Corrosion behaviour of the amorphous Mg65Y10Cu15Ag10 alloy

Corrosion behaviour of the amorphousMg65Y10Cu15Ag10 alloy

R.V. Subba Rao a,b, U. Wolff a,1, S. Baunack a, J. Eckert a,A. Gebert a,*

a IFW Dresden, P.O. Box 270016, D-01171 Dresden, Germanyb Indira Gandhi Centre for Atomic Research, Kalpakkam 603 102, Tamil Nadu, India

Received 10 March 2002; accepted 10 June 2002

Abstract

The corrosion behaviour of the amorphous Mg65Y10Cu15Ag10 alloy as well as of its crys-

talline multiphase counterpart was studied in alkaline electrolytes and compared with that of

the amorphous Mg65Y10Cu25 alloy. Electrochemical investigations were carried out in 0.3 M

H3BO3/Na2B4O7 buffer solution with pH ¼ 8:4 and in 0.1 M NaOH solution with pH ¼ 13.Tafel plots were recorded and cyclic potentiodynamic polarisation tests were conducted,

transients were measured at anodic potentials. Potentiostatically formed surface layers were

characterised by Auger electron spectroscopy and atomic force microscopy. Changes in the

corrosion behaviour were noticed which are attributed to the presence of silver. The passive

layers formed in the two electrolytes were quite different in the composition as well as in

morphology. The layer growth mechanisms also showed some variation presumably mainly

due to the presence of silver, though copper still seems to play a dominant role in the passi-

vation of this alloy in the weakly alkaline solution. The amorphous alloys displayed superior

corrosion behaviour compared to the crystalline alloy, because of the absence of the hete-

rogeneties existing in crystalline alloys.

� 2002 Elsevier Science Ltd. All rights reserved.

Keywords: Magnesium; Metallic glass; AES; Cyclic voltammetry; Passive films

www.elsevier.com/locate/corsci

Corrosion Science 45 (2003) 817–832

*Corresponding author. Tel.: +49-351-4659275; fax: +49-351-4659541.

E-mail address: [email protected] (A. Gebert).1 Present address: Materials Research Department, Risoe National Laboratory, Roskilde, Denmark.

0010-938X/03/$ - see front matter � 2002 Elsevier Science Ltd. All rights reserved.

PII: S0010-938X(02 )00131-2

mail to: [email protected]

1. Introduction

Magnesium-based alloys already find applications in varied fields such as auto-

mobile industries, sporting goods fabrication and to some extent aerospace industry

owing to their low specific weight combined with high strength [1]. Recent advances

in the area of amorphous bulk alloys [2] have sparked the interest in many Mg-basedalloy systems such as Mg–Y–Cu and Mg–Y–Ni [3,4]. Emphasis on bulk production

with larger diameters has motivated the search for alloying additions that fur-

ther enhance the glass-forming ability. Partial replacement of Cu with Ag in the

Mg65Y10Cu25 alloy was found to increase the glass-forming ability as silver reduces

the melting temperature without affecting the other criteria for bulk metallic glass

preparation [2]. Thus, a quaternary amorphous Mg65Y10Cu15Ag10 alloy was devel-

oped which could be prepared with larger diameters (up to 10 mm diameter) com-

pared to the ternary system [5,6]. Corrosion studies were carried out on the ternaryalloy [7], highlighting the superiority of corrosion properties of these amorphous

bulk alloy systems over those of magnesium and also conventional Mg-based alloys.

The current study on the quaternary alloy with the addition of 10 at.% of silver is

a sequel to the reported ones [7,8] and is aimed at characterising its corrosion

behaviour in comparison to that of the ternary alloy in alkaline solutions.

2. Experimental

Amorphous ribbons of the Mg65Y10Cu15Ag10 alloy were prepared by melt spin-

ning as reported in [8]. The microstructure was checked by means of X-ray dif-

fraction, optical microscopy and electron microscopy. Electrochemical investigationswere carried out in a glass cell with a platinum counterelectrode and a standard

saturated calomel electrode (SCE) with 240 mV(SHE). Potentiodynamic polarisation

tests in both, linear and cyclic modes, and potentiostatic current transient mea-

surements were conducted using a Potentiostat/Galvanostat Meinsberg PS6 and a

Solartron Analytical 1287 Electrochemical Interface. The electrolytes were 0.3 M

borate buffer solution (H3BO3/Na2B4O7) with a pH of 8.4 and 0.1 M sodium hy-

droxide solution (NaOH) with a pH of 13. These solutions were nitrogen purged for

1 h before each measurement. The specimens were subjected to a pre-reduction at)1.5 V (SCE) for 1 min to ensure a similar surface condition before each measure-ment. The electrode behaviour was investigated by recording Tafel characteristics at

a scan rate of 0.5 mV/s, by cyclic potentiodynamic polarisation measurements at a

scan rate of 5 mV/s (three cycles) and by current transient measurements at anodic

potentials in the passive region. Anodically treated surfaces were characterised by

Auger electron spectroscopy (AES) for obtaining elemental as well as oxide (using

factor analysis approach) depth profiles and by atomic force microscopy (AFM) for

surface topographical information. AES depth profiles were obtained in a PHI 660Auger microprobe (Physical Electronics, USA) with primary electrons of 10 keV,

100 nA and an energy resolution of 0.6%. The samples were sputtered by a ra-

stered beam of argon ions at 1.5 keV. The resulting sputter rate in silicon oxide was

818 R.V. Subba Rao et al. / Corrosion Science 45 (2003) 817–832

3.1 nm/min. The chemical composition of the electrolyte after the potentiostatic

current transient measurement was analysed and determined by inductively coupled

plasma mass spectrometry (ICPMS).

A cylindrical bulk sample of 7 mm diameter, prepared by die-casting, was used as

multiphase crystalline counterpart material to compare the results with those ob-

tained for the amorphous Mg65Y10Cu15Ag10 alloy ribbons. All the electrochemicaltests were carried out on the cross-sectional area of this rod material.

3. Results and discussion

Microstructural investigations on the melt spun ribbons and on the die cast rods

revealed similar features to those observed for the ternary Mg65Y10Cu25 alloy [8]. The

rapidly quenched material is fully amorphous, except for some few single inclusions

identified as Y2O3 particles which are formed during the preparation process. Fromthe DSC trace, recorded at 20 K/min, supercooled liquid region of DTx ¼ 40 K wasdetermined which correlates well with data reported in literature [5]. The cast rod

consists of a variety of intermetallic phases, some of which could be identified as the

Mg2Cu main phase and MgY. The role of silver in the crystalline phase formation

could not be clearly identified, yet.

Fig. 1 shows the results of cyclic voltammetric measurements on the amorphous

ribbons of the Mg65Y10Cu15Ag10 and the Mg65Y10Cu25 alloys and also on the crys-

talline Mg65Y10Cu15Ag10 sample in the borate buffer solution with pH ¼ 8:4. The

Fig. 1. Cyclic potentiodynamic polarisation tests performed in 0.3 M borate buffer solution with pH ¼ 8:4at a potential sweep rate of 5 mV/s (the first of three cycles).

R.V. Subba Rao et al. / Corrosion Science 45 (2003) 817–832 819

curves of the Mg65Y10Cu15Ag10 alloy samples show certain differences when com-

pared with that of the Mg65Y10Cu25 alloy. The free corrosion potential (OCP) is

shifted considerably to more positive potentials by addition of silver indicating the

alloy tending to be more noble. The characteristic peak (I) of copper oxidation [8]

was found to be shifted slightly in positive direction. In addition, at about 450

mV(SCE) a peak (II) occurs which is attributed to oxidation of silver (Ag2O/AgO)[9]. A corresponding reduction peak (III) is indicated in the reverse scan. It can be

seen that, for the amorphous alloy ribbons the passive current density of the qua-

ternary alloy after the silver oxidation peak (II) is similar to that of the ternary alloy

at these potentials. The current densities in the cathodic regime are considerably

lower and hydrogen evolution occurs at more negative potentials indicating that

cathodic surface reactions are more inhibited at the quaternary alloy surface. Also,

the anodic oxygen evolution reaction appears to be more pronounced and a re-

duction peak close to 800 mV(SCE) in the reverse scan may indicate further oxi-dation of the silver oxide, formed at (II), at higher potentials in the forward scan.

Potential–pH diagrams (Pourbaix diagram) [9] of the silver–water system indicate

that the higher valency oxide (Ag2O3) is unstable in the presence of water and be-

comes a site of evolution of oxygen. The higher current densities in this region could

be explained by the solubility of AgO on the verge of its further oxidation to Ag2O3.

The multiphase crystalline Mg65Y10Cu15Ag10 alloy exhibits higher current densi-

ties throughout the whole cycle (Fig. 1). The silver oxidation peak (II) appears a bit

diffused compared to that of the amorphous alloy and the copper oxidation peak (I)appears a bit more pronounced compared to the corresponding peak of the amor-

phous alloy. This indicates the differences of the activities of the various crystalline

phases in the multiphase system leading to higher current upsurges as the scan is

progressing.

A similar effect of enhanced activity of the crystalline alloy is obvious from the

cyclic voltammograms obtained when the materials are tested in 0.1 M sodium

hydroxide electrolyte of pH 13. Typical results are summarised in Fig. 2. Both, the

copper oxidation peaks (I and II) and the silver oxidation peaks (III and IV) alongwith their reduction counterparts in the reverse scan are prominently seen in the

curve of the multiphase crystalline material. The curves pertaining to the amorphous

alloy ribbons of ternary and quaternary Mg-based alloys significantly differ from

this. But, both exhibit a similar shape which is quite similar to the curve shape of

pure magnesium in this electrolyte [8]. The curve of the quaternary amorphous

alloy, however, shows an oxidation shoulder before the oxygen evolution reaction

occurs, which is a little more pronounced compared to that in the curve of the

ternary alloy. The occurrence of cathodic reduction peaks without the presence ofprominent oxidation peaks (confirmed by reversing the cycle even before the onset of

copper oxidation peak) suggests that thin films of Ag2O formed during the immer-

sion might have been converted to the soluble form of AgO� ions in this highly

alkaline electrolyte [10]. It has also been observed that when amorphous ribbons of

the quaternary alloy were immersed in electrolyte of pH 8.4, the typical brownish

golden film indicative of Ag2O [9] is formed whereas such a colour change is not

visible upon the immersion in the electrolyte of pH 13. Both, the oxygen evolution


reaction and hydrogen evolution reaction are slightly less inhibited for the amor-phous quaternary alloy compared to the ternary one perhaps due to the higher

solubility of the silver oxides compared to the copper oxides. The multiphase crys-

talline quaternary alloy shows considerably higher activity for anodic and cathodic

reactions compared to the amorphous alloy. This effect seems to be much more

pronounced in the strongly alkaline electrolyte (Fig. 2) compared to the weakly al-

kaline electrolyte (Fig. 1).

Potentiodynamic polarisation tests with slower scan rate of 0.5 mV/s were carried

out to estimate the corrosion potential (Ecorr) and corrosion current density (icorr)values from the Tafel plot extrapolations. Fig. 3 shows the typical Tafel character-

istics of the quaternary amorphous Mg65Y10Cu15Ag10 alloy as well as of the multi-

phase crystalline Mg65Y10Cu15Ag10 alloy in both electrolytes. The amorphous alloy

shows clearly more positive corrosion potentials and lower corrosion current densities

than the multiphase crystalline alloy. In both electrolytes, anodic passive current

densities of the amorphous alloy are significantly lower than those of the crystalline

alloy suggesting the formation of more uniform and dense and hence, more protective

surface layers under these conditions. This is a clear indication that the single phase,chemically homogeneous nature of the amorphous Mg65Y10Cu10Ag15 alloy provides

significantly increased corrosion resistance. The results of these investigations on

quaternary alloys as well as parameters for the ternary alloys from [8] are summarised

in Table 1. Comparing the amorphous alloy ribbons, the quaternary alloy clearly

shows a superior behaviour with more noble Ecorr and smaller icorr values compared tothe ternary alloy in both electrolytes. This could be clearly attributed to the noble

Fig. 2. Cyclic potentiodynamic polarisation tests performed in 0.1 M sodium hydroxide solution with

pH ¼ 13 at a potential sweep rate of 5 mV/s (the first of three cycles).


nature of silver. However, this advantageous effect of silver is lost in the case of the

crystalline multiphase alloy where the ternary alloy exhibits more noble Ecorr values.Similarly it is also noticed that the corrosion current densities for the crystalline alloys

Fig. 3. Potentiodynamic polarisation measurements recorded at a potential sweep rate of 0.5 mV/s in

0.3 M borate buffer solution and 0.1 m sodium hydroxide solution.

Table 1

Results of potentiodynamic polarisation (Fig. 3) and potentiostatic polarisation (Fig. 4) tests performed

on quaternary amorphous and crystalline Mg65Y10Cu15Ag10 alloy samples in weakly and strongly alkaline

electrolytes compared with those of ternary Mg65Y10Cu25 alloy Ref. [8]

Measured parameter Mg65Y10Cu25crystalline

multiphase

from [8]

Mg65Y10Cu25amorphous

from [8]

Mg65Y10Cu15-

Ag10crystalline

multiphase

Mg65Y10Cu15-

Ag10 amor-

phous

Ucorr in 0.3 M borate buffer (mV/SCE) )679 )720 )1050 )241icorr in 0.3 M borate buffer (lA/cm2) 7 5 35 1.5

Ucorr in 0.1 N NaOH (mV/SCE) )624 )575 )1130 )318icorr in 0.1 N NaOH (lA/cm2) 2.5 0.5 0.8 0.15

ipassive in 0.3 M borate buffer, potentio-

dynamic polarisation (lA/cm2)270 49 150 8

ipassive in 0.1 N NaOH, potentiodynamicpolarisation (lA/cm2)

64 2 40 0.65

ipassive in 0.3 M borate buffer, potentio-

static control (60 min) (lA/cm2) Fig. 455 14 100 140

ipassive in 0.1 N NaOH, potentiostaticcontrol (60 min) (lA/cm2) Fig. 4

149 2 6 0.6


are lower for the ternary alloy compared to the quaternary alloy in the weakly al-kaline electrolyte of pH 8.4 whereas the reverse is true for the strongly alkaline

electrolyte with pH 13. The higher icorr values for the quaternary alloy in the boratebuffer electrolyte are assumed to be the result of different activities between differ-

ently composed adjacent phases. The passive current densities during the potentio-

dynamic polarisation measurements at a very slow scan rate (Fig. 3) were seen to be

consistently lower for the quaternary alloy compared to the ternary alloy (Table 1).

Fig. 4 shows the results of current transient measurements carried out for the

amorphous Mg65Y10Cu15Ag10 alloy ribbons at selected anodic potentials in thepassive region, i.e., at 300 mV(SCE) for the NaOH solution and at 800 mV(SCE) for

the borate buffer electrolyte. These potentials are indicated by arrows (#) in Fig. 3. Inthe double logarithmic plot, the current density transient falls monotonously in the

strongly alkaline electrolyte whereas it shows a specific pattern in the weakly alkaline

electrolyte. The overall current density level in this electrolyte with pH ¼ 8:4 ishigher by about two orders of magnitude than that in the solution with pH ¼ 13.The anodic current density transient at pH ¼ 8:4 shows a slope close to )0.5 indi-cating a process of diffusion-controlled layer growth [8] in the initial part up to about2 s, then falls rapidly with a slope close to )1 indicating a high field controlled oxidelayer growth [8] till about 20 s, and then shows the trend of more or less constant

level in the range of about 150 lA/cm2. In the strongly alkaline electrolyte, thecurrent density transient falls with a slope of about )0.7 upto a passive currentdensity of 0.6 lA/cm2 after 1 hour (Table 1). The passive layer formation mechanisms

Fig. 4. Anodic potentiostatic polarisation tests (current transient measurements) performed at 800

mV(SCE) in 0.3 M borate buffer solution and at 300 mV(SCE) in 0.1 M sodium hydroxide solution for the

amorphous Mg65Y10Cu15Ag10 alloy.


seem to be different in the quaternary alloy when compared with those reported for

the ternary alloy [8] in both electrolytes. The presence of silver could have some

influence in the variation observed. However, as is seen from the values reported in

Table 1, the passive current density of the amorphous alloy after potentiostatic

polarization for 1 h in borate buffer at the anodic potential is higher than those of the

ternary and even crystalline quaternary alloys. This is in contrast to results of po-tentiodynamic measurements and could be attributed partly to the altering processes

of the passive layer over a larger time scale.

From the electrochemical investigations so far, it is possible to state that the

quaternary amorphous alloy displays an improved corrosion behaviour in both the

chosen alkaline electrolytes of pH 8.4 and 13 with more noble corrosion potentials

and lower corrosion current densities compared to that of ternary amorphous alloy.

The crystalline quaternary alloy, however, shows a different behaviour i.e. a less

noble corrosion potentials compared to that of the ternary crystalline alloy mostlydue to galvanic activities of adjacent phases. The corrosion current densities as well

as the passive current densities are higher in the weakly alkaline electrolyte compared

to the strongly alkaline electrolyte for both amorphous as well as crystalline qua-

ternary alloys. The passive layer growth mechanisms seem to be different for the

quaternary alloy compared to those for the ternary alloy.

The AFM has revealed different topographies of passive layers formed in different

electrolytes on the amorphous Mg65Y10Cu15Ag10 alloy surface. Typical AFM images

are shown in Fig. 5. The reference specimen (as-quenched ribbon) (Fig. 5a) shows amainly homogeneous flat surface with only a few inhomogeneities, certainly, the

inclusions in the ribbon i.e., the yttrium oxide particles. As compared to this, the

ribbon passivated in the borate buffer solution (Fig. 5b) shows a quite heterogeneous

rough surface topography. The growth was uneven with mounds and deeper valleys.

For the ribbons passivated in sodium hydroxide electrolyte (Fig. 5c), a less rough

oxide layer topography was observed.

The passive films grown for 2 and 20 h under potentiostatic polarisation in both

electrolytes of pH 8.4 and 13 were subjected to surface chemical analysis by AES andthe results are shown in Figs. 6 and 8. In Fig. 6, the net concentrations obtained from

Auger depth profiles for the as-quenched amorphous ribbon and for ribbon samples

polarised at anodic potentials in the passive region under potentiostatic control in

both electrolytes (800 mV(SCE) in case of the borate buffer electrolyte and 300

mV(SCE) in case of the NaOH electrolyte) for time durations of 2 and 20 h are

displayed. It can be seen that in the borate buffer solution thick oxides containing

remarkable amounts of boron were formed. They displayed rough surfaces when

observed using the secondary electron and absorbed current images in the Augermicroprobe. This roughness will cause a non-homogeneous sputter removal. Due to

this and also due to the fact that the sputter rate depends on the chemical compo-

sition, the concentrations of the elements are displayed with respect to the equivalent

sputter depth in silicon oxide. The decay of the oxygen concentration to 50% of its

maximum is used as the measure for the thickness. For the thin oxides formed in the

NaOH solution this thickness is estimated to be 16 and 29 nm (Fig. 6d and e), re-

spectively, for 2 and 20 h of exposure times. For the oxides formed in the borate


Fig. 5. AFM images of (a) as-quenched amorphous Mg65Y10Cu15Ag10 ribbon surface, (b) passive layer

grown in 0.3 M borate buffer solution at pH 8.4 and 800 mV(SCE) for 20 h (c) passive layer grown in

0.1 M sodium hydroxide solution at pH 13 and 300 mV(SCE).


buffer solution the oxide layer removal times correspond to thicknesses of 90 and 260

nm (Fig. 6b and c), respectively, which may not be accurate from the reasonsmentioned above. For the as-quenched specimen, the native oxide thickness was

estimated to be about 6 nm (Fig. 6a).

After sputter removal of the oxide layers the concentrations of the metals reach a

steady state. It is seen here that even for the as-quenched specimen, the nominal

composition of the Mg65Y10Cu15Ag10 ribbon material is not found to be in the steady

Fig. 6. AES concentration depth profiles of the elements with the depth scale as equivalent sputter depth

in silicon dioxide. The horizontal dashed lines indicate the nominal compositions of the components: (a)

as-quenched amorphous Mg65Y10Cu15Ag10 alloy ribbon surface, (b) passivated in borate buffer solution at

pH 8.4 and 800 mV(SCE) for 2 h, (c) passivated in borate buffer solution at pH 8.4 and 800 mV(SCE) for

20 h, (d) passivated in sodium hydroxide solution at pH 13 and 300 mV(SCE) for 2 h, (e) passivated in

sodium hydroxide solution at pH 13 and 300 mV(SCE) for 20 h.


state of the depth profiles. The measurement (Fig. 6a) shows a depletion in Mg and

enrichments in Cu and Y, whereas the Ag concentration is only slightly affected. This

is attributed to the preferential sputtering of the light Mg atoms. The uppermost

surface is depleted in the atoms with the higher sputter yields until a new steady state

of sputter removal is reached. Using the stopping and range of ions in materials

(SRIM) code [11], the sputter yields of the elements in the alloy were calculated to be8.7 (Mg), 5.9 (Cu), 4.1 (Y) and 5.4 (Ag), i.e. the observed depletion in Mg with

respect to Cu and Y is in agreement with the calculation. It is known that prefer-

ential sputtering depends on mass and surface binding energy of the sputtered atoms

and that effects like segregation also can alter the surface composition further.

However, the measurements show that the steady state reached during sputter depth

profiling differs considerably with the electrochemical treatment, also. It could be

deduced by observing the magnesium profile in all the five cases presented. While we

notice that the specimen exposed to the NaOH solution for 20 h (Fig. 6e) has athicker oxide layer than the one exposed for 2 h (Fig. 6d) as well as as-quenched

specimen (Fig. 6a), the steady state Mg concentration also is somewhat lower in this

specimen. This could be attributed to the depletion due to sputtering as the anodic

current transient (Fig. 4) does not suggest any active dissolution in the NaOH so-

lution. In case of samples passivated in the borate buffer solution the magnesium

profile shows further depletion. The current transient in this case (Fig. 4), however,

suggests some active dissolution while the passive layer is formed. We carried out the

solution analysis, using ICPMS after the experiment, to identify the elements pre-sent. It was confirmed that the solution contains significant amounts of magnesium

compared to the reference solution. In addition, the composition in the steady state

here (in the Auger analysis) for these specimens may be affected by the roughening.

The shape of the Auger peaks can vary with the chemical state. The changes in the

peak shape have been analysed by factor analysis [12–15] of the spectra obtained

during depth profiling. The contributions of the different peak shapes to the spectra

can be calculated and show how the chemical state changes within the specimens. It

was found that, the spectra of the alloy components are dominated by two peakshapes. While Mg(KLL) and Y(MNN) showed pronounced changes in shape and

position, the Cu(LVV) and Ag(MNN) peaks showed only small changes. Fig. 7

shows typical spectra of elements and their oxides observed by Auger spectrometer

and resolved by factor analysis in the current study. Because these changes in peak

shape and position can be due to variations in binding energy, shape of the valence

band and different relaxation effects, the peaks were broadly classified as metallic and

in the oxidized state with no specific identification assigned. Boron was only in the

oxidized state with two slightly different peak positions. For oxygen, three differentpeak shapes caused by different metal–oxide bonds were found. Fig. 8 shows the

result of the factor analysis and the corresponding display of the so called chemical

depth profiles. Due to the large number of components, Fig. 8 shows only the

chemically resolved depth profiles of the oxidized components and Ag, the effects of

which are especially focussed upon and presumed to be the reason for the differ-

ences in the ternary and quaternary alloys. It should be noted that Ag and oxidized

state of Ag profiles were shown with the appropriate multiplication factor for easy


identification on the graph. The three oxygen species identified by factor analysis

differ mainly in the peak position (Fig. 7). The peak designated O1 is located at 512

eV and its profile is connected to that of Mg oxide and O2 (at 510 eV) follows the Yoxide and boron profiles. The profile of O3 (at 514 eV) is similar to a second boron

(oxide) profile.

From the Auger depth profiles one can see that the outermost layer of the passive

oxide film for the specimen exposed to borate buffer solution (Figs. 6b and 8b) is

enriched in copper oxide and to a smaller extent in silver oxide. The subsequent

layers are rich in yttrium and boron and at the passive layer–alloy substrate interface

there is an enrichment of magnesium oxide. The profile of the sample passivated for

20 h duration (Figs. 6c and 8c) displays a similar pattern but appears broadened andmore diffuse. The corresponding current density transient in this electrolyte (Fig. 4)

exhibits an initial stage with a slope of about )0.5 indicating a diffusion-controlledgrowth, then a stage with a mainly high field-driven growth rate (slope close to )1).It is well known that the kinetics of anodic oxide film formation is governed by the

thermodynamics at the surface and the mass transfer across the oxide layer as it

forms. It could be presumed here that a thin film of magnesium hydroxide and silver

oxide form initially due to the presence of silver that could be slightly enriched at the

surface in the as-quenched specimen. Mass transfer across this native film perhapsdictates the growth rate of other stable oxides at the interface initially. As the native

Fig. 7. Characteristic spectra obtained by factor analysis of the transitions: (a) Mg(KL23L23), (b) O(KLL),

(c) Y(M45N23N45) and (d) Ag(M45N45N45). For the metals solid and broken lines are for the peak in the

alloy and the oxide, respectively. The oxygen peak shapes labelled O1, O2 and O3 are discussed in the text.


film is not stable enough and dissolves in the electrolyte, the stable copper oxidestarts forming and the growth kinetics of this oxide is high field-driven [8,16,17].

Yttrium oxide also starts forming. As the layer of these combined oxides grow, the

current density rapidly decreases. It is assumed that, due to solubilities of silver oxide

and magnesium hydroxide in this electrolyte there could be some pores formed that

allow the contact of electrolyte with the interface in select places. The AFM images

of these surfaces display deeper valleys suggestive of such a possibility though actual

confirmation is possible only through a cross-sectional view. We assume that the

current starts rising again and that at the interface there is a growth of magnesiumoxide due to the local concentration of oxygen and due to the availability of mag-

nesium from the bulk. As this oxide grows the current again shows a decreasing

Fig. 8. Chemically resolved depth profiles of the oxidized components of Fig. 6 in the same order along

with silver elemental profile (magnified): (a) as-quenched, (b) passivated in borate buffer for 2 h, (c)

passivated in borate buffer for 20 h, (d) passivated in sodium hydroxide for 2 h, (e) passivated in sodium

hydroxide for 20 h. Silver oxide also is shown magnified with the multiplication factor in the brackets.


trend. This is perhaps the reason for the ‘‘waved’’ shape of the curve at the times

larger than 1 h. The reason for the higher current density in this electrolyte after

passivation for 1 h as compared to that of the ternary alloy, shown in Table 1, may

be due to this dissolution process of magnesium oxide/hydroxide at the interface

inside the pores of the passivating layers. This, however, needs further experimental

confirmation. Other studies using TEM are underway to obtain evidence to thisassumption.

The passive layers anodically formed on the ribbons exposed to the strongly al-

kaline electrolyte, however, show a different composition (Fig. 6d and e and Fig. 8d

and e). The outer layer is enriched with magnesium oxide/hydroxide whereas the

interfacial layers show an enrichment of yttrium oxide (Fig. 8d and e). The layers are

thin compared to the layers anodically formed on ribbons in weakly alkaline elec-

trolyte. Very small enrichment of silver oxide is seen upon careful observation on the

outermost layer. The stability diagrams [9] indicate that magnesium hydroxide isquite stable at these pH values. Silver oxide has a soluble component of AgO� and

so, the small enrichment of this oxide noticed on the surface is perhaps due to what

remained in the film as most of the soluble part remains in the solution. Since there is

no barrier from the silver oxide as in the case of borate buffer electrolyte, the thin

adherent magnesium oxide and the subsequent stable yttrium oxide grow by a

mechanism which is in between the diffusion-controlled growth and the high field-

driven growth, as indicated by the current density transient, which has a slope of

about �0:7. These oxides quickly passivate the surface. In the ternary alloy, thegrowth of the passive layer in this electrolyte of pH 13 was found [8] to be mainly by

diffusion-controlled mechanism. Here we observe a mechanism which is in between

the diffusion-controlled and the high field-driven growth mechanisms. This again

perhaps could be attributed to the presence of silver, a valve metal.

From these investigations we conclude that the mechanisms of passivation of the

amorphous Mg65Y10Cu15Ag10 alloy are quite different in these alkaline electrolytes

and that the film growth kinetics also largely differs. The presence of silver does seem

to cause the observed changes in the corrosion potential and the corrosion currentdensities as compared to the ternary Mg65Y10Cu25 alloy, i.e. the corrosion potentials

shift in the noble direction and the corrosion current densities tend to be lower

values. However, due to the higher solubilities of the different oxides of silver formed

at different anodic potentials, the film characteristics, passivating nature and the

growth mechanisms get altered. Ag2O is soluble in acid media with the formation of

Agþ ions and in very alkaline media with the formation of AgO� ions. The reported

solubility of this oxide in pure water is 10�3:85 mol Agþ/l [9]. It is worthwhile to note

that the solubilities of the oxides (e.g., 10�7 mol Cuþþ/l for CuO and 10�3:95 molMgþþ/l for Mg(OH)2 in pure water) and the stabilities of those oxides in different

electrolytes vary considerably even for the simple parametric combination of po-

tentials (300 and 800 mV(SCE)) and pH values (8.4 and 13) chosen here. This leads

to the difference in the passivation mechanisms where copper plays a dominant role

in the weakly alkaline electrolyte and magnesium in the strongly alkaline electrolyte

for this quaternary amorphous Mg65Y10Cu15Ag10 alloy in addition to the above

discussed role of silver.


4. Conclusions

• The corrosion behaviour of the quaternary Mg65Y10Cu15Ag10 alloy containing sil-

ver is altered compared to ternary Mg65Y10Cu25 alloy without silver. Due to the

presence of silver in the matrix, the quaternary alloy in the amorphous state dis-

plays more noble and stable behaviour in both electrolytes of pH 8.4 and 13.• The multiphase crystalline quaternary alloy exhibits a poorer corrosion resistance

compared to that of the amorphous alloy. The difference in the corrosion behav-

iour between the amorphous and crystalline states is much more pronounced for

the quaternary alloy than for the ternary alloy.

• The passive layers formed on the amorphous Mg65Y10Cu15Ag10 alloy in the

weakly and strongly alkaline electrolytes are quite different in the composition

as well as in their morphology.

• In the weakly alkaline electrolyte, copper plays a dominant role in the passivationof the quaternary alloy along with yttrium. Silver is supposed to play an impor-

tant role in stabilising the passive films at lower potentials. The growth mecha-

nism has a distinctive pattern starting with an initially diffusion-controlled

growth, followed by a high field-driven growth mainly attributed to copper oxide

formation which is then followed by a steady state corrosion.

• Magnesium hydroxide along with yttrium hydroxide and silver oxides form the

thin homogenous passive layer on the amorphous quaternary alloy surface in

the highly alkaline electrolyte. The growth mechanism, which was mainly a diffu-sion-controlled mechanism for the ternary alloy [8], gets altered with a tendency

towards a high field-driven growth mechanism for the quaternary alloy, attributed

to the participation of silver in the surface layer formation.

• The beneficial effect of the amorphous single phase nature of a metallic material in

terms of its improved corrosion resistance is once again demonstrated for this

quaternary Mg65Y10Cu15Ag10 alloy system.

Acknowledgements

R.V. Subba Rao gratefully acknowledges the International Office of the BMBF at

the DLR Bonn, Germany, for financial support. We thank profoundly M. Johne and

M. Frey for technical assistance and E. Fangh€aahnel for ICPMS analysis.

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Corrosion behaviour of the amorphous Mg65Y10Cu15Ag10 alloy

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