled magnesium electrode also presents an active–passive transition in the same potential region but the active dissolution current peak is drastically decreased (idiss = 1.7 mA cm−2 compared to 4.7 mA cm−2 for unmilled Mg). The reduced passive current density measured at 0.2 V versus SCE (ipass = 0.8 mA cm−2 for milled Mg compared to 1.7 mA cm−2 for unmilled Mg) and the positive shift of the corrosion potential (Ecorr = −1.59 V for milled Mg and −1.71 V versus SCE for unmilled Mg) also confirm the corrosion resistance improvement of Mg in borate solution through the milling process. However, the corrosion current density value determined using the first Stern method is almost unmodified by the milling (icorr ≈ 0.7–0.8 mA cm−2 ) and the polarization resistance increases only slightly with milling (Rp = 63  cm2 for milled Mg compared to 42  cm2 for unmilled Mg). This may be related to the fact that the corrosion of Mg in borate solution is controlled by the cathodic reaction (i.e. reduction of water causing hydrogen evolution) which is unmodified by the milling process.

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Fig. 9. Mg 2p XPS peak of unmilled Mg, (a) exposed to air, (b) immersed in water for 3 h and dried in air and (c) immersed in 1 M KOH for 3 h, rinsed with water and dried in air.

• As seen in Figs. 9b and 10b, the Mg 2p lines for Mg exposed to water are combinations of Mg oxide and hydroxide peaks. No apparent difference was observed between milled and unmilled Mg. • For Mg exposed to 1 M KOH (Figs. 9c and 10c), the mid position of the Mg 2p lines is at lower binding energies compared to water, which is indicative that the MgO fraction is larger in the KOH case. No significant difference was observed between milled and unmilled Mg. The comparative Auger depth profiles of unmilled and 10 h milled Mg are represented in Fig. 11 as a function of the exposure media (air, water and 1 M KOH). When Mg and O atomic concentrations are equal to 50%, MgO is assumed to be formed and for Mg and O atomic concentrations close to 66 and 33%, respectively, Mg(OH)2 is considered to be present.

3.5. Surface analysis XPS surface analysis focused on the Mg 2p spectrum is represented in Fig. 9 for unmilled and in Fig. 10 for 10 h milled Mg as a function of the exposure media (air, water and 1 M KOH). These results are summarized as follows: • For unmilled and 10 h milled Mg exposed to air, the major compound present at the surface of the sample is MgO. However, the Mg 2p spectrum exhibits also a significant contribution from the metallic state (at 49.6 eV), indicating that the MgO layer is thin. In addition, the intensity ratio between the XPS peak of metallic Mg and MgO is higher in the case of the 10 h milled Mg (Fig. 10a) compared to the unmilled Mg (Fig. 9a), suggesting a thinner MgO layer onto milled Mg.

Fig. 10. Mg 2p XPS peak of Mg milled for 10 h, (a) exposed to air, (b) immersed in water for 3 h and dried in air and (c) immersed in 1 M KOH for 3 h, rinsed with water and dried in air.

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Fig. 11. Auger depth profiles of (a) unmilled Mg exposed to air, (b) Mg milled for 10 h and exposed to air, (c) unmilled Mg immersed in water for 3 h and dried in air, (d) 10 h milled Mg immersed in water for 3 h and then dried in air, (e) unmilled Mg immersed in 1 M KOH for 3 h, rinsed with water and dried in air and, (f) 10 h milled Mg immersed in 1 M KOH for 3 h, rinsed with water and dried in air.

Fig. 11a and b shows the Auger depth profiles for the Mg and oxygen peaks for unmilled and 10 h milled Mg exposed to air, respectively. Fig. 11a shows the surface of the unmilled Mg is composed primarily of MgO. The thickness of the oxide layer is about 50 nm. In the case of the 10 h milled sample, Fig. 11b, the concentration of metallic Mg right at

the surface is higher than in the previous case. It may be related to the fact that the milling process creates fresh surfaces by shear deformations and these oxide free surfaces and interfaces are brought to the surface in higher concentration by the sample preparation procedure. However, Fig. 11b shows also that, despite the fact that the oxygen content is

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lower right at the surface, decrease in O concentration with sputtering depth is more gradual and its overall concentration in the bulk of the 10 h milled sample is higher than that in the case of the unmilled Mg. This is due to the fact that all oxygen on the surface of the initial particles and the additional oxygen picked during the milling process by contamination are dispersed throughout the sample during the intensive milling and, as a result, the bulk value of oxygen in the milled sample is much higher than in the unmilled case. This is in accordance with previous discussion in the powder contamination section. However, it must be noted that the O bulk values determined from AES measurements is higher than measured from O contamination analysis. It may be due to surface re-oxidation during sputtering, considering that Mg oxidizes rapidly even under ultra-high vacuum conditions [23,24]. Fig. 11c and d shows the corresponding Auger profiles for the unmilled and milled Mg after immersion in water for 3 h. The figures show surfaces primarily composed of Mg(OH)2 and MgO, magnesium hydroxide being located closer to the surface. The total thickness of the surface layers is about 500 nm in the case of unmilled Mg and 350 nm in the case of the 10 h milled Mg. More than half of these layers is composed of hydroxide. This difference in thickness is probably associated with the better corrosion resistance of the milled sample. It is interesting to point out that the oxygen content in the bulk of the milled Mg is, as before, significantly higher that that of the unmilled Mg. The same is true for the samples immersed in 1 M KOH for 3 h. Fig. 11e and f shows the Auger depth profiles. The results are similar to the previous case but the thickness of the layers are about half that of the previous case. The fact that the oxy-hydroxide layer is thinner in KOH solution than in water is not surprising considering the higher tendency of Mg to passivate in strong alkaline solution as indicated in the electrochemical tests.

4. Summary and conclusion Our observations indicate clearly that mechanical milling improves the corrosion resistance of magnesium in passive conditions (KOH solution) as well as in more active corrosion conditions (borate solution). This statement is firstly illustrated by the corrosion potential which is shifted to more positive value with milling (see Table 2 and Fig. 5). Secondly, the variation of the polarization resistance and related corrosion current with milling time (see Table 2 and Fig. 7) is also an indication of the improvement of the Mg corrosion resistance due to the milling. Thirdly, the passive current is significantly lower for milled Mg (see Table 2, Figs. 6 and 8), reflecting the formation of a more protective passive film. In addition, our recent electrochemical impedance investigations of unmilled and 10 h milled Mg [25] show that the variations with immersion time of the interfacial resistance and capacitance resulting from charge

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transfer and film effect are less accentuated and tend more rapidly to a steady state with milled Mg suggesting a greater kinetic for the formation of the passive layer. Chemical analysis indicates a MgO enrichment in the milled powders (see Fig. 4), which seems to be essentially related to the powder oxidation during the milling. Even if MgO reacts with water to form Mg(OH)2 [26], it could contribute to the improved corrosion resistance of milled Mg since MgO is thermodynamically much less reactive than Mg in the presence of water (the standard enthalpy of reaction of MgO with water to form Mg(OH)2 is −37 kJ mol−1 compared to −353 kJ mol−1 for the standard enthalpy of reaction of Mg with water to form Mg(OH)2 + H2 [27]). However, XPS and AES data indicate that MgO enrichment occurs in the bulk of the material and no MgO enrichment is detected at the electrode surface (see Figs. 9–11). Moreover, a significant corrosion improvement appears after only 30 min of milling (see Fig. 7) despite a non-apparent increase of the O content in the sample (see Fig. 4). Thus, an additional positive effect of the ball milling on the Mg corrosion resistance must be involved. Few studies have reported the improved corrosion resistance of nanostructured materials. For example, nanocrystalline Ni made by electrodeposition [28] or by ball milling [29], nanocrystalline Ru prepared by ball milling [30] and nanocrystalline Fe produced by ball milling-hot compaction [31] exhibit a better corrosion resistance than their polycrystalline counterpart. Inturi and Szklarska-Smialowska [32] have also observed superior localized corrosion resistance in NaCl for nanocrystalline sputtered stainless steel films compared to the conventional material with the same composition. The better corrosion behavior of the nanocrystalline sputtered film is explained by the presence of a large number of uniformly distributed defects (primarily located at the grain boundaries), which results in a high degree of distribution of chloride ions on the metal surface preventing localized chloride enrichment and subsequent acidification. In addition, studies on stainless steel modified by sandblasting-annealing treatment demonstrate that the nanocrystalline surface has higher resistance to corrosion, greater capability to repassivation and higher chemical stability [33]. This is supported by Grake et al. [34] who argue that the grain boundaries and dislocations (provided by a fine-grain microstructure and/or surface working) act as easy-diffusion paths and thus, lead to rapid formation of a protective layer. Additionally, the study of the interaction of water vapour with magnesium indicated that the nucleation of oxide is faster on the more open-packed faces, where the surface atom coordination number is smaller [35]. On the other hand, one must note that in the case of non-passivating materials such as Ni–P alloys in H2 SO4 solution [36] or Co–Cu alloys in borate solution [37], the nanocrystalline character of the material has rather a negative effect on the corrosion resistance, probably for the same reason, i.e. faster atomic diffusion at the grain boundaries leading to a higher anodic dissolution of the metal. One must also point out the

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works of Burke et al. [38] which indicate the importance of the surface defects, e.g. adatoms, vacancies, microclusters, grains boundaries, etc., which constitute superactive sites for various surface electrochemical processes. So, in the present study, we suggest that the increase through the milling process of the density of surface defects and grain boundaries with atoms having low coordination numbers may increase the number of nucleation sites for Mg hydroxylation in addition to accelerate the atomic diffusion for the growing of the Mg(OH)2 film. This is likely to facilitate the rapid formation of a dense protective passive layer. This assumption is in accordance with our observations indicating that the oxy-hydroxide passive film formed onto the milled Mg electrode has a lower thickness than the one formed on unmilled Mg but a higher corrosion resistance; this apparent contradiction being explained by the denser (i.e. less permeable) structure of the passive layer formed on milled Mg. Additional investigations such as TEM or AFM/STM characterizations of the passive layer will be useful to confirm this hypothesis. At last, our observations are consistent with recent studies showing the improved corrosion behavior in alkaline media of nanocrystalline Zn deposit [39] and nanocrystalline (Ni70 Mo30 )90 B10 alloy [40] due to the formation of a more uniform/complete and thus more effective passive film.

Acknowledgements We thank E. Irissou for support in the Rietveld refinement analyses. This work has been financially supported by HERA Hydrogen Storage Systems Inc. and the National Sciences and Engineering Research Council (NSERC) of Canada through a Collaborative R&D grant.

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