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Enhancing wear resistance of Cu–Al alloy by controlling subsurfacedynamic recrystallization

X. Chen, Z. Han⇑

and K. Lu⇑

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences,

72 Wenhua Road, Shenyang 110016, China

Received 27 November 2014; revised 21 January 2015; accepted 21 January 2015Available online 11 February 2015

The sliding wear resistance of Cu-2.2 wt.% Al alloy can be remarkably enhanced by controlling subsurface dynamic recrystallization (DRX) pro-cess. The wear volume decreases from 3.1 � 107 to 0.9 � 107 lm3 when the DRX grain size increases from 0.28 to 0.62 lm induced by dynamic plasticdeformation and subsequent annealing. The enhanced wear resistance stems from a wear mechanism transition from cracking and peeling-off of thefine-grained DRX layer to that of the topmost nanostructured mixing layer as DRX grains become coarser.� 2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Copper alloys; Nanostructure; Worn subsurface microstructure; Dynamic recrystallization; Wear

Sliding of metals generally results in a sophisticatedmicrostructural evolution beneath the worn surface [1–3].The formation of nanostructured tribolayer was alwaysreported in sliding of metals [4], and atomistic level mixingwas observed at the interfaces in hard-soft tribo-pairs byatomistic simulations [5]. Among various sliding-inducedprocesses, dynamic recrystallization (DRX) of thedeformed structure, frequently observed in the worn sub-surface layer of copper [6,7], aluminum, and their alloys[8], is found to be crucial in determining the wear behaviorof the metals. For instance, in pure copper a subsurface lay-er of DRX structure is generated beneath the topmostnanostructured mixing layer (NML) during dry sliding[9]. A pronounced correlation is identified that the wearresistance increases monotonically with a decreasingDRX grain size.

By adding Al (0.1–2.2 wt.%) into Cu, the recrystalliza-tion kinetics can be adjusted so that the recrystallized grainsizes are changed, which corresponds to an obvious changein the wear resistance. A wear mechanism transition hasbeen found when the recrystallized grain sizes are below acritical value [10]. As the recrystallized grain sizes are largerthan 0.7 lm, the wear volume decreases monotonically witha decreasing grain size and the wear mechanism is dominat-ed by peeling-off of the NML. While for the recrystallizedgrains smaller than 0.7 lm, cracking and peeling-off ofthe DRX layer becomes a dominating mechanism,corresponding to increased wear volume. Due to such a

http://dx.doi.org/10.1016/j.scriptamat.2015.01.0231359-6462/� 2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights

⇑Corresponding authors. Tel.: +86 24 23971891; fax: +86 2423971215; e-mail addresses: zhonghan@imr.ac.cn; lu@imr.ac.cn

transition, the Cu-2.2 wt.% Al alloy with very fine recrystal-lized grains exhibits a rather poor wear resistance comparedto that of pure copper, in consistent with the other investi-gations of Cu–Al alloys [11–13].

The objective of the present study is to explore if thewear resistance of a Cu–Al alloy with a fixed compositioncan be enhanced by controlling the subsurface recrystalliza-tion process via pre-treatments such as plastic deformationand/or heat treatment.

A solid solution Cu-2.2 wt.% Al alloy was preparedfrom 99.999 wt.% Cu and 99.995 wt.% Al. A homogeneousmicrostructure with coarse grains (CG, �200 lm in size)was obtained, which was subjected to dynamic plasticdeformation (DPD) with a strain about 2.0 at liquid nitro-gen temperature. The setup and processing parameters ofthe DPD facility have been described in detail elsewhere[14]. The deformation strain is defined as e = ln (L0/Lf),where L0 and Lf are the initial and the final thickness ofthe deformed sample, respectively. Microstructure of theas-DPD Cu–Al alloy is similar to that in the DPD Cu[14], consisting of about 25 vol.% nano-scale twin/matrix(T/M) lamellae and 75 vol.% nano-sized grains with anaverage size of about 30 nm [15]. The as-DPD samples wereannealed at 100–460 �C for 10 min for modifyingmicrostructure via recovery and/or recrystallization.

Sliding wear tests of the Cu–Al samples were performedon an Optimal SRVIII oscillating friction and wear tester ina ball-on-plate contact configuration under dry conditionat room temperature (25 �C) in air with a relative humidityof 45%. Plates were cut from the Cu–Al specimens to adimension of 6 � 6 � 3 mm3 with an electro-polished sur-

reserved.

X. Chen et al. / Scripta Materialia 101 (2015) 76–79 77

face. WC-Co balls of 10 mm in diameter with a micro-hard-ness of 17.5 GPa were used as the counterface. The frictionand wear tests were carried out at a sliding amplitude of500 lm, a normal load of 30 N, a frequency of 5 Hz, anda duration of 60 min. Profiles of the worn surfaces weremeasured by using a MicroXAM 3 dimensional (3D) sur-face profilometer system so as to determine the wearvolume.

Before the worn subsurface structure characterization,the wear scars were ultrasonically cleaned with acetoneand alcohol respectively, on which a protective layer ofcopper was electrodeposited. Cross-sectional samples werecut parallel to the sliding direction and grinded to the cen-ter of the wear scars. The worn subsurface structure of theCu–Al specimens was characterized by using scanning elec-tron microscopy (SEM) in electron channeling contrast(ECC) mode with a thermal field emission gun on a FEINano-SEM Nova 430 system operated at 15 kV. Detailedmicrostructure characterization was performed on a TecnaiG2 F-20 transmission electron microscope (TEM) operatedat a voltage of 200 kV on a sectional wear scar. Thin foilsfor TEM observations were prepared on a Nova 200 Nano-lab focused ion beam (FIB) system. A thin layer of plat-inum was deposited on the worn surface to protect thesurface features against beam damage.

The measured wear volume of Cu-2.2 wt.% Al alloy is3.1 � 107 lm3 after sliding against WC-Co ball for 60 minunder a normal load of 30 N, which is higher than that ofcopper sample (1.8 � 107 lm3) [9]. After the DPD process-ing, the wear volume of the Cu–Al sample decreases toabout 0.9 � 107 lm3 under the same sliding condition.The measured steady-state friction coefficient of the DPDCu–Al sample is about 0.7, comparable to that of theCu–Al sample.

For both the CG and the as-DPD Cu–Al samples, thesteady state subsurface structures are formed under theworn surface during sliding. Extremely fine structure isfound in the subsurface layer of the CG sample (Fig. 1a).TEM observations show a thin and continuous nanostruc-tured mixing layer, consisting of nano-sized grains with

a

20 µm

c

b

NML

DRX

NML

Figure 1. SEM-ECC images of the cross-sectional worn subsurface structurNMLs for the CG (a and d) and the DPD (c and e) Cu-2.2% Al alloy after simage of the NML with the DRX layer for the Cu-2.2% Al alloy. The worn syellow dashed lines. (For interpretation of the references to color in this figu

random crystallographic orientations, is developed on thetop of the DRX layer with extremely fine structure(Fig. 1b). The detailed microstructures in the NML andDRX layer were investigated in previous work [10]. Forthe as-DPD sample, a subsurface layer of obvious DRXstructure is generated beneath the topmost NML, as shownin Figure 1c. There is a sharp boundary between the NMLand the DRX layer (as outlined by dashed lines). Sliding onthe as-DPD sample leads to the formation of much largerrecrystallized grains compared to that in the CGcounterpart.

Cracking and peeling-off in the DRX layer wereobserved (Fig. 1a), which is the dominating wear processfor the Cu-2.2% Al sample [10]. While for the as-DPD sam-ple, micro-cracks are found either inside the NMLs or atthe NML/DRX interfaces, similar to that observed in Cu[9]. Propagation and coalescence of these cracks result infracture and peeling-off of the topmost NML consequently,which is quite different from the material removal mechan-ism exists in the Cu–Al sample.

Careful examinations reveal an obvious difference in theDRX grain size between the two samples (see Fig. 1d ande): the average size is about 0.28 and 0.62 lm for the CGand the as-DPD sample, respectively. It seems that theenhanced wear resistance of the as-DPD sample is attribut-ed to larger recrystallization grain size. To verify recrystal-lization grain size effect on the wear resistance, different Cu-2.2% Al samples were prepared by controlling recrystallizedkinetics via DPD and subsequent heat treatment.

Microstructural evolution of the DPD samples duringthermal annealing has been well-studied in copper [16]and in other materials [17]. Annealing induced static recrys-tallization (SRX) process of the DPD Cu-2.2% Al sample(above 260 �C), analogous to that in Cu [16], occurs initial-ly in the nano-grained regions due to their higher storedenergy compared with that of the nano-twined structure.The partially recrystallized samples are structurally charac-terized by a mixed structure of micro-sizes SRX grains,nano-sized grains and nano-twin bundles. Microhardnessdrops sharply when the annealing temperature (Ta) exceeds

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e and corresponding distributions of DRX grain size underneath theliding against a WC-Co ball at a load of 30 N. (b) A bright field TEMurfaces are outlined by red dashed lines and the NMLs are outlined byre legend, the reader is referred to the web version of this article.)

Figure 2. Variations of hardness (a) and corresponding wear volume(b) for the CG, the DPD and the annealed DPD Cu–Al alloy atdifferent temperatures. The inserts in (a) show the SEM-ECC images ofthe structure in the DPD samples annealed at 280 �C (left) and 460 �C(right).

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Figure 3. SEM micrographs of worn subsurface structures andcorresponding distributions of DRX grain size for DPD samplesannealed at (a and b) 160 �C; (c and d) 340 �C; (e and f) 460 �C for10 min.

78 X. Chen et al. / Scripta Materialia 101 (2015) 76–79

280 �C due to the occurrence of more obvious SRX grains(constituting �50% in volume, left insert in Fig. 2a), asdepicted in Figure 2a. The mean SRX grain size is about2.5 lm. The volume of fraction of SRX grains increasesat higher annealing temperatures. Hardness decreases from2.3 GPa of the as-DPD sample to 0.9 GPa for the sampleannealed at 460 �C, which corresponds to an increasing vol-ume fraction of SRX grains to �98% (right insert inFig. 2a). The mean SRX grain size is about 3 lm. However,the corresponding wear volumes of the annealed samplesdo not follow the hardness variation linearly, as expressedin Archard’s law (see Fig. 2b). For instance, there existsno obvious decrease in the wear volume when Ta is below340 �C, although the hardness decreases to a lower levelof 1.2 GPa when annealing at 340 �C. But the wear volumeincreases from 1.0 � 107 to 1.8 � 107 lm3 as Ta increasesfrom 340 to 460 �C, accompanied by a further decrease ofhardness from 1.2 to 0.9 GPa. Obviously, initialmicrostructure and hardness may not have a straightfor-ward effect on the wear resistance of the samples.

For the DPD samples annealed at 160 and 340 �C(Fig. 3a and c), similar DRX structures are formed beneaththe top NML. Micro-cracks are frequently observed insidethe NMLs and/or at the NML/DRX interfaces, analogousto that in the as-DPD sample. However, for the DPD sam-ple annealed at 460 �C, the subsurface DRX grains aremuch refined (Fig. 3e) and cracks are often found in theDRX layer, as observed in the Cu–Al sample. Quantitativemeasurements show that the grain size distribution of thesample annealed at 160 �C is identical to that of the sample

annealed at 340 �C (see Fig. 3b and d). Back to the wearvolume data (Fig. 2b), no obvious difference exists betweentwo samples despite of a remarkably difference in the initialhardness. However, the average grain size of the sampleannealed at 460 �C is 0.42 lm (Fig. 3f), which is obviouslysmaller than that of the samples annealed at 160 �C(0.58 lm) and 340 �C (0.6 lm). This is consistent with ourprevious observation [10] that cracking and peeling-off ofthe DRX layer is the dominating process when DRX grainsize is less than 0.5 lm.

Figure 4 plots a relationship between the wear volumeand the average DRX grain size for the Cu-2.2% Al sam-ples with different initial microstructures. The wear volumedecreases significantly with an increasing grain size from0.28 to 0.62 lm. It indicates that the wear resistance canbe elevated when the recrystallized process is enhancedfor the Cu-2.2% Al alloy. Combined with the data for theCu–Al alloys with different Al concentrations [10], thesedata follow two different variation tendencies, dependingon the DRX grain size. The wear volume decreases mono-tonically with a decreasing DRX grain size from 0.92 to0.7 lm, and increases with a further decrease from 0.7 to0.28 lm.

The present work reveals that the recrystallized kineticsduring sliding is controlled by adjusting the initialmicrostructure of Cu-2.2% Al sample. DRX grain size isdetermined by the stored energy and grain boundary mobi-lity during the DRX process [18]. For the Cu-2.2% Al sam-ple, extremely small DRX grains are derived from a muchreduced boundary mobility due to alloying with Al, com-pared with that in Cu. However, after the DPD treatment,the driving force for dynamic recrystallization during slid-ing is enlarged owing to very high stored energy in thedeformed matrix [19]. The DRX kinetics is enhanced by

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Figure 4. Correlation of the wear volume with the average DRX grainsize for a series of Cu–Al alloys [10], the DPD and the annealed DPDCu-2.2% Al alloy at different temperatures.

X. Chen et al. / Scripta Materialia 101 (2015) 76–79 79

the high-energy state of the plastically deformed sample inwhich there exist nano-sized grains with a very largepopulation of grain boundaries and other defects [20]. Asthe samples were annealed at higher temperatures, thestored energy is released so that the DRX grain sizedecreases during sliding.

There exists a wear mechanism transition for the Cu-2.2% Al alloy accompanied with the DPD and subsequentannealing. For the Cu–Al alloy, cracking and peeling-off ofthe brittle DRX layer is the dominating wear process.While for the as-DPD and the annealed DPD samples atlower temperatures, peeling-off of the topmost NML dom-inates the material removal process. The latter mechanismcorresponds to higher wear resistance. For the Cu-2.2%Al alloy, the wear resistance increases by nearly two times,with an increasing DRX grain size. When the recrystallizedgrains are much refined, it is difficult to accommodate plas-tic deformation during sliding, resulting in cracking withinthe recrystallized layer. Once the vortical NML is formedwhen the recrystallized grains are larger enough, slidingmay induce cracking and peeling-off of the NML, accom-panied with energy-consuming process of transformationfrom the DRX structure to the NML. For the Cu–Al alloys[10], with an increasing Al concentration, there exists a cri-tical DRX grain size of 0.7 lm determining the minimumwear volume when Al concentration is 0.5 wt.%. TheDPD processing seems effective in improving the wear resis-tance of Cu–Al alloy by adjusting the recrystallizedstructure.

It is interesting that a lower wear rate of the DPD Cu–Alalloy can be achieved by increasing DRX grain size, which

corresponds to the subsurface softening process. Althoughthe DPD Cu–Al alloy exhibits very limited uniform tensileelongation (�1–2%) [15], the recrystallized structure offersplenty rooms to release accumulated plastic strain withoutcracking in the recrystallized layer during sliding.

In summary, the wear resistance of Cu-2.2 wt.% Al alloycan be remarkably enhanced by increasing DRX grain sizein the worn subsurface layer. It is attributed to a wearmechanism transition from cracking and peeling-off of thebrittle fine-grained DRX layer to peeling-off of the topmostNML. The present finding provides an effective approachto elevate the wear resistance of the Cu–Al alloy by control-ling subsurface DRX process.

The authors are grateful for the financial supports fromthe National Natural Science Foundation of China (Grant51231006), the MOST 973 Project (Grant 2012CB932201), theDanish-Chinese Center for Nano-metals (Grant 51261130091),the Key Research Program of Chinese Academy of Sciences(Grant KGZD-EW-T06) and Shenyang National Laboratory forMaterials Science.

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