Mechanical and tribological behavior of nanostructured copper–alumina cermets obtained by Pulsed Electric Current Sintering

Wear 297 (2013) 762–773

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Mechanical and tribological behavior of nanostructured copper–alumina cermets obtained by Pulsed Electric Current Sintering Wassim Zein Eddine a,n, Paolo Matteazzi b, Jean-Pierre Celis a a b

KU Leuven, Department of Metallurgy and Materials Engineering (MTM), Kasteelpark Arenberg 44-Bus 2450, 3001 Leuven, Belgium CSGI and MBN Nanomaterialia spa, Via Bortolan 42, 31050 Vascon di Carbonera, Treviso, Italy

a r t i c l e i n f o

a b s t r a c t

Article history: Received 24 March 2012 Received in revised form 10 August 2012 Accepted 15 October 2012 Available online 26 October 2012

Nanostructured Cu–Al2O3 powders obtained by the reduction of CuO with Al in a high energy ball mill were successfully consolidated by Pulsed Electric Current Sintering (PECS). The effect of the composition and microstructure of these PECS Cu–Al2O3 cermets on their strength was investigated. The friction and wear behavior of these cermets were studied under reciprocating sliding against corundum at 23 1C and 50% RH, and compared to the behavior of coarse grained PECS sintered pure copper. The effect of grain size on the coefficient of friction was masked by the formation of a surface tribolayer. The wear depth recorded on Cu–Al2O3 is lesser than half the one on coarse grained copper. Surface and subsurface deformation studied through FIB cross-sections showed that delamination and oxidative wear were active on Cu and Cu–Al2O3 cermets respectively under the current sliding test conditions. PECS Cu–Al2O3 cermets showed a good thermal stability even at 600 1C. & 2012 Elsevier B.V. All rights reserved.

Keywords: Nanostructured copper–alumina Pulsed electric current sintering Dispersion strengthening Tribology Mild oxidational wear Sub-surface deformation

1. Introduction Materials with ultrafine/nanostructured grains offer some advantages over their micron-sized counterparts when it comes to mechanical properties. Nanostructured powders can be obtained by several techniques including inert gas condensation and mechanical alloying [1], the latter showing a large versatility in producing homogeneous metallic alloys and cermets. Reduction reactions can also be achieved by mechanical activation during ballmilling, such as the reduction of various metal oxides with aluminum to obtain nanometer-sized a-Al2O3-metal (Fe, Cr, Cu, and Ni) composites [2]. The presence of an inert second phase, e.g. nanosized oxides, in single phase metallic materials can enhance the strength of materials in line with the Orowan strengthening mechanism [3]. Kupcis et al. [4] reported that Cu–Al2O3 containing Al2O3 particles with a size between 40 and 80 nm exhibits enhanced tensile properties in line with the Orowan mechanism while coherent particles smaller than 15 nm are sheared or fractured at yield. These nano-sized oxides pin not only the movement of dislocations but also the grain boundaries. This last process helps to prevent grains from growing when exposed to high temperatures according to the Zener drag process [5]. Nowadays, bulk nanostructured materials can be synthesized in different ways like pressureless sintering and hot pressing of

n

Corresponding author: Tel.: þ 32 16 32 1292; fax: þ 32 16 32 1991. E-mail addresses: [email protected] (W. Zein Eddine), [email protected] (P. Matteazzi), [email protected] (J.-P. Celis). 0043-1648/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.wear.2012.10.011

nanostructured powders. To harvest improvements in engineering properties, care must be taken on exposing these nanostructured materials to high temperatures during consolidation. The duration of exposure to high temperatures is a critical factor in inducing a grain growth. Pulsed electric current sintering (PECS) is a recent technique to sinter powders to high densities in a relatively short time at low or high temperatures and pressures. It is based on passing a current through a powder compact at high heating rates up to 400 1C/ min [6]. Near-full density ceramics and cermets (SiC, NbC, WC–Co etc.) were achieved with that technique [7,8] as well as pure metals and metallic alloys such as FeMo, Cu and Ni [9–11]. Zhang et al. [10,12] have shown that the PECS sintering temperature of copper powders has to be below 750 1C (at 50 MPa) to avoid a substantial grain growth, but above 250 1C (at 600 MPa) to achieve a relative density better than 95%. They achieved a grain size of 120 nm after sintering of 50 nm copper powders at 300 1C and 600 MPa. Ultrafine/nanostructured materials have been reported to have an improved wear resistance over micron-sized ones. Studies on Cu, Ti, Al, and Ni [13–15] have shown an improved wear resistance by at least a factor of two when moving from coarse grained to nanostructured materials. The relation between the coefficient of friction and grain size was however found to be less straight forward. Studies on Cu [16] and Al [14] showed that the peak coefficient of friction (the highest value before steady state) decreases by 40% on nanostructured Cu and Al as compared to the coefficient of friction recorded on similar but coarse grained Cu and Al. Considering that an increase in hardness and E-modulus reduces the adhesive component of friction, one would indeed expect a drop in the coefficient of friction with nanostructuring. On the contrary,

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Table 1 Powder sintering conditions and some properties of the PECS sintered samples. Material tested

CA5n-450 CA5n-900 Cu

PECS processing conditions

Characteristics of PECS treated samples

Die/punch material

T (1C)

Pressure (MPa)

Time (min)

r/rtheo  100 (%)

Grain size (nm)

HV300 (kgf/mm2)

E-modulus (GPa)

Steel Graphite Graphite

450 900 700

285 60 60

7 7 7

95.8 96.1 99.3

160 280 20000

304 7 9 207 7 4 73 7 3

1417 4 1507 7 1337 6

the steady state coefficient of friction recorded on these nanostructured and coarse grained Cu and Al materials was the same. On comparing nanostructured and coarse grained Ti and Cu materials, no difference in the steady state coefficient of friction was recorded [13]. However, research by Wang et al. [17] on Ni showed a decrease of the coefficient of friction with nanostructuring while their work on Co showed an increase of the coefficient of friction with nanostructuring. A study of the friction recorded on equal channel angular pressed (ECAP) Cu with repeated scratching tests against a spherical sapphire indenter, where material strength was varied at a relatively constant grain size, revealed that the strength rather than the grain size is a friction determining parameter [18]. Zhou et al. [19] studied Cu–Al2O3 sliding against a AISI440C counterbody at low loads (0.1–1 N) and found the least wear rate at 2 wt% Al2O3 with a wear rate 3 times lower than the one of pure copper, while the coefficient of friction for Cu–Al2O3 varied between 0.45 and 0.55. To the authors’ best knowledge, no literature is available on the tribological behavior of PECS sintered nanostructured Cu–Al2O3 cermets. This study reports on the possibility to retain the ultrafine structure of high energy ball milled copper powders reinforced with nano-sized alumina after densification by PECS. The effect of grain size and the presence of Al2O3 particles on hardness, friction, wear behavior, and thermal stability of these nanostructured Cu-based cermets are investigated and discussed in terms of material-related mechanisms.

2. Experimental procedure 2.1. Powder preparation Nanostructured Cu–Al2O3 composite powders containing 5 vol% alumina particles were produced by high energy ball milling under vacuum (MBN Nanomaterialia S.p.a.). Hereto the starting powders used were Cu powders (99.5% purity, mesh 45–350 mm), CuO powders (99.6% purity, mesh 45–350 mm), and Al powders (99.4% purity, mesh 75–350 mm). They were milled together for 8 h in a steel container at room temperature. Hardened steel milling balls were used at a ball-to-powder ratio of 15:1. During the high energy milling, alumina was formed due to the highly exothermic redox reaction between stoichiometric quantities of CuO and Al as follows [20]: Cuþ 3CuOþ2Al-4CuþAl2O3

(1)

At the same time, a powder agglomeration took place ending up in a composite powder that can be handled without health or environmental risks. The size distribution of the agglomerated powders that were further processed by pulsed electric current sintering (PECS) was mesh þ 38–75 mm. For comparison purposes, Cu powders (mesh 45–350) were also processed by pulsed electric current sintering. 2.2. Pulsed electric current sintering of pure Cu and agglomerated Cu–Al2O3 powders In the current study, Cu and Cu–Al2O3 powders were PECS sintered under 0.05 Pa vacuum to achieve compacted samples of

Fig. 1. Selected sets of parameters used in this study for the consolidation of Cu and Cu–Al2O3 powders based on results by Zhang et al. [10,12] and limitations in the used PECS equipment.

30 mm diameter and 3 mm height using a die/punch setup made of either steel for high pressure processing or graphite for high temperature processing. The PECS sintering conditions are summarized in Table 1. This processing was aimed to obtain a relative density above 95% for all samples without altering the nanostructure. Based on the results of Zhang et al. [10,12] (see Section 1) and on the technical limitations of the used PECS equipment, a processing parameters window was worked out (Fig. 1). Within that window, a sintering temperature of 450 1C was selected higher than the 250 1C proposed by Zhang et al. to compensate for the lower pressure reachable in our PECS equipment, and to account for a potential decrease in densification due to the presence of Al2O3 particles. A sintering temperature of 900 1C was selected as close as possible to the melting temperature of Cu but below it to account for any possible overshoot on heating up. A sintering temperature of 700 1C was selected based on previous experience with PECS sintering of pure Cu. The holding time of 7 min was the time after which the relative displacement of the piston stabilized at 710 mm/min. For simplicity, ‘‘CA5n-450’’ is used as a code name for Cu–Al2O3 PECS sintered at 450 1C while ‘‘CA5n-900’’ refers to the one sintered at 900 1C and ‘‘Cu’’ refers to the pure copper PECS sample. The resulting PECS sintered cylindrical samples were subsequently sand blasted, ground, and finally polished on one of their flat circular sides with 0.25 mm colloidal silica, cleaned ultrasonically first with acetone and then with ethanol before further testing. 2.3. Structural characterization of PECS sintered samples The density of the sintered samples was measured by the Archimedes method and is reported in Table 1 as the ratio of the measured density, r, and the theoretical one, rtheo multiplied by 100. The constituent phases were determined by X-ray diffraction (Seifert 3003 T/T with Cu-Ka radiation). The microstructure and chemical composition of the PECS sintered bulks were characterized by

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ultra-high resolution dual beam microscopy SEM/FIB (FEI Nova 600 NanoLab) operated in SE and BSE modes and equipped with an energy dispersive X-ray spectroscope (EDXS). The grain size was determined by the electron backscattered diffraction (EBSD) technique. The hardness is reported as an averaged value of 5 measurements on polished samples using a Vickers hardness tester operated at 300 g load applied for 10 s. The nanohardness and E-modulus were derived from nanoindentation loading/unloading curves. These measured values are also reported in Table 1. 2.4. Tribological tests The tribological properties were determined in a dry reciprocating sliding ball-on-flat tester [21] operated at 5 N normal force, 500 mm peak-to-peak displacement amplitude, and 5 Hz sliding frequency for up to 100,000 cycles. The displacement amplitude selected allows to over-ride several sintered powders during each sliding cycle, and to achieve a gross slip regime. Corundum balls of 10 mm diameter were used as counterbody (CERATEC, 2000 HV, E¼300 GPa). All the tests were performed in ambient air at 23 1C and 50% relative humidity. The coefficient of friction was derived from the tangential force, Ft, versus displacement, d, loops recorded during the reciprocating sliding tests. After the sliding tests, the maximum wear depth on the PECS samples was determined by non-contact profilometry (WYKO NT 3300). The debris, the wear track morphology, and its composition as well as the subsurface microstructure were investigated by SEM/FIB.

3. Results and discussion 3.1. (Micro)structural characteristics of PECS Cu–Al2O3 samples The structure of the Cu–Al2O3 starting agglomerated powders and the PECS samples were derived from XRD spectra (Fig. 2). The only visible peaks appearing in both CA5n-450 and CA5n-900 XRD spectra were those of copper. Al2O3 peaks did not clearly appear which might be due to the low weight percent and the nanometric size of the alumina particles and the fact that they are embedded in a Cu matrix which has high density. A sharpening in the Cu peaks is noticed when one goes from the starting CA5n-powder to powders sintered at 450 1C and 900 1C. This reveals that grain growth took place during the PECS processing.

The microstructure of both CA5n-450 and CA5n-900 reveals two major areas, namely Cu–Al2O3 zones surrounded by a pure copper network which was confirmed by EDXS analysis (Fig. 3a). The overall microstructure of CA5n-450 is finer than the one of CA5n-900, and the network of pure copper surrounding the Cu–Al2O3 zones is thinner in the case of CA5n-450. A closer look (Fig. 3b) reveals in the case of PECS Cu–Al2O3 the presence of porosity at the inter-powder boundaries. High resolution imaging inside the Cu–Al2O3 (Fig. 3c) area reveals a homogeneous dispersion of Al2O3 particles with sizes below 50 nm. The two zones formation is explained in literature on Cu–Al2O3 by a preferential flow of copper during sintering [22]. Published data on Ag–SnO4 [23,24] showed also a similar microstructure which was explained by exo-diffusion where gases retained in the powder after milling desorb upon heating, and probably condense at nano-pores located in the boundaries (zones of low crystallinity) inducing stresses in the surrounding matrix. This results in a gradient of chemical potential for vacancies which diffuse toward the stress sources. As heating proceeds, a counter flux of Ag atoms toward the free surface results in pores and in a surface Ag-layer. It is interesting to note that at the sintered powder boundaries in CA5n-450 and CA5n-900 where a pure copper network is detected after PECS sintering, the absence of Al2O3 particles has caused during the 7 min treatment a faster grain growth up to several hundreds nanometers or even few microns in size as visible in the EBSD image Fig.3d. This is in line with the Zener drag process which shows the positive effect of dispersoids on the pinning of grain boundaries and on slowing down the grain growth of a metal matrix when exposed to high temperatures [5]. The grains determined by EBSD in the starting Cu–Al2O3 powders and in CA5n-450 and CA5n-900 samples are shown in Fig. 4a, b and c, respectively. The average grain size determined by the intercept length method (30 lines) is at about 120 nm, 160 nm, and 280 nm for starting powders, CA5n-450, and CA5n-900 samples respectively. Only a limited grain growth happened during the PECS processing at 450 1C for 7 min but the grain size doubled after a PECS processing at 900 1C for 7 min. The EBSD patterns allow thus to conclude that even for a short sintering time of only 7 min grain growth takes place to a certain extent in this rapid sintering technique. Zhang et al. [10] investigated grain growth during PECS in copper with an average starting grain size of 1 mm. After a PECS processing at 400 1C for 6 min the grain size was 2.5 mm and after processing at 600 1C it was 6 mm. They also reported [12] that a

Fig. 2. XRD of the Cu–Al2O3 starting agglomerated powders and the two PECS processed samples.

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Fig. 3. (a) BSE image showing the microstructure of PECS processed CA5n-900. EDX analysis (inset) reveals zones of Cu–Al2O3 surrounded by a network of pure Cu at the powder boundaries, (b) Porosity at the inter-powder boundaries, (c) High resolution imaging showing the nano-Al2O3 particles and (d) EBSD image showing large grains inside the Cu network as opposed to finer grains in the Cu–Al2O3 which emphasizes the importance of grain boundaries pinning by nano-Al2O3.

copper grain growth from 50 nm to 260 nm took place during PECS when the processing conditions were 500 1C for 5 min. 3.2. Strengthening mechanisms Ductile metallic powders undergo during ball milling a severe mechanical and plastic deformation leading to a grain refinement. That process is however accompanied by a large amount of work hardening [25]. Through the addition of a nanosized inert second phase like Al2O3-particles to a single phase nanostructured matrix, e.g. Cu, four major strengthening mechanisms can add up resulting in a global strengthening effect, namely: a- a strengthening resulting from the grain size known as grain boundary strengthening, b- a strengthening resulting from an oxide particle reinforcement, known as dispersion strengthening, c- a strengthening resulting from the residual plastic strain due to the thermal mismatch between particles and matrix during post-processing cooling, and d- a strengthening resulting from the accumulation of dislocations due to plastic deformation during ball milling, known as work hardening. On the contrary, the effect of 5 vol % Al2O3 particles on the load bearing capacity is small due to the low content fraction of hard particles. The grain boundary strengthening, DsHP, is given by the HallPetch relation [26,27]: k

ffiffiffi DsHP ¼ s0 þ pHP

d

ð2Þ

with s0 the Peierls stress equal to 25 MPa for Cu, kHP a strengthening coefficient constant equal to 0.14 MPa m1/2 for Cu, and d the grain size of Cu expressed in m. This shows that as the grain size decreases the strength of the material increases. In the case of the presently investigated PECS samples, this will account for an

increase in yield stress by 370 MPa and 284 MPa for CA5n-450 and CA5n-900 respectively. The dispersion strengthening, Dsorowan, is given by the Orowan-Ashby equation [28, 29]:

Dsorowan ¼

0:13Gm b

l

r  ln b

ð3Þ

with Gm the shear modulus of the copper matrix equal to 42 GPa, b the Burgers vector for copper equal to 2.56  10  10 m, r the particle radius, and l the interparticle distance given by " 1 # 1 3 l ¼ dp 1 ð4Þ 2V p with dp the particle diameter of average 35 nm, and Vp the volume fraction of dispersed Al2O3 particles. This strengthening is due to the resistance of closely spaced hard particles to the passing of dislocations. Orowan bowing is necessary for dislocations to bypass particles [3]. As a result, an additional increase in yield strength by 146 MPa is anticipated for both CA5n-450 and CA5n-900. The strengthening due to particles-matrix post-processing thermal mismatch, Dsthermal, is given by [30]: pffiffiffiffi Dsthermal ¼ kGm b r ð5Þ with k a constant equal to 1.25, and r the dislocation density induced by plastic strain due to a thermal mismatch. This r is given by [29]:

r ¼ 12

DaDTV p   bdp 1V p

ð6Þ

with Da the difference in the coefficient of thermal expansion between the copper matrix and the Al2O3 particles which is equal to 7.8  10  6 K  1, and DT the difference between the processing and the test temperatures. An increase in yield strength due to this mechanism by 206 MPa and 295 MPa is anticipated for CA5-450 and CA5n-900 respectively.

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Fig. 4. EBSD images showing grains of different orientations in (a) Cu–Al2O3 powder (b) CA5n-450 and (c) CA5n-900. Note: One color is not linked to a given crystallographic orientation, but different colors are used simply to differentiate grains. A grain boundary is defined once a 51 minimum orientation difference exists.

Table 2 Contribution of the various strengthening mechanisms to the yield strength of PECS processed Cu–Al2O3. The difference in these contributions is listed in the last row.

CA5n-450 CA5n-900

Ds(CA5n-450-CA5n-900)

DsHP

DsOrowan

Dsthermal

SsTotal

(MPa)

(MPa)

(MPa)

(MPa)

370 284 86

146 146 0

206 295  89

722 725 3

In total, these three major strengthening mechanisms should provide an increase in yield strength by 722 MPa for CA5n-450 and by 725 MPa for CA5n-900 as shown in Table 2. Note that according to Tabor [31], the Vickers hardness of large indents is around 3 times the yield strength. Taking that into consideration, an approximation of the strength should be around 993 MPa and 676 MPa for CA5n-450 and CA5n-900 respectively. In the case of CA5n-900, the measured and calculated values are 49 MPa apart which could be accommodated by approximations made in the theoretical equations. However, the difference in yield strength in the case of CA5n-450 is much larger, namely around 271 MPa. In order to understand that difference, the following assumption can be made: the CA5n-900 sample was processed at a temperature above 0.8 Tm of copper, it is thus in an annealed state, and therefore the yield strength contribution of the increase in dislocation density due to plastic strain during ball milling, also known as work hardening (DsWH), can be neglected. Since the only difference in the processing of CA5n-900 and CA5n-450 was the sintering temperature and since phase transformations or chemical interactions between Cu and inert Al2O3 particles are not expected, the additional strength due to work hardening is a logical addition. If the only microstructural difference between CA5n-450 and CA5n-900 are additional dislocations from work hardening, DsWH can be estimated by subtracting the measured yield strength of CA5n-900 from that of CA5n-450. However, the contributions of Dsthermal and DsHP differ for both materials, and they must be taken into account in the subtraction. These differences are mentioned in the last row of

Table 2. The contribution of DsWH to the strength of CA5n-450 is therefore estimated to be DsWH ¼Measured s(CA5n-450)–Measured s(CA5n-900)–Ds(CA5n-450  CA5n-900) ¼993–676–(3)¼320 MPa. This brings the difference between the calculated and the measured yield strength of CA5n-450 down to 49 MPa. This indicates thus that PECS of Cu–Al2O3 performed at temperatures below 0.5 Tm of Cu for short holding times (7 min in this case) preserves a higher dislocation density due to work hardening achieved during the ball milling process compared to the more annealed state of CA5n-900 sintered at 900 1C. 3.3. Tribological behavior of PECS sintered Cu and Cu–Al2O3 (5 vol%) The tribological behavior of PECS sintered nanostructured Cu–Al2O3 (5 vol%), namely CA5n-450 and CA5n-900, was compared to the one of PECS sintered coarse grained Cu. The initial contact pressure, Pmax, resulting from the applied loading conditions (ball-on-flat) was estimated from the Hertzian contact pressure approach [32]: Pmax ¼

3 Fn 2 pa2

ð7Þ

with Fn the normal load which is 5 N in this work, and a the resulting contact radius given by  1 3 FnR 3 ð8Þ a¼ 4 En with R the radius of the counterbody equal to 5 mm in this work, and En the effective Young’s modulus given by 1u21 1u22 1 þ n ¼ E1 E2 E

ð9Þ

with E1 the Young modulus of the sample, namely 141 GPa, 150 GPa and 133 GPa for CA5n-450, CA5n-900, and Cu respectively, E2 the Young modulus of corundum, namely 300 GPa, u1 is the Poisson ratio of the sample, namely 0.33 on average for all samples, and u2 the Poisson ratio of corundum, namely 0.24. The calculated maximum Hertzian pressure is equal to 756 MPa, 777 MPa and 736 MPa for CA5n-450, CA5n-900 and Cu respectively. These values are close to or below the yield strength of

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Fig. 5. SE images showing (a) a wear track on PECS Cu after 20,000 sliding cycles at a normal load of 5 N, 5 Hz, 500 mm displacement in ambient air at 23 1C and 50%RH, and (b) plastic deformation slip planes at the edge of this wear track.

Fig. 6. SE images of the wear track after 60,000 cycles on (a) CA5n-900 and (b) Cu. Tests performed in ambient air (23 1C, 50%RH) at 5 N normal load, 5 Hz and 500 mm displacement. Counterbody was a corundum ball (diameter 10 mm).

Fig. 7. SE images taken inside the wear track on CA5n-450 obtained after a reciprocating sliding test done at 5 N normal load, 5 Hz, and 500 mm displacement in ambient air (23 1C, 50%RH) for 20,000 cycles: (a) before; and (b) after ultrasonically cleaning for 5 min in ethanol.

Cu–Al2O3 but 3 times higher than the one of pure Cu. Plastic deformation is thus inevitable for the latter in the sliding contact. This was confirmed by SEM revealing the presence of slip planes at the rim of a wear track on pure copper tested under sliding at a 5 N normal load for 20,000 cycles (Fig. 5). Such plastic deformation was detected neither on CA5n-450 nor on CA5n-900 tested under similar sliding conditions. Due to the fact that the initial normal load exceeds the yield strength of copper, a larger and deeper imprint on the surface of Cu is noticed. At extended sliding cycles, the surface deformation becomes more prominent leading to a rougher appearance (Fig. 6). The wear track surface features on Cu–Al2O3 were basically similar to those on Cu but much less pronounced. A SEM investigation at higher magnification of the wear track on CA5n450 obtained after 20,000 sliding cycles, before (Fig. 7a) and after (Fig. 7b) ultrasonic cleaning in ethanol of the worn surface, shows that the surface features inside these sliding tracks did not result

from a mechanical deformation of the actual surface but are cracked, delaminated tribolayers or compacted debris forming a tribolayer. Some flaked tribolayer surfaces were removed revealing the native surface with its actual roughness. The remaining parts of the tribolayer adhere well to the underlying native material. The coefficient of friction recorded on PECS processed materials in this work is reported in Fig. 8 on sliding against corundum during 100,000 sliding cycles. The steady state coefficient of friction recorded is rather equal, namely, 0.58 for Cu, 0.55 for CA5n-450 and 0.54 for CA5n-900. The running-in phase in the case of Cu–Al2O3 is limited to less than 10,000 cycles while for Cu it extends to around 35,000 cycles. Wear depths were recorded after increasing numbers of cycles (Fig. 9). The evolution of wear depth with the number of sliding cycles is quite similar in the cases of CA5n-450 and CA5n-900, however that wear depth is lesser than half the wear depth recorded on PECS Cu.

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Fig. 8. Evolution of average coefficient of friction during reciprocating sliding tests on PECS processed Cu, CA5n-450, and CA5n-900. Sliding conditions—corundum ball as counterbody, ambient air (23 1C, 50%RH), 5 N normal load, 5 Hz, 500 mm displacement and 100,000 sliding cycles.

Fig. 9. Evolution of wear depth during reciprocating sliding tests on PECS processed Cu, CA5n-450, and CA5n-900. Sliding conditions—corundum ball as counterbody, ambient air (23 1C, 50%RH), 5 N normal load, 5 Hz, and 500 mm displacement.

It is interesting to note that the wear track profiles recorded on CA5n-450 and CA5n-900 reveal a slight swelling up of the surface during the first 5,000 sliding cycles. Such a swelling up appears clearly on the wear depth (Z) profiles recorded on CA5n-450 along the sliding direction (XX’) and perpendicular to it (YY’) (Fig. 10). In order to check whether a change in surface state could be behind that swelling up of the sliding track, the nanohardness inside the sliding track obtained after 500 sliding cycles was measured at a normal load of 20 mN. The nanohardness on CA5n450 remained, inside and outside the wear track, within the standard deviation so that a notable hardening did not occur. SEM and white light interferometry observations of the Al2O3 counterbody did not show any wear on it which excludes any material transfer from ball onto wear track. A material transfer on the Al2O3 counterbody was not detected. Adhesive wear can thus be excluded as an active wear mechanism on PECS processed CA5n-450. However, EDXS analyses show that the oxygen content in the wear track increased compared to the oxygen content on unworn material (from 7 to 39 at% O). The volumetric expansion during the first 5,000 sliding cycles has thus to be linked solely to oxidation that takes place in the sliding tracks on Cu–Al2O3. It was difficult to observe such swelling on Cu since the wear depth due

to plastic deformation was larger than any potential volumetric gain due to oxidation. A debris ejected from the sliding track on CA5n-450 tested for 100,000 cycles is shown in Fig. 11a. A zoom-in (Fig. 11b) reveals that the debris consists of particles with sizes below 100 nm. That size range is characteristic of debris resulting primarily from a triboreaction in a sliding track between a material and the ambient atmosphere as reported by Quinn [33]. In ambient air, surface oxidation during sliding tests and its effects on wear of nanostructured materials need thus to be considered closely. Grain boundaries can act as high diffusion paths for oxygen, and in the case of nanostructured materials, the density of grain boundaries acting as preferential nucleation sites for oxides is higher compared to coarse grained materials [34]. During sliding, the real contact area consists of several patches bearing the load. These patches expand progressively during sliding to become plateaus of contact. The tribolayer consisting of a mixture of Cu and Cu oxide may periodically detach, and generate oxidized debris. This is illustrated in Fig. 12 showing the wear track on CA5n-450 after 20,000 sliding cycles. Smooth areas appear which can be considered as plateaus. These plateaus often contain surface cracks perpendicular to the sliding direction. These cracks result from the repetitive heat and stress cycles suffered by materials in these plateaus during reciprocating sliding tests. When the oxidized tribolayer on the Cu–Al2O3 cermets reaches a critical thickness, they may break up forming flakes (Fig. 7) and eventually form debris. On breaking up, the load bearing action is transferred to plateaus located elsewhere in the sliding area. The virgin material underneath the removed plateaus can then deform and oxidize progressively at the prevailing local surface temperature. FIB cross-sections were made inside the wear tracks on Cu and Cu–Al2O3 along the XX’ and YY’ directions in order to characterize the thickness and morphology of the tribolayer, and to detect a possible subsurface deformation. A thin platinum layer was sputtered on the surface prior to ion milling. This sample preparation protects the original sample surface in comparison to mechanical cutting which alters the outermost layer frequently. In the case of PECS Cu, several subsurface features were observed (Fig. 13) starting underneath the sputtered platinum layer with a tribolayer (compacted debris) after which a nanocrystalline layer found at zones of high deformation is followed by a layer consisting of sub-micron deformed grains that expand into the micron-sized original matrix. The formation of a subsurface gradient structure was attributed in literature to a severe plastic deformation due to friction [35]. Low angle subgrain boundaries may form as a way to decrease the energy stored in the material due to an accumulation of dislocations. The subgrains misorientation increases with increasing sliding contact events while their diameter and boundary thickness decrease leading to the formation of high angle, ultrafine/nano-grains [36]. We have found that such a refinement process can happen even at one sliding pass on coarse grained copper at the current testing conditions. FIB cross-sections in the middle of wear tracks on PECS Cu along the sliding direction (XX’) revealed that the deformed submicron layer becomes deeper, more uniform, and unidirectional after 60,000 cycles as opposed to a sub-layer divided into discrete semi-circular deformation zones observed below 30,000 sliding cycles. Fig. 14 a and b shows the different sub-surface microstructures observed. The development of such a uniform layer of severely deformed grains corresponds to the stage where a stable coefficient of friction is recorded on Cu after about 35,000 sliding cycles (see Fig.8). It corresponds also to the slope change in the evolution of wear track depth on Cu with sliding cycles after 30,000 cycles (see Fig. 9). The formation of this uniform layer

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Fig. 10. (a) Top view of a wear track recorded by white light interferometry on CA5n-450, and wear depth profiles along (b) the sliding direction (XX’), and (c) perpendicular to it (YY’) on CA5n-450 after reciprocating sliding tests in ambient air (23 1C, 50%RH) performed at 5 N, 5 Hz, 500 mm displacement after 4 different numbers of sliding cycles.

Fig. 11. (a) Debris ejected from the sliding track on CA5n-450 tested in ambient air (23 1C, 50%RH) at 5 N normal load, 5 Hz, 500 mm displacement for 100,000 cycles. (b) Zoom-in showing the compacted nature of nanometer size particles forming a debris.

indicates a steady-state in the sub-surface deformation process which determines extrinsic material properties such as friction and wear. The outermost most heavily deformed layer on PECS Cu has an extremely fine grain size (below 100 nm) and plays an important role in the wear process. As sliding progresses, this layer is further refined, becomes brittle, and generates the debris. Heilmann et al. [37] have shown through TEM observations that the debris have the same structure as the highly deformed surface layer. Rough and debris-full surface features correspond to a thicker tribolayer and a deeper deformation level as observed in our work on PECS Cu (Fig. 15). A closer look at the outermost heavily deformed layer (Fig. 16) reveals a discontinuous cracking after 20,000 sliding cycles while a fully interconnected cracking and an outlining of the tribolayer perimeter is noticed after 60,000 sliding cycles. These cracks were found down to depths of around 10 mm. This could be explained by Suh’s delamination theory of wear [38]. That theory is based on nucleation and growth of subsurface cracks in the base material after large plastic strains.

In the case of CA5n-450, even at higher magnifications, the subsurface features found in Cu are not noticed except for the tribolayer (Fig. 17) that is noticed at an early stage of sliding (10,000 cycles) as well as at extended sliding cycles (100,000 cycles). The investigation of the tribolayer on CA5n-450 by FIB revealed an oxide layer of variable thickness and morphology along the cross-section. The layer thickness varied from a few nanometers up to several hundred nanometers in a continuous or discontinuous way. There is however no sign of any deformation underneath this thin tribolayer. It seems thus that the nano-Al2O3 particles pin the nano-sized grains and prevent any visible subsurface deformation. The higher yield strength of CA5n-450 and the pinning effect of the nanoAl2O3 particles result in a deformation limited to the topmost surface leading to a lower rate of debris formation as opposed to PECS Cu. This however has limited our ability to visualize the deformation as in the case of PECS Cu. Although cracks can be seen in cross-sections, these might be compacted and delaminated debris. This seems to be in line with

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Rigney’s [39] statement that crack nucleation can be rate controlling with a crack propagation being so rapid that crosssections will not reveal any real cracks.

Fig. 12. SE image of a wear track on CA5n-450 after a reciprocating sliding test performed at 5 N, 5 Hz, and 500 mm displacement in ambient air (23 1C, 50%RH) for 20,000 sliding cycles. Cracks inside the smooth contact plateau are perpendicular to the sliding direction XX’.

Fig. 13. FIB-cross section on Cu showing the gradient structure formed under the surface of a wear track after 100,000 sliding cycles. Test conditions: 5 N, 5 Hz, 500 mm in ambient air (23 1C, 50%RH).

A FIB cross-section done at a location where both an oxidized plateau (flake) and a scratched zone are present (Fig.18), shows that the underlying virgin material is almost at the same level. The tribolayer is thicker at the flake indicating a preferential oxidation that takes place in the contact zone as expected. The presence of this flaked tribolayer of similar composition (CuþCu oxide) covering the surface of Cu and Cu–Al2O3 cermets alters the original nature of the contacting surfaces, and thus masks any direct effect of nanostructuring on friction. Since the Al2O3 content is relatively low (5 vol%) to have a major influence

Fig. 15. FIB cross-section along YY’, perpendicular to the sliding direction in the middle of a wear track on PECS Cu after 60,000 sliding cycles. A deep tribolayer is visible below the heavily deformed surface area which is full of debris. Test conditions: 5 N, 5 Hz, 500 mm in ambient air (23 1C, 50%RH).

Fig. 16. FIB cross-sections along the sliding direction (XX’) in the middle of the wear tracks on PECS Cu after sliding tests performed for (a) 20,000 cycles exhibiting a discontinuous cracking, and (b) 60,000 cycles showing continuous cracks outlining the perimeter of the tribolayer. Test conditions: 5 N, 5 Hz, 500 mm in ambient air (23 1C, 50%RH).

Fig. 14. FIB cross-sections on Cu along the sliding direction (XX’) showing (a) more discrete semi-circular subsurface deformation zones after 20,000 sliding cycles and (b) more uniform, unidirectional layer after 60,000 sliding cycles. Test conditions: 5 N, 5 Hz, 500 mm in ambient air (23 1C, 50%RH).

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Fig. 17. FIB cross sections perpendicular to the sliding direction showing the tribolayer in the middle a wear track on CA5n-450 tested at 5 N, 5 Hz, and 500 mm displacement in ambient air (23 1C, 50%RH) after (a) 10,000 and (b) 100,000 cycles. Note in (a) the crack located inside the oxide layer and not at its base.

Fig. 18. FIB cross section through both flaked off and scratched zones in the wear track on CA5n-450 tested at 5 N, 5 Hz, and 500 mm displacement in ambient air (23 1C, 50%RH) for 20,000 cycles. Note the thicker tribolayer at the flaked off zone.

Fig. 19. (a) SEM image of the wear track on CA5n-450 after sliding in ambient air (23 1C, 50%RH) against alumina at 15 N normal load, 5 Hz and 500 mm displacement for 20,000 cycles. (b) FIB cross-section showing no sign of deformation aside from the surface tribolayer.

on the coefficient of friction, this tribolayer can be at the origin of the almost equal values of the coefficient of friction recorded on the three tested materials after the development of a steady-state subsurface microstructure. The effect of contact pressure on the subsurface microstructure of CA5-450 was investigated by performing a reciprocating sliding test for 20,000 cycles at 15 N normal load corresponding to a Hertzian pressure of 1094 MPa. This resulted in a wear depth of 2 mm compared to 1 mm at 5 N. On the corresponding wear track and FIB cross section shown in Fig. 19, a grain refinement or grain growth was not noticed in the subsurface material. The microstructural stability of the nano-Al2O3 reinforced nanostructured Cu matrix is thus the main reason behind the decreased wear noticed on CA5n-450 compared to pure Cu. The subsurface refinement in wear tracks performed on PECS Cu is expected to induce hardening in the wear track due to a decrease in grain size. Nanoindentation measurements across the wear track (YY’ direction) on Cu after 20,000 sliding cycles indeed show that the nanohardness value doubled inside the wear track as opposed to the native material (Fig. 20). Similar measurements done on a similar wear track on CA5n-450 revealed no conclusive data. Optical microscopy revealed that the lowest hardness values

Fig. 20. Nanohardness obtained along the YY’-direction on wear tracks on Cu and CA5n-450 after 20,000 sliding cycles. Test conditions: 5 N, 5 Hz, and 500 mm displacement in ambient air (23 1C, 50%RH).

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are recorded on pure Cu inter-powder regions. As previously mentioned, it is possible that the stability of the microstructure prevented any detectable subsurface refinement.

below the tribolayer difficult to happen. This was responsible for a wear depth lesser than half that of pure copper. The thermal stability of the structure of Cu–Al2O3 cermets after exposure to 600 1C for 20 h was confirmed.

3.4. Thermal stability Finally, the thermal stability of CA5n-450 was checked by performing a heat treatment in an oven at 600 1C for 20 h in ambient air. The sample surface was then polished by 0.25 mm colloidal silica to remove the oxidized surface layer. The hardness measured as a 5-measurement averaged value at a load of 300 g was 243 HV300 showing a drop of 21% compared to the 307 HV300 recorded on non-heat treated sample. The average grain size measured by EBSD was 250 nm showing an increase by 36% compared to the starting size of 160 nm. The PECS sintered Cubased cermets retained thus a relatively good hardness even after thermal exposure at 600 1C. Nanostructured Cu–Al2O3 cermets show thus an improved wear resistance and thermal stability compared to coarse grained copper. They can thus be considered as good candidates for applications where low wear and good thermal stability is required.

4. Conclusions The original nanostructure of the composite Cu–Al2O3 (5 vol%) powders was maintained with minimal grain growth during PECS sintering at 450 1C and 900 1C for 7 min. The presence of nanoAl2O3 pins the grain boundaries, limits grain growth during thermal processing and contributes to the total strength of the material through a dispersion strengthening mechanism. In addition to that, the residual plastic strain due to particles-matrix post-processing thermal mismatch contributed to the total strength of the cermet. High dislocation density induced by ball milling which was not fully dissipated after thermal processing at 450 1C contributed as well to the strength of the Cu–Al2O3 composite. No major difference was noticed in the steady state coefficient of friction of copper and nanostructured Cu–Al2O3 cermets. Its value is affected by the tribolayer formed in the contact area between the virgin surface and the counterbody what complicates the identification of an effect of nanostructuring on friction. Two processes dictate the wear behavior of the PECS sintered Cu and Cu–Al2O3 materials. On one hand, there is surface oxidation that occurs during sliding. This was revealed by the swellingup of the surface during the early stages of sliding in the case of Cu–Al2O3. The normal pressure applied was below the yield strength of the material what excludes any plastic deformation. Such a swelling up was difficult to observe on Cu due to plastic deformation which was larger than any volumetric gain due to oxidation. On the other hand, there is the evolution of the subsurface microstructure. That evolution was easily visible in the case of pure Cu due to the low yield strength. The grain refinement due to a friction induced severe plastic deformation is visible even at a depth of few tens of micrometers. The sub-micron layer evolves from discrete zones to a more uniform and unidirectional layer at increasing number of sliding cycles. This explains the stability of the coefficient of friction as well as the slope change of the wear track depth recorded on Cu after sliding for more than 30,000 sliding cycles. However, such a grain refinement was hardly observed on nano-Al2O3 reinforced nanostructured Cu. Their microstructure is much more stable and the strength of the material higher. As a result, the deformation is limited to the top few nanometers sub-surface zone making any deformation

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