- Email: [email protected]
Enhancing wear resistance of Cu–Al alloy by controlling subsurface dynamic recrystallization
Available online at www.sciencedirect.com
ScienceDirect Scripta Materialia 101 (2015) 76–79 www.elsevier.com/locate/scriptamat
Enhancing wear resistance of Cu–Al alloy by controlling subsurface dynamic 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 2015 Available 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) process. 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 plastic deformation and subsequent annealing. The enhanced wear resistance stems from a wear mechanism transition from cracking and peeling-off of the fine-grained DRX layer to that of the topmost nanostructured mixing layer as DRX grains become coarser. Ó 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 sophisticated microstructural evolution beneath the worn surface [1–3]. The formation of nanostructured tribolayer was always reported in sliding of metals [4], and atomistic level mixing was observed at the interfaces in hard-soft tribo-pairs by atomistic simulations [5]. Among various sliding-induced processes, dynamic recrystallization (DRX) of the deformed structure, frequently observed in the worn subsurface layer of copper [6,7], aluminum, and their alloys [8], is found to be crucial in determining the wear behavior of the metals. For instance, in pure copper a subsurface layer of DRX structure is generated beneath the topmost nanostructured mixing layer (NML) during dry sliding [9]. A pronounced correlation is identified that the wear resistance increases monotonically with a decreasing DRX grain size. By adding Al (0.1–2.2 wt.%) into Cu, the recrystallization kinetics can be adjusted so that the recrystallized grain sizes are changed, which corresponds to an obvious change in the wear resistance. A wear mechanism transition has been found when the recrystallized grain sizes are below a critical value [10]. As the recrystallized grain sizes are larger than 0.7 lm, the wear volume decreases monotonically with a decreasing grain size and the wear mechanism is dominated by peeling-off of the NML. While for the recrystallized grains smaller than 0.7 lm, cracking and peeling-off of the DRX layer becomes a dominating mechanism, corresponding to increased wear volume. Due to such a
⇑ Corresponding
authors. Tel.: +86 24 23971891; fax: +86 24 23971215; e-mail addresses: [email protected]; [email protected]
transition, the Cu-2.2 wt.% Al alloy with very fine recrystallized grains exhibits a rather poor wear resistance compared to that of pure copper, in consistent with the other investigations of Cu–Al alloys [11–13]. The objective of the present study is to explore if the wear resistance of a Cu–Al alloy with a fixed composition can be enhanced by controlling the subsurface recrystallization process via pre-treatments such as plastic deformation and/or heat treatment. A solid solution Cu-2.2 wt.% Al alloy was prepared from 99.999 wt.% Cu and 99.995 wt.% Al. A homogeneous microstructure with coarse grains (CG, 200 lm in size) was obtained, which was subjected to dynamic plastic deformation (DPD) with a strain about 2.0 at liquid nitrogen temperature. The setup and processing parameters of the 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 of the deformed sample, respectively. Microstructure of the as-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 an average size of about 30 nm [15]. The as-DPD samples were annealed at 100–460 °C for 10 min for modifying microstructure via recovery and/or recrystallization. Sliding wear tests of the Cu–Al samples were performed on an Optimal SRVIII oscillating friction and wear tester in a ball-on-plate contact configuration under dry condition at room temperature (25 °C) in air with a relative humidity of 45%. Plates were cut from the Cu–Al specimens to a dimension of 6 6 3 mm3 with an electro-polished sur-
http://dx.doi.org/10.1016/j.scriptamat.2015.01.023 1359-6462/Ó 2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
X. Chen et al. / Scripta Materialia 101 (2015) 76–79
face. WC-Co balls of 10 mm in diameter with a micro-hardness of 17.5 GPa were used as the counterface. The friction and wear tests were carried out at a sliding amplitude of 500 lm, a normal load of 30 N, a frequency of 5 Hz, and a duration of 60 min. Profiles of the worn surfaces were measured by using a MicroXAM 3 dimensional (3D) surface profilometer system so as to determine the wear volume. Before the worn subsurface structure characterization, the wear scars were ultrasonically cleaned with acetone and alcohol respectively, on which a protective layer of copper was electrodeposited. Cross-sectional samples were cut parallel to the sliding direction and grinded to the center of the wear scars. The worn subsurface structure of the Cu–Al specimens was characterized by using scanning electron microscopy (SEM) in electron channeling contrast (ECC) mode with a thermal field emission gun on a FEI Nano-SEM Nova 430 system operated at 15 kV. Detailed microstructure characterization was performed on a Tecnai G2 F-20 transmission electron microscope (TEM) operated at a voltage of 200 kV on a sectional wear scar. Thin foils for TEM observations were prepared on a Nova 200 Nanolab focused ion beam (FIB) system. A thin layer of platinum was deposited on the worn surface to protect the surface features against beam damage. The measured wear volume of Cu-2.2 wt.% Al alloy is 3.1 107 lm3 after sliding against WC-Co ball for 60 min under a normal load of 30 N, which is higher than that of copper sample (1.8 107 lm3) [9]. After the DPD processing, the wear volume of the Cu–Al sample decreases to about 0.9 107 lm3 under the same sliding condition. The measured steady-state friction coefficient of the DPD Cu–Al sample is about 0.7, comparable to that of the Cu–Al sample. For both the CG and the as-DPD Cu–Al samples, the steady state subsurface structures are formed under the worn surface during sliding. Extremely fine structure is found in the subsurface layer of the CG sample (Fig. 1a). TEM observations show a thin and continuous nanostructured mixing layer, consisting of nano-sized grains with
random crystallographic orientations, is developed on the top of the DRX layer with extremely fine structure (Fig. 1b). The detailed microstructures in the NML and DRX layer were investigated in previous work [10]. For the as-DPD sample, a subsurface layer of obvious DRX structure is generated beneath the topmost NML, as shown in Figure 1c. There is a sharp boundary between the NML and the DRX layer (as outlined by dashed lines). Sliding on the as-DPD sample leads to the formation of much larger recrystallized grains compared to that in the CG counterpart. Cracking and peeling-off in the DRX layer were observed (Fig. 1a), which is the dominating wear process for the Cu-2.2% Al sample [10]. While for the as-DPD sample, micro-cracks are found either inside the NMLs or at the NML/DRX interfaces, similar to that observed in Cu [9]. Propagation and coalescence of these cracks result in fracture and peeling-off of the topmost NML consequently, which is quite different from the material removal mechanism exists in the Cu–Al sample. Careful examinations reveal an obvious difference in the DRX grain size between the two samples (see Fig. 1d and e): the average size is about 0.28 and 0.62 lm for the CG and the as-DPD sample, respectively. It seems that the enhanced wear resistance of the as-DPD sample is attributed to larger recrystallization grain size. To verify recrystallization grain size effect on the wear resistance, different Cu2.2% Al samples were prepared by controlling recrystallized kinetics via DPD and subsequent heat treatment. Microstructural evolution of the DPD samples during thermal annealing has been well-studied in copper [16] and in other materials [17]. Annealing induced static recrystallization (SRX) process of the DPD Cu-2.2% Al sample (above 260 °C), analogous to that in Cu [16], occurs initially in the nano-grained regions due to their higher stored energy compared with that of the nano-twined structure. The partially recrystallized samples are structurally characterized by a mixed structure of micro-sizes SRX grains, nano-sized grains and nano-twin bundles. Microhardness drops sharply when the annealing temperature (Ta) exceeds
b
NML
b
0.20
d
0.16
Volume Fraction
a
77
DRX
0.12 0.08 0.04 0.00 0.0
500 nm
0.2
0.4
0.6
0.8
1.0
1.2
DRX grain size
c
0.20
NML
20 µm
Volume Fraction
DRX
e
0.16 0.12 0.08 0.04 0.00 0.0
0.2
0.4
0.6
0.8
1.0
1.2
DRX grain size
Figure 1. SEM-ECC images of the cross-sectional worn subsurface structure and corresponding distributions of DRX grain size underneath the NMLs for the CG (a and d) and the DPD (c and e) Cu-2.2% Al alloy after sliding against a WC-Co ball at a load of 30 N. (b) A bright field TEM image of the NML with the DRX layer for the Cu-2.2% Al alloy. The worn surfaces are outlined by red dashed lines and the NMLs are outlined by yellow dashed lines. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
78
X. Chen et al. / Scripta Materialia 101 (2015) 76–79
a
0.24
b
0.16 0.08
Volume Fraction
0.00
c
d
0.20 0.12 0.04 0.24
f
0.16
e
0.08 0.00
10 m
0.2
0.4
0.6
0.8
1.0
DRX grain size (µm)
Figure 3. SEM micrographs of worn subsurface structures and corresponding distributions of DRX grain size for DPD samples annealed at (a and b) 160 °C; (c and d) 340 °C; (e and f) 460 °C for 10 min. Figure 2. Variations of hardness (a) and corresponding wear volume (b) for the CG, the DPD and the annealed DPD Cu–Al alloy at different temperatures. The inserts in (a) show the SEM-ECC images of the structure in the DPD samples annealed at 280 °C (left) and 460 °C (right).
280 °C due to the occurrence of more obvious SRX grains (constituting 50% in volume, left insert in Fig. 2a), as depicted in Figure 2a. The mean SRX grain size is about 2.5 lm. The volume of fraction of SRX grains increases at higher annealing temperatures. Hardness decreases from 2.3 GPa of the as-DPD sample to 0.9 GPa for the sample annealed at 460 °C, which corresponds to an increasing volume fraction of SRX grains to 98% (right insert in Fig. 2a). The mean SRX grain size is about 3 lm. However, the corresponding wear volumes of the annealed samples do not follow the hardness variation linearly, as expressed in Archard’s law (see Fig. 2b). For instance, there exists no obvious decrease in the wear volume when Ta is below 340 °C, although the hardness decreases to a lower level of 1.2 GPa when annealing at 340 °C. But the wear volume increases from 1.0 107 to 1.8 107 lm3 as Ta increases from 340 to 460 °C, accompanied by a further decrease of hardness from 1.2 to 0.9 GPa. Obviously, initial microstructure and hardness may not have a straightforward 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 beneath the top NML. Micro-cracks are frequently observed inside the NMLs and/or at the NML/DRX interfaces, analogous to that in the as-DPD sample. However, for the DPD sample annealed at 460 °C, the subsurface DRX grains are much refined (Fig. 3e) and cracks are often found in the DRX layer, as observed in the Cu–Al sample. Quantitative measurements show that the grain size distribution of the sample annealed at 160 °C is identical to that of the sample
annealed at 340 °C (see Fig. 3b and d). Back to the wear volume data (Fig. 2b), no obvious difference exists between two samples despite of a remarkably difference in the initial hardness. However, the average grain size of the sample annealed at 460 °C is 0.42 lm (Fig. 3f), which is obviously smaller than that of the samples annealed at 160 °C (0.58 lm) and 340 °C (0.6 lm). This is consistent with our previous observation [10] that cracking and peeling-off of the DRX layer is the dominating process when DRX grain size is less than 0.5 lm. Figure 4 plots a relationship between the wear volume and the average DRX grain size for the Cu-2.2% Al samples with different initial microstructures. The wear volume decreases significantly with an increasing grain size from 0.28 to 0.62 lm. It indicates that the wear resistance can be elevated when the recrystallized process is enhanced for the Cu-2.2% Al alloy. Combined with the data for the Cu–Al alloys with different Al concentrations [10], these data follow two different variation tendencies, depending on the DRX grain size. The wear volume decreases monotonically with a decreasing DRX grain size from 0.92 to 0.7 lm, and increases with a further decrease from 0.7 to 0.28 lm. The present work reveals that the recrystallized kinetics during sliding is controlled by adjusting the initial microstructure of Cu-2.2% Al sample. DRX grain size is determined by the stored energy and grain boundary mobility during the DRX process [18]. For the Cu-2.2% Al sample, extremely small DRX grains are derived from a much reduced boundary mobility due to alloying with Al, compared with that in Cu. However, after the DPD treatment, the driving force for dynamic recrystallization during sliding is enlarged owing to very high stored energy in the deformed matrix [19]. The DRX kinetics is enhanced by
X. Chen et al. / Scripta Materialia 101 (2015) 76–79 4.0
Wear volume (107µm3)
3.5
corresponds to the subsurface softening process. Although the DPD Cu–Al alloy exhibits very limited uniform tensile elongation (1–2%) [15], the recrystallized structure offers plenty rooms to release accumulated plastic strain without cracking in the recrystallized layer during sliding. In summary, the wear resistance of Cu-2.2 wt.% Al alloy can be remarkably enhanced by increasing DRX grain size in the worn subsurface layer. It is attributed to a wear mechanism transition from cracking and peeling-off of the brittle fine-grained DRX layer to peeling-off of the topmost NML. The present finding provides an effective approach to elevate the wear resistance of the Cu–Al alloy by controlling subsurface DRX process.
CG Cu-Al alloys [10] Cu-2.2 wt.% Al alloy
3.0 2.5 2.0 1.5 1.0 0.5
0.2
0.4
0.6
0.8
1.0
79
1.2
DRX grain size (µm) Figure 4. Correlation of the wear volume with the average DRX grain size for a series of Cu–Al alloys [10], the DPD and the annealed DPD Cu-2.2% Al alloy at different temperatures.
the high-energy state of the plastically deformed sample in which there exist nano-sized grains with a very large population of grain boundaries and other defects [20]. As the samples were annealed at higher temperatures, the stored energy is released so that the DRX grain size decreases during sliding. There exists a wear mechanism transition for the Cu2.2% Al alloy accompanied with the DPD and subsequent annealing. For the Cu–Al alloy, cracking and peeling-off of the brittle DRX layer is the dominating wear process. While for the as-DPD and the annealed DPD samples at lower temperatures, peeling-off of the topmost NML dominates the material removal process. The latter mechanism corresponds 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 recrystallized grains are much refined, it is difficult to accommodate plastic deformation during sliding, resulting in cracking within the recrystallized layer. Once the vortical NML is formed when the recrystallized grains are larger enough, sliding may induce cracking and peeling-off of the NML, accompanied with energy-consuming process of transformation from the DRX structure to the NML. For the Cu–Al alloys [10], with an increasing Al concentration, there exists a critical DRX grain size of 0.7 lm determining the minimum wear volume when Al concentration is 0.5 wt.%. The DPD processing seems effective in improving the wear resistance of Cu–Al alloy by adjusting the recrystallized structure. It is interesting that a lower wear rate of the DPD Cu–Al alloy can be achieved by increasing DRX grain size, which
The authors are grateful for the financial supports from the National Natural Science Foundation of China (Grant 51231006), the MOST 973 Project (Grant 2012CB932201), the Danish-Chinese Center for Nano-metals (Grant 51261130091), the Key Research Program of Chinese Academy of Sciences (Grant KGZD-EW-T06) and Shenyang National Laboratory for Materials Science.
[1] D.A. Rigney, X.Y. Fu, J.E. Hammerberg, B.L. Holian, M.L. Falk, Scripta Mater. 49 (2003) 977–983. [2] A. Mahato, N. Verma, V. Jayaram, S.K. Biswas, Acta Mater. 59 (2011) 6069–6082. [3] F. Ren, S.N. Arshad, P. Bellon, R.S. Averback, M. Pouryazdan, H. Hahn, Acta Mater. 72 (2014) 148–158. [4] D. Rigney, S. Karthikeyan, Tribol. Lett. 39 (2010) 3–7. [5] S. Karthikeyan, A. Agrawal, D.A. Rigney, Wear 267 (2009) 1166–1176. [6] J.H. Dautzenberg, Wear 60 (1980) 401–411. [7] A. Emge, S. Karthikeyan, D.A. Rigney, Wear 267 (2009) 562–567. [8] S.M. Kuo, D.A. Rigney, Mater. Sci. Eng., A 157 (1992) 131– 143. [9] B. Yao, Z. Han, K. Lu, Wear 294–295 (2012) 438–445. [10] X. Chen, Z. Han, K. Lu, Wear 320 (2014) 41–50. [11] R.A. Poggie, J.J. Wert, Wear 156 (1992) 315–326. [12] J.J. Wert, W.M. Cook, Wear 123 (1988) 171–192. [13] J.J. Wert, S.A. Singerman, S.G. Caldwell, D.K. Chaudhuri, Wear 92 (1983) 213–229. [14] Y.S. Li, N.R. Tao, K. Lu, Acta Mater. 56 (2008) 230–241. [15] Y. Zhang, N.R. Tao, K. Lu, Acta Mater. 59 (2011) 6048– 6058. [16] Y.S. Li, Y. Zhang, N.R. Tao, K. Lu, Scripta Mater. 59 (2008) 475–478. [17] F.K. Yan, G.Z. Liu, N.R. Tao, K. Lu, Acta Mater. 60 (2012) 1059–1071. [18] C.C. Koch, R.O. Scattergood, K.A. Darling, J.E. Semones, J. Mater. Sci. 43 (2008) 7264–7272. [19] F. Yan, H.W. Zhang, J. Mater. Sci. Technol. 28 (2012) 289– 293. [20] L. Yang, N.R. Tao, K. Lu, L. Lu, Scripta Mater. 68 (2013) 801–804.
ScienceDirect Scripta Materialia 101 (2015) 76–79 www.elsevier.com/locate/scriptamat
Enhancing wear resistance of Cu–Al alloy by controlling subsurface dynamic 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 2015 Available 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) process. 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 plastic deformation and subsequent annealing. The enhanced wear resistance stems from a wear mechanism transition from cracking and peeling-off of the fine-grained DRX layer to that of the topmost nanostructured mixing layer as DRX grains become coarser. Ó 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 sophisticated microstructural evolution beneath the worn surface [1–3]. The formation of nanostructured tribolayer was always reported in sliding of metals [4], and atomistic level mixing was observed at the interfaces in hard-soft tribo-pairs by atomistic simulations [5]. Among various sliding-induced processes, dynamic recrystallization (DRX) of the deformed structure, frequently observed in the worn subsurface layer of copper [6,7], aluminum, and their alloys [8], is found to be crucial in determining the wear behavior of the metals. For instance, in pure copper a subsurface layer of DRX structure is generated beneath the topmost nanostructured mixing layer (NML) during dry sliding [9]. A pronounced correlation is identified that the wear resistance increases monotonically with a decreasing DRX grain size. By adding Al (0.1–2.2 wt.%) into Cu, the recrystallization kinetics can be adjusted so that the recrystallized grain sizes are changed, which corresponds to an obvious change in the wear resistance. A wear mechanism transition has been found when the recrystallized grain sizes are below a critical value [10]. As the recrystallized grain sizes are larger than 0.7 lm, the wear volume decreases monotonically with a decreasing grain size and the wear mechanism is dominated by peeling-off of the NML. While for the recrystallized grains smaller than 0.7 lm, cracking and peeling-off of the DRX layer becomes a dominating mechanism, corresponding to increased wear volume. Due to such a
⇑ Corresponding
authors. Tel.: +86 24 23971891; fax: +86 24 23971215; e-mail addresses: [email protected]; [email protected]
transition, the Cu-2.2 wt.% Al alloy with very fine recrystallized grains exhibits a rather poor wear resistance compared to that of pure copper, in consistent with the other investigations of Cu–Al alloys [11–13]. The objective of the present study is to explore if the wear resistance of a Cu–Al alloy with a fixed composition can be enhanced by controlling the subsurface recrystallization process via pre-treatments such as plastic deformation and/or heat treatment. A solid solution Cu-2.2 wt.% Al alloy was prepared from 99.999 wt.% Cu and 99.995 wt.% Al. A homogeneous microstructure with coarse grains (CG, 200 lm in size) was obtained, which was subjected to dynamic plastic deformation (DPD) with a strain about 2.0 at liquid nitrogen temperature. The setup and processing parameters of the 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 of the deformed sample, respectively. Microstructure of the as-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 an average size of about 30 nm [15]. The as-DPD samples were annealed at 100–460 °C for 10 min for modifying microstructure via recovery and/or recrystallization. Sliding wear tests of the Cu–Al samples were performed on an Optimal SRVIII oscillating friction and wear tester in a ball-on-plate contact configuration under dry condition at room temperature (25 °C) in air with a relative humidity of 45%. Plates were cut from the Cu–Al specimens to a dimension of 6 6 3 mm3 with an electro-polished sur-
http://dx.doi.org/10.1016/j.scriptamat.2015.01.023 1359-6462/Ó 2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
X. Chen et al. / Scripta Materialia 101 (2015) 76–79
face. WC-Co balls of 10 mm in diameter with a micro-hardness of 17.5 GPa were used as the counterface. The friction and wear tests were carried out at a sliding amplitude of 500 lm, a normal load of 30 N, a frequency of 5 Hz, and a duration of 60 min. Profiles of the worn surfaces were measured by using a MicroXAM 3 dimensional (3D) surface profilometer system so as to determine the wear volume. Before the worn subsurface structure characterization, the wear scars were ultrasonically cleaned with acetone and alcohol respectively, on which a protective layer of copper was electrodeposited. Cross-sectional samples were cut parallel to the sliding direction and grinded to the center of the wear scars. The worn subsurface structure of the Cu–Al specimens was characterized by using scanning electron microscopy (SEM) in electron channeling contrast (ECC) mode with a thermal field emission gun on a FEI Nano-SEM Nova 430 system operated at 15 kV. Detailed microstructure characterization was performed on a Tecnai G2 F-20 transmission electron microscope (TEM) operated at a voltage of 200 kV on a sectional wear scar. Thin foils for TEM observations were prepared on a Nova 200 Nanolab focused ion beam (FIB) system. A thin layer of platinum was deposited on the worn surface to protect the surface features against beam damage. The measured wear volume of Cu-2.2 wt.% Al alloy is 3.1 107 lm3 after sliding against WC-Co ball for 60 min under a normal load of 30 N, which is higher than that of copper sample (1.8 107 lm3) [9]. After the DPD processing, the wear volume of the Cu–Al sample decreases to about 0.9 107 lm3 under the same sliding condition. The measured steady-state friction coefficient of the DPD Cu–Al sample is about 0.7, comparable to that of the Cu–Al sample. For both the CG and the as-DPD Cu–Al samples, the steady state subsurface structures are formed under the worn surface during sliding. Extremely fine structure is found in the subsurface layer of the CG sample (Fig. 1a). TEM observations show a thin and continuous nanostructured mixing layer, consisting of nano-sized grains with
random crystallographic orientations, is developed on the top of the DRX layer with extremely fine structure (Fig. 1b). The detailed microstructures in the NML and DRX layer were investigated in previous work [10]. For the as-DPD sample, a subsurface layer of obvious DRX structure is generated beneath the topmost NML, as shown in Figure 1c. There is a sharp boundary between the NML and the DRX layer (as outlined by dashed lines). Sliding on the as-DPD sample leads to the formation of much larger recrystallized grains compared to that in the CG counterpart. Cracking and peeling-off in the DRX layer were observed (Fig. 1a), which is the dominating wear process for the Cu-2.2% Al sample [10]. While for the as-DPD sample, micro-cracks are found either inside the NMLs or at the NML/DRX interfaces, similar to that observed in Cu [9]. Propagation and coalescence of these cracks result in fracture and peeling-off of the topmost NML consequently, which is quite different from the material removal mechanism exists in the Cu–Al sample. Careful examinations reveal an obvious difference in the DRX grain size between the two samples (see Fig. 1d and e): the average size is about 0.28 and 0.62 lm for the CG and the as-DPD sample, respectively. It seems that the enhanced wear resistance of the as-DPD sample is attributed to larger recrystallization grain size. To verify recrystallization grain size effect on the wear resistance, different Cu2.2% Al samples were prepared by controlling recrystallized kinetics via DPD and subsequent heat treatment. Microstructural evolution of the DPD samples during thermal annealing has been well-studied in copper [16] and in other materials [17]. Annealing induced static recrystallization (SRX) process of the DPD Cu-2.2% Al sample (above 260 °C), analogous to that in Cu [16], occurs initially in the nano-grained regions due to their higher stored energy compared with that of the nano-twined structure. The partially recrystallized samples are structurally characterized by a mixed structure of micro-sizes SRX grains, nano-sized grains and nano-twin bundles. Microhardness drops sharply when the annealing temperature (Ta) exceeds
b
NML
b
0.20
d
0.16
Volume Fraction
a
77
DRX
0.12 0.08 0.04 0.00 0.0
500 nm
0.2
0.4
0.6
0.8
1.0
1.2
DRX grain size
c
0.20
NML
20 µm
Volume Fraction
DRX
e
0.16 0.12 0.08 0.04 0.00 0.0
0.2
0.4
0.6
0.8
1.0
1.2
DRX grain size
Figure 1. SEM-ECC images of the cross-sectional worn subsurface structure and corresponding distributions of DRX grain size underneath the NMLs for the CG (a and d) and the DPD (c and e) Cu-2.2% Al alloy after sliding against a WC-Co ball at a load of 30 N. (b) A bright field TEM image of the NML with the DRX layer for the Cu-2.2% Al alloy. The worn surfaces are outlined by red dashed lines and the NMLs are outlined by yellow dashed lines. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
78
X. Chen et al. / Scripta Materialia 101 (2015) 76–79
a
0.24
b
0.16 0.08
Volume Fraction
0.00
c
d
0.20 0.12 0.04 0.24
f
0.16
e
0.08 0.00
10 m
0.2
0.4
0.6
0.8
1.0
DRX grain size (µm)
Figure 3. SEM micrographs of worn subsurface structures and corresponding distributions of DRX grain size for DPD samples annealed at (a and b) 160 °C; (c and d) 340 °C; (e and f) 460 °C for 10 min. Figure 2. Variations of hardness (a) and corresponding wear volume (b) for the CG, the DPD and the annealed DPD Cu–Al alloy at different temperatures. The inserts in (a) show the SEM-ECC images of the structure in the DPD samples annealed at 280 °C (left) and 460 °C (right).
280 °C due to the occurrence of more obvious SRX grains (constituting 50% in volume, left insert in Fig. 2a), as depicted in Figure 2a. The mean SRX grain size is about 2.5 lm. The volume of fraction of SRX grains increases at higher annealing temperatures. Hardness decreases from 2.3 GPa of the as-DPD sample to 0.9 GPa for the sample annealed at 460 °C, which corresponds to an increasing volume fraction of SRX grains to 98% (right insert in Fig. 2a). The mean SRX grain size is about 3 lm. However, the corresponding wear volumes of the annealed samples do not follow the hardness variation linearly, as expressed in Archard’s law (see Fig. 2b). For instance, there exists no obvious decrease in the wear volume when Ta is below 340 °C, although the hardness decreases to a lower level of 1.2 GPa when annealing at 340 °C. But the wear volume increases from 1.0 107 to 1.8 107 lm3 as Ta increases from 340 to 460 °C, accompanied by a further decrease of hardness from 1.2 to 0.9 GPa. Obviously, initial microstructure and hardness may not have a straightforward 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 beneath the top NML. Micro-cracks are frequently observed inside the NMLs and/or at the NML/DRX interfaces, analogous to that in the as-DPD sample. However, for the DPD sample annealed at 460 °C, the subsurface DRX grains are much refined (Fig. 3e) and cracks are often found in the DRX layer, as observed in the Cu–Al sample. Quantitative measurements show that the grain size distribution of the sample annealed at 160 °C is identical to that of the sample
annealed at 340 °C (see Fig. 3b and d). Back to the wear volume data (Fig. 2b), no obvious difference exists between two samples despite of a remarkably difference in the initial hardness. However, the average grain size of the sample annealed at 460 °C is 0.42 lm (Fig. 3f), which is obviously smaller than that of the samples annealed at 160 °C (0.58 lm) and 340 °C (0.6 lm). This is consistent with our previous observation [10] that cracking and peeling-off of the DRX layer is the dominating process when DRX grain size is less than 0.5 lm. Figure 4 plots a relationship between the wear volume and the average DRX grain size for the Cu-2.2% Al samples with different initial microstructures. The wear volume decreases significantly with an increasing grain size from 0.28 to 0.62 lm. It indicates that the wear resistance can be elevated when the recrystallized process is enhanced for the Cu-2.2% Al alloy. Combined with the data for the Cu–Al alloys with different Al concentrations [10], these data follow two different variation tendencies, depending on the DRX grain size. The wear volume decreases monotonically with a decreasing DRX grain size from 0.92 to 0.7 lm, and increases with a further decrease from 0.7 to 0.28 lm. The present work reveals that the recrystallized kinetics during sliding is controlled by adjusting the initial microstructure of Cu-2.2% Al sample. DRX grain size is determined by the stored energy and grain boundary mobility during the DRX process [18]. For the Cu-2.2% Al sample, extremely small DRX grains are derived from a much reduced boundary mobility due to alloying with Al, compared with that in Cu. However, after the DPD treatment, the driving force for dynamic recrystallization during sliding is enlarged owing to very high stored energy in the deformed matrix [19]. The DRX kinetics is enhanced by
X. Chen et al. / Scripta Materialia 101 (2015) 76–79 4.0
Wear volume (107µm3)
3.5
corresponds to the subsurface softening process. Although the DPD Cu–Al alloy exhibits very limited uniform tensile elongation (1–2%) [15], the recrystallized structure offers plenty rooms to release accumulated plastic strain without cracking in the recrystallized layer during sliding. In summary, the wear resistance of Cu-2.2 wt.% Al alloy can be remarkably enhanced by increasing DRX grain size in the worn subsurface layer. It is attributed to a wear mechanism transition from cracking and peeling-off of the brittle fine-grained DRX layer to peeling-off of the topmost NML. The present finding provides an effective approach to elevate the wear resistance of the Cu–Al alloy by controlling subsurface DRX process.
CG Cu-Al alloys [10] Cu-2.2 wt.% Al alloy
3.0 2.5 2.0 1.5 1.0 0.5
0.2
0.4
0.6
0.8
1.0
79
1.2
DRX grain size (µm) Figure 4. Correlation of the wear volume with the average DRX grain size for a series of Cu–Al alloys [10], the DPD and the annealed DPD Cu-2.2% Al alloy at different temperatures.
the high-energy state of the plastically deformed sample in which there exist nano-sized grains with a very large population of grain boundaries and other defects [20]. As the samples were annealed at higher temperatures, the stored energy is released so that the DRX grain size decreases during sliding. There exists a wear mechanism transition for the Cu2.2% Al alloy accompanied with the DPD and subsequent annealing. For the Cu–Al alloy, cracking and peeling-off of the brittle DRX layer is the dominating wear process. While for the as-DPD and the annealed DPD samples at lower temperatures, peeling-off of the topmost NML dominates the material removal process. The latter mechanism corresponds 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 recrystallized grains are much refined, it is difficult to accommodate plastic deformation during sliding, resulting in cracking within the recrystallized layer. Once the vortical NML is formed when the recrystallized grains are larger enough, sliding may induce cracking and peeling-off of the NML, accompanied with energy-consuming process of transformation from the DRX structure to the NML. For the Cu–Al alloys [10], with an increasing Al concentration, there exists a critical DRX grain size of 0.7 lm determining the minimum wear volume when Al concentration is 0.5 wt.%. The DPD processing seems effective in improving the wear resistance of Cu–Al alloy by adjusting the recrystallized structure. It is interesting that a lower wear rate of the DPD Cu–Al alloy can be achieved by increasing DRX grain size, which
The authors are grateful for the financial supports from the National Natural Science Foundation of China (Grant 51231006), the MOST 973 Project (Grant 2012CB932201), the Danish-Chinese Center for Nano-metals (Grant 51261130091), the Key Research Program of Chinese Academy of Sciences (Grant KGZD-EW-T06) and Shenyang National Laboratory for Materials Science.
[1] D.A. Rigney, X.Y. Fu, J.E. Hammerberg, B.L. Holian, M.L. Falk, Scripta Mater. 49 (2003) 977–983. [2] A. Mahato, N. Verma, V. Jayaram, S.K. Biswas, Acta Mater. 59 (2011) 6069–6082. [3] F. Ren, S.N. Arshad, P. Bellon, R.S. Averback, M. Pouryazdan, H. Hahn, Acta Mater. 72 (2014) 148–158. [4] D. Rigney, S. Karthikeyan, Tribol. Lett. 39 (2010) 3–7. [5] S. Karthikeyan, A. Agrawal, D.A. Rigney, Wear 267 (2009) 1166–1176. [6] J.H. Dautzenberg, Wear 60 (1980) 401–411. [7] A. Emge, S. Karthikeyan, D.A. Rigney, Wear 267 (2009) 562–567. [8] S.M. Kuo, D.A. Rigney, Mater. Sci. Eng., A 157 (1992) 131– 143. [9] B. Yao, Z. Han, K. Lu, Wear 294–295 (2012) 438–445. [10] X. Chen, Z. Han, K. Lu, Wear 320 (2014) 41–50. [11] R.A. Poggie, J.J. Wert, Wear 156 (1992) 315–326. [12] J.J. Wert, W.M. Cook, Wear 123 (1988) 171–192. [13] J.J. Wert, S.A. Singerman, S.G. Caldwell, D.K. Chaudhuri, Wear 92 (1983) 213–229. [14] Y.S. Li, N.R. Tao, K. Lu, Acta Mater. 56 (2008) 230–241. [15] Y. Zhang, N.R. Tao, K. Lu, Acta Mater. 59 (2011) 6048– 6058. [16] Y.S. Li, Y. Zhang, N.R. Tao, K. Lu, Scripta Mater. 59 (2008) 475–478. [17] F.K. Yan, G.Z. Liu, N.R. Tao, K. Lu, Acta Mater. 60 (2012) 1059–1071. [18] C.C. Koch, R.O. Scattergood, K.A. Darling, J.E. Semones, J. Mater. Sci. 43 (2008) 7264–7272. [19] F. Yan, H.W. Zhang, J. Mater. Sci. Technol. 28 (2012) 289– 293. [20] L. Yang, N.R. Tao, K. Lu, L. Lu, Scripta Mater. 68 (2013) 801–804.