Friction and Wear Reduction in Copper with a Gradient Nano-grained

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Functional Nanostructured Materials (including low-D carbon)

Friction and wear reduction in copper with a gradient nano-grained surface layer Xiang Chen, Zhong Han, and Ke Lu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b01205 • Publication Date (Web): 04 Apr 2018 Downloaded from http://pubs.acs.org on April 5, 2018

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ACS Applied Materials & Interfaces

Friction and wear reduction in copper with a gradient nano-grained surface layer X. Chen a,b, Z. Han a,*, K. Lu a a

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China

b

Institute for Applied Materials (IAM), Karlsruhe Institute of Technology (KIT), 12 Kaiserstrasse, Karlsruhe 76131, Germany

* Corresponding author. E-mail address: [email protected] (Z. Han) Abstract A gradient nano-grained (GNG) surface layer is fabricated on a commercial-purity Cu sample, in which a significant reduction in the coefficient of friction and the wear loss is obtained compared to the coarse-grained and the nano-grained counterparts. A novel mild ploughing mechanism without subsurface damage has been identified in the GNG sample, giving rise to a much reduced wear rate. Sliding induced surface deformation brings about the unique inhomogeneous substructure in the GNG Cu: the topmost layer persists with nano-grains without being oxidized, underneath which deformation is well accommodated by grain coarsening adjacent to the dynamic recrystallization layer. Both subsurface structural evolution and stress field model confirm that sliding-induced strain localization is suppressed, which is responsible for the superior friction and wear behaviors of the GNG Cu. Keywords: Copper; Gradient nano-grained (GNG) surface layer; Friction; Wear; Surface deformation mechanism

1. Introduction Reducing macroscopic friction and wear losses has always been one of the challenges for metallic components operating under sliding contact conditions, which will significantly prolong machinery lifetime and improve energy efficiency.1-3 According to the classical adhesive theory of Bowden and Tabor, the coefficient of friction (COF) in a tribo-system can be expressed as the ratio of the ACS Paragon Plus Environment

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interfacial shear strength to the hardness of softer component during sliding.4 Therefore, lowering the interfacial shear strength and/or elevating the hardness of material itself seem to be two possible strategies to achieve the low COFs. Generally, extrinsic low friction (0.01-0.3) and wear loss has been successfully approached by applying all sorts of solid and liquid lubricants between the contacting metallic surfaces to reduce the interfacial shear strength.5-9 While for metallic lubricants such as Ag, In and Sn widely utilized in precision bearings under high loading, the COFs can be as low as 0.2-0.4, but accompanied by considerably large wear rates (2×10−5 mm3/Nm) of the lubricants at room temperature.10 Over the past decades, intensive investigations aimed at lowering the intrinsic COFs and enhancing wear resistance have been carried out in various nano-grained (NG) metals and alloys with much higher hardness due to the grain size effect.11-17 Unfortunately, reduced COFs were observed only in a few NG metals at low loads/sliding speeds,14-15 where the applied contact stress is far below the flow stress of the metals. For many NG metals and alloys, no obvious reduction in COFs was observed under high loads compared with their coarse-grained (CG) counterparts.12-15 This arises from the fact that the NG metals possess very limited plastic deformation ability so that strain localization occurs in the sliding surface,18 triggering surface roughening and formation of delaminating tribo-layers.19 The roughening sites unavoidably increase the component of interfacial shear strength as a result of incidental adhesion, which will make it difficult to exhibit the low COFs for the NG metals although hardness is elevated greatly. According to Archard equation, the appreciably positive effect of high hardness on the wear resistance of the NG metals has been anticipated compared to the CG counterparts. However, due to the high COFs, NG metals showed very limited enhancement in the wear resistance under dry sliding,16 for instance, resulting in the wear rate of the NG copper in the order of 10−5 mm3/Nm (CG Cu: 10-4-10−5 mm3/Nm). In some NG cases, controversial results of reduced wear resistance were obtained, which were attributed to a lack of strain hardening capability.20-21 Above all, it is hard to reduce the intrinsic COFs and the wear rates simultaneously for materials under high loading. ACS Paragon Plus Environment

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Tribological loading induces large plastic strain and strain gradients in the subsurface layer of ductile metals, and mechanical mixing with the counterparts or the environment. These processes lead to formation of a sharp demarcation between the nanostructured tribo-layer and adjacent subgrains and dislocation cells along the depth.22-23 The distinguishing feature of the tribo-layers is hard but brittle under repeated sliding, which well corresponds to measured high COF during sliding.14-16 Similarly, finishing or machining processes generate a very thin nanostructured layer on the metallic surface due to plastic deformation induced the strain and the strain gradient below the surface.24-25 This submicrometer-thick layer may enable the COF and wear reduction in running-in stage under oillubricated conditions,24 but their reduction disappeared at high sliding speeds.25 A gradient nanograined (GNG) surface layer can be produced on various metals and alloys by using a newly developed surface mechanical grinding treatment (SMGT), in which grain sizes increase gradually from tens of nanometers in the top surface to micrometers in the interior.19, 26 This GNG layer with a thickness in the millimeter scale exhibited an extraordinary tensile ductility as a result of effectively suppressing strain localization, as well as high mechanical stability.27 Our recent work28 firstly demonstrated that the GNG surface layer of a Cu-Ag alloy yielded a significant reduction in COF under a load of 50 N, from 0.64 (CG) to 0.29. Importantly, a distinct sliding-induced microstructural feature was found: there existed a stable submicrometer-thick topmost layer consisting of nano-sized grains, beneath which grain coarsening occurred. A gradient change in grain sizes was developed with an increasing depth from the surface. The low COF originated from effective suppression of surface roughening and formation of delaminating tribo-layer (e.g. Cu oxides) induced by plastic deformation during sliding. Moreover, for pure Cu samples with relatively low recrystallization temperature, dynamic recrystallization (DRX) structures were commonly observed in the very early sliding stage either in the NG or the CG samples.29 Transformation mechanism from softer DRX structures into the topmost brittle tribo-layers corresponded to surface roughening and the high COFs (~ 0.7) in both Cu samples. In the present study, we utilized the SMGT to prepare the GNG surface layer on the Cu sample. The ACS Paragon Plus Environment


objective is to explore if low COF and wear loss can be achieved in the GNG Cu which is more thermodynamically unstable, and to reveal the involved surface deformation mechanisms of the GNG structure under tribological loading.

2. Experimental 2.1. Sample preparation Commercial-purity Cu (99.97 wt.%) rods were annealed at 723 K for 2 h to obtain a homogeneous CG structure with an average grain size of 20 µm. A GNG surface layer was prepared on the Cu rod by means of the SMGT. A Cu rod with a diameter of 10 mm and a length of 150 mm was processed by using a hemi-spherical WC/Co tool tip with a radius of 8 mm at cryogenic temperature (~173 K). The SMGT processing parameters can be found in the reference.28 For each sample, the SMGT process was repeated six times with the same parameters, in order to achieve a thick and uniform GNG surface layer. No cracks and material removal were identified on the treated surface with a small surface roughness Ra of ~0.2 µm. Bulk NG Cu samples were prepared by using the dynamic plastic deformation (DPD) technique.30 A CG copper cylinder (16 mm in diameter and 25 mm in height) cooled by liquid nitrogen was compressed at a high strain rate of 103 s-1. Multiple impacts were applied to deform a cylinder sample eventually to a disc with a thickness of 3 mm. The total accumulative strain is about 2.0. The microstructure of the DPD Cu sample consists of nano-sized grains with an average size of about 75 nm mixed with about 33 vol.% of nano-twinned regions (twin thickness λ = 49 nm).30 It exhibits a micro-hardness of 1.53 ± 0.08 GPa. 2.2. Friction and wear tests Sliding tests of the Cu samples were performed on an Optimol SRVIII oscillating friction and wear tester in a ball-on-plate contact configuration under dry condition at room temperature (25oC) in air with a relative humidity of 45±5%. Balls of 10 mm in diameter were made of WC-Co with a micro-

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hardness of 17.5 GPa. Before sliding tests, the three kinds of Cu samples were electrochemically polished to the same level of surface roughness (Ra~0.03 µm). The tests were carried out along the length direction of the SMGT rod at a sliding stroke of 1 mm, a sliding velocity of 0.01 m/s and normal loads of 10, 30 and 50 N. One sliding cycle is defined as two strokes. COF, µ = F/P, where F is frictional force and P is normal applied load. COF values were recorded automatically. With a Young’s modulus of 117 GPa for Cu and 680 GPa for WC and a Poisson ratio of 0.3 for Cu and 0.24 for WC, under normal loads of 10, 30 and 50 N, the corresponding Hertzian contact stress was 0.97, 1.40 and 1.67 GPa, respectively. Profiles of the worn surfaces were measured by using a MicroXAM 3 dimensional (3D) surface profilometer system so as to determine the wear volumes after different sliding cycles. A reference surface is determined to quantify the volume of a wear scar, and the volume of material below the reference surface is taken as the wear volume, including the small pile-ups volume above the reference surface at the edge of the wear scar. An effective length of 1 mm of the wear scar is used for determining the wear volume. The wear rate is calculated as wear volume per unit distance and unit load after sliding. 2.3. Surface morphologies observation Surface morphologies and roughness of the Cu samples were measured by using an Olympus 4000 confocal laser scanning microscope (CISM) with a height resolution in the Z axis of 10 nm. In addition, micro-hardness measurements were carried out on a Mitutoyo MVK-H3 micro-hardness tester with a load of 25 g and a loading duration of 10 s. 2.4. Structural characterization Cross-sectional structural characterization of the Cu sample was performed on an emission gun on a FEI Nano-SEM Nova 430 system operated at 15 kV and a JEM-2010 TEM operated at a voltage of 200 kV, respectively. The cross-sectional TEM foils of the as-prepared GNG samples were made by first electrodepositing a Cu coating (about 1.5 mm in thickness) on the GNG surface, then cutting the ACS Paragon Plus Environment


cross-sectional foils, followed by mechanical polishing, and ion milling with a Gatan precision ion polishing system. For the worn Cu samples, thin foils for TEM characterization were prepared by cutting in the middle of the wear scars parallel to the sliding direction in a focused ion beam (FIB) system (FEI Helios NanoLab DualBeam 650). A thin layer of platinum was deposited on the worn surface to protect against beam damage. For deep subsurface positions, the cross-sectional samples were cut perpendicular to the sliding direction along the center of the wear scars. TEM foils were accurately positioned and prepared by using FIB lift-out method. Similarly, a thin platinum layer was deposited to protect the sample surface. 3. Results 3.1 Microstructural characterization A depth-dependent gradient nano-structure was fabricated in a CG Cu rod by using the SMGT at cryogenic temperature (Figure 1a). TEM observations (Figure 1b-e) show a gradually increasing grain size with an increasing depth below the surface. The topmost layer (0-3 µm) of the SMGT sample consists of roughly equiaxed nano-grains with an average transversal size of 30 nm (Figure 1b and c). The grain size gradually increases to be about 60 nm at a depth of 15 µm (Figure 1d). The average grain sizes are less than 100 nm in the top 25 µm-thick layer from the surface and increase to about 300 nm at a depth of 25-150 µm (in the ultra-fine grained (UFG) regime, Figure 2a). UFG structures are separated by sharp grain boundaries (Figure 1e). At a depth greater than 150 µm, typical deformation structures are characterized by dislocation cells with sizes ranging from sub-micrometers to micrometers. The total thickness of the deformed layer is about 0.6-0.7 mm. Corresponding to the grain size gradient is a micro-hardness gradient, from 1.8 GPa in the surface to 1.1 GPa at a depth of 150 nm (Figure 2b). 3.2 Friction and wear behavior


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COFs were tested for the GNG Cu sample sliding against a WC-Co ball with a load of 30 N and a speed of 10 mm/s. For comparison, COFs of the CG and the NG samples were also measured under the same conditions. The measured COF of the CG sample increases from about 0.4 to a steady state value at ~0.74 after sliding for several thousand cycles (Figure 3a), a typical dry frictional response of ductile metals in the literatures.31 A similar COF variation was observed in the NG sample: COF rises from 0.37 to a steady state value of 0.67, slightly lower than that of the CG sample. This friction behavior is in account with that of the other NG materials under high loads.14 However, for the GNG sample, COF remains unchanged at 0.37 during all the sliding cycles, much lower than the steady state COFs of the NG and the CG samples. Repeated measurements with at least three tests show that the steady state COF of the GNG sample changes slightly from 0.35 to 0.4 with an increasing load from 10 to 50 N (average 0.37, Figure 3b), much lower than that of the NG and the CG samples (0.62-0.75). Simultaneously, the measured wear volumes of the GNG sample are much smaller than that in the NG and the CG samples under different sliding cycle at a load of 30 N (Figure 4a). For instance, the wear volume of the GNG sample is 1.6 × 106 µm3 after sliding for 6000 cycles, much smaller than that in the NG (3.7 × 106 µm3) and the CG (1.4 × 107 µm3) samples. Correspondingly, the GNG sample also exhibits much reduced wear rates in the load range from 10 to 50 N (Figure 4b), compared to the CG and the NG counterparts. Under a load of 30 N, the wear rate for the GNG sample (4.4 × 10−6 mm3/Nm) is much smaller than that of the NG (1.0 × 10−5 mm3/Nm) and the CG (3.9 × 10−5 mm3/Nm) samples. 3.3 Surface damage feature COFs of metals are sensitive to the surface morphology and its variation during sliding.32 As illustrated in Figure 5 (a, b, d, e), the surfaces of the CG and the NG samples become rather rough in the steady state of sliding, wherein plenty of cracks and plate-like debris are formed. The measured surface roughness Ra is about 0.12 and 0.18 µm for the CG and the NG samples after sliding for 6000 cycles, respectively. For the softer CG Cu sample, surface cracks easily appear as a result of sliding



induced severe plastic deformation. And for the NG sample crack formation and propagation in the surface is attributed to strain localization for its limited plastic deformability. This type of sliding induced damage was also reported in other tribo-systems,14 which corresponded to high surface roughness and COFs. However, for the GNG sample sliding after 6000 cycles, the sliding surface (Figure 5c and f) remains as smooth as the original one, with a steady surface roughness of Ra = 0.03 µm. The GNG surface exhibits high resistance to surface roughening, which is quite distinct from the CG and the NG samples. 3.4 Worn subsurface structure For the CG and the NG samples, microstructure evolution and sliding-induced damage accumulation have been intensively investigated in our previous publications.29,33 DRX structure was always generated beneath the topmost tribo-layer under sliding contact.29 The tribolayer is a nanostructured layer with a grain size as small as 15 nm (Figure S1), containing a substantial amount of oxygen mixing with the environment. Micro-cracks were frequently observed within the tribo-layer (Figure S1) and at the interfaces between the tribo-layer and DRX layer.29 The formation of the brittle tribo-layer and its delamination during sliding corresponded to high surface roughness and COFs.14 However, in the GNG sample sliding after 6000 cycles, no obvious structure change is observed in the topmost layer ( 1 GPa, see Figure 9b), which far exceeds its yield stress (~585 MPa), implying that fracture and delamination inevitably occur in the surface layer. The stress field analysis is in well accord with high COF in the NG sample as a result of strain localization. However, an inherent gradient in the elastic limit is present in the GNG sample. During sliding, plastic deformation is accommodated along a thick subsurface layer when comparing the applied stress to the yield stress in the GNG sample, which will definitely generate a smaller strain gradient than that in the NG sample (Figure 9c). The proposed mechanism can also be verified from the above experimental results of rather smooth sliding surfaces and deeper subsurface structural evolution layer that has undergone plastic deformation in the GNG Cu sample. Interestingly, the estimated layer depth that can endure plastic deformation is over 60 µm in the leading edge of contact from the stress field analysis (Figure 9c), which is comparable to that determined from the SEM images for the subsurface ACS Paragon Plus Environment

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structural observations (60-90 µm, Figure 6a). Hence, strain localization in the sliding surface layer is released effectively, which may suppress cracking and pile-ups induced surface roughening, finally result in the obviously reduced COF. Apparently, the very low and stable COFs observed in the GNG Cu sample originate from the unique GNG structure that can accommodate massive plastic deformation effectively during sliding. Actually, the low COF was also achieved in a CG Cu sample subjected to mild sliding under a load of 20 mN, which was ascribed to the intermediate formation of a very thin ultra-nanocrystalline (10-20 nm) films in the surface layer.35 The corresponding contact stress is calculated to be about 178 MPa in the low friction regime, much less than the yield stress of the ultra-nanocrystalline films. However, this low friction regime disappeared while the contact stress exceeded its yield stress, which agrees well with our results in the CG and NG samples. Therefore, the low friction behavior in the GNG sample cannot be simply ascribed to grain boundary sliding and dislocation activity, due to distinctly different deformation mechanisms identified in the surface layer under high tribological stress. 4.1.2 Gradient effect To generate gradient distribution of grain sizes from the nano to the macro-scale can be regarded as a new and effective strategy for achieving low COFs of Cu and CuAg alloy. Then, a quick question turns to whether gradient grain sizes from the submicron to the macro-scale can still be effective in lowering the COF of the Cu sample. In order to address the above question, the gradient-structured (GS) sample was prepared by removing the topmost 20-30 µm nano-grained layer in the GNG sample. As shown in Figure 10a, micro-hardness of the GS sample decreases from 1.4 GPa in the surface to 0.7 GPa in the CG matrix, which confirms the removal of nano-grained layer. Measured COF of the GS sample increases from 0.3 to 0.54 after sliding for 6000 cycles (Figure 10b) and tends to be a steady state at 0.7 after 12000 cycles (not shown in Figure 10 for brevity). For the GS sample, DRX structure is generated beneath the topmost delaminating tribo-layer in the steady-state sliding, without the GC layer formed in the GNG sample (Figure 10c). Micro-cracks are frequently observed within the tribo-layer and at the interfaces between the tribo-layer and DRX layer, which is similar to that in ACS Paragon Plus Environment


the NG sample.29 Both COF elevation and subsurface structural feature indirectly confirm that the nano-scale grains in the topmost layer indeed play a crucial role in maintaining the low friction state in the GNG sample. Nevertheless, the GS sample shows lower COF value compared to the NG sample, although the maximum micro-hardness in the surface (1.4 GPa) is less than that in the NG sample (1.6 GPa). In this respect, the gradient distribution of grain sizes can somehow lower the COFs of metals, which merits further investigations. 4.2 Wear mechanism of GNG Cu sample Accompanied by COF reduction, the wear loss of the GNG Cu sample is markedly lowered under high tribological stress. Previous results of Cu samples with various initial homogeneous structures showed a DRX dominated wear mechanism, underlying that the wear volume decreases monotonically with a decreasing DRX grain size in the subsurface layer.29 It is worth mentioning that the straightforward correlation applies to the NG, the UFG and the CG structures prepared by using different deformation techniques, including DPD, quasi-static compression and multiple cold rolling. And the wear debris originates from cracking and peeling-off of the topmost tribo-layers transformed from the DRX structure (Figure S1). Obviously, the very low wear volume of the GNG Cu deviates from above monotonic relationship, although DRX grain sizes in the NG and the GNG samples are nearly the same (~700 nm). It can be seen that the wear mechanism of the GNG Cu is not dominated by subsurface DRX structure anymore. Superior wear resistance of the GNG Cu sample benefits from a sort of low friction state without surface cracking and peeling-off. The underlying wear mechanism of the GNG Cu sample is mild ploughing without cracking in the surface and the subsurface layer (see Figure 5c and 5f), which corresponds to extra low COF and wear loss. The mild ploughing mechanism was only observed in a few NG metals under low tribological stress,14-15 that is, at low sliding speeds and/or forces or with lubrication. However, most NG metals suffered from high wear rates (in the order of 10−5 mm3/Nm) because of rapid peeling-off of the tribo-layers and even cracking in the adjacent deformation layer under high-load sliding.16, 19, 36 To the best of our knowledge, a novel mild wear mechanism has been identified under high load sliding and a very low wear rate has been


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achieved in the GNG Cu in comparison with that of homogeneous NG structures under comparable sliding conditions in the literatures.16, 19-21, 36 4.3 Subsurface structure stability In comparison with the GNG Cu-Ag alloy, the GNG Cu sample exhibits different subsurface microstructural stability during sliding contact. Apart from grain coarsening in the GNG Cu-Ag alloy, DRX occurs adjacent to the GC layer in the GNG Cu sample. Firstly, DRX structure with the grain size of ~700 nm is commonly observed beneath the GC layer in the GNG Cu sample during sliding, while no DRX grains are detected below the surface in the GNG Cu-Ag alloy. The main reason is that DRX kinetics is extremely suppressed by adding Ag into Cu, and recrystallization temperature is much lower for the Cu sample (120oC,37) relative to the Cu-Ag alloy (375oC,38). Secondly, the thickness of the GC layer (~20 µm) in the GNG Cu is much larger than that in the Cu-Ag alloy (~2 µm). Recent investigation39 revealed that mechanically-induced grain coarsening in the GNG Cu is not only controlled by the stress, but also thermally activated. Hence, more stable GC layer in the GNG Cu-Ag alloy may also originate from the pinning effect of grain boundary migration by alloying Ag. Sliding induced plastic deformation was found to be dominated by grain coarsening and DRX, indicating that massive plastic strain can be accommodated on the GNG surface layer of Cu sample. Therefore, strain localization induced surface deformation instability can be effectively suppressed, which ultimately achieves a sort of low friction state as a result of the cooperative deformation mechanisms of the GNG Cu sample. The present stress distribution model clearly demonstrates that the shear stress underneath the surface makes a critical contribution to the driving force for DRX process. At the leading edge of contact, one may roughly define DRX activated region by comparing the applied shear stress to the yield stress of the GNG Cu. The predicted DRX zone during sliding is consistent with the measured thickness of the DRX layer from microstructure observations (which is usually 40-50 µm in thickness, see Figure 6a). A minimum shear stress for activating DRX is estimated to be about 420 MPa from the



stress distribution (Figure 9c). It also applies for DRX activation in the NG sample when COF increases to a steady state at 0.66. There is an excellent qualitative agreement between simple model and experimental results. As a final point of discussion, we note that the top NG layer (0-25 µm in thickness) in the GNG Cu sample exhibits higher stability with a smaller saturated grain size (Figure 8f) after sliding for several thousand cycles, compared to that in the initial UFG layer (25-70 µm). As described in the axial tension of the GNG Cu, the grain coarsening is dominated by a mechanically driven grain boundary process,26 resulting in a saturated microstructural size approaching to the sub-micrometer level. For the GNG Cu sample, the grain coarsening in the NG layer can be prohibited by increasing the strain rate of the tensile test, underlying that grain coarsening kinetics is stress-driven.39 It is also well known that the strain rate is very high in the topmost layer and decreases gradually with the depth below the sliding surface,40-41 which implies that high strain rate may play a crucial role in stabilizing the fine NG structures in the topmost layer. Besides, the thicker GC layer and DRX layer in the GNG Cu relative to the GNG Cu-Ag alloy lead to more pronounced strain softening that competes with dislocation hardening. Different strain gradient below the sliding surface may bring a different stability of the GNG structure, which will determine the persistence of low friction state. 5. Conclusions A GNG surface layer is developed on a commercial-purity Cu sample by means of the SMGT. The mean grain size is approximately 30 nm at the topmost layer and increases gradually with the depth from the surface. A significant reduction in the COF and the wear loss was achieved in the GNG Cu sample under high tribological stress, compared to the CG and the NG counterparts. Interestingly, a novel mild ploughing mechanism without subsurface damage has been identified in the GNG Cu sample, leading to much reduced wear rates. During all sliding cycles, the topmost layer in the GNG Cu remains nano-grained structure without being oxidized, underneath which the nano-grains coarsen and DRX structures develop in the initial ACS Paragon Plus Environment

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UFG region. Both subsurface structural evolution and stress field model confirm that plastic deformation is well accommodated within the GNG layer and surface deformation localization is suppressed. The stable gradient nano-structures enable significantly reduced COF and wear loss simultaneously in pure Cu, which suggests the potential engineering applications of metals under high tribological stress.

Acknowledgements We thank J. Tan for assistance in FIB experiments. We are grateful for the financial supports of the National Key R&D Program of China (2017YFA0204401 and 2017YFA0204403), the National Natural Science Foundation (51231006) and the Key Research Program of Chinese Academy of Sciences (KGZD-EW-T06). Thanks for the elastic model guidance from Prof. Lars Pastewka in University of Freiburg.

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(9) Klemenz, A.; Pastewka, L.; Balakrishna, S. G.; Caron, A.; Bennewitz, R.; Moseler, M. Atomic Scale Mechanisms of Friction Reduction and Wear Protection by Graphene. Nano Lett. 2014, 14 (12), 7145-7152. (10) Aouadi, S. M.; Paudel, Y.; Simonson, W. J.; Ge, Q.; Kohli, P.; Muratore, C.; Voevodin, A. A. Tribological investigation of adaptive Mo2N/MoS2/Ag coatings with high sulfur content. Surf. Coat. Technol. 2009, 203 (10), 1304-1309. (11) Farhat, Z. N.; Ding, Y.; Northwood, D. O.; Alpas, A. T. Effect of grain size on friction and wear of nanocrystalline aluminum. Mater. Sci. Eng. A 1996, 206 (2), 302-313. (12) Rupert, T. J.; Schuh, C. A. Sliding wear of nanocrystalline Ni–W: Structural evolution and the apparent breakdown of Archard scaling. Acta Mater. 2010, 58 (12), 4137-4148. (13) Argibay, N.; Furnish, T.; Boyce, B.; Clark, B.; Chandross, M. Stress-dependent grain size evolution of nanocrystalline Ni-W and its impact on friction behavior. Scr. Mater. 2016, 123, 26-29. (14) Padilla, H. A.; Boyce, B. L.; Battaile, C. C.; Prasad, S. V. Frictional performance and nearsurface evolution of nanocrystalline Ni–Fe as governed by contact stress and sliding velocity. Wear 2013, 297 (1), 860-871. (15) Prasad, S. V.; Battaile, C. C.; Kotula, P. G. Friction transitions in nanocrystalline nickel. Scri. Mater. 2011, 64 (8), 729-732. (16) Zhang, Y.; Han, Z.; Wang, K.; Lu, K. Friction and wear behaviors of nanocrystalline surface layer of pure copper. Wear 2006, 260 (9), 942-948. (17) Li, A.; Szlufarska, I. How grain size controls friction and wear in nanocrystalline metals. Phy. Rev. B 2015, 92 (7), 075418. (18) Meyers, M. A.; Mishra, A.; Benson, D. J. Mechanical properties of nanocrystalline materials. Prog. Mater. Sci. 2006, 51 (4), 427-556. (19) Chen, X.; Han, Z.; Lu, K. Wear mechanism transition dominated by subsurface recrystallization structure in Cu–Al alloys. Wear 2014, 320, 41-50. (20) Zhilyaev, A.; Morozova, A.; Cabrera, J.; Kaibyshev, R.; Langdon, T. Wear resistance and electroconductivity in a Cu–0.3 Cr–0.5 Zr alloy processed by ECAP. J. Mater. Sci. 2017, 52 (1), 305313. (21) Zhilyaev, A. P.; Shakhova, I.; Belyakov, A.; Kaibyshev, R.; Langdon, T. G. Wear resistance and electroconductivity in copper processed by severe plastic deformation. Wear 2013, 305 (1), 89-99. (22) Garbar, I. I.; Skorinin, J. V. Metal surface layer structure formation under sliding friction. Wear 1978, 51 (2), 327-336. (23) Rigney, D. A. Transfer, mixing and associated chemical and mechanical processes during the sliding of ductile materials. Wear 2000, 245 (1-2), 1-9. (24) Linsler, D.; Schröckert, F.; Scherge, M. Influence of subsurface plastic deformation on the running-in behavior of a hypoeutectic AlSi alloy. Tribo. Inter. 2016, 100, 224-230. (25) Shakhvorostov, D.; Pöhlmann, K.; Scherge, M. Structure and mechanical properties of tribologically induced nanolayers. Wear 2006, 260 (4), 433-437.


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(26) Fang, T. H.; Li, W. L.; Tao, N. R.; Lu, K. Revealing Extraordinary Intrinsic Tensile Plasticity in Gradient Nano-Grained Copper. Science 2011, 331 (6024), 1587-1590. (27) Li, X.; Lu, K. Playing with defects in metals. Nat. Mater. 2017, 16, 700-701. (28) Chen, X.; Han, Z.; Li, X.; Lu, K. Lowering coefficient of friction in Cu alloys with stable gradient nanostructures. Sci. Adv. 2016, 2 (12), e1601942. (29) Yao, B.; Han, Z.; Lu, K. Correlation between wear resistance and subsurface recrystallization structure in copper. Wear 2012, 294, 438-445. (30) Li, Y. S.; Tao, N. R.; Lu, K. Microstructural evolution and nanostructure formation in copper during dynamic plastic deformation at cryogenic temperatures. Acta Mater. 2008, 56 (2), 230-241. (31) Jungk, J. M.; Michael, J. R.; Prasad, S. V. The role of substrate plasticity on the tribological behavior of diamond-like nanocomposite coatings. Acta Mater. 2008, 56 (9), 1956-1966. (32) Stoyanov, P.; Stemmer, P.; Järvi, T. T.; Merz, R.; Romero, P. A.; Scherge, M.; Kopnarski, M.; Moseler, M.; Fischer, A.; Dienwiebel, M. Friction and Wear Mechanisms of Tungsten–Carbon Systems: A Comparison of Dry and Lubricated Conditions. ACS Appl. Mater. Interfaces 2013, 5 (13), 6123–6135. (33) Yao, B.; Han, Z.; Li, Y. S.; Tao, N. R.; Lu, K. Dry sliding tribological properties of nanostructured copper subjected to dynamic plastic deformation. Wear 2011, 271 (9), 1609-1616. (34) Hamilton, G. M. Explicit equations for the stresses beneath a sliding spherical contact. Proc. Inst. Mech. Eng., Part C 1983, 197 (1), 53-59. (35) Argibay, N.; Chandross, M.; Cheng, S.; Michael, J. R. Linking microstructural evolution and macro-scale friction behavior in metals. J. Mater. Sci. 2017, 52 (5), 2780-2799. (36) Chen, X.; Han, Z.; Lu, K. Enhancing wear resistance of Cu–Al alloy by controlling subsurface dynamic recrystallization. Scr. Mater. 2015, 101, 76-79. (37) Li, Y. S.; Zhang, Y.; Tao, N. R.; Lu, K. Effect of thermal annealing on mechanical properties of a nanostructured copper prepared by means of dynamic plastic deformation. Scr. Mater. 2008, 59 (4), 475-478. (38) Sitarama Raju, K.; Subramanya Sarma, V.; Kauffmann, A.; Hegedűs, Z.; Gubicza, J.; Peterlechner, M.; Freudenberger, J.; Wilde, G. High strength and ductile ultrafine-grained Cu–Ag alloy through bimodal grain size, dislocation density and solute distribution. Acta Mater. 2013, 61 (1), 228-238. (39) Chen, W.; You, Z. S.; Tao, N. R.; Jin, Z. H.; Lu, L. Mechanically-induced grain coarsening in gradient nano-grained copper. Acta Mater. 2017, 125, 255-264. (40) Karthikeyan, S.; Kim, H. J.; Rigney, D. A. Velocity and Strain-Rate Profiles in Materials Subjected to Unlubricated Sliding. Phys. Rev. Lett. 2005, 95 (10), 106001. (41) Sundaram, N. K.; Guo, Y.; Chandrasekar, S. Mesoscale Folding, Instability, and Disruption of Laminar Flow in Metal Surfaces. Phys. Rev. Lett. 2012, 109 (10), 106001.



surface

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200 nm

200 nm

200 nm

200 nm

Fig. 1. (a) Typical longitudinal-sectional SEM image of the as-prepared GNG Cu sample. Typical TEM images showing the microstructures at different depths below the treated surface: (b) 0 μm, (c) 3 μm, (d) 15 μm, (e) 30 μm.


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Grain size (nm)

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Fig. 2. Variations of longitudinal (dl) and transverse (dt) grain sizes (a) and micro-hardness (b) along depth from the surface in the GNG Cu sample.



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Fig. 4. (a) Variations of wear volumes with number of cycles for the CG, NG, and GNG Cu samples sliding against WC-Co balls under a load of 30 N, a slide stroke of 1 mm, and a velocity of 10 mm/s. (b) Variations of wear rates with the sliding load for the three samples.


ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56

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Fig. 5. SEM images (a-c) and 3D confocal laser microscopy profiles (d-f) for surface morphologies of the CG (a, d), NG (b, e), and GNG (c, f) Cu samples after sliding for 6000 cycles. A load of 30 N, a slide stroke of 1 mm, and a velocity of 10 mm/s were applied for each sample. Double-ended arrows indicate the sliding directions. Insets show the surface roughness (Ra) values.


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a GC layer

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Fig. 6. (a) Typical cross-sectional SEM image of the subsurface layer in the GNG sample after sliding for 6000 cycles. The yellow dotted lines outline the grain coarsening (GC)/DRX layer and DRX/deformation layer boundaries. The bright field TEM images show the microstructures indicated in (a) at different depth spans: (b) 0-1 mm, (c) 5-10 mm, (d) 10-15 mm. Double-ended arrows indicate the sliding directions.


ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56

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Fig. 7. Typical worn subsurface structure of topmost layer in the GNG Cu sample after sliding for (a) 3000 and (b) 4500 cycles against a WC-Co ball under a load of 30 N, a slide stroke of 1 mm and a velocity of 10 mm/s, respectively. Insets show the corresponding electron diffraction patterns of topmost layer.


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GNG Cu-Ag (ref. [23]) GNG Cu

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Fig. 8. (a) A typical TEM image shows the microstructure of original GNG Cu sample at a depth of 10 mm below the treated surface. (b-d) Typical cross-sectional TEM images of the GNG Cu sample after sliding for (b) 3000 cycles, (c) 4500 cycles, (d) 6000 cycles in a depth span of 10 to 15 mm. (e) Variation of the average grain size along the depth determined from TEM images before and after sliding for 3000, 4500, and 6000 cycles, respectively. (f) Variation of absolute grain growth along the depth from the surface for the GNG Cu-Ag (ref. 23) and Cu samples. A load of 30 N, a slide stroke of 1 mm, and a velocity of 10 mm/s were applied for each Cu sample. ACS Paragon Plus Environment


a

c

600 Shear stress xz (MPa)

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center of contact front of contact yield, GNG yield, NG

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Fig. 9. The xz-component of the stress field in the middle of the wear scar along the sliding direction for different friction coefficients: (a) COF= 0.37 and (b) COF= 0.66. The arrow indicates the sliding direction. Solid lines show contours of constant stress. (c) Variation of applied stress at the center and the front of contact along depth from the sliding surface in the GNG (COF= 0.37) and NG samples. Measured variations of yield strength along depth for the GNG (taken as approximately 1/3 micro-hardness) and NG samples are included. The positions of the center and the front of contact in (c) are indicated in (a) with the dash dotted and dashed lines, respectively.


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10 mm Fig. 10. (a) Variation of measured micro-hardness with depth from the treated surface for the GS, NG, GNG Cu samples. The GS sample was prepared after removal of the topmost 20-30 mm NG layer in the GNG sample. (b) Variations of COFs with sliding cycles after sliding against WC-Co balls under a load of 30 N for the three samples. (c) Typical steady-state worn subsurface structure in the GS sample.


Friction and Wear Reduction in Copper with a Gradient Nano-grained

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