Available online at www.sciencedirect.com

www.elsevier.com/locate/actamat

Acta Materialia 60 (2012) 851–859

Effects of microstructure on native oxide scale developmentand electrical characteristics of eutectic Cu–Cu6La alloys

B.S. Senturk a, Y. Liu a, J.V. Mantese b, S.P. Alpay a, M. Aindow a,⇑

a Institute of Materials Science and Department of Chemical Materials and Biomolecular Engineering, University of Connecticut,

97 North Eagleville Road, Storrs, CT 06269-3136, USAb United Technologies Research Center, Silver Lane, East Hartford, CT 06108, USA

Received 25 May 2011; received in revised form 4 November 2011; accepted 6 November 2011

Abstract

A combination of electron microscopy, focused ion beam and conductive atomic force microscopy techniques have been used to studythe microstructure, oxide scale development, and electrical behavior of a Cu–9 at.% La alloy. The as-cast alloy exhibits a eutectic micro-structure comprising 30 vol.% Cu rods in a Cu6La matrix. The eutectic colonies exhibit a singular orientation relationship with [010]Cu6La parallel to h011i Cu along the rod axis, and it is shown that this corresponds to lattice matching of the two phases along thisdirection (�0.02% misfit). Oxidation of the alloy at 100 �C to accelerate formation of a native oxide scale led to the development ofa Cu2O layer less than 25 nm thick on the Cu rods and a La-doped Cu2O scale up to 1 lm thick on the Cu6La matrix. The La-dopedregions of the scale are more conductive despite being much thicker, which is consistent with previous contact resistance data obtainedfor this alloy. The mechanisms responsible for the formation of this non-equilibrium oxide scale structure and for the enhanced electricalconductivity of the La-doped regions are discussed.� 2011 Published by Elsevier Ltd. on behalf of Acta Materialia Inc.

Keywords: Copper alloys; Electrical resistivity/conductivity; Electron microscopy; Microstructure; Oxidation

1. Introduction

Metallic contacts are a ubiquitous method of connectingelectrical and electronic components/systems. Over the lastfew decades the number of metallic contacts used in indus-trial applications such as transportation, electronic andtelecommunication systems have increased dramaticallysuch that approximately one trillion metallic connectorsare now produced annually. These contacts are usually fab-ricated from base metals because they are inexpensive, havehigh bulk electrical conductivities and exhibit excellentformability [1,2]. Unfortunately, such base metals oxidizein air under ambient conditions, and the characteristicsof the native oxide scales lead to contact resistances ordersof magnitude higher than those for mating bare metal

1359-6454/$36.00 � 2011 Published by Elsevier Ltd. on behalf of Acta Mater

doi:10.1016/j.actamat.2011.11.013

⇑ Corresponding author. Tel.: +1 860 486 2644; fax: +1 860 486 4745.E-mail address: m.aindow@uconn.edu (M. Aindow).

surfaces [3,4]. This is a critical technological issue, sincethe development of unacceptably high contact resistancesover time is now by far the most common cause of failurein electrical/electronic devices and systems. One commonapproach to this problem is the application of a noblemetal coating to a base metal contact to prevent oxide scalegrowth; this is viable in situations where the improved reli-ability of the contact outweighs the penalty of added cost[5]. While precious metal coatings can dramaticallyincrease reliability and signal/power transfer efficiency,such plated systems can still fail through abrasive wearand fretting, which results in the formation of high contactresistance via oxidation of bare base metal beneath thecoating [6,7].

In our work we have approached this problem by devel-oping strategies for alloying base metals to promote the for-mation of inherently conductive native oxide scales. The useof such alloys for electrical contacts would obviate the need

ialia Inc.

852 B.S. Senturk et al. / Acta Materialia 60 (2012) 851–859

for noble metal plating. Moreover, since the native scaleswould be self-healing, contacts produced from these alloyswould be far more reliable in situations where abrasive wearor fretting might occur. In a recent paper [8] several distinctapproaches to enhancing the conductivity of oxide scaleswere identified, including: doping to increase mobile carrierconcentration; inducing mixed oxidation states resulting inelectron/polaron hopping; and phase separation givingconducting pathways. Preliminary data obtained frombinary alloys were used to demonstrate the feasibility of theseapproaches and contact resistances significantly lower thanthose for the respective oxidized base metals were measured [8].

In this paper we consider one of these binary systems,Cu–9 at.% La (Cu–9La), in greater detail. Copper and itsalloys are widely used for electrical contact components,but copper oxidizes even under ambient conditions, forminga non-protective Cu2O (cuprous oxide) scale, whichincreases electrical contact resistance and dramaticallydegrades contact reliability [9–11]. Results obtained in ourpreliminary study on oxidized Cu–9La [8] revealed that thecontact resistance was approximately three times lower thanthat for oxidized pure Cu, with only a modest decrease in thebulk electrical conductivity (25.1 vs. 83.1 nX m for pure Cu[12]). Here we present a combined electron and scannedprobe microscopy study of the microstructures in the alloyand in the native oxide scale. These observations are usedto relate the morphology of the scale to both microscopicand macroscopic measurements of the electrical characteris-tics, and the mechanisms responsible for the formation of amore conductive native scale on the alloy are discussed.

2. Materials and methods

The Cu–La system was selected because of the potentialto form La cuprates such as La2CuO4, which are known toshow metallic conductivity under ambient conditions andto exhibit superconducting properties at low temperatures[13–15]. Thus alloys in this system are good candidates forthe formation of conducting oxides via phase separation ofthe scale whereby the high conductivity second oxide phaseforms percolating pathways through the low conductivitybase metal oxide. In the standard binary Cu–La equilibriumphase diagram of Cirafici and Palenzona [16] the Cu–9Lacomposition corresponds to a binary eutectic between Cuand Cu6La (Fig. 1a); this composition was chosen for thepresent study because the microstructure should be moreuniform than those of hypo- or hyper-eutectic alloys. Wenote that later studies on the Cu–La system by Bloch et al.[17] and Meyer-Liautaud et al. [18] identified a new Cu13La“X” phase with the structure NaZn13. It was proposed thatthis Cu13La phase forms peritectically at 873 �C via a reac-tion between Cu and (Cu,La) liquid and then decomposesto Cu + Cu6La on cooling to 820 �C [19]. In subsequent hightemperature diffraction studies by Bolmgren and Lundstrom[20], however, it was shown that the X phase peaks could beaccounted for on the basis of a mixture of the high tempera-ture orthorhombic b and low temperature monoclinic a

polymorphs of Cu6La. Since the evidence for the Cu13LaX phase is equivocal, the microstructural data obtained inthis study have been interpreted on the basis of the simplifiedCu–La phase diagram [16] shown in Fig. 1a. From the equi-librium oxide phase diagram (Fig. 1b) one would anticipateobserving the growth of a copper oxide scale on the Cu phaseand a mixed copper oxide/La2CuO4 scale on the Cu6Laphase [21].

An experimental alloy ingot 25 mm in diameter and50–100 mm in length was produced by arc melting high pur-ity (>99.99%) metal precursors under an inert atmosphere.The overall composition of the as-cast alloy and the compo-sitional homogeneity were verified using energy dispersiveX-ray spectrometry (EDXS) in a scanning electron micro-scope. For oxidation experiments circular slices of �5 mmthickness were cut from the ingot. These slices were roughground using SiC papers, polished to a 0.3 lm finish usingalumina paste, and then rinsed in ethanol. The samples wereoxidized in air at 100 �C using a muffle furnace with a tem-perature control of ±3 �C.

X-ray diffraction (XRD) data were acquired with a Bru-ker AXS D5005 X-ray diffractometer using a CuKa sourcewith glancing incidence (tube angle of 0.5–1�) to enhancethe signals from the thin oxide scale. The XRD spectra wereobtained by scanning over the angular range 2h = 25–80� ata scan speed of 5� min�1 for a total of 24 h. The peaks wereidentified by comparison with the appropriate JCPDS files.For the as-cast alloy metallographic sections were preparedby mechanical grinding, polishing, and then etching with anaqueous ferric chloride solution (10 g FeCl3, 20 ml HCl,80 ml H2O). Oxidized samples were observed directly with-out further preparation. Secondary electron (SE) imageswere obtained from these samples in a JEOL JSM-6335Ffield emission scanning electron microscope operating atan accelerating voltage of 15 keV.

Cross-sections were obtained from both as-cast and oxi-dized samples by focused ion beam (FIB) sectioning in aFEI Strata 400S Dual-Beam instrument, equipped with highresolution electron and Ga+ ion beam columns, and anEDAX Phoenix STD-S EDXS system. Prior to sectioninga thin strap of platinum was deposited over the area of inter-est to minimize ion beam damage. A trench was then milledinto the surface and SE images were obtained from the wallof the trench using the electron column in the dual beaminstrument. The accelerating voltage used was 5–10 keV,and the SE signals were detected using an Everhardt Thorn-ley detector.

Thin foil samples for transmission electron microscopy(TEM) were prepared by twin jet electropolishing to perfo-ration using a solution consisting of 67% methanol and33% nitric acid at �30 �C with an applied voltage of10 V. The samples were then Ar+ ion milled briefly toremove any residual contamination using a Gatan Duomill600 equipped with a liquid nitrogen cooled stage and oper-ating at an accelerating voltage of 2 kV. To avoid specimenpreparation artifacts TEM samples were not producedfrom the oxidized alloy discs, but instead thin foil alloy

Fig. 1. (a) Partial Cu–La phase diagram with the alloy composition indicated. (b) Partial La2O3–CuO phase diagram with the composition of the Cu6Laphase indicated to show the oxides CuO and La2CuO4 that should form in the scale on this phase at equilibrium.

B.S. Senturk et al. / Acta Materialia 60 (2012) 851–859 853

samples were oxidized under the same conditions as thoseused for the oxidation experiments and these oxidized thinfoils were then examined directly by TEM. The TEMexperiments were performed in an FEI Tecnai T12 scan-ning/transmission electron microscope operating at120 keV and a JEOL JEM-2010 FasTEM operating at200 keV.

Conductive atomic force microscopy (C-AFM) experi-ments were performed using an Asylum Research MFP-3dinstrument operated in contact mode at room temperaturein an ambient atmosphere. The Standard Asylum ORCAmodule (http://www.asylumresearch.com/Applications/Orca/Orca.shtml), which includes a specially designed canti-lever holder with a transimpedance amplifier, was employedto obtain topographic and current images of the surfacesimultaneously. In such C-AFM experiments the surface isscanned in contact mode with a d.c. voltage in the range�10 to 10 V applied between the sample and the groundedconductive tip. The resultant current is amplified, digitizedusing an auxiliary 100 kHz analog to digital converter(ADC), and then filtered digitally at 1 kHz [22]. For the datapresented here, a standard ORCA cantilever holder with asensitivity of 2 nA V�1 and a gain of 5 � 108 V A�1 (�1pA to 20 nA) was used with nitrogen-doped diamond-coatedconductive tips and a d.c. voltage of 1 V applied to theoxidized alloy surface.

3. Experimental results

3.1. Microstructure of the as-cast alloy

The overall microstructure of the as-cast Cu–9La alloy isrevealed in SE scanning electron microscopy (SEM) images,such as those shown in Fig. 2a, obtained from the etchedmetallographic sections. The lower magnification mainimage shows that the alloy structure consists of uniform,randomly oriented eutectic colonies several tens of micronsin diameter. The morphology of the phases within an indi-vidual eutectic colony can be seen more clearly in the insethigher magnification image. There is a rod-like phase withinthe eutectic structure with an average rod spacing of 0.75 lm

and an average rod diameter of 0.3 lm. EDXS point analy-sis measurements from the minority rod-like phaseconfirmed that they corresponded to almost pure Cu. Theferric chloride solution used to etch the metallographicsample preferentially dissolves the Cu6La phase and thusreveals the distribution of Cu phase. Further metallographicinvestigations were performed on the as-cast alloy by pre-paring a cross-sectional sample using a dual beam FIBinstrument to observe the morphology and distribution ofthe Cu6La phase. Fig. 2b is a typical SE SEM imageobtained from a FIB cut section through the alloy showingthe distribution of the Cu rods within the Cu6La matrix. Noevidence for a pro-eutectic phase was observed in eitheretched metallographic sections or FIB cut cross-sections ofthe alloy, indicating that the true alloy composition is extre-mely close to the nominal eutectic value. The area fractionsof the two phases were determined using standard imageanalysis techniques on several such images and mean valuesof 29%Cu, 71%Cu6La were obtained. These values corre-spond closely to the volume fractions one would expectfor a eutectic mixture of pure Cu and stoichiometric Cu6Labased upon published values of the lattice parameters(31.25 vol.% Cu, 68.75 vol.% Cu6La).

The details of the alloy microstructure were revealedmore clearly in data obtained from thin TEM foils.Fig. 3a is a typical bright field (BF) TEM image of theas-cast alloy. As expected from the SEM data, a two phasemicrostructure was observed. The EDXS spectra (notshown here) and selected area diffraction patterns (SADP)obtained from such regions showed that the darker rod-likephase is fcc Cu and the lighter matrix is the monoclinicintermetallic phase a-Cu6La, as expected. An example ofa SADP recorded from the interface between the Cu andCu6La phases obtained with the beam direction parallelto [001] in Cu6La is shown in Fig. 3b. The [011] zone axisin the Cu phase is almost parallel to the beam direction,indicating that there is a well-defined orientation relation-ship (OR) between the two phases in the Cu–Cu6La eutec-tic. An analysis of SADP from several different coloniesrevealed that the phases adopt the OR: [001] Cu6La //[011] Cu; ½0�10� Cu6La // ½0�11� Cu.

Fig. 2. Secondary electron SEM images of the as-cast alloy. (a) Deep etched metallographic sample with a high magnification detail inset showing thedistribution of the etch-resistant Cu phase. (b) FIB cut cross-sectional sample showing the protective Pt strap and the distribution of the Cu and Cu6Laphases.

Fig. 3. TEM data obtained from an electropolished Cu–9La thin foil. (a) BF image showing the distribution of the phases. (b) SADP obtained with thebeam direction parallel to [001] for Cu6La showing the OR between the phases.

854 B.S. Senturk et al. / Acta Materialia 60 (2012) 851–859

In each case the Cu rods were aligned with the latterdirections, i.e. the rod axes were h011i type and the rodslay parallel to the b-axis of the Cu6La matrix phase.

3.2. Characteristics of the oxide scale

The morphology of the oxide scale forming on the alloyafter exposure to air at 100 �C was investigated using a com-bination of SEM, FIB cross-sectional imaging and AFM. Aselection of the SE SEM images obtained from the oxidizedsamples is presented in Fig. 4. A typical plan view imageobtained from a sample oxidized for 200 h is shown inFig. 4a. Such images reveal that the base alloy microstruc-ture has a profound influence on the morphology of theoxide scale. There is a thicker oxide scale with a more com-plex morphology on one phase and a much thinner oxide onthe other phase. Higher magnification images, such asFig. 4b, show that the protruding thicker regions of thescale are comprised of spherical particles with diametersof 20–50 nm. In contrast, the thinner regions are very uni-form, so that scratches on the alloy surface from metallo-graphic specimen preparation are clearly visible. Fig. 4c isa typical SE SEM image obtained from a FIB cut section

through the same sample. In such images there is a porousoxide scale between the Pt layer and the Cu6La matrix phasein the underlying alloy, but the scale on the Cu rods is thin-ner and denser, which can be seen more clearly in Fig. 4d,which is a higher magnification detail of the region indi-cated by the box in Fig. 4c. Moreover, there is a band ofbright contrast in the matrix phase that extends to a depthof approximately 1 lm. In separate elemental mappingexperiments it was revealed that this band of contrast corre-sponds to a region of enhanced oxygen content in thematrix phase only, indicating that the Cu6La phase oxidizesinwards very rapidly, while oxidation of the Cu phase isnegligible. The topography revealed in the plan view SESEM images was quantified by obtaining contact modeAFM images such as Fig. 5a. This is a gray scale image ofthe oxide scale topography on a eutectic colony orientedwith the Cu rods almost perpendicular to the surface.Fig. 5b is a topographic trace across the horizontal line indi-cated in Fig. 5a. Such traces reveal that the oxide on theCu6La phase protrudes above the surface by up to 100 nm.

Further experiments were conducted to reveal the struc-ture of the oxide scale. Fig. 6a is an XRD spectrum takenfrom a bulk alloy sample oxidized for 400 h. All of the

Fig. 4. Secondary electron SEM images showing the morphology of the oxide scale after oxidation for 200 h. (a) Plan view image of the oxide scalesurface. (b) High magnification detail of the region in (a) showing the morphological differences in the scale on the two alloy phases. (c) Cross-sectionalimage of the wall for a FIB cut trench through the scale showing penetration of the oxidation front into the Cu6La matrix. (d) High magnification detail ofthe boxed region in (c) showing the very thin scale on the Cu phase.

Fig. 5. Contact mode AFM data from the surface of the oxide scale afteroxidation for 200 h. (a) Topographic image. (b) Height profile across thehorizontal line indicated in (a).

B.S. Senturk et al. / Acta Materialia 60 (2012) 851–859 855

peaks in such spectra corresponded to those expected forthe Cu, Cu6La, and Cu2O phases. Since XRD is relativelyinsensitive to the presence of minority phases, particularlyin thin oxide scales, further analysis was performed usingTEM techniques. Fig. 6b is a BF TEM image obtained

from a region at the edge of a hole in an electropolishedalloy foil that had been oxidized for 10 h. There is a bound-ary running across the field of view and EDXS datarevealed that there is no La in the region on the right ofthe boundary but the Cu:La ratio in the region on the leftis 6:1. Thus the boundary corresponds to an interfacebetween the matrix phase and a Cu rod in the alloy priorto oxidation. Diffraction data obtained from such regionsrevealed that the extent of oxidation in the Cu6La matrixphase was far greater than that in the Cu phase, as expectedfrom the SEM and AFM data (Figs. 4 and 5). Fig. 6c is atypical SADP obtained from the oxidized matrix phase; allof the rings in such patterns could be indexed unambigu-ously as arising from the Cu2O structure. It is interestingto note that, although the formation of metastable Cu2Oscales on Cu is well known, the formation of a single phaseCu2O scale on the Cu6La was not expected [23]. Since the6:1 Cu:La ratio is maintained in these regions the La couldeither be incorporated substitutionally into the Cu2O lat-tice or be present as an amorphous oxide film along theCu2O grain boundaries. To distinguish between these pos-sibilities, high resolution TEM (HR-TEM) phase contrastlattice images were obtained from the thinnest regions ofthe oxidized TEM foils. One example of a HR-TEM imagefrom such a region is shown in Fig. 6d. The lattice andmoire fringes across the field of view in such images reveal

Fig. 6. Structural data obtained from the oxide scales. (a) XRD data from a sample oxidized for 400 h. (b–d) TEM data from a sample oxidized for 10 h.(b) BF image showing the differences between the oxidized phases. (c) SADP of the oxidized Cu6La phase showing rings corresponding to untexturedpolycrystalline Cu2O. (d) HR-TEM lattice image showing the fully crystalline scale on the Cu6La phase comprising Cu2O grains �10 nm in diameter.

856 B.S. Senturk et al. / Acta Materialia 60 (2012) 851–859

that the scale on the Cu6La phase comprises Cu2O grains>10 nm in diameter, with no evidence for amorphous grainboundary films. Thus the La must be incorporated into theCu2O lattice as a substitutional cation.

C-AFM experiments were performed to investigate therelationship between the electrical characteristics of theoxidized surfaces and the oxide scale structure/morphol-ogy. Examples of the data obtained are presented inFig. 7. Fig. 7a and c are topographic images obtained fromtwo different areas on a sample oxidized for 200 h, andFig. 7b and d are the corresponding current images. Thereis a clear correlation between the bright regions in Fig. 7aand 7, indicating that there is a higher current flowingthrough the tip when it is in contact with the oxide onthe Cu6La matrix phase. Thus the oxide scale on the Cu6Laphase has a higher conductance than that on the Cu phase,despite being much thicker. To demonstrate that these fea-tures are not an artifact of the sample topography considerFig. 7c and d. The very bright feature arrowed in Fig. 7ccorresponds to a darker region in Fig. 7d, i.e. the contrastfor this feature is different from that of the scale on theCu6La phase. This contrast is what one would expect fora low conductivity dust particle �0.2 lm in diameter lyingon top of the oxide scale.

4. Discussion

Binary transition metal/rare earth systems haveattracted a great deal of interest, largely because of their

potential as permanent magnet and/or hydrogen storagematerials. In the Cu–La system, however, the primaryinterest has been in their oxide derivatives, which formthe basis for many of the high Tc superconducting oxides[24–26]. As such, there has been relatively little publishedwork on the microstructures exhibited by binary Cu–Laalloys. In this work we firstly investigated the microstruc-ture of the eutectic Cu–9La alloy. It was shown that themicrostructure was comprised of randomly oriented eutec-tic colonies with diameters of several tens of microns. Eacheutectic colony consisted of parallel fcc Cu rods within amonoclinic a-Cu6La matrix; these phases and their volumefractions were what one would expect on the basis of thesimplified phase diagram presented in Fig. 1a [16]. Therewas no evidence for the presence of the so-called X Cu13Laphase in this microstructure. Moreover, if the alloy were tosolidify by the formation of a Cu13La + Cu6La eutecticand the Cu13La phase then underwent eutectoid decompo-sition to a mixture of Cu + Cu6La, as implied by the cur-rently accepted Cu–La phase diagram [19], then onewould expect to see some microstructural evidence for thistransformation. No such evidence was observed.

The morphologies adopted by eutectic microstructuresare usually determined by some combination of the volumefractions of the phases present, the available interfacialconfigurations and, where appropriate, the elastic straininduced in the phases by coherency at the interfaces. Inthe current study a singular OR was observed within theCu–Cu6La eutectic colonies, even though the colonies were

Fig. 7. C-AFM data from a sample oxidized for 200 h. (a and c) Topographic images obtained in contact mode. (b and d) Current images obtained at asurface bias of 1 V from the regions in (a) and (c), respectively.

B.S. Senturk et al. / Acta Materialia 60 (2012) 851–859 857

oriented randomly with respect to the external axes of thesample. Thus this OR probably arises because it is an ener-getically favorable configuration for the phases. We recallthat the Cu rods within the colonies have their major axesparallel to h011i, and that this lies parallel to [010] in theCu6La matrix. Since solubility of La in Cu is negligible andCu6La is a line compound we can estimate the lattice mis-match from the published lattice parameters of the purephases. The lattice parameter a0 for Cu is 0.3607 nm andso the {01 1} planes in Cu have a spacing of 0.25505 nm.The b0 parameter for Cu6La is 1.0204 nm, giving a spacingfor the (040) planes of 0.25510 nm. Since the (040) planesof Cu6La lie parallel to {011} of Cu along the rod axis, themisfit in this direction will be remarkably low (�0.02%). Assuch, the magnitude of the elastic strain required for theformation of a coherent interface along the rod axis willbe extremely small, which is presumably why the eutecticadopts this combination of OR and phase morphology.We note that while the same degree of lattice matchingcould not be found in any other direction for the OR mea-sured, there is some evidence for the development of facetsin the rod cross-sections (see Fig. 2), suggesting that theremay also be some anisotropy in the energy of interfaceslying in the h011i Cu/[01 0] Cu6La zone.

The character of the oxide formed on the Cu rods isbroadly consistent with what one might expect based uponthe literature. Cu is known to oxidize readily in air even atroom temperature and forms a non-protective Cu2O layer

on the surface at oxidation temperatures below 200 �C.Many studies have shown that at low temperatures oxida-tion of Cu follows an inverse logarithmic rate law wherebythe oxidation reaction proceeds by tunneling of electronsand inward migration of cation vacancies under the influ-ence of a strong electric field formed between the copper/oxide and oxide/oxygen interfaces. Thus the oxidation rateis mostly controlled by the number of mobile vacancies inthe oxide scale [23,27].

The oxide formed on the Cu6La phase was, however,very different from that anticipated. One might haveexpected the formation of a phase separated scale compris-ing a mixture of copper oxide and copper lanthanum oxidein the manner suggested by the equilibrium La2O3–CuOphase diagram (Fig. 2b) or even a pure La2O3 scale viaselective oxidation of La. Instead, a scale with Cu2O asthe only crystalline phase was formed and the thicknessof this scale was up to 1 lm (Fig. 4c), i.e. nearly two ordersof magnitude thicker than that on Cu. Since the Cu:Laratio in these regions of the scale was approximately equalto that in the Cu6La phase, and no significant amorphousmaterial was observed at the boundaries of the Cu2O grains(Fig. 5d), the La must be incorporated into the Cu2O crys-tal structure. It is implausible that La could be accommo-dated substitutionally in the twofold coordination sitesadopted by Cu in the Cu2O crystal structure and thus weinfer that the La must lie in the octahedral interstices.The distortion that one would expect for such interstitial

858 B.S. Senturk et al. / Acta Materialia 60 (2012) 851–859

ions is fairly large: La3+ in sixfold coordination with oxy-gen has a radius of 114 pm, whereas the correspondingvalue for Cu+ would be 96 pm [28]. Indeed, it seems likelythat it is this ionic radius mismatch that inhibits the out-ward diffusion of La at such low temperatures and preventsselective oxidation. It is not clear how the incorporation ofLa into Cu2O leads to such a dramatic difference in thegrowth rate of the scale, but possible factors include anincrease in the density and/or mobility of the cation vacan-cies. Systematic studies of oxidation kinetics would berequired to resolve this issue.

In our preliminary report on this system [8] it was shownthat the contact resistance of Cu–9La is several times lowerthan that of pure Cu oxidized under the same conditions,even though the bulk conductivity, which affects the con-tact resistance, is over three times higher (83.1 vs.25.1 nX m for pure Cu). This is particularly remarkablesince the scale formed on the Cu6La phase is much thickerthan that formed on the Cu phase (see Fig. 4c). Indeed, theC-AFM data presented in Fig. 7 confirm that the scale onthe Cu6La phase is much more conductive than that on theCu rods, despite being much thicker. It is difficult to extractmeaningful conductivity measurements from the contactresistance or C-AFM data for this system due to a combi-nation of uncertainty in the contact area due to the probeand oxide scale topographies and uncertainty in the currentpath length from the point of contact on the La-rich scaleto the interdigitated Cu rods that will carry the currentfrom the scale to the underlying alloy. What is clear is thatthe conductivity must be orders of magnitude higher thanthat of pure Cu2O.

Cuprous oxide is a natural p-type semi-conductor with aroom temperature band gap of approximately 2.2 eV. Thep-type behavior is usually ascribed to an excess of cationvacancies resulting in delocalized hole states [29,30]. Therehave been various attempts to enhance the conductivity ofCu2O by doping, predominantly with non-metallic accep-tors such as N, Si and Cl [31–33]. The introduction of alio-valent metallic impurities should lead to n-type behavior,since these dopants will act as donors, but there is a paucityof experimental evidence for this and density functionaltheory (DFT) approaches have been used to show thatthese dopants may be compensated by binding to Cuvacancies [34]. Subsequent DFT studies by Nolan andElliot [35] on a range of aliovalent doping additions toCu2O showed that the formation of compensating Cuvacancies was thermodynamically favorable for most ofthe potential n-type dopants, including La3+. We note thatthe situation in the native oxide scales on Cu–9La consid-ered here is very different from that envisaged in these pre-vious studies. Firstly, these are thin scales that form on abulk alloy, which serves as a large reservoir of Cu and willthereby inhibit the formation of excess Cu vacancies. Sec-ondly, we note that the system is far from equilibriumand thus the validity of predictions based on thermody-

namic criteria is questionable. Thirdly, the concentrationof La is so high in these scales that the formation of thetwo Cu vacancies required to compensate each La3+ dop-ant would be sufficient to transform the oxide structurefrom Cu2O to a mixture of Cu2O + CuO. Since no evidencefor CuO was found in any of these samples, it is our con-tention that the most likely explanation for the enhancedconductivity in the La-rich regions of the Cu2O scale is thatthese are heavily doped n-type semi-conducting regions.

Finally, we note that while the contact resistance forCu–9La is not sufficiently low for this to constitute a viablealternative to coated base metal alloys for contact applica-tions, the data do suggest that very significant improve-ments in oxide scale conductivity could be achieved byincorporating suitable dopants. Thus if an alloy could beidentified that contained an appropriate oxide dopant butwhich did not accelerate the oxidation kinetics in the man-ner observed for La in Cu, then this could form the basisfor a new class of contact materials.

5. Conclusions

A eutectic Cu–9La alloy has been studied in the as-castcondition and after oxidation in air at 100 �C. Analysis bySEM, FIB, TEM and C-AFM techniques has shown thefollowing.

1. The alloy structure consists of eutectic colonies tens ofmicrons in diameter wherein a rod-like Cu phase lieswithin Cu6La matrix phase. There is no evidence for apro-eutectic phase or a Cu13La phase in this alloymicrostructure.

2. The Cu rods adopt a well-defined singular orientationrelationship with respect to the Cu6La matrix giving a0.02% lattice mismatch parallel to the rod axes.

3. The alloy microstructure has a profound influence onthe morphology of the oxide scale. There is a thickeroxide scale with a more complex morphology on theCu6La phase and a much thinner oxide on the Cu phase.

4. The thin scale formed on the Cu phase was Cu2O asexpected, while the thicker scale formed on the Cu6Laphase was a polycrystalline La-rich Cu2O layer.

5. C-AFM data revealed that the conductance of the La-rich regions in the scale is much higher than that ofthe pure Cu2O on the Cu rods, despite the former scalebeing two orders of magnitude thicker. It is inferred thatthis enhanced conductivity arises from heavy n-typedoping of the Cu2O lattice by La3+.

Acknowledgement

The authors gratefully acknowledge support by the USArmy Research Office through Grant No. W-911-NF0710388.

B.S. Senturk et al. / Acta Materialia 60 (2012) 851–859 859

References

[1] Braunovic M, Konchits VV, Myshkin NK. Electrical contacts:fundamentals, applications and technology. New York: CRC Press;2007.

[2] Holm R. Electric contacts: theory and applications. Berlin: Springer-Verlag; 1967.

[3] Stott FH, Glascott J, Wood GC. J Phys D Appl Phys 1985;18:541.[4] Tamai T. Electron Commun Jpn Part 2 1989;72:1349.[5] Patton ST, Zabinski JS. Tribol Lett 2005;18:215.[6] Hannel S, Fouvry S, Kapsa Ph, Vincent L. Wear 2001;249:761.[7] Sankara Narayanan TSN, Park YW, Lee KY. Tribol Int 2008;41:87.[8] Aindow M, Alpay SP, Liu Y, Mantese JV, Senturk BS. Appl Phys

Lett 2010;97:152103.[9] Iijima J, Lim JW, Hong SH, Suzuki S, Mimura K, Isshiki M. Appl

Surf Sci 2006;253:2825.[10] Tamai T. IEICE Trans Electron 2007;E90-C:1391.[11] Tsuda N, Nasu K, Fujimori A, Siratori K. Electronic conduction in

oxides. 2nd ed. New York: Springer; 2000.[12] Dyos GT, Farrell T. Electrical resistivity handbook: IEE materials

and devices series. 1st ed. London: Institution of Electrical Engineers;1992.

[13] Logvenov G, Gozar A, Bozovic I. Science 2009;326:699.[14] Attfield JP, Kharlanov AL, McAllister JA. Nature 1998;394:157.[15] Mohottala HE, Wells BO, Budnick JI, Hines WA, Niedermayer C,

Udby L, et al. Nat Mater 2006;5:377.[16] Cirafici S, Palenzona A. J Less-Common Met 1977;53:199.

[17] Bloch JM, Shaltiel D, Davidov D. J Less-Common Met 1981;79:323.[18] Meyer-Liautaud F, Allibert C, Moreau JM. J Less-Common Met

1985;110:81.[19] Okamoto H. J Phase Equilib 1991;12:504.[20] Bolmgren H, Lundstrom T. J Less-Common Met 1990;163:79.[21] Oka K, Unoki H. Jpn J Appl Phys 1987;26:L1590.[22] Lee JC, Chuang JH. Microelectron Reliab 2003;43:1687.[23] Honkanen M, Vippola M, Lepisto T. J Mater Sci 2007;42:4684.[24] Du Z, Xu Y, Zhang W. J Alloy Compd 1999;289:88.[25] Zhan Y, Xie H, Jiang J, Xu Y, Wang Y, Zhuang Y. J Alloy Compd

2008;461:570.[26] Logvenov G, Gozar A, Butko VY, Bollinger AT, Bozovic N, Radovic

Z, et al. J Phys Chem Solids 2010;71:1098.[27] O’Reilly M, Jiang X, Beechinor JT, Lynch S, Ni Dheasuna C,

Patterson JC, et al. Appl Surf Sci 1995;91:152.[28] Shannon RD, Prewitt CT. Acta Crystallogr 1969;B25:925.[29] Nolan M, Elliott SD. Phys Chem Chem Phys 2006;8:5350.[30] Raebiger H, Lany S, Zunger A. Phys Rev B 2007;76:045209.[31] Ishizuka S, Kato S, Maruyama T, Akimoto K. Jpn J Appl Phys

2001;40:2765.[32] Ishizuka S, Kato S, Okamoto Y, Akimoto K. Appl Phys Lett

2002;80:950.[33] Biccari F, Malerba C, Mittiga A. Solar Energy Mater Solar C

2010;94:1947.[34] Wright AF, Nelson JS. J Appl Phys 2002;92:5849.[35] Nolan M, Elliott SD. Chem Mater 2008;20:5522.