Wear resistance of Cu-Ag multilayers: A microscopic study - ACS

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Wear resistance of Cu-Ag multilayers: A microscopic study Madhavan R, Pascal Bellon, and Robert Averback ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b03015 • Publication Date (Web): 07 Apr 2018 Downloaded from http://pubs.acs.org on April 7, 2018

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ACS Applied Materials & Interfaces

Wear resistance of Cu-Ag multilayers: A microscopic study R Madhavan*, P Bellon and RS Averback Department of Materials Science and Engineering, University of Illinois at Urbana-Champaign, Urbana, IL 61801, USA *Corresponding author. Email: [email protected]

Abstract The microscopic wear behavior of copper-silver multilayer samples was studied by performing sliding wear tests using a tribo-indenter. Multilayers with an average composition of Cu90Ag10 and Ag layer thicknesses ranging from 2 to 20 nm were grown by magnetron sputtering. For reference, a homogeneous Cu90Ag10 solid solution film was similarly grown. The thin films were subjected to two-dimensional wear tests by rastering a cono-spherical diamond indenter under loads of 100 to 400 µN, for 1 to 20 consecutive passes, or cycles. The wear volumes were determined by atomic force microscopy. Characterization of the specimens employed nanoindentation, nano-scratch and transmission electron microscopy (TEM). Wear rates were found to reach steady state after 5 cycles or less. The hardness values of as-grown and worn samples both increased with decreasing thickness of the Cu and Ag layers, while the steady state wear rates decreased. Notably, the wear resistance increased faster than the corresponding increase in indentation hardness, indicating a deviation from Archard’s law. An inverse relationship between wear rate and hardness was, however, recovered when using scratch hardness, suggesting that scratch hardness is a better predictor of wear resistance. Characterization of sub-surface wear microstructures by TEM revealed that forced chemical mixing and dissolution of layers occur to a depth of ≈ 40 to 50 nm, stabilizing a chemically homogeneous solid solution below the wear surface. Comparative wear tests on thicker multilayers revealed that Cu/Ag interfaces reduced wear rate significantly, thus helping to rationalize the high wear resistance of thin multilayers.

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Keywords: Cu-Ag multilayers; Microscopic wear; Chemical mixing; Scratch hardness; Archard law.

1. Introduction Materials for wear applications are traditionally selected based on key mechanical and chemical properties. In the case of sliding wear of metallic materials, the focus of the present work, it is in particular common to select materials with high hardness, high toughness, and a low coefficient of friction.1-2 These criteria, however, are insufficient or inadequate at high loads when contacting asperities from material surfaces undergo plastic deformation, resulting in significant modifications of the microstructure of these materials. At the interface of two materials in sliding contact, for instance, deformation-induced chemical mixing can favor galling, and it has been suggested that in order to suppress this effect the main elements in the two materials sliding past each other should be chemically immiscible.3 This approach, however, is complicated by the fact that tribolayers, or third bodies, with distinct composition can form at the interface between the first bodies.4 Furthermore, below the sliding interface, plastic deformation takes place by dislocation generation and accumulation, leading to broad microstructural reorganization, such as grain refinement and texturing.5-7 In elemental systems such as Cu, Ni, or Al, for instance, it has been reported that plastic deformation can induce the formation of layers with alternating crystallographic orientation,8 with layer thicknesses decreasing with increasing strain, and reaching thicknesses as low as 12 nm in Cu subjected to sliding friction. More recently, it was observed in macroscopic pin-on-disk wear tests that Cu90Ag10 alloys, with an initial microstructure comprised of equiaxed Ag nanoprecipitates in Cu, spontaneously form chemical nanolayers, with alternating Ag-rich and Cu-rich layers.9 The higher the strain, the thinner are the layers, and close to the wear surface, where the plastic strain is the largest, the thickness of these layers can reach ≈ 5-10 nm. Remarkably, this nanolayering was found to increase the wear resistance of Cu-Ag alloys by a factor 2 to 15 compared to solid solution


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alloys with the same composition. This effect is even more surprising when considering the fact that those solid solutions have higher hardness than the Ag/Cu nanolayered structures. In order to fully assess and realize the potential of this self-organization reaction for improving the wear resistance of metallic materials, it is essential to understand the effect of the thickness of alternating Cu and Ag layers on wear resistance. It is, however, very challenging to address this question using macroscopic samples, where wear induces the chemical nanolayering reaction. First, the layer thickness cannot be directly and simply varied as it results from a complex selforganization reaction. Second, owing to surface roughness and microstructural heterogeneities, loading conditions and plastic deformation are not uniform at the sliding surface, resulting in a spectrum of layer thickness values rather than one that is well defined. In this work, we use an alternative approach, where Cu/Ag thin film multilayers are grown by physical vapor deposition, thus affording complete control over layer thickness and composition, and we employ a nano-tribo-indenter to perform two-dimensional wear tests with well-defined tip-surface geometry and applied load. We note that metallic nanolaminates have been studied for applications that demand exceptional tribological performances owing to their high hardness and tunable toughness properties, in comparison to the individual constituents of the multilayer.10-11 For layer thickness λ > 50 nm, the flow strength of multilayers follows a conventional Hall-Petch (H-P) model, with the strength proportional to λ/. As the layer thickness is reduced below ≈ 50 nm, the strength of multilayers is often found to increase faster than predicted by the H-P model, an effect that is attributed to confined layer slip12-13. As a consequence, very high hardness values can be achieved using multilayers, as will be also illustrated in the present study. There has been, however, limited studies on the tribological performances of metallic multilayers. Studies related to one dimensional scratch experiments on metallic multilayers such as Ag/Ni14 and Ag/Cu15 showed that the coefficient of friction reduces with decreasing bi-layer period, a property that is attributed to higher hardness/modulus ratio for



thin layers which decreases the material pile-up. Luo et al.16 reported on the microstructural evolution under sliding scratch tests on Cu/Au multilayer. The study suggested that the formation of local microstructural inhomogeneities at the interface act as microstructural precursors for vortices and subsequent mechanical mixing, but the impact of the layered structure on wear was not studied. Two-dimensional microscopic wear testing has been emerged as a technique to determine the wear properties of thin films, since it mimics the dry sliding wear condition under single-asperity contact in small-scale devices. Recently, Schultz et al.17 studied the tribological properties of monolithic copper film by nanowear tests under varying contact pressures and line spacings between parallel passes of the sliding probe. The study primarily focused on determining the hardness variation within the wear pattern as a function of wear track and stress-field interactions caused by applied load and line spacing between adjacent scratch passes. A softening effect was reported at lower loads due to the presence of residual stresses in the films and hardening was observed at higher loads due to larger plastic zones and their interactions. The study, however, did not address the correlation between mechanical properties and evolution of the sub-surface microstructure. There are, however, very few studies on the wear resistance of metallic multilayers. One exception is the recent study of Al/Ti multilayers with either 2.5 or 30 nm individual layer thickness, subjected to 2D microscopic wear testing18. It was reported that the wear resistance scales linearly with applied load, as predicted by Archard’s law, and that the thinner multilayer, which was ≈ 60% harder than the thicker one, had improved resistance. The wear resistance improvement, however, was 2-3 times larger than predicted by Archard’s law. Based on transmission electron microscopy characterization, the authors suggested that this may be due to the relaxation of interfacial dislocations in the thinner multilayer. The study, however, considered only two bi-layer periods, so it is difficult to quantify the effect of layer thickness on wear resistance. Furthermore, Al and Ti have a negative heat of mixing and are very reactive, in contrast to Cu and Ag, which have positive heat of mixing of ≈ 6 kJ /mol.19 This is important since it has been established that the


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heat of mixing can play an important role on alloy evolution during sustained plastic deformation.19-22 In this work, we thus investigate the effect of layer thickness on wear resistance in Cu/Ag multilayers by growing samples with Ag layer thickness ranging from 2 to 20 nm. The overall composition is chosen to be Cu90Ag10 to match that used in the previous macroscopic pin-on-disk studies.23 The thin films are worn using a tribo-indenter. The 2D wear patterns are analyzed by atomic force microscopy as a function of load and number of cycles, and the multilayers are characterized before and after wear testing using nanoindentation, nanoscratch, and transmission electron microscopy.

2. Experimental procedures Copper-silver multilayers having a nominal chemical composition equivalent to 90 at.% Cu and 10 at.% Ag, and a reference solid solution film of the same composition were synthesized by magnetron sputtering. Pure copper and silver targets for sputtering Cu/Ag films were supplied by ESPI metals Inc. and the purity of targets were 99.99%. The films were grown on Si [100] substrates having 1 µm SiO2 layer on top. For the multilayers, both copper and silver were sputtered at a rate of 10 nm/min, by varying the power from 110 W to 40 W, respectively, and the argon pressure was maintained at 2.1 mTorr. Throughout this paper, multilayered samples are designated in A/B format, where A is the thickness of a single Cu layer in nm, B is the thickness of a single Ag layer in nm and A+B is defined as the bi-layer period. Though the composition ratio of Cu and Ag is 9:1, the thickness ratio of Cu and Ag layers is ≈ 6:1 due to the difference in atomic number densities of the elements. Four multilayers with an equivalent alloy composition of Cu90Ag10 were sputtered with different individual layer thicknesses and, using the definition introduced above, are referred to as 12/2, 30/5, 60/10 and 120/20. The Cu90Ag10 solid solution sample was obtained by co-sputtering from separate Ag and Cu targets, also onto an oxidized Si substrate at room temperature. While Cu and Ag are immiscible at that temperature,



past works have shown that homogeneous solid solutions can be achieved across the whole composition range by sputtering at RT24-25. The total thicknesses of the sputtered films were at least 1 µm. All multilayer samples terminated in a silver layer to protect against oxidation. Microscopic wear testing of multilayers and the solid solution was performed at room temperature using a Hysitron TI-950 Triboindenter, using a conospherical tip with a nominal radius of curvature of ~0.7 µm and cone angle of 92° (Fig. S1), purchased from Micro Star Technologies. The radius of curvature calculated by fitting the tip contour from an SEM image is found to be 0.695 µm. The tip was rastered over 10 µm by 10 µm areas using a saw-tooth pattern, as illustrated schematically in Fig. 1a. The tip moves in the X-direction for 10 µm and displaces laterally in the Y-direction to make a two-dimensional wear pattern. The wear tests were carried out in reciprocating wear condition, and a single wear cycle comprises 256 back and forth scratch lines spread over the scan area. Therefore, the successive scratches starting from X=0 are displaced by ~ 40 nm. Wear testing was carried out in dry sliding condition. Before each wear experiment, the tip is cleaned with isopropanol to remove any attached debris and the characterization of the tip before and after wear testing indicated that there was no change in the tip geometry. Samples were tested under 100, 200 and 400 µN loads, and the tip sliding velocity was maintained at 5 µm/s. Using Hertz’s contact model, the contact radius values (a) for these loads were estimated to be 75, 100 and 120 nm. The overall contact (2a) is, therefore, 4-6 times larger than the separation between adjacent scratch grooves, i.e. 40 nm. These calculations imply a complete overlap of adjacent scratch lines. Although the above wear test could be referred to a scanning scratch test since it consists of multiple scratch grooves made in scanning mode, we use here the more common terminology of micro-wear test18, 26. Samples were subjected to successive wear between 1 and 10 cycles. Wear tracks were imaged after wear testing using an atomic force microscope (AFM) (Asylum Research MFP-3D) in tapping mode (Fig. 1b). The average residual wear depths were determined using 10


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horizontal profiles across the track, along the X direction. Since the profile is curved at both ends of the track, only the flat portion of the profiles is considered for wear rate calculations (Fig. 1c). The volume of material removed by the sliding tip is calculated from average wear depth and lateral dimensions of the full wear pattern (10 x 10 µm2). Three tracks were made on each multilayer sample and the solid solution sample to check for reproducibility and to obtain statistics. Note that for each wear test for a given number of cycles, a new wear track was created, rather than adding cycles to a previous wear track. This avoids errors introduced by tip re-positioning when attempting to accumulate wear cycles on the same track. In the present work, wear rates are expressed as the ratio of wear volume (in mm3) normalized by the total sliding distance (in meters) of the tip. In contrast to the macroscopic pin-on-disk wear test process, the micro-wear test does not cause any physical change to the dimensions and weight of samples. So, the term ‘wear volume’ here refers to the volume of material that was displaced as wear debris under the action of the sliding tip. In order to investigate the relationship between microscopic wear rate and mechanical properties of the sputtered films, nanoindentation and scratch hardness values of the sputtered films were measured using the TI-950 system. Nanoindentation experiments were carried out using a Berkovich indenter on as-sputtered films and locally within wear patterns. Hardness and elastic modulus values were determined by fitting the initial slope of the unloading segment of the load-displacement curves using Oliver-Pharr method27. Hardness measurements of thin films by nanoindentation are in general influenced by the presence of the substrate, indentation size effects and residual stresses in films. Indentation loads were chosen to maintain the indentation depths to less than 1/10th of the total film thickness to minimize possible substrate effects. To avoid artefacts from indentation size effects, hardness values were measured under a range of loads between 1000 µN and 2500 µN to determine a threshold load value beyond which the hardness value is not influenced by the penetration depth of the tip, and ensuring that



at least one bi-layer period of the multilayers is indented (Fig. S2). A load of 1500 µN met these criteria for all samples, and this load was thus used for all nanoindentation data reported in Section 3. Regarding the residual stress in thin films, past studies on thin metallic films have shown that this effect is generally small, typically less than 10%28. Moreover, the films are subjected to large stresses under the indent and localized plastic deformation, which should reduce further any influence of the residual stresses. A minimum of 9 indents were made on each as-deposited sample to obtain a mean value and standard deviation. For hardness measurements within the wear track, the tests were planned such that at least 4-5 indents fell on the flat surface of the wear pattern. For the scratch hardness, scratch tests were performed using the same conospherical tip that was used for the micro-wear tests. Scratches of length 10 µm were made with constant 200 µN normal force for multiple scratch cycles (one cycle comprises a back and forth scratch pass). The widths of scratch grooves were determined after 1, 5, 25, 50, 100 scratch cycles using AFM. Similar to micro-wear test, new scratch track was made for each predefined number of scratch cycles. The scratch hardness was calculated from the scratch width using the following simplified expression,29

H =

(1)

where Hs is the scratch hardness in GPa, FN is the normal load, w is the scratch width (Fig. S3). Scratch width is defined as the width between points of inflection across the residual scratch depth profile30, here measured by AFM. In order to investigate the influence of wear properties by the presence of single buried Ag layer, an additional multilayered structure was synthesized by sputtering under identical conditions. The multilayer is designated as 100/10/800, as the thickness of top and bottom Cu layers are 100 nm and 800 nm respectively, and the Ag layer thickness is 10 nm. Micro-wear


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tests, similar to the one outlined above, were carried out to determine the wear rates after each incremental wear cycle. Microstructural characterization of the wear tracks was performed by transmission electron microscopy (TEM), using a JEOL 2010 LaB6 and 2010F field emission TEMs both operated at 200KV. Chemical analysis was performed using energy dispersive spectroscopy (EDS) in the JEOL 2010F in scanning mode, using a probe size of 0.5 nm. Cross-sectional samples were lifted out from the mid-section of the wear patterns, along the Y-direction as shown in Fig. 1b, and thinned in a FEI Helios 600i dual beam focussed ion microscope. A platinum coating was applied on the wear area before the lift-out for protection from ion-beam damage during FIB milling.

Fig. 1: (a) Schematic of stylus motion on the sample surface during micro-wear test. (The individual scratch grooves are shown for representative purpose only and not up to scale). (b) Representative two-dimensional wear pattern made with a spherical tip after 10 consecutive wear cycles. Tip motion is along the X-direction. (c) Depth profile across the wear track shown as line AB in (b), corresponding to 120nm/20nm multilayer film.

3. Results 3.1 Characterization of as-grown films Fig. 2 shows cross-sectional TEM microstructures of the as-grown solid solution and a representative multilayer sample, here the 12/2 sample. Both types of film show columnar



grains in the as-grown state, and in the multilayer sample, the columnar grain growth in Cu layers does not appear to be affected by the presence of the Ag layers. The average grain size for Cu90Ag10 solid solution is 20 nm. Since the grain shape in solid solution film is columnar, the length of smaller dimension, i.e. the length perpendicular to the growth direction is taken as the grain size. For all multilayers, the grain size parallel to the growth direction (out-of-plane) is limited by the layer thickness. Therefore, the out-of-plane grain sizes in Cu layer are same as the layer thickness. The average in-plane grain size of Cu in 12/2 multilayer is about 20 nm, and in 60/10 multilayer is about 60 nm. The interface roughness in multilayers, measured from STEM-HAADF images, vary with individual layer thicknesses and location of the interface from substrate. In 12/2 multilayer, the RMS roughness of the interfaces increases from 2 nm (near the substrate) to 6 nm (near the surface), and in 60/10 multilayer, it increases from 2 nm (near the substrate) to only 3nm (near the surface). Selected area diffraction patterns from the solid solution film confirm the presence of a single-phase FCC phase and that grains exhibit a nearrandom texture (Fig. 2a). On the other hand, patterns from multilayers show diffraction rings corresponding to Cu and Ag layers. The presence of brighter arcs on the {111} rings along the direction normal to the films in the 12/2 multilayer indicate that both Cu and Ag layers are weakly textured (Fig. 4b), whereas in the 60/10 multilayer, the diffracted intensities of the {111} rings are equally distributed suggesting no preferential texture formation during growth (Fig. S4). Fig. 3 shows the nanoindentation hardness of sputtered multilayers plotted as a function of inverse square root of bi-layer period (Cu+Ag). As mentioned above, the bi-layer period correlates with the out-of-plane grain sizes in both Cu and Ag layers. The linear relationship suggests a conventional Hall-Petch-like behavior with respect to bi-layer period, with = 2.5 GPa for Cu/Ag multilayers, and a Hall-Petch coefficient of 4.6 GPa√nm, or 0.145 MPa√m. These values agree very well with values reported for Cu-Ag eutectic alloys and multilayers with layer thickness in the range, 10-100 nm.31-33 It should be noted that the hardness of Cu90Ag10


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solid solution film is, ≈ 4.2 GPa. This hardness, which is consistent with the value of 4.1 GPa reported elsewhere34, is greater than in any of the multilayers, or in nanocrystalline copper, which was found to be ~2.06 GPa, both here and in the literature35. Such a high hardness of the Cu90Ag10 solid solution over that of pure Cu cannot be accounted for by existing models of solid solution hardening in nanocrystalline alloys36 since the relative change of lattice parameter37, and that of shear modulus (measured during our nanoindentation experiments) are small. However, molecular dynamics simulations38 and nanoindentation experiments39 on several Cubase alloy systems suggest that such a large hardening in nanocrystalline solid solutions can result from solute segregation at the grain boundaries, which reduces grain boundary energy, suppresses grain boundary sliding and strengthens the grain boundaries.

Fig. 2: As-sputtered microstructures of Cu90Ag10 alloy in different morphologies. (a) solid solution, (b) multilayered structure with ~12nm thick Cu and ~2nm thick Ag layers. Insets show corresponding diffraction pattern. Arrows in the images denote growth direction.



Fig. 3: Indentation hardness of as-sputtered Cu/Ag multilayers as a function of inverse square-root of bi-layer period. Dotted line denotes the Hall-Petch fit. Hardness of solid solution film, with 20 nm grain size, is shown for comparison. (Multilayered samples are designated as A/B in the top X-axis, where A is the thickness of Cu layer in nm and B is the thickness of Ag layer in nm.) 3.2 Micro-Wear behavior Fig. S5 shows the effect of applied normal load and number of wear cycles on the micro-wear rates obtained from representative Cu/Ag multilayers as well as from the solid solution sample. The wear response as a function of load is obtained by testing under three different applied loads, namely 100 µN, 200 µN, and 400 µN, here measured after 5 cycles (Fig. S5a). The wear rate is observed to increase with the load. A power law fit of the wear rate vs. load plot yields exponent values greater than unity, which suggests a deviation from Archard’s equation as the latter predicts a linear relation between the wear volume and applied load. In light of this observation, the wear rates are expressed in terms of wear volume normalized by the total sliding distance but not by load (i.e. mm3/m). The wear rates of multilayers plotted as a function of number of cycles show that the rates decrease with increasing number of cycles before approaching a steady state value after ≈ 5 wear cycles or less, as seen in Fig. S5b.


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Consequently, the wear rates for the multilayers and solid solution in the present work are reported for a load of 200 µN (to limit the volume of wear debris observed at higher loads) and 10 wear cycles (i.e. in the steady-state regime). The steady-state wear rates for Cu90Ag10 multilayers are reported in Fig. 4. In all the multilayer samples, at least one bi-layer period is worn, and to avoid substrate effects, the maximum wear depth is restricted to 10% of the film thickness. The wear rate of the Cu90Ag10 solid solution is shown for comparison. It is observed that wear rates decrease monotonously with decreasing Cu and Ag layer thicknesses. Note that the wear rates observed in these micro-wear tests are about two orders of magnitude larger than those reported in conventional pin-on-disc wear tests,9 but this is due to use of a sharp indenter here, resulting in much higher contact stresses, ≈ 10 GPa, compared to less than 1 GPa in macroscopic tests.

Fig. 4: Wear rates plotted against bi-layer periods of Cu/Ag multilayers (with an equivalent nominal composition of Cu90Ag10) for a normal load of 200 µN and after 10 wear cycles. Wear rate of Cu90Ag10 solid solution is shown as dotted horizontal line for comparison.



The above results indicate that wear resistance improves with decreasing layer thickness. In order to investigate the relationship between wear rate and hardness, it is useful to measure the hardness of the wear tracks, as opposed to the as-prepared specimens, since metallic multilayers are expected to work harden during the wear test due to plastic deformation. The hardness values were thus measured locally from within the wear pattern (Fig. 5). It should be noted that the RMS values of surface roughness within the wear patterns range between 2 and 4 nm, which are less than 5% of the penetration depths of hardness indents in all the samples. It is, therefore, expected that the surface roughness of the wear patterns has a negligible effect on hardness measurements. The wear induced hardening, assessed here from the difference in hardness before and after wear, increases as the bi-layer period decreases. We note that the worn region hardness from the solid solution sample is higher than those for the multilayer samples. This could be caused by two possible factors: one is the difference in pile-up around the indents in different multilayers, which can underestimate the contact area and leads to error in hardness values. Second is the difference in the evolution of wear microstructures beneath the wear surface. Bolshakov and Pharr40 predicted that the changes in hardness values due to pile-up can be identified by the ratio of depth of residual impression of the indent (hf) to the depth of indent at maximum load (hmax). When hf/hmax ≤ 0.7, the effects of pile-up are not significant and the Oliver-Pharr method provides a good estimate of hardness, whereas if hf/hmax > 0.7, pile-up causes an overestimation of hardness. The hf and hmax values were obtained from the experimental loading-unloading curves of all multilayers and solid solution films before and after wear tests and the ratio has been plotted (Fig. S6). We observe that all values lie within the range of 0.6 to 0.8, which indicate a slight effect on the hardness values from the pile-ups. In order to estimate the actual contribution from the pile-up on the hardness, imaging of residual indents was carried out by AFM to determine the extent of pile-up and estimate the projected contact area (Fig. S7). It is found that no pile-up formed in the solid solution samples, but some pile-ups were present in the multilayers. The hardness values


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calculated using projected contact area indicate an overestimation of hardness in all the multilayers in the range of 7-10%. These corrections are therefore rather small, and do not affect significantly the trends observed in Fig. 5. Regarding the wear microstructure, it will be shown in the next section that the microstructures are only inhomogeneous over the top ≈ 40-50 nm, so the measured hardness of the worn regions should be considered a lower bound for the true hardness of the top surface materials. Lastly, we note that although the solid solution film showed greater indentation hardness values than the multilayer samples, see Fig. 5, its wear rate is, nevertheless, ~25% higher than that of the 12/2 multilayer, see Fig. 4.

Fig. 5: Post-wear indentation hardness values of Cu/Ag multilayers, shown in red circles, compared with the as-sputtered values. The measurements were taken from within wear patterns subjected to 10 wear cycles at 200µN. Hardness values of solid solution sample are shown for comparison, as horizontal lines at corresponding values. 3.3 Wear microstructures In order to shed light on the effect of multilayers on wear processes, we compare and contrast the wear microstructures of the various Cu90Ag10 samples. Fig. 6 shows cross-sectional TEM



microstructures of wear tracks from solid solution and multilayered films after 10 wear cycles. The cross sections were prepared in such a way that they contain the normal direction (ND) of the thin film and the transverse direction (TD) (i.e. direction perpendicular to sliding direction), which we refer to as ND-TD cross sections. In the solid solution film, the wear process alters the specimen microstructure to a depth of ∼ 40 nm; below this depth the microstructure is similar to the as-sputtered condition with columnar grains. The affected surface region shows coarser grains, with grain dimensions of ∼ 50 nm normal to the film, and up to ∼150 nm in-plane (Fig. 6a). Nano-beam diffraction (NBD) patterns indicate this region remains a single phase solid solution of Ag in Cu (Fig. S8). In the case of the multilayers, two different morphologies are observed near the wear surface, depending on layer thicknesses. The microstructures for the 12/2 and 60/10 multilayers are shown as representative cases for small and large bi-layer periods. For the thinner layer thickness, i.e. 12/2, a chemically mixed, homogeneous layer is present at the surface, and it extends to a depth of ∼40 nm (Fig. 6b). This chemically mixed layer is very similar to the surface layer that forms below the surface in the solid solution film after wear. The chemical composition, as determined using STEM-EDS from multiple spots within the mixed layer, is similar in both these cases and is close to the nominal composition, Cu90Ag10 (Fig. S9). The presence of this mixed layer in the steady state wear regime implies that as material is continuously removed from the wear surface new Cu and Ag layers are mixed when they approach the surface. In the multilayers with larger layer thickness, i.e., the 60/10 film, the mixing is incomplete, with the interfaces becoming more diffuse as the surface is approached; the bi-layer period is, however unchanged (Fig. 6c). Further below the wear surface, the microstructure is again similar to the undeformed condition, but evidence of plastic deformation can be seen, for example, the presence of deformation twins to a depth up to ~270 nm (Fig. 6d).


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Fig. 6: Cross-section TEM and STEM microstructures from wear tracks subjected to 10 wear cycles under 200µN normal load. (a) Cu90Ag10 solid solution, (b) 12/2 multilayer. (c) 60/10 multilayer. Inset is the magnified view of diffuse Ag layer near the wear surface. (d) Bright field TEM image from 60/10 multilayer. Inset shows the magnified view of twin features. Sliding direction of the tip is into the plane of paper. Boundaries between surface and sub-surface regions are highlighted using dashed-lines in (a) and (b). By comparing the worn microstructures with the wear-induced hardening in multilayers (Fig. 5), it can be inferred that differences in worn microstructures contributed to the pronounced hardening in thin multilayers, e.g., 12/2, owing to solid solution strengthening. The coarsening of sub-surface microstructure observed in solid solution and 12/2 multilayer could be due to two



possible causes: thermal grain growth due to frictional heating or deformation-induced grain growth. The local temperature rise on the sample surface by the sliding contact, estimated using the equation proposed by Kannel and Barber41, is less than 0.5°C. (The parameters related to the sample are: thermal conductivity of nanocrystalline Cu-Ag alloy - ~200 W/m/k; specific heat 0.385 J/g/K; density - 9.11 g/cc; heat transfer coefficient for convection in air - 20 W/m2/K). This suggests that the observed grain growth is entirely attributed to plastic deformation. It is interesting to note that the wear-induced hardening observed in the solid solution films and in the thin multilayer films is associated with coarsening of the microstructure below the wear surface. While this may appear as surprising at first since grain growth is generally associated with reduction in hardness, Vo et al. showed using molecular dynamics simulations for pure nanocrystalline Cu, that grain boundary energies are reduced during plastic deformation and that this leads to strengthening.42 Similar findings have been reported for sliding wear in Ni-W by Rupert and Schuh43 and attributed to grain boundary relaxation phenomena. We further note that when the initial layer thickness exceeds the thickness of the larger grains induced by wear, ≈ 40-50 nm, the near surface microstructure is very different, as the chemical mixing forced by deformation is not sufficient to stabilize the formation of a solid solution at the wear surface. The large period multilayers did nevertheless yield improvement in wear resistance over pure Cu or pure Ag. The reason for this point will be discussed in the next section. 3.4 Effect of single buried Ag layer Despite the absence of full mixing at the wear surface in thicker multilayers, an increase in wear resistance is measured for the samples with more interfaces, e.g., 60/10 compared to 120/20. This suggests that buried Cu/Ag interfaces might impact the wear response. In order to investigate the possibility of such an effect, we designed a multilayered structure with one thin Ag layer inserted in between two thick Cu layers, the sample described as 100/10/800 in section 2, and tested this sample using the same parameters as for the periodic multilayers. The wear


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rate for this buried layer sample is plotted as a function of increasing number of wear cycles in Fig. 7. It is observed that the wear rate is nearly constant for the first 4 wear cycles, i.e. up to a depth of ≈ 60 nm, and then drops steadily as the tip approaches the interface. In order to understand the morphology of Ag layer as the wear surface nears the interface, cross-section microstructural characterization of the wear track after 5 wear cycles were carried out. It is clear that the silver layer, which is ~25 nm below the worn surface after 5 wear cycles, is affected by the wear process (Figs. 8, S10). The layer becomes very rough and wavy, with a RMS roughness value of the interface of ~18 nm, compared to ~3.5 nm in the undeformed region. With continuous sliding, the concentration of Ag increases in the thick Cu layer present beneath the Ag layer due to wear induced mixing (Fig. S11). The wear rate reaches a minimum value after 7 cycles, which corresponds to a depth close to the Cu/Ag interface (~96 nm). At depths greater than the initial depth of the interface, i.e., after 7 cycles in Fig. 7, the wear rate begins to recover, returning to its original value after 10 cycles, equivalent to a depth of ≈ 140 nm, i.e., ~ 30 nm beyond the original depth of the buried layer. The consequences of this overall sequence for wear resistance will be discussed in Section 4.3.

Fig. 7: Variation in cumulative wear rate as a function of wear cycles for a Cu/Ag/Cu multi-layered structure. Numbers on the plot corresponds to average depth of wear (in [19] ACS Paragon Plus Environment


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nm) after corresponding wear cycle. Inset shows the schematic of the multilayer configuration marked with the thickness of individual layer.

Fig. 8: Z-contrast TEM images from the cross-section of 100nm Cu/10nm Ag/ 800nm Cu multilayer, after 5 wear cycles. (a) as-deposited condition, (b) below wear surface. 4. Discussion 4.1 Analysis of the wear-hardness relationship The first main result of this work is that the wear rates decreased continuously and significantly with decreasing bi-layer period, as shown in Fig. 4. In order to determine whether this improvement in wear resistance can be accounted for by the higher hardness of the smaller bilayer periods (Fig. 3), we checked whether the wear rate and hardness were related through Archard’s law2, i.e.,:

=

.

(2)

here V is the wear volume, k is the wear coefficient, FN is the normal load, s is the total sliding distance of the stylus and H is the hardness. Though this relation was empirically observed to be widely applicable at the macroscopic scale, there has been contrasting reports on its validity


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at smaller length scales.18,

26, 44

Here we find that Archard’s law does not satisfactorily fit the

wear data as seen in Fig. 9, regardless if one uses the as-sputtered or the worn-region hardness values for scaling. This observation is applicable to hardness values estimated by both Oliver-Pharr method and AFM method described above. Specifically, the wear rates fall off too rapidly with decreasing bi-layer period. A similar effect is also present in the data recently reported for Cu/Nb multilayers subjected to micro-wear tests similar to ours,45 although this effect was not directly identified or discussed in that work.

Fig. 9: Steady state wear rate corresponding to 10 wear cycles at 200µN load, plotted as a function of inverse of as-sputtered hardness (filled square symbols) and the corresponding worn region hardness (filled disc symbols) for the Cu90Ag10 multilayered films. Data points of solid solution (open symbols) are shown for comparison. Dotted lines are best fits to Archard’s equation for the multilayer data. Solid solution data are not used in the fit.



Fig. 10: Average wear rate plotted as a function of inverse of scratch hardness for the multilayered films. Fitted dotted line is the expected wear rate according to Archard’s equation. A key limitation of indentation hardness for predicting the microscopic wear as measured here is that it does not take into account the effects of ploughing and lateral force. These effects are, however, present in scratch and micro-wear tests.46-47 While the models of Bowden et al48 and Szlufarska et al49 consider the evolution of friction coefficients due plowing and lateral forces, they do not address the generation of wear volume under the sliding contact. We investigate whether for the present testing conditions a physical correlation might exist between wear rate and scratch hardness. We used the steady-state scratch hardness, specifically reached after 50 passes, for comparison with steady-state wear rates. The remarkable result is that the wear rate of multilayers now follows very well a linear scaling with the inverse of scratch hardness (Fig. 10), in contrast to the poor agreement obtained when using indentation hardness (Fig. 9). The main reason for this difference is that as the bi-layer period decreases the hardness increase is larger for scratch hardness than for indentation hardness, thus better capturing the high wear resistance of the thin multilayers. Moreover, unlike the study of Wen et


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al. on nanoscratch testing of Cu/Ag multilayers,50 there was no change in the scratch morphology between the thinner and thicker multilayers. The use of a Berkovich tip in that study, instead of a conospherical tip, might have influenced the relative role of micro-cutting versus ploughing. It is also interesting to note that in Fig. 10, the linear scaling between wear rate and inverse scratch hardness is even well obeyed for pure Cu and pure Ag thin films. For the Cu90Ag10 solid solution sample, however, the wear resistance is lower than that expected based on its scratch hardness, when compared to the multilayered films. Although the reasons for this difference are not yet fully elucidated, results presented above in Section 3.4 suggest that in multilayered samples, Cu/Ag interfaces below the wear surface provide an additional wear resistance mechanism. In summary, we showed here that a modified Archard law, one relying on scratch hardness instead of indentation hardness, could be used to rationalize, and predict, the wear resistance of Cu/Ag metallic multilayers subjected to two-dimensional wear patterns.

4.2 Sub-surface plastic deformation Multilayers with shorter bi-layer periods performed better than those with thicker layers as seen in Fig. 4, and it is thus interesting to assess the role of interfaces on wear-induced microstructural evolution and on wear resistance. We first consider the extent of plastic deformation below the wear surface. Analysis of the cross-sectional images of the microstructures indicates that the plastic deformation is concentrated at or near the sliding surface. This can be rationalized by considering contact stresses generated by a single point of contact. Using the Hertz’s contact model approach,51 the contact radius (a) and peak stress (p ) at the point of contact for a 0.7 µm spherical tip and 200 µN normal load are found to be 96 nm and 10.3 GPa respectively. An elastic modulus value of ~125 GPa, estimated from nanoindentation test, and Poisson ratio of 0.35 were used for calculating the contact radius. The peak shear stress is located at a depth equal to ∼ 0.49 times the contact radius, i.e. ≈ 47 nm



below the surface, and the distribution of stress is symmetrical along the loading axis (refer to stress field contours for maximum shear stress in Ref.51). However, Hertz’s model assumes a frictionless contact. The effect of frictional stresses can be assessed by considering the solution proposed by Liu52 for a cylinder-plane contact. The depth of the maximum shear stress moves toward the sliding surface as the coefficient of friction µ is increased, e.g., ≈ 0.80a for µ = 0, to 0.70a for µ = 0.15, and ≈ 0 (i.e., at the sliding surface) when µ > 1/3. The coefficient of friction measured during the nanoscratch tests near steady-state was ≈ 0.15, suggesting that for a sphere-plane contact the maximum shear stress peaks at a depth of ∼ 40 nm. This depth is comparable to grain size induced by wear ≈ 30 to 50 nm, as seen in Figs. 6(a, b). The above models, however, neglect plastic deformation of the contact and therefore not well suited for determining the extent of the plastic zone. A better estimate of the plastic zone depth can be obtained using Johnson’s spherical cavity model,53 which yields values here ranging from 3.2a to 2.6a, thus 310 nm to 250 nm for film hardness values ranging from 3 GPa to 4 GPa, respectively. This estimate is consistent with our previous observation of mechanical twins between 100 to 300 nm below the sliding surface (Fig. 6d), and also from the cross-section TEM images obtained from pure copper film and 100/10/800 multilayered film after 5 wear cycles (Fig. S12). Although plastic deformation is concentrated within 300 nm from the surface, some dislocation features were visible even up to ~900 nm. This is not too surprising since once dislocations are nucleated, much lower stresses are sufficient to cause glide in large grain metallic materials. The formation of a solid solution just below the wear surfaces of the multilayer samples provides an estimate the plastic strains resulting from sliding. It has indeed been observed that near and below room temperature, severe plastic deformation of Cu-Ag mixtures by high-energy ball milling54 or by high pressure torsion (HPT)55-57 can force the stabilization of solid solutions across the full composition range, despite the fact that Cu and Ag are thermodynamically


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immiscible at these temperatures. This was rationalized by the chemical mixing forced by dislocation glide across interfaces, combined with the absence of thermally activated diffusion at these low homologous temperatures, thus preventing the alloy system from relaxing toward its two-phase thermodynamic equilibrium.55,

58-59

Furthermore, it was determined by atomistic

simulations59-60 and validated by experiments on Cu-Ag mixtures61 that the strain required to homogenize a two-phase structure would scale linearly with the initial length scale. In the case of Cu/Ag multilayers obtained by eutectic solidification, it was found in particular that the homogenization by HPT of Cu and Ag layers with bi-layer period of ≈ 30 nm would require a plastic shear strain of ≈ 350.55 The fact that the 12/2 and 30/5 multilayers formed solid solutions just below the sliding surface, but not the thicker multilayers, thus suggests that the plastic strains generated in the present 2D microscopic wear test are ∼ 300 – 400 in the top 50 nm. In the case of multilayers with layer thicknesses exceeding 50 nm, the near-surface plastic strains are insufficient to fully mix the Cu and Ag layers, but partial mixing is nevertheless expected. This is consistent with diffuse and rough interfaces at the top Ag layer observed in the 60/10 sample, as seen in Fig. 6c.

4.3 Impact of interfaces on wear response Wear tests on a multilayer sample having a single buried layer of Ag inserted between two thick Cu layers (100nm Cu/10nm Ag/800nm Cu) show that the wear rate is at first nearly constant, at a value typically expected for pure Cu, then dropping as the Cu-Ag interface is approached, and finally rising back to values near the original ones as the sample is worn past the Ag layer. Two key conclusions can be implied from these observations: (i) multilayers with thick layers show periodic oscillation in wear rates as the wear surface moves from a thick layer to interfaces; (ii) the buried interface influences the wear behavior at distances up to ≈ 40 nm. The progressive return of the wear rate to the initial value is a consequence of two factors. First, the roughness of the buried Ag layer, which is estimated to be ≈ 18 nm from TEM observations (Figs. 8b, S10),



results in an averaging of the wear rate over a similar depth. Second, some amount of Ag mixed into Cu is still present even after the tip is sliding into the bottom Cu layer (Fig. S11). This mixing, which is estimated to extend over ≈ 10-20 nm from Ag concentration profiles, increases the hardness of the thick Cu layer and thus decreases its wear rate. These additional experiments and analyses thus suggest that in order to maximize the beneficial effect of interfaces on wear resistance, the bi-layer period should be small, typically below ≈ 20-30 nm in Cu/Ag. This conclusion is in good agreement with recent results reported on Al/Ti18 and Cu/Nb multilayers45.

5. Conclusion The microscopic wear response of Cu90Ag10 multilayer samples was studied and compared with a homogeneous solid solution of the same composition. The wear tests were carried out in a tribo-indenter by two-dimensional, back and forth, wear patterning to multiple wear cycles. The major findings of this work are: 1. As the layer thickness decreased, the indentation hardness of multilayers increased, following a Hall-Petch relationship. The hardness of the thinnest multilayer, 3.7 GPa for 12nm Cu/2nm Ag, was lower than the hardness of the solid solution film, 4.2 GPa, although much harder than pure Cu with comparable grain size ≈ 2 GPa. 2. Irrespective of the layer thickness, multilayers reached the steady-state wear regime after 5 wear cycles or less. Decreasing the layer thicknesses led to a reduction in the steady-state wear rates. The decrease in wear rate was, however, faster than the increase in indentation hardness, thus representing a deviation from Archard’s wear law. Furthermore, the thinnest multilayers displayed a wear resistance that is 25% better than the reference solid solution, despite their lower hardness, illustrating the role of Cu/Ag interfaces in promoting wear resistance.


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3. Scratch hardness was shown to be a better predictor of wear rates than indentation hardness, as wear rates were found to vary linearly with the inverse of scratch hardness, thus suggesting a modified Archard law for metallic multilayers at small length scales. 4. In thin multilayers, sliding led to the mixing of the top layers and the stabilization of a supersaturated solid solution over ≈ 40-50 nm below the sliding surface. This depth corresponds approximately to the region experiencing maximum shear stress. In multilayers with a bi-layer period exceeding the mixing threshold depth, sliding could no longer force the stabilization of a solid solution, but instead made the interfaces diffuse and rough. Additional tests on samples with a single buried Ag layer helped quantify the distance over which interfaces can affect wear rates, this distance is ≈ 40 nm.

Acknowledgment The authors acknowledge the funding and technical support from BP through the BPInternational Centre for Advanced Materials (BP-ICAM), which made this research possible. The work was carried out in part in the Frederick Seitz Materials Research Laboratory Central Facilities, University of Illinois at Urbana-Champaign. RM and PB wish to thank Prof. Greg Sawyer, University of Florida, for stimulating discussions.

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Supporting information SEM of indenter tip used in micro-wear test, Variation of hardness with indentation load and depth, Wear rate as a function of load and number of wear cycles, AFM image of wear pattern showing indent impression and depth profile of indents, nano-beam diffraction pattern of subsurface region below worn surface in solid solution, Z-contrast TEM image showing wearinduced waviness and roughness in buried Ag layer, EDS line scan showing localized mixing of Ag and Cu near the interface, extent of plastic zone induced by micro-wear in pure copper and multilayer film.

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Wear resistance of Cu-Ag multilayers: A microscopic study - ACS

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