Electrochemical response of amorphous and devitrified Al-Ni-La-X (X = Ag, Cu) alloys

Materials and Corrosion 2009, 60, No. 6 DOI: 10.1002/maco.200805118 431

Electrochemical response of amorphous anddevitrified Al-Ni-La-X (X¼Ag, Cu) alloys

A. Roy, A. K. Mandhyan, K. L. Sahoo,J. Banhart and I. Chattoraj*

The electrochemical response of melt-spun Al-Ni-La alloys withpartial substitution of Ni after different stages of devitrification wasstudied. The base alloy was found to have the best corrosionresistance. It was observed that primary crystallization causedminimal deterioration in the corrosion resistance of the base alloyas compared to its amorphous state. The substitute alloys haddifferent corrosion resistance dependent on the substituting element

� I. Chattoraj, A. Roy, A. K. Mandhyan, K. L. SahooNational Metallurgical Laboratory, Jamshedpur 831007 (India)E-mail: [email protected]

J. BanhartDepartment of Materials, Hahn-Meitner Institute, GlienickerStrasse 100, 14109 Berlin (Germany)

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with the Ag containing alloy having the least resistance. This couldbe attributed to the operation of local galvanic cells, enhanced bychemical heterogeneities in the alloys. Secondary crystallizationcaused a reduction in the corrosion resistance of all the alloys due tothe creation of intermetallic phases that increased the galvanicactivity.

1 Introduction

A number of Al-based amorphous alloys synthesized byrapid solidification [1–4] demonstrate a favorable combina-tion of high strength and acceptable ductility. Among thesealloys, Al-Ni-RE (RE¼La, Ce, Gd, etc.) have generatedconsiderable attention owing to their promising mechanicalproperties. Partial crystallization of these alloy systems hasbeen shown to significantly enhance their strength due tohardening of the amorphous matrix by the primary crystal-lization product [4]. During partial crystallization the size ofthe precipitates should be restricted to the range of a fewnanometers, beyond which the favorable properties start todeteriorate. Unfortunately, these alloys undergo a detri-mental secondary crystallization at higher temperatures withsignificant reduction of physical and chemical properties [4].It is therefore, desirable to have a separation between the firstand second crystallization temperature which is as wide aspossible. Partial substitution of the rare earth or transitionmetal component by other transition metals has beenattempted with the aim to enhance the working range(difference between the first and second crystallizationtemperatures), to cause changes in structure and mechanicalstrength and to positively influence primary crystallization ofthe amorphous alloys [5–7]. It is reported [5] that theaddition of Cu to amorphous Al85Y8Ni5Co2 alloys causes analmost complete deprivation of glass transition due to thedecrement of primary crystallization temperature andformation of quenched-in-nuclei in the as-melt-spun state.It is reported that Cu and Ag refine the particle size of Alphase and significantly increase the number of Al particlesduring primary crystallization [8]. Partial substitution in Al-

Ni-RE alloys by other elements could be a promising way toachieve the goals described above, but at present literature onthis is scarce.An assessment of the corrosion resistance of Al-Ni-La

alloys, especially those with Ni partially substituted by otherelements, is scarce in literature [9,10]. Commercialacceptance of substitute alloys will depend on their corrosionresistance in addition to their stability. The first of these twopapers [9] showed that Cu had a positive influence on thecorrosion resistance of the base alloy, with regards to thepassivation current in alkaline NaCl. This study wasconducted for the amorphous alloys only. The second study[10] reported the effect of devitrification on the base alloyonly. This study provides the effect of devitrification as wellas alloy substitution along with the effect of chlorideconcentration on the corrosion behavior. It is remarkablydifferent from the earlier studies that the presence of Cu andAg is found to be detrimental to corrosion resistance,although the solution pH in our case is near neutral. Thestudy also compares the effect of different substitutionelements (Ag, Cu) on the electrochemical response of thealloys before and after devitrification.

2 Experimental procedure

Ingots with the compositions Al87Ni6La7, Al87Ni5La7Cu1,and Al87Ni5La7Ag1 were prepared by alloying purecomponents (purities: 4N/3N) by induction melting undera purified argon atmosphere. The ingots were inductively re-melted (temperature ranges from 1280 to 1300 K) in analumina-coated quartz crucible, after which melt-spunribbons were prepared by ejecting the melt onto a rotatingcopper wheel (tangential speed 40 m/s) in a He atmosphere.The structure of the as-melt-spun as well as annealed

ribbons was investigated by X-ray diffractometry (XRD) andtransmission electron microscopy (TEM). The crystalliza-tion behavior of the ribbons was studied by differentialscanning calorimetry (DSC) under a high purity argonatmosphere. The DSCwas calibrated by using pure In and Zn

bH & Co. KGaA, Weinheim

432 Roy, Mandhyan, Sahoo, Banhart and Chattoraj Materials and Corrosion 2009, 60, No. 6

ig. 2. Continuous heating DSC curve of amorphous Al-Ni-La-XX¼Cu/Ag) alloys at a heating rate of 40 K/min

standards, giving an accuracy of �0.3 K for temperature and�0.02 mW for heat flow. From the XRD patterns thestructure of the crystalline phases was determined. The TEMsamples were prepared by jet thinning using a solution of25% nitric acid and 75% methanol at 243 K and 12 V.Small angle neutron scattering (SANS) investigations

were carried out at the BER-II reactor of the Hahn-Meitner-Institut Berlin, Germany, using instrument V4. The meas-urements were performed at two sample-detector distancesof 1 and 4 m at a neutron wavelength l¼ 0.605 nm. Detailsof the measurements are given in ref. [10].The electrochemical studies were conducted in Na2SO4

(0.1 N) and NaCl solutions with three different concentra-tions (0.001 N, 0.01 N, and 0.1 N). Potentiodynamic meas-urements were obtained starting from a potential of �1 Vwith respect to a saturated calomel electrode (SCE) and werecontinued either up to 1.6 VSCE or to a voltage at which theanodic current density attained a value of around 1 mA/cm2.A slow scan rate of 1 mV/s was used. The bright side of theribbons was exposed while the dull side was masked off. Thealloy studied was made the working electrode, a platinumwire was used as the counter electrode, and a saturatedcalomel electrode served as the reference electrode.Corrosion products were generated by holding the samplesfor 30 min at a potential ofþ50 mV relative to the respectiveopen circuit potential in the corroding medium, followed byrinsing in water and acetone and drying. The corrosionproducts formed were studied under a scanning electronmicroscope (SEM) with energy dispersive spectroscopic(EDS) facilities.

3 Results

3.1 Crystallization

The ribbons prepared by melt spinning were 30� 5 mmthick, �2 mm wide, and several meters long. XRD patternsrevealed that the ribbons obtained from the basealloy Al87Ni6La7, Cu- and Ag-containing alloys are fullyamorphous on the substrate side as well as on the air-cooledside. Figures 1(a)–(c) shows the XRD patterns of the ribbonson the air-cooled side. These patterns showed a distinct broadhalo centered on 37.58, which is characteristic of the glassyphase. The continuous heating DSC curve (Fig. 2) of thealloys revealed two stages of crystallization, correspondingto the formation of different phases. The first stage, i.e.,

Fig. 1. XRD pattern of as-melt-spun alloys, air-cooled side(a) Al87Ni6La7, (b) Al87Ni5La7Ag1, and (c) Al87Ni5La7Cu1

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primary crystallization, is due to the formation of ametastable phase that transforms to a stable phase at highertemperatures. After primary crystallization the structurecontains a uniform distribution of the metastable phase in anamorphous matrix. The second stage, i.e., secondarycrystallization, is due to the formation of stable intermetalliccompounds and fcc Al-phase from the amorphous matrix.Therewere definite differences in the peak temperature of thefirst crystallization (Tx1p) as well as second crystallization(Tx2p) stages. It was observed that Tx1p of the base alloy ishigher and Tx2p is lower than that of the Cu- or Ag-containingalloy. From Fig. 2 it can be calculated that DTx (Tx2p�Tx1p)is greatest for the Ag-containing alloy (Ag: DTx¼ 86 K, Cu:DTx¼ 81 K, base alloy: DTx¼ 60 K).The nature of the initial melts-spun ribbons and

the evolutions of crystallinity were investigated by TEM.Figure 3a shows a bright field TEM and selected areaelectron diffraction (SAED) pattern of the as-melt-spun basealloy. In accordance with the XRD results no crystallineareas are found in this sample. The as-melt-spun Ag- and Cu-containing alloys show similar results though the bright fieldimage of the Cu containing alloy (Fig. 3b) shows somecontrast. This appears due to mass-thickness contrast. TheSAED pattern in the deep contrast areas also showsamorphous nature and no crystallinity was noted in thesesamples through SAED (Fig. 3b).In order to identify the phases formed in the first

crystallization stage the samples were heated at a rate of20 K/min up to the end of their crystallization. Correspond-ing representative TEM images are shown in Fig. 4. Theprimary crystalline product of the base alloy is a metastablebcc phase of lattice parameter 0.663 nm [10]. Addition of Agor Cu does not change the nature of the primary crystal-lization product. In all cases crystallization of metastable bccphases was observed. Major changes were observedregarding the number and growth of the metastableprecipitates. Figures 4(a)–(c) show that after completionof the first crystallization stage the bcc metastable phases aremore or less uniformly distributed in the remainingamorphous matrix. The number of precipitates in the Agcontaining samples was much higher than that in the basealloy but their size was much smaller. Similarly, the size ofthe precipitates in Cu-containing samples is smaller than inthe base alloy. Figure 5 shows the results of small angleneutron scattering experiments and displays the sizedistribution of concentration fluctuations, or, more generally,

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Materials and Corrosion 2009, 60, No. 6 Electrochemical response of amorphous and devitrified Al-Ni-La-X alloys 433

ig. 3. TEM bright field image (inset SAED pattern) of as-melt-pun ribbons (a) base alloy and (b) Cu containing alloy
Fs
heterogeneities, in the as-melt-spun ribbons obtained by theanalysis of the scattering cross-section curve versusmomentum transfer measured ex situ after heat treatment.According to Fig. 5, the number of heterogeneities in the Ag-containing sample is much higher than in the base alloy buttheir size distribution is narrower.

ig. 4. TEM images of samples after primary crystallizationa) base alloy, (b) Ag-containing alloy, and (c) Cu-containing alloy

3.2 Electrochemical behavior

3.2.1 Polarization studies

The electrochemical responses of the alloys weremeasured after different degrees of crystallization, includingthe amorphous alloy, the alloy after primary crystallizationand the alloy after secondary crystallization. For this, melt-spun samples were heated applying the same heating profileas used for DSC and were then quenched. Figures 6–8provide these responses in different media for the threealloys.All the alloys in their amorphous state as well as after

crystallization demonstrated passivation in the Na2SO4

solution, see Figs. 6(a), 7(a), and 8(a), expressed by alimitation of the current density to<10�4 A/cm2 even for thehighest applied positive voltages at 1.6 V.In the 0.001 N NaCl medium, Fig. 6(b), all the amorphous

alloys demonstrated passivity, but the passive range wassignificantly reduced compared to the Na2SO4 solution,expressed by the much lower positive voltage (0.2 V), atwhich the current density starts to rise to high values. The


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base alloy showed the maximum and the most stable passiverange (of the order of �400 mV), while the Ag-containingalloy exhibited the smallest passive range (of the order of�150 mV). The passive region of the base alloy appearedstable, while all the substitute alloys demonstrate electro-chemical noise, possibly indicating instability of the passivefilm. After primary crystallization, all the alloys showed adeterioration of corrosion resistance in the same medium


ig. 5. Fitted size distributions (volume weighted) of the as-melt-pun ribbons obtained by SANS experiments
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Fig. 6. Electrochemical response of the different alloys in the as-cast state measured through potentiodynamic polarization in themedia indicated. Alloy designations: 1, base; 2, Cu-containing; 3,Ag-containing

(0.001 N NaCl), see Fig. 7(b). There is an evident lack ofpassivation for the Ag-containing alloy and transientpassivity in the Cu-containing alloy. Only the base alloysshowed a significant passive range. This ability of the basealloy to passivate was considerably reduced after secondarycrystallization. All the other substitute alloys showed activedissolution when tested in 0.001 N NaCl, see Fig. 8(b).On increasing the NaCl concentration to 0.01 N and then

to 0.1 N, a progressive decrease in the corrosion resistance ofall the alloys was observed. In the as-cast state, the base alloyexhibits passivity in all the media tested, although the passiverange decreased significantly with increase in the NaClconcentration. This tendency of the base alloy to passivatewas seen even after the primary crystallization. In fact, forthe base alloy there was little difference in the electro-chemical response between the as-cast and first-crystal-lization samples. The base alloy was the only one todemonstrate transient, weak passivity after secondarycrystallization in the 0.001 N, see Fig. 8(b), and 0.01 N,see Fig. 8(c), NaCl solutions, although it showed activedissolution in the higher concentration, 0.1 N NaCl, seeFig. 8(d). The Cu-containing alloy showed a small passiveregion in 0.01 N NaCl and none in 0.1 N NaCl. Of thesubstitute alloys the performance of the Ag-containing alloywas the worst of all the alloys when compared to the as-caststate. Primary crystallization caused a general deteriorationin corrosion resistance of the substitute alloys. Theelectrochemical response of the substitute alloys after firstcrystallization followed a trend similar to that for their as-cast state. The Cu-containing alloy showed transientpassivity only in 0.001 N NaCl, whereas the Ag-containingalloy actively dissolved in all the solutions. None of the othersubstitute alloys showed passivation after secondary crystal-lization in any medium and actively dissolved.

3.2.2 Corrosion product analysis

All alloys after primary crystallization were exposedto anodic polarization to obtain hints on the character ofanodic activities. Potentiostatic treatment was carried out atþ50 mVwith respect to the respective open circuit potentialsfor 30 min in 0.001 N NaCl. This galvanic treatment andsolution combination were chosen as it provided definite

differences in the electrochemical response of the threealloys. After potentiostatic holding the samples wereimmediately observed in a SEM equipped with EDS.Changes in current density during potentiostatic holding

was continuously monitored. The anodic activity in all thealloys increased with the duration of potentiostatic holdingwith the exception of the base alloy, indicating that the latterdeveloped passivation while the substitute alloys did not.This is in accordance with our earlier observation that afterprimary crystallization only the base alloy showed sig-nificant passivity over a range of potentials. The passivationtendency of the base alloy is evident in the lack of corrosionon the surface of this alloy; see Fig. 9(a). The rippledstructure seen here is the normal morphology of rapidlysolidified ribbons. These shapes are absent in the substitutealloys, indicating anodic activity that smoothened out thesurface irregularities by anodic dissolution. The Cu-contain-ing alloy shows regions of corrosion (dark areas in the SEMmicrograph) interspersed in relatively unaffected areas,



Fig. 7. Electrochemical response of the different alloys afterprimary crystallization in the media indicated. Alloy designationsame as in Fig. 6

Fig. 8. Electrochemical response of the different alloys aftersecondary crystallization in the media indicated. Alloy designationsame as in Fig. 6

Fig. 9(b). The extent of corrosion was the most for the Ag-containing alloy, Fig. 9(c). To further elucidate the extent ofcorrosion, elemental mappings of the exposed samples werecarried out. Significant differences between the exposedsamples were found with regard to the oxygen distributionwhich was found to increase with the extent of corrosion, i.e.,it was the highest in the Ag-containing alloys and least in thebase alloy, see Fig. 10. The corrosion products developedwere therefore rich in oxygen, presumably representingoxides of the constituent elements.

4 Discussion

Partial substitution of Ni by Ag or Cu leads to adeterioration of corrosion resistance of the alloys. The levelof corrosion resistance was found to strongly depend on thesubstituting element. The formation of a passive film on thebase alloy was observed even after primary crystallization,indicating that crystallization per se does not cause the


deterioration of corrosion resistance. The substitute alloysform similar phases during crystallization. Therefore, thedifference in the corrosion response between the four alloyscannot be attributed to the primary crystals. In fact, structuralheterogeneities seemed less important for the as-cast andprimary crystallized alloys in determining the corrosionresponse as compared to chemical heterogeneities. In spite ofthe structural homogeneity of the Ag-containing alloy thehigher corrosion of this alloy suggests a detrimental role ofAg for the corrosion resistance of this alloy. In this regard thecluster distribution indicated by SANS (Fig. 5) is instructive.Amongst the amorphous alloys, the Ag-containing alloy hadthe highest density of clusters and the average size of theseclusters was the smallest. The chemical fluctuations and thepresence of Ag result in a low corrosion resistance. Localgalvanic activity is responsible for this enhanced corrosion inthe as-cast alloy containing Ag. On the other hand, the basealloy has the dual benefits of lower number, larger chemicalclusters, and structural homogeneity. A lower number andlarger size of clusters implies a lower interfacial area


Fig. 9. Surface morphology after exposure to 0.001 N NaCl for thealloys after primary crystallization (a) base; (b) Cu-containingalloy; (c) Ag-containing alloy

Fig. 10. Oxygen distribution on the exposed surfaces indicated inFig. 9 measured by EDS: (a) base; (b) Cu-containing alloy; (c) Ag-containing alloy

between chemically different clusters. Ag- and Cu-contain-ing alloys have local fluctuations even in the as-castconditions. This could create interfaces and enhancecorrosive activity. However, the local galvanic cell actioninvolving chemical clusters (as opposed to those involvingdifferent phases) shows a more dominant effect. Ag or Ag-containing clusters would have a higher galvanic potentialdifferential by virtue of the element’s position in the galvanic

series. The concentration fluctuations present in the Cu- andAg-containing alloy in the as-cast condition and in all thealloys after primary crystallization are Al-rich bcc phases,the compositions of which are not very different from thecomposition of the remaining amorphous matrix [11].Therefore, galvanic action between these crystals and theamorphous matrix is minimal. This also explains why afterprimary crystallization there is hardly any change in



corrosion resistance for the base alloy. For the substitutealloys too, primary crystallization causes only a slightdeterioration in corrosion resistance in any medium,compared to that for the as-cast alloy, due to the presenceof additional structural and chemical heterogeneities asdiscussed earlier. Secondary crystallization results in a groupof intermetallics such as Al11La3, Al3Ni, and Al4La. Nointermetallic or other phase containing the substitute elementcould be found in any of the three alloys. Either such phasesdo not exist or are formed in very small amounts. Theequilibrium solid solubility of Cu and Ag in Al is 2.48 and23.8 (at%), respectively [12] which would account for theirnon-participation in any intermetallic phase. Intermetallicphases largely increase galvanic activity as they are presentin significant amounts and are electrochemically differentfrom the primary bcc phase. Thus, after secondary crystal-lization there is a tremendous decrease in the corrosionresistance of all the alloys. Even so, the base alloy is superiorto the substitutes and the Ag-containing alloy continued tohave the lowest corrosion resistance after secondarycrystallization. The higher amount of Ni in the base alloyis responsible for this. Ni is known to improve passivationand prevent pitting in various conventional crystalline alloys.The two substitute alloys had similar Ni contents, so thedifference in their response is due to the substitute element.The process of primary crystallization leads to a concentra-tion of these substitutes in the amorphous matrix (as theprimary phase is devoid of these elements) therebyincreasing galvanic activity. After secondary crystallizationthese elements would either form intermetallics, remain insolution in the primary phase, or form elemental precipitates.Unfortunately, due to the very low concentration of thesubstituting elements, none of these structures could beindividually found. Irrespective of the nature of occurrence,these would all enhance galvanic activity and add to thedetrimental effect of the intermetallics detected aftersecondary crystallization.

5 Conclusions

In the amorphous state, the base alloy Al87Ni6La7 wasfound to have the best, and the Ag-containing substitute alloythe least corrosion resistance. SANS studies showing aprofuse number of small clusters in the Ag-containing alloysindicate that local galvanic action was responsible for thisenhanced corrosion.


Structural heterogeneities in the form of nanocrystals orconcentration fluctuations present in the as-cast Cu alloywere less important in determining the corrosion resistancethan the galvanic activity of local clusters.Primary crystallization resulting in a bcc-Al phase in all

the alloys caused varying degrees of deterioration in thecorrosion resistance of the alloys as compared to theamorphous alloys. For the base alloy Al87Ni6La7 thisdeterioration was negligible.The onset of secondary crystallization that resulted in the

generation of intermetallic phases caused a pronouncedreduction in corrosion resistance for all the alloys. This is dueto increased galvanic activity as well as the loss of theamorphous phase.

Acknowledgements: The help provided by Ms. SeemaKumari in heat treating the samples is gratefully acknowl-edged.

6 References

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[11] K. L. Sahoo, M. Wollgarten, J. Haug, J. Banhart, Acta Mater.2005, 53, 3861.

[12] R. H. Brown, L. A. Willey, in: J. E. Hatch (Ed.), Aluminium,Properties and Physical Metallurgy, ASM Publication, USA,

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(Received: July 21, 2008) W5118(Accepted: July 28, 2008)

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Electrochemical response of amorphous and devitrified Al-Ni-La-X (X = Ag, Cu) alloys