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Materials Chemistry and Physics 132 (2012) 109– 118

Contents lists available at SciVerse ScienceDirect

Materials Chemistry and Physics

j ourna l ho me pag e: www.elsev ier .com/ locate /matchemphys

ynthesis of CuCr and CuCrAg alloy with nano-ceramic dispersion by mechanicallloying and consolidation by laser assisted sintering

. Bera, I. Manna ∗

etallurgical and Materials Engineering Department, Indian Institute of Technology Kharagpur, WB 721302, India

r t i c l e i n f o

rticle history:eceived 2 February 2010eceived in revised form 31 July 2011ccepted 11 November 2011

a b s t r a c t

Cu–4.5Cr and Cu–4.5Cr–3Ag (in wt%) alloys without or with 10 wt% nanocrystalline Al2O3 and ZrO2 disper-sion have been synthesized by mechanical alloying or milling and consolidated by laser assisted sinteringin Ar atmosphere. Microstructural characterization by scanning and transmission electron microscopyand phase analysis by X-ray diffraction suggest that the alloyed matrix undergoes significant grain growth

eywords:. Composite materials. Powder metallurgy. Electron microscopy. Wear. Electrical conductivity

after sintering while the dispersoids retain their ultrafine size and uniform distribution in the matrix. Thedispersion of nano-Al2O3 is more effective than that of nano-ZrO2 in enhancing the mechanical propertiesdue to the smaller initial particle size of Al2O3 than that of ZrO2. In general, laser sintering of mechani-cally alloyed Cu–4.5Cr and Cu–4.5Cr–3Ag alloys with 10 wt% nanocrystalline Al2O3 at 100 W laser powerand 1–2 mm s−1 scan speed yields the optimum combination of high density (7.1–7.5 mg m−3), hardness(165–225 VHN), wear resistance and electrical conductivity (13–20% IACS).

. Introduction

Copper and its alloys are universally used as electrical con-uctors and devices due to their excellent electrical conductivity.owever, mechanical properties (hardness, wear or erosion resis-

ance) of Cu-alloys are often found inadequate. Solid solution andrecipitation strengthening through addition of solute elementshowing decrease in solid solubility with temperature (e.g. Cr,g, Nb and Zr) without adversely affecting electrical conductivityre well established strategies to enhance mechanical propertiesf Cu [1–8]. Furthermore, strength and hardness of the matrixhase may increase with decrease in grain size following theall–Petch relationship [9]. A number of investigations with Cu–Crlloys have studied the crystal structure of the metastable coherentrecipitates [10,11] and their influence on tribological perfor-ance [11–13], oxidation behavior [14] and mechanical properties

12–15] of these alloys. Besides alloying, dispersion of ceramicxide in Cu-matrix is also known to improve mechanical proper-ies of Cu alloys even at elevated temperature generated during use16–20].

Nanometric materials, with ultrafine grain or domain size, areharacterized by a very high specific surface or interfacial area of

he crystallites, and correspondingly, a host of novel properties noteasible in their coarse-grained counterparts [21,22]. Among thearious possible routes applied to produce nanostructured alloys

∗ Corresponding author. Fax: +91 3222 255303.E-mail address: imanna@metal.iitkgp.ernet.in (I. Manna).

254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.matchemphys.2011.11.005

© 2011 Elsevier B.V. All rights reserved.

and composites, mechanical alloying is one of most convenient,flexible and inexpensive methods of producing nanostructuredalloys from elemental powder blends with significant extensionof solid solubility [23–25]. However, the scope of developing nano-ceramic or intermetallic phase dispersed extended solid solution ofCu–Cr and Cu–Cr–Ag by mechanical alloying has not been explored.

Laser assisted sintering is an advanced method of sinteringpowders of different composition, nature and type and size intocomponents of precise shape, size and geometry [26–30]. In par-ticular, laser sintering has been widely exploited for producingfunctional components, prototypes and tools based on polymericand ceramic powders [31–34]. A recent improvisation of laser sin-tering, called selective laser sintering, offers the unique advantageof fabricating solid components of desired geometry and dimen-sion by computer-aided layer-by-layer consolidation of powders[33,34]. However, sintering of green compacts by laser irradiationis a more widely practiced method of producing solid componentsfrom metallic, ceramic or polymeric powders [32,35–37].

The advantages of laser sintering lie in selectivity of location,time and energy economy, flexibility, versatility, automation-worthiness, and faster heating and cooling rate including transientheating to elevated temperature [32,33]. The main parameters tocontrol the sintering process are laser power density and scanspeed. Though laser sintering of polymers, ceramics and coarsemetallic powders has been extensively studied, the scope of laser

assisted consolidation of mechanically alloyed nanocrystallinepowders has not been explored.

The present investigation aims to (a) synthesize Cu–4.5Cr andCu–4.5Cr–3Ag (in wt%) alloy powders by mechanical alloying,

1 istry and Physics 132 (2012) 109– 118

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10 S. Bera, I. Manna / Materials Chem

b) simultaneously disperse 10 wt% nano-Al2O3 or ZrO2 particlesomogeneously in the alloyed matrix by mechanical milling, andhen (c) consolidate the product into bulk solid component by laserintering for microstructural characterization and assessment ofhysical, mechanical and electrical properties of interest.

. Experimental

Pure Cu, Cr and Ag powders (each with > 99.5 wt% purity and about 50–100 �marticle size) were mixed in the weight proportions of Cu–4.5Cr and Cu–4.5Cr–3Agnd milled in a Retsch planetary mill for 25 h of cumulative time at 300 rpm atoom temperature. The milling operation was performed using hardened steel vialsnd balls. About 50 ml toluene was added to the vials prior to milling to avoid coldelding of powders to the surface of milling media (vials and balls) and changing

he milling dynamics. 10 wt% nanometric Al2O3 (with 10–40 nm particle size) orrO2 (with 20–50 nm particle size) powders were blended with 25 h mechanicallylloyed powders by an extra hour of milling to ensure homogeneous dispersion. Thelloyed powder was cold compacted at ambient temperature using 100 MPa pres-ure within a cylindrical graphite die of 10 mm diameter followed by laser irradiationssisted consolidation and sintering using a 1.5 kW continuous wave (Gaussian) CO2

aser at two different incident power levels (P = 100 and 150 W) and two differentcanning speeds (v = 1 and 2 mm s−1) at each incident laser power level. The corre-ponding power density level was 14.1 and 21.2 W mm−2 for a beam diameter of.5 mm for P = 100 and 150 W, respectively. For sintering a large surface area of theomponent (compared to the small beam diameter), about 10–15% overlap betweenuccessive passes was adopted. To avoid oxidation, the laser beam was shrouded byehumidified argon gas flowing at a rate of 6 l min−1.

The laser sintered samples were polished using 800, 1200 and 2400 grit SiCaper. The identity and sequence of phases were monitored by X-ray diffractionXRD) using a Panalytical X’Pert Pro diffractometer with Cu K� (0.1542 nm) radi-tion. Microstructure of the compacted and sintered samples was studied usingcanning electron microscope (SEM) coupled with energy dispersion spectroscopyEDS). Selected samples in as-milled power and high pressure sintered conditionsere studied by a JEOL transmission electron microscope (TEM) in the bright andark field and high resolution lattice imaging mode.

Average hardness of the composite samples was measured with the help of aickers microhardness tester at 0.5 or 1 N load from about 10 indentations at equiv-lent locations on the same sample. Density of the compacted and sintered samplesas measured using Archimedes principle after weighing in air and water sepa-

ately using an electronic balance with a precision of 0.1 mg. Wear resistance of theellets was studied by a fretting wear tester applying up to 20 N load within a wearpan of 1 mm at 10 Hz frequency for a cumulative duration of up to 1 h. The volumef the worn or fretted material was calculated using the empirical equation usedy Qu and Truhan [38]. Electrical resistivity of the sintered pellets was measured atmbient temperature using the van der Pauw four-probe method with more than5% accuracy [39].

. Results and discussion

Mechanical alloying for 25 h produces an alloy with extendedolid solubility of Cr (in Cu–4.5Cr and Cu–4.5Cr–3Ag alloy) andg (in Cu–4.5Cr–3Ag alloy), respectively. Subsequent mechani-al milling of 1 h with nanometric Al2O3 or ZrO2 (10 wt%) yieldsomogeneous dispersion of these ultrafine ceramic particles in thelloyed matrix. It may be noted that dissolution of Cr and Ag haseen partial. The details of the synthesis procedure and mechani-al alloying or milling parameters were reported in an earlier study40]. Fig. 1a and b shows the XRD patterns of laser sintered Cu–4.5Crlloy with different laser power (P) and scan speed (v) withoutnd with 10 wt% nanometric Al2O3 dispersion, respectively. TheRD patterns of the cold compacted pellets with identical compo-ition prior to laser irradiation are included for a ready referencend comparison. The XRD patterns of the laser sintered compos-tes show no separate peak of elemental Cr or Ag except that of theu-rich matrix phase. Therefore, it can be concluded that the phaseggregate after laser sintering is identical as that after mechani-al alloying or milling. Though laser irradiation generates adequateeat for sintering, limited time of exposure to high temperature dueo high scan speed and efficient self quenching of the narrow irra-

iation zone by the rest of the sample instantaneously to ambientemperature prevent precipitation of considerable amount of Cr org rich phase from the supersaturated matrix. However, the XRDeaks of the laser sintered alloy are sharper than those of the same

Fig. 1. XRD patterns of cold compacted and laser sintered pellets of (a) Cu–4.5Cr and(b) 10 wt% Al2O3 dispersed Cu–4.5Cr extended solid solutions (SS). Laser sinteringwas conducted at different levels of power (P) and scan speed (v).

mechanically alloyed powder compacted at ambient temperature.This suggests that sufficient heat or thermal activation is generatedduring laser sintering to cause grain coarsening of the matrix. Aftersintering some peaks corresponding to Cu2O are observed in theXRD patterns suggesting that the argon shroud during laser sin-tering could not totally prevent oxidation. However, intensity ofthese peaks and hence the volume percent of the oxide phase isinsignificant.

Fig. 2a and b shows the XRD patterns of laser sintered (withdifferent combination of P and v) Cu–4.5Cr–3Ag alloy with 10 wt%nanometric Al2O3 and ZrO2 dispersion, respectively. For the con-venience of comparison, XRD patterns of the same cold compactedalloy of identical composition have been appended. The XRD pat-terns show identical trend as that of the 10 wt% Al2O3 dispersedCu–4.5Cr alloy composite (Fig. 1b). Absence of significant peaks ofelemental Cr or Ag suggests that the sintered product remains apredominantly single phase extended solid solution albeit a notice-able coarsening of the crystallite size of the matrix due to the heatgeneration by laser sintering. Formation of small amount of Cu2Ocould not be prevented despite the continuous argon flow duringlaser sintering.

The precise lattice parameter (a0) of the face centered cubic(FCC) matrix after laser sintering (100 W and 1 mm s−1) of Cu–4.5Crand Cu–4.5Cr–3Ag + 10Al2O3 compacts has been determined bysuitable extrapolation of the variation of lattice parameter of

Cu-base solid solution (aCu) as a function of the Nelson–Riley(N–R) parameter (cos2 �/sin � + cos2 �/�) (Fig. 3a and b, respec-tively). The amount of Cr dissolved in Cu–4.5Cr is calculated fromthe increment in precise lattice parameter [40] and is ∼1.5 wt%.

S. Bera, I. Manna / Materials Chemistry and Physics 132 (2012) 109– 118 111

Fas

EbtTitCtCcp

Fig. 3. Variation of lattice parameter as a function of the Nelson–Riley (N–R) param-

TE

ig. 2. XRD patterns of cold compacted and laser sintered pellets of (a) 10 wt% Al2O3

nd (b) 10 wt% ZrO2 dispersed Cu–4.5Cr–3Ag extended solid solutions (SS). Laserintering was conducted at different levels of power (P) and scan speed (v).

vidently, the precise lattice parameter of Cu–4.5Cr–3Ag powderlend (0.3621 nm) is higher than that of the precise lattice parame-er of mechanically alloyed Cu–4.5Cr powder blend (0.3619 nm).his increase is attributed to the dissolution of both Cr and Agnto Cu. However, it is interesting to note that these precise lat-ice parameter values (0.3619 nm and 0.3621 nm for Cu–4.5Cr andu–4.5Cr–3Ag + 10Al2O3 alloy, respectively) are slightly smaller

han the lattice parameter values (0.3622 nm and 0.3624 nm foru–4.5Cr and Cu–4.5Cr–3Ag alloy, respectively) of the mechani-ally alloyed powder blend [40]. This marginal decrease in latticearameter in the present alloys could arise due to small degree of

able 1DS (spot or large area) analysis of laser sintered alloys of different initial compositions.

Alloy material Cu (wt%) Cr (wt%

Cu–4.5Cr(100 W–1 mm s−1)

Bulk 91.6 4.4

Spot 1 63.2 29.0

Spot 2 58.6 34.7

Spot 3 96.0 1.7

Spot 4 95.5 1.8

Cu–4.5Cr(100 W–2 mm s−1)

Bulk 92.1 4.8

Spot 1 22.5 70.5

Spot 2 62.6 31.7

Spot 3 94.4 2.1

Spot 4 95.1 2.3

Cu–4.5Cr–10Al2O3 (100 W–1 mm s−1) Bulk 83.7 4.6

Cu–4.5Cr–3Ag–10Al2O3 (100 W–1 mm s−1) Bulk 78.8 4.3

Cu–4.5Cr–3Ag–10ZrO2 (100 W–1 mm s−1) Bulk 82.2 4.1

eter (cos2 �/sin � + cos2 �/�) of laser sintered (100 W, 1 mm s−1) (a) Cu–4.5Cr and (b)10 wt% Al2O3 dispersed Cu–4.5Cr–3Ag alloy.

precipitation of Cr or Cr-rich phase from the supersaturated matrixduring laser sintering. In fact, Fig. 5a and b along with Table 1indeed confirm that Cr or Cr-rich precipitates are present in themicrostructure after laser sintering.

A thermodynamic analysis has been carried out to explain thesolubility of Cr in Cu. Primarily Gibbs energy change (�G) dueto mixing of Cr in Cu at 300 K was calculated with the equation�G = �H − T�S, where �H is enthalpy of mixing, �S is the entropychange and T is temperature. �S is calculated assuming ideal solidsolution formation. �H can be calculated by as per Miedema’sapproach [41] and details of the calculation is described elsewhere[42]. Fig. 4 shows the predicted values of �G as a function of Crcontent in Cu. Essentially positive free energy of mixing restricts

the solubility of Cr in Cu. At this point, it is well established thatmechanical alloying stores enthalpy in the materials by producinglarge dislocation density and decreasing the crystallite size in

) Ag (wt%) Al (wt%) Zr (wt%) O2 (wt%) Fe (wt%)

– – – 3.6 0.4– – – 7.8 –– – – 6.7 –– – – 2.3 –– – – 2.7 –

– – – 3.1 –– – – 7.0 –– – – 5.7 –– – – 3.5 –– – – 2.6 –

– 5.2 – 6.9 0.62.9 5.9 – 7.7 0.42.8 – 4.8 6.1 –

112 S. Bera, I. Manna / Materials Chemistry

FC

nheas(ofse�viwCe�o�e1ot�f(

wv

the oxide component (alumina) added to the powder charge, pres-

ig. 4. Gibbs free energy changes (�G) as a function of Cr content (up to 5 wt%) inu–Cr system.

anometer range [24]. In this context, Sheibani et al. [25,43]ave shown that during mechanical alloying both Gibbs strainnergy (�Gse) due to the presence of dislocation in the materialnd the Gibbs free energy (�Ggb) due to decrease in crystalliteize increase significantly. If the total increase in free energy�Gtotal = �Gse + �Ggb) is sufficiently large so that �Gtotal ∼ �Gf similar composition then it can provide the driving force forormation of an extended solid solution [25,43]. Adopting theimilar approach �Gse (�Gse = � × � × Vm, where � is dislocationlastic energy, � is dislocation density and Vm is molar volume,

= ((Gb2)/4�) × ln((dc)/b), where G is shear modulas, b is burger’sector, dc is crystallite size) and �Ggb (=� × (A/V) × Vm, where �s grain boundary energy and A/V is the surface by volume ratio)

ere calculated with the help of the data from the XRD pattern ofu–4.5wt%Cr alloy (after 25 h mechanical alloying) reported in ourarlier study [40]. The value of �Gtotal (�Gse = 0.55 kJ mol−1 andGgb = 0.56 kJ mol−1) is 1.11 kJ mol−1 which is closer to the value

f �G of Cu–3wt%Cr alloy (1.13 kJ mol−1) but much lower thanG of Cu–4.5wt%Cr alloy (1.71 kJ mol−1) (Fig. 4). This certainly

xplains the fact that about 3 wt% Cr dissolves in Cu and about.5 wt% Cr remains undissolved in Cu–4.5wt%Cr system after 25 hf mechanical alloying [40]. During laser sintering strain energy ofhe material decreases and grain coarsening occurs that decreases

Gtotal and consequently causes some amount of Cr precipitationrom the alloy and decreases the solubility of Cr in Cu to ∼2 wt%Table 1).

Fig. 5 shows the microstructure of laser sintered Cu–4.5Cr alloyith P = 100 W at two different levels of (a) v = 1 mm s−1 and (b)

= 2 mm s−1, respectively. A few isolated pores of sub-micron size

Fig. 5. SEM image of laser sintered Cu–4.5Cr alloy with different laser parame

and Physics 132 (2012) 109– 118

are observed at the surface of the consolidated alloy. Otherwise,the microstructure is homogeneous and dense. The points marked1, 2, 3 and 4 (both in Fig. 5a and b) represent the locations of theEDS micro-analysis to determine the local composition. The resultsof such elemental micro-analysis are summarized in Table 1. Thebulk EDS analysis (with a large area of analysis) shows that thecomposition is similar to that of the initial powder blend prior tomilling. Both in Cu–4.5Cr and Cu–4.5Cr–3Ag alloys the amount ofCr found in spots 1 and 2 is significantly higher (∼25 wt%) thanthe quantity of Cr content (4.5 wt%) in the initial powder blend.This implies that certain regions in the microstructure consist ofCr-rich precipitates which are in sub-micron size range, unlike ourearlier study concerning high pressure sintering of Cu–10Cr–3Agalloy [44] that led to precipitation of coarse Cr particles dur-ing sintering. Although the entire amount of added elemental Crdoes not dissolve in Cu, both large area and spot EDS analysis ofCu–4.5Cr–3Ag alloy suggests that the entire amount of elementalAg in the initial powder blend dissolves in the Cu-rich matrix. Itmay be pointed out that the marginal amount of Fe found in EDSanalysis may arise as impurity picked up from the milling medium(hardened steel) during mechanical alloying. Furthermore, oxy-gen may dissolve into the sintered alloy either from the processcontrol agent (toluene) during wet milling or Ar shroud duringlaser sintering. Presence of oxygen in the sintered product is sub-stantiated by the weak Cu2O peaks observed in the XRD patterns(Figs. 1 and 2). However, the total impurity content (Fe, Cu2O)is not very large. Thus, results of XRD and EDS analysis suggestthat constitution (phase aggregate) and composition of the lasersintered composites compare well with that of the initial powderblend.

Fig. 6a and b shows the SEM images of 10 wt% Al2O3 dispersedCu–4.5Cr alloy pressed in ambient condition at (a) low (2000×) and(b) high (50,000×) magnifications, respectively. The green com-pact contains porosities among loosely bound particles (Fig. 6a).In addition, the nanometric Al2O3 particles are dispersed homo-geneously throughout the alloy matrix (Fig. 6b). In contrast, afterlaser sintering (100 W and 1 mm s−1) the compacts are dense andhomogeneous with very few sub-micrometer sized pores (Fig. 6c).Furthermore, the nanometric ceramic oxide particles (Al2O3) forma strong bond with the matrix (Fig. 6d) due to adequate interdif-fusion aided by laser irradiation during sintering. The results ofEDS micro-analysis of the laser sintered 10 wt% Al2O3 dispersedCu–4.5Cr alloy (Fig. 6c) are summarized in Table 1. The analysisshows that the composition is similar to that of the initial powderblend prior to milling. While oxygen is present either as impu-rity dissolved from milling atmosphere (toluene) or as a part of

ence of negligible amount of Fe (<1 wt%) arises from marginal levelof impurity pick up from the grinding medium during mechanicalalloying or milling.

ters (a) 100 W and 1 mm s−1 and (b) 100 W and 2 mm s−1, respectively.

S. Bera, I. Manna / Materials Chemistry and Physics 132 (2012) 109– 118 113

F ent tea ificati

Cai(cpru

F1

ig. 6. SEM image of 10 wt% nano-Al2O3 dispersed Cu–4.5Cr alloy pressed at ambilloy in cold compacted and laser sintered (100 W and 1 mm s−1) condition at magn

Fig. 7a shows the SEM image of 10 wt% Al2O3 dispersedu–4.5Cr–3Ag alloy pressed at ambient temperature. The powdersre loosely bonded with a large distribution of pores as that seenn the green compact of 10 wt% Al2O3 dispersed Cu–4.5Cr alloyFig. 6a). Laser sintering (100 W and 1 mm s−1) converts this green

ompact into a dense product with only a few ultrafine residualores (Fig. 7b). A close observation at higher magnification (Fig. 7c)eveals that ultrafine Al2O3 of average particle size < 100 nm areniformly dispersed in the Cu-rich matrix. The large window EDS

ig. 7. SEM image of 10 wt% nano-Al2O3 dispersed Cu–4.5Cr–3Ag alloy (a) cold pressed a mm s−1 scan speed at magnifications of (b) 5000× and (c) 150,000×, respectively.

mperature at different magnifications of (a) 2000× and (b) 50,000× and the sameons of (c) 2000× and (d) 50,000×, respectively.

analysis shows that the composition is similar to that of the initialpowder blend prior to milling with very negligible amount of impu-rity (Table 1). Therefore, the present routine of mechanical alloyingand laser sintering has been effective to produce a homogeneousnano-Al2O3 dispersed Cu-matrix composite. It is interesting to note

that the interface of the nanometric Al2O3 particles with the matrixis free from micro-cracks and porosities suggesting that laser sin-tering has been effective to firmly integrate ceramic nano-particleswith the metallic matrix. This kind of strong bond between the

t ambient temperature and cold pressed and laser sintered with 100 W power and

114 S. Bera, I. Manna / Materials Chemistry and Physics 132 (2012) 109– 118

F bient temperature and viewed at magnifications of (a) 2000× and (b) 100,000×, and thes eed viewed at magnifications of (c) 2000× and (d) 100,000×, respectively.

dpm

Cacddlp1vAdfet

apAttol(s8

ttac

T

Fig. 9. (a) Bright field TEM image of laser sintered (100 W and 2 mm s−1) pellets of10 wt% nano-Al2O3 dispersed Cu–4.5Cr–3Ag alloy, (b) the corresponding SAD pat-

ig. 8. SEM image of 10 wt% nano-ZrO2 dispersed Cu–4.5Cr–3Ag alloy pressed at amame alloy cold pressed and laser sintered with 150 W power and 1 mm s−1 scan sp

ispersoid and matrix is essential in metal matrix composites torevent decohesion or fracture during loading and use and ensureaximum level of dispersion strengthening.Fig. 8a and b shows the SEM images of the 10 wt% ZrO2 dispersed

u–4.5Cr3Ag alloy pressed at ambient temperature at (a) 2000×nd (b) 50,000× magnifications, respectively. The green compactontains pores and inhomogeneities among the loosely bond pow-er particles (Fig. 8a). In addition, nanometric ZrO2 particles areispersed homogeneously but loosely over the alloy matrix (Fig. 8b)

ike the microstructure of the green compact of 10 wt% Al2O3 dis-ersed Cu–4.5Cr alloy (Fig. 6b). After laser sintering (150 W and

mm s−1), the compacts appear dense and homogeneous with aery few isolated sub-micrometer sized pores (Fig. 8c). Similar tol2O3 particles in laser sintered (100 W and 1 mm s−1) 10 wt% Al2O3ispersed Cu–4.5Cr alloy (Fig. 6d), the nanometric ZrO2 particlesorm strong bond or grain bridge with the matrix (Fig. 8d). How-ver, the ZrO2 particles are not distributed homogeneously all overhe matrix and the particles are also not of equal size.

Fig. 9a shows the bright field TEM image of laser sintered (100 Wnd 1 mm s−1) 10 wt% Al2O3 dispersed Cu–4.5Cr–3Ag alloy com-osite. The microstructure shows uniformly distributed ultrafinel2O3 particles of <100 nm size in the Cu-base extended solid solu-

ion matrix. The corresponding SAD ring pattern (Fig. 9b) confirmshat the matrix phase is nanocrystalline FCC solid solution. A closebservation of a 30 nm size Al2O3 particle (Fig. 9a and c) reveals theattice image of the particle and the calculated interplanar spacing0.351 nm) is comparable to the interplanar spacing of equilibriumynthetic corundum Al2O3 phase (0.348 nm, JCPDS data file, PDF no3-2080).

An attempt has been made to calculate the maximum tempera-ure (T) reached at the compact surface during laser sintering usinghe analytical model of Ashby and Easterling [45] model and thepproach of solution adopted by Singh et al. [46]. Accordingly, T

an be predicted from the following expression as:

− T0 = AP/v

2��[t(t + t0)]1/2exp

[− 1

4˛

(z2

t+ y2

t + t0

)](1)

tern and (c) higher magnification lattice image of an isolated nano-Al2O3 dispersoidshowing (0 1 2) planes (d = 0.351 nm).

where T0 is the initial temperature, A is the absorptivity at thesample surface, P is the incident laser power, � is the thermal con-ductivity of the sample, is the thermal diffusivity (= �/c), is thedensity, c is the specific heat, v is the laser scan speed, r is the radiusof laser beam and t0 (= r2/4˛) is the time taken for heat to diffuseover half the beam-width.

The product of density and specific heat of the compact is cal-culated using the rule of mixture [47]:

ccc = fCuCucCu + fCrCrcCr (2)

Here f represents the volume fraction of the corresponding con-stituents in the compact indicated as subscripts in Eq. (2). The

S. Bera, I. Manna / Materials Chemistry

Fd2

tt

tocdfpeitblanq3tm

TS

ig. 10. Variation of temperature (T) as a function of time (t) at two different inci-ent power levels (P = 100 and 150 W) and two different scanning speeds (v = 1 and

mm s−1) at each incident laser power level for Cu–4.5Cr alloy.

hermal conductivity of the compact is related to thermal conduc-ivity of the individual constituents as follows [33].

1�

= fCu

�Cu+ fCr

�Cr(3)

Accordingly, from using Eq. (1) the thermal profile as a func-ion of time has been determined for z = 0 and y = 0 using valuesf absorptivity of the compact as 0.67 [35]. Fig. 10 shows the cal-ulated thermal profile generated during laser sintering at twoifferent incident power levels (P = 100 and 150 W) and two dif-erent scanning speeds (v = 1 and 2 mm s−1) at each incident laserower level for Cu–4.5Cr alloy. The maximum temperature gen-rated during laser sintering is about 6000–18,000 ◦C for differentncident laser power and scanning speed. However, it is observedhat within few ms the temperature comes down to 1000 ◦C orelow in all the cases. Therefore, in can be concluded that though

aser sintering initially produces very high temperature only for transient period (few ms, inadequate for surface melting toucleate), faster quenching (for example P = 100 W, v = 1 mm s−1,

uenching rate is 4 × 105 ◦C s−1 up to 30 ms and 104 ◦C s−1 between0 and 100 ms) to temperature well below the melting or liquidusemperature of the alloy rules out the possibility of surface or bulk

elting of the pellets. However, the overall heating and cooling

able 2ummary of physical, mechanical and electrical properties of compacts consolidated und

Alloys Laserpower (W)

Scan speed(mm s−1)

Density(mg m−3)

Cu–4.5Cr 100 2 7.6

100 1 7.5

Cu–4.5Cr–10Al2O3 Green compact 6.4

100 2 7.3

100 1 7.3

150 2 7.1

150 1 7.2

Cu–4.5Cr–3Ag–10Al2O3 Green compact 5.8

100 2 7.3

100 1 7.4

150 2 7.5

150 1 7.4

Cu–4.5Cr–3Ag–10ZrO2 Green compact 6.3

100 2 7.0

100 1 7.3

150 2 7.3

150 1 7.1

and Physics 132 (2012) 109– 118 115

cycle appears sufficient to form grain bridges or necks required forsintering of the powder mass.

The physical, mechanical and electrical properties of interest ofthe laser sintered composites are summarized in Table 2. While thedensity of the laser sintered alloys with different laser power andlaser scan speed is almost comparable and about 7.1–7.5 mg m−3,the hardness values of the laser sintered compacts differ signif-icantly with laser parameter. In addition the hardness of lasersintered Cu–4.5Cr alloy without Al2O3 dispersion is about 130VHN, whereas laser sintered 10 wt% nanometric Al2O3 dispersedCu–4.5Cr alloy records a hardness of 170–225 VHN in samples sin-tered with different combination of laser power and scan speed.Evidently the hardness value increases significantly due to homo-geneous dispersion of nanometric Al2O3 particles in the Cu-richalloy matrix. In addition, the hardness increases further as 3 wt% Agdissolves into the solid solution. Consequently, the 10 wt% Al2O3dispersed Cu–4.5Cr–3Ag alloy registers the highest hardness of275 VHN. The ZrO2 dispersed Cu–4.5Cr–3Ag alloy shows non-uniformity and lower hardness value than that of Al2O3 dispersedalloy composite of similar composition, laser sintered with identicallaser parameters. This is perhaps due to higher initial particle sizeof ZrO2 than Al2O3 during mixing by ball milling. Therefore, local-ized melting during laser sintering might cause the coalescence ofZrO2 particles which further create non-uniformity in the disper-sion of ZrO2 particles (Fig. 8d) in the Cu-rich alloy matrix and resultin inhomogeneity in hardness. In addition to this, it is observedthat laser sintering at higher laser power produces compacts thatshow lower hardness. This may be due to the fact that sinteringwith higher laser power causes more grain coarsening than sinter-ing with lower laser power. It is interesting to note that the densityand hardness values of these laser sintered compacts are lower thanthat of the consolidated compacts by equi-channel angular pressing[40] or high pressure sintering [44] of mechanically alloyed pow-der. This difference is due to the porosity and other defects presentin the compact consolidated by laser sintering. The porosity arisesduring compaction at room temperature and at relatively lowerpressure (100 MPa). However, it may be noted that these hardnessvalues with nano alumina dispersion are significantly higher thanthat in the earlier reported studies on Cu–Cr alloys [48], Cu–Cr–Agor Cu–Cr–Zr alloys [4,11,48,49] and nano Al2O3 dispersed Cu matrix

composite [16].

The average hardness and modulus of elasticity of some selectedcomposite samples were measured with the help of an MTS nano-indenter operated at 50 g force. The values were determined from

er different conditions of laser sintering.

Theoreticaldensity (%)

Hardness(VHN)

Wearscar (�m×106)

Conductivity(% IACS)

86 131 ± 5 6.03 22 ± 185 135 ± 5 5.39 21 ± 177 115 ± 5 10.53 <187 225 ± 5 1.16 14 ± 189 205 ± 5 1.05 14 ± 185 170 ± 5 1.73 15 ± 186 185 ± 5 1.42 18 ± 169 120 ± 5 9.96 <187 275 ± 5 0.94 13 ± 188 215 ± 5 1.04 14 ± 189 185 ± 5 1.35 14 ± 188 165 ± 5 1.70 14 ± 174 121 ± 5 10.68 <182 209 ± 15 1.52 13 ± 185 172 ± 15 1.94 14 ± 185 147 ± 15 2.53 15 ± 183 152 ± 15 1.98 15 ± 1

116 S. Bera, I. Manna / Materials Chemistry and Physics 132 (2012) 109– 118

Table 3Summary of mechanical properties obtained from nano-indentation test.

Alloy Laser power (W) Beam speed (mm s−1) Hardness (GPa) Young’s modulus (GPa)

Cu–4.5Cr 100 1 1.2 39Cu–4.5Cr + 10Al2O3 100 1 1.7 60Cu–4.5Cr–3Ag + 10Al O 100 1 1.8 64

1dhfpa(

tbacErtAditftd(trnta

tAawAtC

Fgp

conductivity slightly decreases with the addition or dispersion of

2 3

Cu–4.5Cr–3Ag + 10ZrO2 100 1

5 to 20 indentations at equivalent locations (ensuring a standardeviation within 1%) and are enlisted in Table 3. The trend of theardness values is almost similar to that of the values obtained

rom Vickers micro-hardness test (Table 2). The 10 wt% Al2O3 dis-ersed Cu–4.5Cr–3Ag alloy laser sintered with 100 W laser powernd a scan speed of 1 mm s−1 shows the highest Young’s modulus64 GPa).

Fig. 11 shows the extent of wear damage calculated fromhe measured dimensions of wear scar (depth/width) sufferedy mechanically alloyed Cu–4.5Cr or Cu–4.5Cr–10Al2O3 samplesfter consolidation by laser assisted sintering using appropriateombinations of laser power (P) and scan speed (v), respectively.vidently, the laser sintered Cu–4.5Cr compacts with nanomet-ic Al2O3 dispersion show considerably higher resistance to wearhan that of the compacts without any ceramic oxide dispersion.pparently the hard and ultrafine ceramic oxide particles embed-ed in the alloy matrix resist the material from wear. In addition,

t seems that the ceramic oxide particles form a strong bond withhe matrix material that prevents the ceramic oxide to wear outrom the matrix and consequently increases the wear resistant ofhe alloy composite. Similarly, Fig. 12 shows the extent of wearamage (calculated from the measured dimensions of wear scardepth/width)) suffered by mechanically alloyed and laser sin-ered Cu–4.5Cr–3Ag–10Al2O3 and Cu–4.5Cr–3Ag–10ZrO2 samples,espectively. Evidently, the wear volume in green compacts is sig-ificantly higher than that in laser sintered samples. It is obvioushat the particles in the green compacts are loosely bound particlesnd very much prone to wear.

Fig. 13a and b shows typical wear scar developed in laser sin-ered (100 W and 1 mm s−1) Cu–4.5Cr alloy without or with 10 wt%l2O3 dispersion after fretting with 20 N load, 1 mm wear spannd 10 Hz frequency, respectively. It is obvious that Cu–4.5Cr alloyithout oxide dispersion shows wider wear scar than that in nano-

l2O3 dispersed Cu–4.5Cr alloy. Likewise, Fig. 13c and d shows

ypical wear scar developed in laser sintered (100 W and 1 mm s−1)u–4.5Cr–3Ag alloy with 10 wt% Al2O3 and ZrO2 dispersion,

ig. 11. Calculated wear volume Cu–4.5Cr and Cu–4.5Cr–10Al2O3 alloys in cold orreen compacted and laser sintered conditions using different combinations of laserower (P) and scan speed (v), respectively.

1.6 57

respectively. Evidently, the laser sintered compacts with nanomet-ric Al2O3 dispersion show rather higher resistance to wear thanthat of the compacts with nanometric ZrO2 dispersion. The trend ofresistance to wear in different laser sintered samples is almost sim-ilar to that of the trend of hardness variation in the same set of sam-ples. Among the laser sintered samples, the Cu–4.5Cr alloy laser sin-tered with laser power of 100 W and scan speed of 2 mm s−1 recordsthe maximum wear (Fig. 13a), while 10 wt% nano-Al2O3 dispersedCu–4.5Cr–3Ag composite laser sintered with laser power of 100 Wand scan speed of 2 mm s−1 registers the least wear volume undercomparable conditions of fretting wear (Fig. 13c). In addition, itis observed from the micrograph that abrasive wear occurs inCu–4.5Cr alloy (Fig. 13a) as the materials come out in forms ofchunk. There is no evidence of oxidation of the debris or base mate-rial during fretting wear test (from EDS analysis). The alloy withultrafine Al2O3 or ZrO2 dispersion shows comparatively adhesivenature of wear (Fig. 13b–d). There is no evidence of wear occursdue to third body in any of the Cu–4.5Cr and Cu–4.5Cr–3Ag alloys(with and without nanometric Al2O3 or ZrO2 dispersion) as well.

Electrical conductivity of the sintered composites has been mea-sured using direct current van der Pauw four-probe method andreported in Table 2. The conductivity values of these composites areslightly lower than that of similar type of cast and wrought alloysreported in the literature [10,48,50,51]. This is perhaps due to theporosity and other defects that are generated during compactionat moderate pressure and room temperature and retained afterlaser sintering. While the Cu–4.5Cr alloy compacts laser sinteredwith 100 W power and 1 mm s−1 or 2 mm s−1 scan speeds recordcomparable as well as the highest electrical conductivity of about20% IACS, nano-Al2O3 or ZrO2 dispersed Cu–4.5Cr or Cu–4.5Cr–3Agalloy composites possess slightly lower conductivity of about 15%IACS albeit greater hardness and elastic modulus. Evidently the

nanometric Al2O3 or ZrO2. However, the decrease in conductivityis not significant (merely 5% IACS). Therefore, it appears that laser

Fig. 12. Calculated wear volume Cu–4.5Cr–3Ag–10Al2O3 andCu–4.5Cr–3Ag–10ZrO2 alloys in cold or green compacted and laser sinteredconditions using different combinations of laser power (P) and scan speed (v),respectively.

S. Bera, I. Manna / Materials Chemistry and Physics 132 (2012) 109– 118 117

F −1 Cr allo( e that

sombtabd

4

wfaadioudsro1tetZtdce

[[

[[[

[[[[

[[[[[

[

ig. 13. Typical wear scar developed in laser sintered (100 W and 1 mm s ) Cu–4.5c) with 10 wt% nano-Al2O3 and (d) 10 wt% nano-ZrO2 dispersion, respectively. Not

intering is capable of producing a homogeneous nano-Al2O3r ZrO2 dispersed Cu-matrix composite with excellent density,echanical properties and adequate electrical conductivity. It may

e noted that laser sintering is effective for compacts of limitedhickness or dimension. However, larger size or complex shapend form of laser sintered products can be developed through layery layer or continuous sintering of powder using an appropriateirect laser assisted manufacturing technique [32].

. Conclusion

Mechanical alloying of elemental Cu, Cr and Ag powder in theeight ratio of Cu–4.5Cr and Cu–4.5Cr–3Ag in a planetary ball mill

or 25 h leads to formation of an extended solid solution with smallmount of undissolved Cr in finely dispersed state. Cold compactiont room temperature (100 MPa) followed by laser sintering withifferent laser power (100 or 200 W) and scan speed (1 or 2 mm s−1)

s effective to form dense (∼7.5 mg m−3) products with and with-ut nanometric Al2O3 or ZrO2 dispersion. While the matrix phasendergoes significant grain growth, the isolated oxide dispersoidso not undergo noticeable change in the particle size after laserintering. Hardness value of the composite without nano-ceramiceinforcement is about 130 VHN. Addition of 10 wt% nano-Al2O3r ZrO2 particle results in an increase in hardness by more than00% (up to 275 VHN) and a significant improvement in wear resis-ance (reduction by a factor of 4–5) without adversely affecting thelectrical conductivity (15–20% IACS) of the compacts. It appearshat dispersion of nano-Al2O3 is more effective than that of nano-rO2, perhaps, due to the smaller initial particle size of Al2O3 than

hat of ZrO2. Thus, laser sintering could be an effective strategy foreveloping nano-oxide dispersed Cu-based composites for electri-al components requiring high wear resistance without sacrificinglectrical conductivity.

[[

[

y (a) without and with (b) 10 wt% nano-Al2O3 dispersion, and Cu–4.5Cr–3Ag alloy the wear damage in (b), (c) or (d) is significantly lower than that in (a).

Acknowledgements

Partial financial support from the International Copper Associa-tion (USA) and Department of Science and Technology, New Delhi(NSTI Grant No. SR/S5/NM-04/2005) is gratefully acknowledged.

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