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Materials Science and Engineering A 529 (2011) 388– 392

Contents lists available at SciVerse ScienceDirect

Materials Science and Engineering A

journa l h o me pa ge: www.elsev ier .com/ locate /msea

ffect of homogenization treatment on microstructure and properties foru–Fe–Ag in situ composites

hixiong Xie, Haiyan Gao ∗, Jun Wang, Baode Suntate Key Laboratory of Metal Matrix Composites, Shanghai JiaoTong University, Shanghai 200240, PR China

r t i c l e i n f o

rticle history:eceived 18 March 2011eceived in revised form 26 July 2011ccepted 7 September 2011vailable online 17 September 2011

a b s t r a c t

The effects of homogenization treatment on the microstructure and properties of Cu–8Fe–0.1Agand Cu–8Fe–0.5Ag (wt.%, similarly hereinafter) in situ composites were investigated. The compositeswere prepared by inductive melting, casting, homogenization treatment and subsequent cold work.Microstructures of the composites were observed and the tensile strength plus electrical conductivity was

eywords:icrostructure

omposite materialsechanical properties

measured respectively. The results show that the prior homogenization treatment results in refinementof the primary Fe dendrites in Cu matrix and promotes the precipitation of the secondary Fe particlesfrom the Cu matrix, which leads to increase in the strength and conductivity of the composites. Duringthe following cold deformation, the homogenized alloys are readily elongated into fine and continuousfilamentary composites compared with the as-cast, leading to high strength due to the increased interfacedensity.

. Introduction

Due to outstanding combinations of high strength and goodlectrical conductivity, in situ deformation processed Cu-baseetal matrix composites (MMCs) consisting of Cu as the matrix

nd a body centered cubic transition metal with low solubility inu (e.g. Nb, Fe or Cr) or face centered cubic Ag as second phase haveeen the subject of extensive research [1–14]. These in situ compos-

tes are promising candidates for material conductor applicationsn robotics and high field magnet design. In the past two decades

ost studies concentrated on the investigations of binary Cu–Nbnd Cu–Ag in situ composites or ternary Cu–Nb–Ag composites tobtain better properties. Bevk et al. [7] manufactured the Cu–18.2%b (vol.%) in situ composites with extremely high strength of.8 GPa. Mattissen and co-workers [2,5] developed ternary in situu–8.2 wt.% Nb–4 wt.% Ag composites and the strength and conduc-ivity reached 1.8 GPa and 46%IACS respectively. Raabe et al. [14]abricated ternary in situ Cu–5 at.% Ag–3 at.% Nb composites usingire drawing. However, Nb and Ag are costly metals, which impede

he large scale applications of such new materials in devices otherhan high field magnets. For that reason, extensive studies aimedt reducing the use of expensive metal.

The Cu–Fe system has attracted considerable attention duringhe past two decades due to relatively low cost of iron comparedo the other possible alloying elements. However, the relatively

∗ Corresponding author. Tel.: +86 21 54742661; fax: +86 21 54742683.E-mail address: gaohaiyan@sjtu.edu.cn (H. Gao).

921-5093/$ – see front matter. Crown Copyright © 2011 Published by Elsevier B.V. All rioi:10.1016/j.msea.2011.09.047

Crown Copyright © 2011 Published by Elsevier B.V. All rights reserved.

higher solubility of iron in Cu at high temperature, coupled withsluggish kinetics of iron precipitation at low temperature increasesthe solute scattering and limits the development of Cu–Fe in situcomposites. Extensive investigations were conducted to optimizethe combination of strength and conductivity of Cu–Fe in situ com-posites through alloying and intermediate heat treatment. Songet al. [15,16] obtained a better combination of strength and con-ductivity of Cu–9 wt.% Fe in situ composites through introducing1.2 wt.% Ag or Cr element with intermediate heat treatments. Theconductivity of the composite increased from 24%IACS to 56.4%IACSwithout severely deteriorating their mechanical properties. Gaoet al. [17–19] produced Cu–11 wt.% Fe–6 wt.% Ag in situ compositeswith three intermediate heat treatments and the tensile strengthand conductivity of reached 1020 MPa/70.5%IACS respectively at astrain of 10. Verhoeven et al. [20] observed that the conductivity ofCu–(10–30) wt.% Fe in situ composites was improved with the pre-cipitation of more Fe through the optimized heat treatments andsubsequent cooling process.

Recent research has reported that homogenization treatmentplayed an important role in the property improvement of Cu–Fein situ composites besides intermediate heat treatments. Forinstance, Wu et al. [21] reported that prior homogenization treat-ment (950 ◦C × 3 h) with furnace cooling significantly increasedthe strength and electrical conductivity of Cu–12 wt.% Fe in situcomposites. The ultimate tensile strength increased from 750 MPa

to 1000 MPa and corresponding electrical conductivity increaseddramatically from 24%IACS to 56%IACS at strain 8. Gaganov et al.[22] also reported that homogenization treatment could sup-press the discontinuous precipitation reaction in subsequent aging

ghts reserved.

Z. Xie et al. / Materials Science and Engineering A 529 (2011) 388– 392 389

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Fig. 1. Microstructure of as-cast (a) Cu–8Fe–0.1Ag, (b) Cu–8Fe

nd improve the strength of Cu–7Ag. However, few studies wereeported to investigate the behavior of homogenization treatmentn ternary Cu–Fe–Ag alloys without intermediate heat treatments.n present work, the effect of homogenization treatment on the

icrostructure and properties of Cu–8 wt.% Fe–0.1 wt.% Ag andu–8 wt.% Fe–0.5 wt.% Ag composites was studied.

. Experimental procedures

Two compositions were chosen for the investigation, namely,u–8 wt.% Fe–0.1 wt.% Ag and Cu–8 wt.% Fe–0.5 wt.% Ag. The meltsere prepared using electrolytic copper, commercial pure iron andg with at least 99.9 wt.% purity. The two alloys were separatelyelted in a vacuum induction furnace using an alumina crucible.

ngots of 14 mm in diameter were cast under an Ar atmospheret pressure of 0.4 × 105 Pa. The cylindrical ingots were swaged to

mm at room temperature, and then drawn into wires to a min-mum diameter of 0.2 mm using successively smaller dies. Somengots were reheated to 900 ◦C and held for 3 h followed by fur-ace cooling before rotary swaging. The cold drawing strain � isescribed by � = ln(A0/A), where A0 and A are the original and finalross-section areas respectively.

Microstructure observation of the composite was carried outsing optical microscopy (OM) and JSM-6460 scanning electronicroscope (SEM) equipped with an energy dispersive X-ray spec-

rometer (EDS). The ultimate tensile strength of the compositesires was tested on a united machine equipped with an exten-

ometer with a strain rate of 5.0 × 10−4 s−1 at room temperature.he electrical resistivity of the samples was measured using a dig-tal micro-ohmmeter with precision of 1 �� at room temperature.

. Results and discussion

.1. Microstructure

Fig. 1 shows the microstructure of as-cast and homogenizedu–8Fe–0.1Ag and Cu–8Fe–0.5Ag alloys. For the as-cast alloys,

g and homogenized (c) Cu–8Fe–0.1Ag, and (d) Cu–8Fe–0.5Ag.

coarse primary Fe dendrites are found distributing in the Cu-richmatrix with low Ag content (Fig. 1(a)). With increasing Ag content,the primary Fe dendrites are refined, as show in Fig. 1(b). The aver-age diameter of primary dendrite of Cu–8Fe–0.5Ag is about 3 �m,less than that of Cu–8Fe–0.1Ag (6.2 �m), which is in agreementwith our previous work [3,17,23]. The refinement of the dendritesmay be ascribed to the reduction of interface energy between �-Feand liquid copper with increasing Ag atoms during solidification.Fig. 1(c) and (d) represents the microstructure of homogenizedCu–8Fe–0.1Ag and Cu–8Fe–0.5Ag alloys. Compared with the as cast,the primary Fe dendrites of both alloys are smaller and distributedmore uniformly in the Cu-rich matrix. The average diameter ofdendrite is approximately 2.8 �m. During homogenization treat-ment process, dissolution of the primary Fe dendrites happenedand coarse dendrites broke up into smaller ones. Wu et al. [21] alsoreported that the homogenization treatment reduced the primaryFe dendritic arms in Cu–6Fe alloys. With increasing Ag content, fewchanges were observed in primary Fe dendrites size. The reason isnot clear yet. In addition, many small secondary Fe particles pre-cipitated from Cu matrix due to the small solubility of Fe in Cu atroom temperature and furnace cooling. The morphology of the sec-ondary Fe particles is shown in Fig. 1(c). The average diameter isabout 0.4 �m, far less than that of the primary dendrites.

Fig. 2 shows the longitudinal microstructure of Cu–8Fe–0.1Agand Cu–8Fe–0.5Ag alloys deformed at � = 3.6. The Cu-rich phaseand primary Fe dendrites are elongated and rotated to the draw-ing direction during the cold deformation. Most of the primarydendrites in the as-cast Cu–8Fe–0.1Ag develop into blocks dis-tributing in the Cu matrix (Fig. 2(a)). With increasing Ag, the as-castCu–8Fe–0.5Ag transforms into elongated continuously filamen-tary structure composites with the average spacing size betweenthe filaments about 10 �m (Fig. 2(b)). The changes in morphol-ogy of filaments are attributed to size of the primary Fe dendrites.

Hong’s and Sinclair’s work [16,24] suggested that small Fe dendritescould develop into continuous filaments while coarse dendriteswere broken into small ones or evolved into elongated particlesunder large deformation. Fig. 2(c) and (d) shows the longitudinal

390 Z. Xie et al. / Materials Science and Engineering A 529 (2011) 388– 392

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composites can be considered to be responsible for the behav-ior of strength change with strain [28,29]. Usually, in the in situcomposite, the dislocation multiplication and interface obstacleare responsible for the strengthening behaviors, and at the low

Fig. 2. As-drawn microstructure (� = 3.6): (a) Cu–8Fe–0.1Ag, (b) Cu

icrostructure of the homogenized alloys. The dense and finelaments with dispersed secondary Fe particles are uniformlyistributed in the matrix. The average spacing size between thelaments is about 6 �m, which is less than that of the as-cast alloys∼10 �m). The initially finer and uniform distribution of the pri-

ary Fe dendrites caused by homogenization treatment may beesponsible for the refinement of filaments. In addition, the homog-nization treatments can remove many dislocation defects in thelloys. Previous literatures [8,25] reported that the annealing pro-ess at higher temperature led to at least a two orders of magnitudeeduction of dislocation density. Therefore, the homogenized alloysre easy to be elongated into continuous filamentary structure com-osites compared to the as-cast alloys.

.2. Mechanical properties

Fig. 3 represents stress–strain curves of cold drawnu–8Fe–0.1Ag in situ composites with drawing strain 5.4. Therere similar stress–strain curves in the experimental strain range.he ultimate tensile strength and the elongation to fracture ofhe homogenized Cu–8Fe–0.1Ag in situ composites are greaterhan that of the as-cast. The higher strength is attributed tohe fine-scale filamentary microstructure in the homogenizedlloys, as mentioned above. Fig. 4 reveals the fracture surface ofensile samples at strain of 5.4. The fractures of the homogenizedomposite exhibits ductile cup-cone type fractures, while that ofhe as-cast takes on shear characters. The transition from cup-coneype to shear type of Cu–Fe–Ag composites is related to the prioromogenization treatment. The in-depth investigation will bearried out later. Higher magnification observation reveals thathe fracture surface of the homogenized composites is coveredith a population of voids of various size and fine dimples, within

hich Fe filaments can be observed (Fig. 4(b)), which is consistentith the smaller filament size and narrower inter-filament spacinghile dimples in the fracture of the as-cast is relatively coarser

Fig. 4(d)). The finer dimple size of the homogenized should be due

0.5Ag and homogenized (c) Cu–8Fe–0.1Ag, and (d) Cu–8Fe–0.5Ag.

to the high ductility of the composite. Liu et al. [26] also reportedthat the ductility of Cu–14Fe–0.06Ag in situ composites was higherthat of the Cu–14Fe for fine dimple sizes.

Fig. 5 shows the dependence of the ultimate tensile strengthon the cumulative cold drawing strain for the cold workedCu–8Fe–0.1Ag and Cu–8Fe–0.5Ag in situ composites. The strengthof the composites increases with increasing drawing strain, whichis in accordance with previous work [23,26,27]. At low strain level,the strength of homogenized Cu–8Fe–0.1Ag is a bit lower comparedwith the as-cast, while it increases rapidly with increasing strain,especially at high strain level. The strengthening mechanism of dis-location multiplication and interface obstacle for Cu-based in situ

Fig. 3. Typical stress–strain curves.

Z. Xie et al. / Materials Science and Engineering A 529 (2011) 388– 392 391

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3.3. Electrical conductivity

Fig. 4. SEM images of fracture surface with � = 5.4, (a and b)

train level, the dislocation multiplication contributes much. Theomogenization treatment removes large amount of defects inhe composites, leading to more obvious strengthening effect inhe as-cast alloys. With the increase of deformation, the dynamicecrystallization and recovery will take place in the compositesuring drawing deformation and differences in dislocation densityetween the two kinds of composites can be neglected, leading toame dislocation contribution to the strength [8]. And at the sameime, the interface density of Cu and Fe filaments increases withncreasing strain, strengthening effect from the interface obstacle ishus predominant [21]. The homogenized alloys have higher inter-ace density due to much finer filamentary structure, resulting inast increment in the strength. In addition, the small secondary

e particles in the Cu matrix may develop into nano-scale fila-ents during large cold drawing and increases significantly the

nterface density of Cu matrix and Fe filaments [21]. Therefore,

Fig. 5. Dependence of the strength on the strain.

genized Cu–8Fe–0.1Ag and (c and d) as-cast Cu–8Fe–0.1Ag.

the homogenized alloys have even higher strength compared tothe as-cast alloys at higher strain. The strength of the homoge-nized Cu–8Fe–0.5Ag microcomposites is 70 MPa higher than thatof Cu–8Fe–0.1Ag. Solution strengthening of the matrix and nano-scale Ag filaments strengthening may contribute to the difference.The solute Ag atoms will precipitate from the Cu matrix duringsluggish furnace cooling and then be elongated into nano-scale fila-ments at heavy deformation due to good ductility, giving additionalstrengthening effect to the composites.

Fig. 6 represents the electrical conductivity dependence on con-venience against cumulative cold drawing strain for Cu–8Fe–0.1Ag

Fig. 6. Dependence of the electrical conductivity on the strain.

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nd Cu–8Fe–0.5Ag in situ composites. The conductivity of in situomposites decreases gradually with increasing drawing deforma-ion. The electrical conductivity of the homogenized compositess about 5–7%IACS higher than that of the as-cast at each draw-ng strain, and the increment in Cu–8Fe–0.5Ag is a bit higher thanhat in Cu–8Fe–0.1Ag. Extensive investigations showed that thelectrical conductivity of ternary Cu–Fe–Ag in situ composites wasrimarily controlled by the resistivity of the Cu matrix, which coulde portioned into the contribution of four scattering mechanismollows [15,16,21,26,28]:

M = �pho + �dis + �int + �imp

here �pho is the resistivity contribution from phonon scattering,dis the dislocation scattering, �int the interface scattering, and �imp

he impurity scattering. At room temperature, the phonon scat-ering can be ignored. Verhoeven’s work [8,30] reported that theislocation density was constant at strains greater than 4, result-

ng from the occurrence of dynamic recovery/recrystallization ofhe Cu matrix. Therefore, the resistivity in heavily drawn Cu–Fen situ composites should be mainly determined by the interfacecattering and impurity scattering. Wu et al. [21] reported that thenterface scattering in Cu–Fe microcomposites would be predom-nant if the filaments approached nanoscale. Generally speaking,he contributions of both �int and �dis of the in situ compositeshould be roughly the same for Cu–Fe–Ag with same percentagef Fe and the same drawing strain. The impurity scattering is moreredominant in Cu–Fe–Ag in situ composites. Previous investiga-ions [9–11,15–19] suggested that the solution of minority Ag intou had less effect on the electrical conductivity of Cu matrix dueo the similar electronic structure. Therefore, the resistivity of Cu

atrix depends on the amount of solute Fe atoms. For all in situomposites, with increasing strain, the amount of vacancy increasesignificantly and then deteriorates the conductivity of the compos-tes. On the other hand, Raabe et al. [31] reported that multiphasen situ composites with limited mutual solid-state solubility under-

ent plasticity-stimulated chemical mixing to levels far beyondquilibrium solubility, which indicated that increasing deforma-ion resulted in the dissolution of more Fe into the Cu matrix andeduced the conductivity of the composites. The increase in thelectrical conductivity of the homogenized Cu–Fe–Ag is relevanto the resistivity of Cu matrix. The prior homogenization treat-

ent with slow furnace cooling decreases the amount of solutee in Cu at room temperature, resulting in the reduction of thempurity scattering level due to the precipitation of secondary Fearticles from the Cu matrix. Gao’s work [17] suggested that theresence of Ag could inhibit the re-dissolution of Fe atoms intou matrix during the homogenization treatment process. As a con-equence, compared to the as-cast alloys the impurity scatteringevel of homogenized Cu–Fe–Ag in situ composites will decrease

ignificantly, which leads to higher electrical conductivity. The con-uctivity of homogenized Cu–8Fe–0.5Ag is a bit higher than that ofu–8Fe–0.1Ag due to more precipitation of Fe from Cu matrix withresence of Ag.

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ineering A 529 (2011) 388– 392

4. Conclusion

The effect of homogenization treatment on microstructure andproperties of Cu–8 wt.% Fe–0.1 wt.% Ag and Cu–8 wt.% Fe–0.5 wt.%Ag microcomposites was investigated. The homogenization treat-ment leads to a uniform distribution of the Fe dendrites andpromotes the precipitation of the secondary Fe particles. Thehomogenized Cu–8Fe–0.1Ag and Cu–8Fe–0.5Ag alloys are read-ily elongated into filamental structure microcomposites due to therefinement of primary Fe dendrites. The strength and conductivityof the homogenized composites is higher than that of the untreatedat same strain level. The increase in the conductivity is attributed tothe precipitation of secondary Fe particles. The increase in strengthis owed to the increased interfaces density of the homogenizedcomposites.

Acknowledgements

The project is supported by the National Natural Science Foun-dation of China (Grant Nos. 50801046 and 11005143), ResearchFund for the Doctoral Program of Higher Education of China (GrantNo. 200802481138) and Shanghai Science and Technology Devel-opment Funds (Grant No. 10dz1203600).

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