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CrN–Ag nanocomposite coatings: Tribology at room temperature and during a temperature ramp
Surface & Coatings Technology 204 (2010) 1388–1394
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
CrN–Ag nanocomposite coatings: Tribology at room temperature and during a temperature ramp C.P. Mulligan a,c,⁎, T.A. Blanchet b, D. Gall a a b c
Department of Materials Science and Engineering, Rensselaer Polytechnic Institute, Troy, NY 12180, United States Department of Mechanical, Aerospace, and Nuclear Engineering, Rensselaer Polytechnic Institute, Troy, NY 12180, United States Benet Laboratories, U.S. Army Armament Research Development & Engineering Center, Watervliet, NY 12189, United States
a r t i c l e
i n f o
Article history: Received 1 June 2009 Accepted in revised form 8 September 2009 Available online 18 September 2009 Keywords: Nanocomposite coating Friction Wear Solid lubrication High-temperature lubrication CrN–Ag
a b s t r a c t 5 µm thick CrN–Ag composite layers with 22 at.% Ag were deposited by reactive magnetron co-sputtering on 440C stainless steel substrates. Increasing the growth temperature from Ts = 500 to 600 to 700 °C leads to Ag segregation within the CrN matrix and the formation of embedded lamellar Ag aggregates with increasing size, < 105, 9 × 106, and 7 × 107 nm3, respectively. Ball-on-disk tests against 100Cr6 steel, followed by optical profilometry and energy dispersive spectroscopy, indicate that the Ag grains for Ts = 500 °C are too small to facilitate an effective lubricious surface layer, resulting in a friction coefficient μ = 0.58 and a composite coating wear rate of 3.8 × 10− 6 mm3/Nm that are nearly identical to those measured for pure CrN with μ = 0.64 and 3.6 × 10− 6 mm3/Nm. The Ts = 600 °C coating exhibits a Ag concentration which is 15% higher within than outside the wear track, and acts as a lubricious layer that reduces μ to 0.47 and yields a 16× and 2.4× lower wear rate for coating and counterface, respectively. Ts = 700 °C leads to a dramatic increase in surface roughness and an associated increase in friction, μ = 0.85, and wear, 9.9 × 10− 6 mm3/Nm. Replacing the steel counterface with an alumina ball results in the lowest μ = 0.50 for Ts = 500 °C, attributed to the presence of Ag and the relatively low hardness of 16.5 GPa for this particular coating. In contrast, friction and wear increase dramatically for Ts = 600 °C, which is attributed to a breakdown of the lubricious Ag layer by the harder counterface. The transient friction coefficient μt during experiments with continuously increasing testing temperature Tt = 25–700 °C initially decreases for all samples, attributed to drying of the environment and an effective softening of both coating and counterface. For the Ts = 500 °C coating, a temperature activated solid lubricant transport yields a lubricious Ag surface layer and a very low μt = 0.05 at Tt ~ 500 °C. All coatings exhibit an increasing μt for Tt > 500 °C, which is attributed to oxidative degradation. Published by Elsevier B.V.
1. Introduction Multi-functional coatings using soft lubricating phases within a hard wear-resistant matrix have been studied extensively in recent years due to their promise for tribological performance in transient and cyclic environments [1–10]. In particular, inclusions of noble metals such as Ag and Au show promise as solid lubricating phases in carbide [3], oxide [4–6], and nitride [7–10] matrices, as they possess sufficiently low shear strength over a suitably wide temperature range as well as stable thermochemistry allowing them to be used in both ambient air and vacuum [11–14]. Most studies on co-deposited nanocomposites have focused on the tribological properties as a function of noble metal content, with for example CrAlN–Ag having a promising room temperature friction coefficient as low as 0.23 and wear rate as low as 3.0×10− 8 mm3/Nm at a 2 N load, which depends on the solid lubricant content [7]. For this study, ⁎ Corresponding author. Department of Materials Science and Engineering, Rensselaer Polytechnic Institute, Troy, NY 12180, United States. E-mail address: [email protected] (C.P. Mulligan). 0257-8972/$ – see front matter. Published by Elsevier B.V. doi:10.1016/j.surfcoat.2009.09.018
we focus on the effect of the composite microstructure on the tribological properties for a fixed composition. We expect that the microstructure, including lubricant phase, grain size and shape, as well as matrix porosity strongly affect both solid lubricant transport and the development of an effective solid lubricating interface. Consequently, understanding the relationship between microstructure and tribological properties is critical to the design of appropriate coating architectures that provide lubrication and wear protection for a given application. In this report, the tribological performance of a prototype CrN–Ag nanocomposite coating is investigated as a function of coating microstructure and mechanical properties, which are controlled by the substrate growth temperature Ts. Our previous work has shown that Ag is highly mobile within the CrN matrix at elevated temperatures [9] and offers the potential for significant friction and wear reduction [10]. We choose for the present study a relatively high solid lubricant concentration of 22 at.%, corresponding to 34 vol.%, since this leads to Ag segregation at the growing surface and the formation of lamellar Ag aggregate inclusions of increasing size as a function of Ts [15]. We find that the development of an effective solid lubricating interface at room temperature depends
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strongly on both (i) the microstructure, described in terms of the Ag aggregate size and shape, and (ii) by the counterface material which is either 100Cr6 steel or alumina. Additionally, our initial tribological measurements during temperature ramps from 25 to 700 °C indicate that the Ag aggregate size and shape have a significant impact on the transient friction coefficient, which drops to <0.1 at 500 °C. 2. Experimental procedure CrN–Ag composite coatings were grown in a load-locked multichamber ultra-high vacuum (UHV) stainless steel dc dual magnetron sputter deposition system with a base pressure of 1.3 × 10− 7 Pa (1 × 10− 9 Torr). Water cooled 5 cm diameter Cr and Ag targets with purities of 99.95% and 99.99%, respectively, were positioned at 10.5 and 21.6 cm from the substrate at an angle of 45° with respect to the substrate surface normal. The substrates are 5 mm thick and 19 mm diameter metallographically polished 440C stainless steel disks (final polish with a 0.03 μm γ-alumina slurry) cleaned with successive rinses in ultrasonic baths of trichloroethane, acetone, and isopropanol and blown dry with dry N2. The substrates were mounted on a molybdenum holder using Pelco colloidal silver paste and inserted into the load-lock chamber for transport to the deposition chamber where they were heated with a resistive heater to the desired growth temperature Ts = 500, 600, or 700 °C. Pure N2 (99.999%) was further purified using a Micro Torr purifier and introduced through metering valves to reach a constant chamber pressure of 0.4 Pa (3 mTorr), which was measured using a capacitance manometer. Just prior to initiating deposition, the targets were sputter cleaned for 5 min while the substrate was covered with a protective shutter. Sputtering was carried out at a constant power of 450 W to the Cr target for pure CrN deposition at Ts = 600 °C. For deposition of the composites, constant powers of 450 W and 160 W were applied to the Cr and Ag targets, respectively, yielding deposition rates of 30 nm/min for CrN and 15 nm/min for Ag, as determined from cross-sectional micrographs of pure CrN and Ag layers. During deposition, the substrates were at a floating potential of −30 V and were continuously rotated about the polar axis with 50 rpm, in order to obtain optimal coating uniformity. The deposition temperature, including the contribution due to plasma heating, was measured using a pyrometer which was cross-calibrated with a thermocouple within the sample stage. The hardness H, modulus E, and resistance to plastic deformation in terms of the H3/E2 ratio [16–18], were measured on metallographically polished cross-sections (final polish with 0.03 µm γ-alumina slurry) using a Wilson–Tukon Microhardness Tester with a Knoop indenter. The elastic modulus was determined using the method developed by Marshall et al. [19], which has previously been applied to both hard [20,21] and soft [22] coating materials, and employs the expression: 0
0
b b b H ≈ = −α ; a E a0 a
ð1Þ
where b and b′ are the short diagonal of the Knoop indentation before and after elastic recovery, a and a′ are the long diagonal before and after elastic recovery, and α is an experimentally determined constant of 0.45. The change in a with elastic recovery is negligible, i.e. a = a′ [19]. Thus, E is directly determined from the difference in b′/a and b/a. Stated uncertainties for H and E correspond to±1 standard deviation. Room temperature tribological properties were measured by sliding 6 mm diameter 100Cr6 bearing steel and alumina balls with 7 and 15 GPa hardness, respectively, for 10,000 cycles in 35–45% relative humidity air, using a Nanovea Series ball-on-disk tribometer. Total sliding distances of 377 and 157 m were obtained for 12 mm and 5 mm diameter wear tracks against steel and alumina, respectively. A normal load of 15 N and sliding speed of 10 cm/s (159 rpm) was used against steel while a 5 N normal load and 5 cm/s (191 rpm) sliding
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speed was used against alumina. The lower load for the alumina counterface reduces the contact stress to a level comparable to that for the steel counterface, facilitating comparison between the two data sets. In addition, prior studies with alumina counterfaces also use a 5 N load [7,10]. These testing parameters, also listed in Table 1, give calculated values of the maximum Hertzian contact stress and initial Hertzian contact radius of 1.6 GPa and 0.07 mm for steel and 1.3 GPa and 0.04 mm for alumina. Friction coefficients were recorded continuously throughout testing. The wear scars were analyzed using a Nanovea ST400 non-contact optical profilometer. The total volume of material removed below the initial plane of the sample was determined using 3-D profilometer surface topography data from two segments of the wear track located 180° from each other. The two segments were 2.5 and 1 mm long for tests with steel and alumina, respectively, and exhibited wear volumes and wear track widths that agreed with each other to within ±5% for all coatings. The reported wear rates are based on the wear volumes obtained from the average of the two segments, extrapolated over the entire circular track. The surface composition and structure of the wear scars were determined by scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS), using a JEOL JSM-840A SEM with a Kevex Instruments model 2003 detector as well as a JEOL JSM-6330F FE-SEM. EDS linescans across the wear tracks were performed to evaluate microcompositional changes associated with tribological testing. These measurements were performed with an incident beam energy of 10 keV to increase surface sensitivity, which was estimated using [23]
1:67
R=
0:0276AE0 Z 0:89 ρ
;
ð2Þ
where R (μm) is the radius of a sphere centered on the specimen surface that represents the sampling volume, Z and A the average atomic number and weight (g/mol), E0 the incident beam energy in keV, and ρ the average density in g/cm3. The estimated maximum electron penetration depth for these coatings is 0.87 μm, with the major fraction calculated from simple geometrical considerations of the measured X-ray intensity originating from the top 300 nm. That is, these measurements provide a method to discern changes in the composition near the sample surfaces. Temperature dependent transient friction coefficients μt for CrN and CrN–Ag coatings were obtained from ball-on-disk tests against alumina with a continually increasing temperature of 10 °C/min from 25–400 °C and 5 °C/min from 400–700 °C. These transient temperature tests were completed with a wear track diameter of 5 mm, a constant sliding speed of 1 cm/s (38 rpm), and a constant load of 15 N to reduce wear track run-in time due to initial surface roughness effects. μt values for each 50 °C temperature interval were determined by averaging over ~ 190 and ~380 cycles for testing temperatures Tt below and above 400 °C, respectively.
Table 1 Tribological testing parameters for room temperature (RT) testing against 100Cr6 and Al2O3 counterfaces, and during a temperature ramp against Al2O3. Parameter
100Cr6, RT
Al2O3, RT
Al2O3, T ramp
Ball radius (mm) Ball hardness (GPa) Wear track radius (mm) Linear velocity (cm/s) Angular velocity (rad/s) Load (N) Hertzian pressure (GPa) Hertzian radius (mm)
6 7 6 10 16.7 15 1.6 0.07
6 15 2.5 5 20.0 5 1.3 0.04
6 15 2.5 1 4.0 15 3.9 0.06
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3. Results and discussion 3.1. Coating composition and structure Post-deposition EDS analyses and Rutherford back scattering (RBS) show that the layers have an average Ag concentration of 22 ± 0.8 at.%, which is in agreement with the expected composition from the deposition rate ratio. The CrN within the composite is stoichiometric, with a N:Cr ratio of 1.0, as determined by RBS and consistent with previous investigations using comparable deposition conditions [24–27]. All coatings are 5 μm thick and are fully dense (100 ± 3%), based on the mass-to-thickness ratio measured by a combination of SEM and RBS. The as-deposited root-mean-square surface roughness Rrms is 70 ± 10 nm for the pure CrN and 100 ± 20, 200 ± 20, and 1190 ± 70 nm for CrN–Ag composites with Ts = 500, 600, and 700 °C, respectively, as measured by profilometry. As previously reported [15], pure CrN exhibits a highly preferred 002 orientation with a columnar morphology, while the composites show a mixed grain orientation with Ag forming lamellae which disrupt the columnar morphology and are elongated parallel to the growing surface, as evident from the micrographs of polished crosssections shown in Fig. 1. These micrographs are obtained in the backscatter electron mode to show strong elemental contrast and are inverted so that the high-atomic mass Ag aggregates appear dark in the lighter CrN matrix. Fig. 1(a) is an image from a layer grown at Ts = 500 °C. It shows Ag aggregates with an average lateral width of 50 nm and an average height along the growth direction of 25 nm, yielding an estimated aggregate volume of ~5 × 104 nm3. We note that, despite the circular symmetric deposition conditions, not all Ag lamellae are necessarily equiaxed in the orthogonal in-plane directions. The micrograph in Fig. 1(b) has a 4× lower resolution and shows the Ag distribution for Ts = 600 °C. The aggregates are 6× wider and 4× higher, 300 and 100 nm, respectively, than for Ts = 500 °C. Raising the growth temperature further to Ts = 700 °C exacerbates Ag segregation and results in the formation of alternating Ag rich and Ag depleted zones within the CrN matrix. This is illustrated in Fig. 1 (c), which shows lamella that are 600 nm wide and 200 nm high on
Fig. 1. Cross-sectional backscatter electron micrographs from CrN–Ag coatings with 22 at.% Ag grown at (a) 500, (b) 600, and (c) 700 °C. The images are inverted so that the high-atomic-mass Ag rich regions appear as dark contrast.
average. The increase in Ag aggregate size and the decrease in homogeneity with Ts are attributed to the increase in surface adatom mobility that allows Ag to nucleate submicron precipitates on the growing layer surface. An additional effect of the increased Ag segregation at Ts ≥ 600 °C is the development of nodular surface defects and non-uniform Ag distributions which are most dominant for the highest growth temperature of 700 °C [15]. As discussed in the following, the dramatic increase in the Ag aggregate size by more than three orders of magnitudes, from 5 × 104 to 9 × 106 to 7 × 107 nm3 for Ts = 500 to 600 to 700 °C, respectively, strongly affects tribological properties at both room and elevated temperatures. 3.2. Tribological properties against 100Cr6 steel Fig. 2 is a plot of the friction coefficient μ, determined by room temperature ball-on-disk measurements, for CrN–Ag nanocomposite coatings deposited on stainless steel substrates at Ts = 500, 600, and 700 °C. The plot also includes the friction data for a pure CrN sample deposited at 600 °C. The pure CrN specimen displays an initial friction coefficient of ~0.7 which drops to an average of 0.64± 0.01 after a 1000 cycle running-in period, in close agreement with reported values of μ = 0.62 for CrN with a 100Cr6 steel counterface [28]. All CrN–Ag composite samples exhibit an initial room temperature friction coefficient of ~0.5. This is followed by an increase and subsequent decrease to μ = 0.58 ± 0.02 for Ts = 500 °C, a monotonic decrease to μ = 0.47 ± 0.02 for Ts = 600 °C, and an increase to μ = 0.85 ± 0.04 for Ts = 700 °C. These μ values correspond to the average over the final 5000 cycles, while the stated error bars are the standard deviation from the fluctuations over the same (5000–10,000) cycle interval. Optical profilometry surface images for samples tested against steel are shown in Fig. 3. The surface morphologies of the pure CrN and the Ts = 500 °C nanocomposite layers are comparable, showing evidence of grooves and intermittent pitting along with associated jagged surface asperities. They also show similar wear rates of 3.6 × 10− 6 and 3.8 × 10− 6 mm3/Nm, respectively, despite the greater H3/E2 ratio of 0.19 GPa for the pure CrN versus 0.09 GPa for the Ts = 500 °C composite, as given in Table 2. In contrast, the Ts = 600 °C coating exhibits a 16× lower wear, 2.3 × 10− 7 mm3/Nm, and only limited roughening which is likely due to material transfer from the steel counterface to the coating surface. This lower wear rate is observed despite the lower H3/E2 for the composite versus the pure CrN, indicating that the presence of Ag plays the primary role in determining friction coefficient and wear rate. The coating deposited at 700 °C shows the largest damage, the highest wear rate of 9.9 × 10− 6 mm3/Nm, despite a higher hardness of 24 GPa and
Fig. 2. Friction coefficient μ vs number of cycles during ball-on-disk testing against 100Cr6 steel for pure CrN deposited at Ts = 600 °C and for CrN–Ag composites with 22 at.% Ag and Ts = 500, 600 and 700 °C.
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Fig. 3. 3-D surface profiles of a pure CrN coating and CrN–Ag composite coatings with Ts = 500, 600 and 700 °C, after testing against a 100Cr6 steel ball.
H3/E2 ratio of 0.19, and the surface exhibits a highly pitted, rough, and irregular surface. The surfaces roughness is attributed to the combination of material transfer from the steel counterface and the as-deposited roughness, which increases significantly from root-mean-square roughness values Rrms = 70 and 100 nm for the pure CrN and the Ts = 500 °C composite, to 200 and 1190 nm for the Ts = 600 and 700 °C composites, respectively. The measured wear rates, also presented in Table 2, should be treated as lower bound values, since they are determined from the wear volume below the initial plane of the sample surface and do not account for negative wear associated with material transfer from the counterface. In addition to the wearing of the coatings, circular wear scars develop on the steel ball counterfaces. The diameters of the wear scars on the steel correspond to the measured track width on the coating surfaces. They are 1.11 mm for pure CrN, and 1.02, 0.89, and 1.25 mm for the nanocomposites with Ts = 500, 600, and 700 °C, respectively. This corresponds to a 60% lower steel wear rate for the Ts = 600 °C composite vs. the pure CrN coating, based on spherical cap volume calculations as
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shown in Table 2. The reduced wear track widths and associated counterface wear scar diameters for Ts = 500 and 600 °C are concurrent with the lower friction coefficients μ = 0.58 and 0.47 in comparison to pure CrN, μCrN = 0.62, respectively, and are attributed to the presence of Ag as well as the lower coating hardness of 17 and 20 GPa, respectively, which are considerably lower than the value for pure CrN of 28 GPa. The counterface wear rate is particularly low for the Ts = 600 °C composite, which exhibits a lower friction than both the pure CrN and Ts = 500 °C composite. As the Ts = 600 °C composite is slightly harder than the Ts = 500 °C composite, this indicates that the friction coefficient is the primary factor in determining wear response. Evidence for the formation of an effective lubricious Ag layer on the Ts = 600 °C composite surface is shown in Fig. 4. The plot shows the Ag concentration, as determined from EDS linescans across the wear tracks. For Ts = 500 °C, the Ag concentration remains essentially constant across the track, fluctuating in the 18–25 at.% range with an average value of 22 at.%. This corresponds to the bulk coating composition and, in turn, indicates that CrN and Ag are worn uniformly. In contrast, the Ag concentration in the wear track of the Ts = 600 °C sample exhibits multiple peaks to 30–40 at.% Ag, resulting in an average concentration of 25 at.% which is higher than the 22 at.% measured outside of the wear track. This may be due the lamellar shaped Ag aggregates, which are well developed at Ts = 600 °C and smear along the surface, yielding a soft Ag rich solid lubricating interface which, in turn, results in the low μ = 0.47 ± 0.02. The sample with Ts = 700 °C exhibits a constant Ag intensity (not shown), comparable to the Ts = 500 °C result. All these line-scan results, and in particular the Ag increase for Ts = 600 °C, are confirmed by wide area EDS measurements in multiple locations within and outside the tracks. The wide area measurements also indicate that all samples exhibit comparable patches of Fe deposits within the wear track due to the counterface wear and material transfer. The effect of Ts on the room temperature friction coefficient and wear rate is attributed to the orientation and structure of the Ag within the nanocomposite, as discussed in the following. The Ts = 500 °C coating exhibits friction and wear that is comparable to pure CrN, indicating that the very fine grained Ag within the matrix fails to form a continuous lubricious interface, resulting in a negligible effect of the Ag and tribological properties that resemble those for pure CrN. In contrast, Ts = 600 °C leads to a nanocomposite with a 25% lower friction and 16× lower wear than pure CrN. This suggests that the distributed lamellar Ag inclusions in this coating are of suitable size and orientation relative to the surface to provide an effective solid lubricating interface. This lubrication is effective despite the 3× higher
Table 2 Microstructural, mechanical, and tribological properties at room temperature. Sample
Pure CrNa
Ts = 500 °C
Ts = 600 °C
Ts = 700 °C
Ag agg. size (w h, nm) Roughness, Rrms (nm) Hardness (GPa) Elastic modulus (GPa) H3/E2 (GPa)
– 70 ± 10 28 ± 2 340 ± 80 0.19
50 × 25 100 ± 20 17 ± 2 240 ± 40 0.09
300 × 100 200 ± 20 20 ± 3 240 ± 60 0.14
600 × 200 1190 ± 70 24 ± 4 270 ± 70 0.19
100Cr6 steel Friction coefficient—μ Coating wear (mm3/Nm) Ball wear (mm3/Nm)
0.64 3.6 × 10− 6 4.4 × 10− 6
0.58 3.8 × 10− 6 3.1 × 10− 6
0.47 2.3 × 10− 7 1.8 × 10− 6
0.85 9.9 × 10− 6 7.1 × 10− 6
Alumina Friction coefficient—μ Coating wear (mm3/Nm) Ball wear (mm3/Nm)
0.59 4.7 × 10− 6 –
0.50 7.0 × 10− 6 –
0.69 8.0 × 10− 6 –
0.8 ~ 10− 4 –
a
Pure CrN deposited with Ts = 600 °C.
Fig. 4. Plots of the Ag concentration versus position, obtained from EDS line scans across the wear tracks for Ts = 500 and 600 °C composites tested against steel.
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initial surface roughness of the composite coating. However, raising the deposition temperature to Ts = 700 °C leads to composites with significantly higher friction and wear than for pure CrN. We attribute this to the large segregated Ag grains that cause structural inhomogeneity and a dramatic 17× increase in surface roughness due to the development of nodular defects [15]. That is, the material is worn too easily, leaving behind a roughened surface and generating a significant amount of wear debris. 3.3. Tribological properties against alumina The plot in Fig. 5 shows the room temperature friction coefficient of samples tested against 6 mm diameter alumina balls. Pure CrN exhibits a μ that increases from 0.4 to 0.5 to reach an average value of 0.59±0.03. This is slightly lower than μ=0.64 for the softer steel counterface, and slightly higher than reported CrN values with alumina counterfaces of 0.37, 0.52, and 0.55 at normal loads of 5, 5, and 2 N for Refs. [7,10,28], respectively. The composite coatings exhibit room temperature friction coefficients that increase from 0.3 to a steady-state value of μ=0.50±0.03 for Ts =500 °C, rise rapidly from an initial μ=0.4 and then fluctuate around an average value of 0.69±0.06 for Ts =600 °C, and oscillate around μ~0.8 prior to coating failure after ~1000cycles for Ts =700 °C. Fig. 6 shows corresponding optical profilometry results. The wear tracks conform to the spherical shape of the alumina ball, indicating relatively little wear of the alumina counterface material. The measured wear rate is lowest for the pure CrN coating, 4.7×10− 6 mm3/Nm, and increases to 7.0×10− 6, 8.0×10− 6, and ~10-4 mm3/Nm for nanocomposites with Ts =500, 600, and 700 °C, respectively. The CrN–Ag coating grown at Ts =500 °C exhibits a slightly lower friction but a higher wear rate than pure CrN. We attribute this to the Ag inclusions which affect the mechanical behavior of the composite coating, with H and H3/E2 decreasing from 28 and 0.19 GPa for CrN to 17 and 0.09 GPa for CrN–Ag. This softening of the coating and relative reduction in resistance to plastic deformation increases wear and only slightly reduces friction without forming an effective and continuous lubricating interface against the hard alumina counterface. The Ts =600 °C coating also exhibits higher room temperature friction and wear against alumina than pure CrN. This is in stark contrast to the results with a steel counterface. The dramatic dependence on the counterface material is attributed primarily to the difference in counterface wear: For the steel counterface, the ball wears and conforms to the surface, forming a flat on flat type geometry with reduced contact stress due to load spread over a much wider area, which is effectively lubricated by the lamellar Ag aggregates that are expected to shear along the surface as discussed in the prior section. In contrast, the considerably harder alumina ball experi-
Fig. 6. 3-D surface profiles of a pure CrN coating and CrN–Ag composite coatings with Ts = 500, 600 and 700 °C, after testing against an alumina counterface.
ences negligible wear and does not conform to the coating surface, resulting in plowing and uniform wear of both CrN and Ag phases. We attribute the higher μ, in comparison to pure CrN, to higher roughness from localized nodular surface asperities [15] and the associated formation of brittle wear debris from deformation and fracture of these asperities. The argument of uniform wear of CrN and Ag is supported by EDS measurements (not shown), which indicate no significant change in Ag concentration in the wear track for sliding against alumina, which is opposite to the result observed for sliding against steel in Fig. 4.The rapid tribological failure for Ts =700 °C is attributed to fracture of the large, 1–5 μm high nodular defects along the rough surface which likely lead to brittle failure of the coating. The differences in behavior for the steel and alumina counterfaces against Ts = 600 and 700 °C composites, wherein the composites exhibit large nodular surface asperities with average asperity tip radii Rs of 3.2 and 2.3 μm and average asperity heights of 0.4 and 2.5 μm, as measured by profilometry, respectively, can be further explained using the plasticity index Ψ, which is a dimensionless quantity that describes deformation behavior independent of normal load and sliding speed. In contrast to H3/E2, in which higher numbers represent an increased ability of a contact to resist plastic deformation on the macroscale of the planar surface and ball counterface, the plasticity index describes the deformation behavior on the microscale of surface asperities in terms of the likelihood of a contact to experience plastic deformation, with higher numbers indicating a greater likelihood of plastic deformation. That is, for Ψ ≪ 1, deformation is expected to be entirely elastic, while Ψ ≫ 1 leads to plastic deformation. The plasticity index is obtained using [29]: Ψ=
Fig. 5. Friction coefficient μ vs number of cycles during ball-on-disk testing against alumina for pure CrN and for CrN–Ag composites with Ts = 500, 600 and 700 °C.
E* H
rffiffiffiffiffi σs ; Rs
ð3Þ
where σs is the composite standard deviation of asperity peak heights which we estimate with the root-mean-square surface roughness of
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the composite surfaces, 200 and 1190 nm for Ts = 600 and 700 °C, respectively. E⁎ is the contact modulus defined by the elastic moduli and Poisson's ratios of the two materials in contact according to
E* =
1−ν21 1−ν22 + E1 E2
!−1 ;
ð4Þ
and is determined using the elastic moduli of the composites given in Table 2, an estimated composite Poisson's ratio of 0.23 [30], and elastic moduli and Poisson's ratios of 210 GPa and 0.3 for steel and 375 GPa and 0.22 for alumina [31], respectively. We calculate Ψ = 4.3 and 1.5 for counterface and coating, respectively, for the case of a steel counterface against the Ts = 600 °C composite, and obtain Ψ = 13.1 and 3.8 for the much rougher Ts = 700 °C composite. For both growth temperatures, all Ψ are above unity and the Ψ values for the steel counterfaces are considerably higher than for the CrN–Ag surfaces, indicating that plastic deformation of the softer steel counterfaces dominate in relieving the contact stresses at asperity tips. This is in agreement with the observed wear of the steel counterface material. For the alumina counterface, the situation is complicated by the closer relative hardness of coating and counterface, resulting in comparable plasticity index values of 2.6:1.9 and 7.9:5.0 for alumina:CrN–Ag at Ts = 600 and 700 °C, respectively. For the alumina counterface, the high plasticity index for Ts = 700 °C is in agreement with the observed severe wear, while the lower but comparable numbers between alumina and CrN–Ag for Ts = 600 °C suggest a likely simultaneous onset of plastic deformation for both coating and counterface. Due to the brittle behavior of ceramics such as CrN, at ambient temperature, onset of yielding would likely lead to fracture of large surface asperities and the generation of wear debris leading to the higher friction and wear observed for Ts = 600 °C and most extensively for Ts = 700 °C, which exhibits sharper asperity radii and larger asperity heights. 3.4. Tribological response as a function of temperature In order to determine temperatures at which the self-lubricating behavior of these films is most prevalent, the transient tribological response as a function of temperature was analyzed. Fig. 7 is a plot of the average transient friction coefficient μt during a continuous temperature ramp from Tt = 25–700 °C. The μt for pure CrN decreases from 0.65 to 0.26 for Tt = 100 to 300 °C, remains relatively constant at
Fig. 7. Transient friction coefficients μt measured against alumina during a continuous temperature ramp from Tt = 25–700 °C, for pure CrN and for CrN–Ag composites deposited at Ts = 500, 600 and 700 °C.
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μ t = 0.25 for 300 < T t < 500 °C, and increases to μ t = 0.55 for Tt ≥ 600 °C. The CrN–Ag coating with Ts = 500 °C exhibits a nearly identical friction as pure CrN for Tt ≤ 300 °C. However, μt continues to drop for Tt > 300 °C to reach a minimum of 0.05 at Tt = 500 °C, followed by a rise to 0.33 for Tt ≥ 600 °C. The Ts = 600 °C coating shows a qualitatively similar curve as for Ts = 500 °C with, however, a consistently higher μt that decreases from 0.64 at Tt = 100 °C to a minimum of 0.34 at Tt = 500 °C and reaches 0.58 for Tt ≥ 600 °C. The Ts = 700 °C coating exhibits rapid failure after ~ 1000 cycles and at Tt = 200–250 °C, as expected from room temperature testing. At low temperatures (Tt < 300 °C), μt decreases with increasing Tt for all samples. This is attributed to a drying of the environment along with an effective softening of the coatings due to an increasing dislocation mobility in CrN and Ag as well as a reduction in shear strength of Ag, and a reduction in the hardness of alumina which decreases from 15 GPa at room temperature to 8.5 GPa at 500 °C [32]. The continued drop in μt at Tt > 300 °C for Ag containing coatings indicates the presence of an increasingly effective lubricious Ag surface layer. The formation of this Ag surface layer with increasing Tt > 300 °C is attributed to a temperature activated solid lubricant transport to the surface, which we have previously demonstrated to occur at elevated temperatures [9]. The very low transient friction (μt = 0.05) for Ts = 500 °C at Tt = 500 °C is comparable to reported values for steady-state μ < 0.05 for ultra-thin (<0.2 μm) Ag films tested at room temperature in vacuum [33] and μ < 0.1 for thin (~2 μm) Ag films tested at 570 °C in an inert atmosphere [14]. Similar to the room temperature results against alumina, the friction for Ts = 600 °C is higher than for Ts = 500 °C. This is attributed to the inability of the Ag to effectively lubricate the 3× rougher Ts = 600 °C surface against the non-conforming alumina ball. The increase in μt above Tt = 500 °C for all samples is likely due to oxidative degradation similar to what has been previously reported for CrN [10,28]. These results suggest, consistent with the room temperature data, that the as-deposited Ag aggregate size and, in turn, the lubricant transport and the surface morphology, strongly affect the high-temperature tribological performance of CrN–Ag nanocomposite coatings. 4. Conclusions Tribological characterization of reactive magnetron sputtered CrN and CrN–Ag nanocomposite coatings against both 100Cr6 steel and alumina indicate that the nanocomposite structure at fixed composition strongly affects friction and wear. For the steel counterface, pure CrN and composites grown at 500 °C show comparable room temperature friction and wear, while Ts = 600 °C composites exhibit a 25% reduced friction and a 16 fold reduced wear rate in comparison to pure CrN. This is attributed to the formation of a lubricious Ag surface layer, as evidenced by EDS analyses, which is facilitated by the orientation and structure of the fine Ag lamellae within the CrN matrix for Ts = 600 °C. The wear increases considerably for the Ts = 700 °C, which is attributed to rough surfaces exacerbated by the formation of large nodules. For tribological tests against alumina, μ is found to increase with surface roughness with the lowest value observed for the CrN–Ag composite deposited at Ts = 500 °C, which exhibits a comparable roughness as the pure CrN coating. Wear rates for the CrN–Ag composites are a factor of 1.5–2 higher than for pure CrN, for both Ts = 500 and 600 °C. This higher wear rate is attributed to a reduction in overall hardness and H3/E2 ratio due to the addition of Ag as well as the inability of the Ag to form an effective lubricating interface against the harder alumina counterface, and for Ts = 600 °C also the higher initial roughness. The most promising benefits for CrN–Ag composite coatings are expected at elevated temperatures. Tribological testing of the coatings over a continuous temperature range of Tt = 25–700 °C indicates significant friction reductions for Ag containing composites with the minimum transient friction coefficient of 0.05 occurring at a transient test temperature of 500 °C for the finest grained composite. The
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Acknowledgments
[12] [13] [14] [15] [16] [17] [18]
This research was supported by the United States Army TEX3 Program, through the Armament Research, Development and Engineering Center (ARDEC), and by the National Science Foundation under grant no. CMMI-0653843.
[19] [20] [21] [22] [23]
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Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
CrN–Ag nanocomposite coatings: Tribology at room temperature and during a temperature ramp C.P. Mulligan a,c,⁎, T.A. Blanchet b, D. Gall a a b c
Department of Materials Science and Engineering, Rensselaer Polytechnic Institute, Troy, NY 12180, United States Department of Mechanical, Aerospace, and Nuclear Engineering, Rensselaer Polytechnic Institute, Troy, NY 12180, United States Benet Laboratories, U.S. Army Armament Research Development & Engineering Center, Watervliet, NY 12189, United States
a r t i c l e
i n f o
Article history: Received 1 June 2009 Accepted in revised form 8 September 2009 Available online 18 September 2009 Keywords: Nanocomposite coating Friction Wear Solid lubrication High-temperature lubrication CrN–Ag
a b s t r a c t 5 µm thick CrN–Ag composite layers with 22 at.% Ag were deposited by reactive magnetron co-sputtering on 440C stainless steel substrates. Increasing the growth temperature from Ts = 500 to 600 to 700 °C leads to Ag segregation within the CrN matrix and the formation of embedded lamellar Ag aggregates with increasing size, < 105, 9 × 106, and 7 × 107 nm3, respectively. Ball-on-disk tests against 100Cr6 steel, followed by optical profilometry and energy dispersive spectroscopy, indicate that the Ag grains for Ts = 500 °C are too small to facilitate an effective lubricious surface layer, resulting in a friction coefficient μ = 0.58 and a composite coating wear rate of 3.8 × 10− 6 mm3/Nm that are nearly identical to those measured for pure CrN with μ = 0.64 and 3.6 × 10− 6 mm3/Nm. The Ts = 600 °C coating exhibits a Ag concentration which is 15% higher within than outside the wear track, and acts as a lubricious layer that reduces μ to 0.47 and yields a 16× and 2.4× lower wear rate for coating and counterface, respectively. Ts = 700 °C leads to a dramatic increase in surface roughness and an associated increase in friction, μ = 0.85, and wear, 9.9 × 10− 6 mm3/Nm. Replacing the steel counterface with an alumina ball results in the lowest μ = 0.50 for Ts = 500 °C, attributed to the presence of Ag and the relatively low hardness of 16.5 GPa for this particular coating. In contrast, friction and wear increase dramatically for Ts = 600 °C, which is attributed to a breakdown of the lubricious Ag layer by the harder counterface. The transient friction coefficient μt during experiments with continuously increasing testing temperature Tt = 25–700 °C initially decreases for all samples, attributed to drying of the environment and an effective softening of both coating and counterface. For the Ts = 500 °C coating, a temperature activated solid lubricant transport yields a lubricious Ag surface layer and a very low μt = 0.05 at Tt ~ 500 °C. All coatings exhibit an increasing μt for Tt > 500 °C, which is attributed to oxidative degradation. Published by Elsevier B.V.
1. Introduction Multi-functional coatings using soft lubricating phases within a hard wear-resistant matrix have been studied extensively in recent years due to their promise for tribological performance in transient and cyclic environments [1–10]. In particular, inclusions of noble metals such as Ag and Au show promise as solid lubricating phases in carbide [3], oxide [4–6], and nitride [7–10] matrices, as they possess sufficiently low shear strength over a suitably wide temperature range as well as stable thermochemistry allowing them to be used in both ambient air and vacuum [11–14]. Most studies on co-deposited nanocomposites have focused on the tribological properties as a function of noble metal content, with for example CrAlN–Ag having a promising room temperature friction coefficient as low as 0.23 and wear rate as low as 3.0×10− 8 mm3/Nm at a 2 N load, which depends on the solid lubricant content [7]. For this study, ⁎ Corresponding author. Department of Materials Science and Engineering, Rensselaer Polytechnic Institute, Troy, NY 12180, United States. E-mail address: [email protected] (C.P. Mulligan). 0257-8972/$ – see front matter. Published by Elsevier B.V. doi:10.1016/j.surfcoat.2009.09.018
we focus on the effect of the composite microstructure on the tribological properties for a fixed composition. We expect that the microstructure, including lubricant phase, grain size and shape, as well as matrix porosity strongly affect both solid lubricant transport and the development of an effective solid lubricating interface. Consequently, understanding the relationship between microstructure and tribological properties is critical to the design of appropriate coating architectures that provide lubrication and wear protection for a given application. In this report, the tribological performance of a prototype CrN–Ag nanocomposite coating is investigated as a function of coating microstructure and mechanical properties, which are controlled by the substrate growth temperature Ts. Our previous work has shown that Ag is highly mobile within the CrN matrix at elevated temperatures [9] and offers the potential for significant friction and wear reduction [10]. We choose for the present study a relatively high solid lubricant concentration of 22 at.%, corresponding to 34 vol.%, since this leads to Ag segregation at the growing surface and the formation of lamellar Ag aggregate inclusions of increasing size as a function of Ts [15]. We find that the development of an effective solid lubricating interface at room temperature depends
C.P. Mulligan et al. / Surface & Coatings Technology 204 (2010) 1388–1394
strongly on both (i) the microstructure, described in terms of the Ag aggregate size and shape, and (ii) by the counterface material which is either 100Cr6 steel or alumina. Additionally, our initial tribological measurements during temperature ramps from 25 to 700 °C indicate that the Ag aggregate size and shape have a significant impact on the transient friction coefficient, which drops to <0.1 at 500 °C. 2. Experimental procedure CrN–Ag composite coatings were grown in a load-locked multichamber ultra-high vacuum (UHV) stainless steel dc dual magnetron sputter deposition system with a base pressure of 1.3 × 10− 7 Pa (1 × 10− 9 Torr). Water cooled 5 cm diameter Cr and Ag targets with purities of 99.95% and 99.99%, respectively, were positioned at 10.5 and 21.6 cm from the substrate at an angle of 45° with respect to the substrate surface normal. The substrates are 5 mm thick and 19 mm diameter metallographically polished 440C stainless steel disks (final polish with a 0.03 μm γ-alumina slurry) cleaned with successive rinses in ultrasonic baths of trichloroethane, acetone, and isopropanol and blown dry with dry N2. The substrates were mounted on a molybdenum holder using Pelco colloidal silver paste and inserted into the load-lock chamber for transport to the deposition chamber where they were heated with a resistive heater to the desired growth temperature Ts = 500, 600, or 700 °C. Pure N2 (99.999%) was further purified using a Micro Torr purifier and introduced through metering valves to reach a constant chamber pressure of 0.4 Pa (3 mTorr), which was measured using a capacitance manometer. Just prior to initiating deposition, the targets were sputter cleaned for 5 min while the substrate was covered with a protective shutter. Sputtering was carried out at a constant power of 450 W to the Cr target for pure CrN deposition at Ts = 600 °C. For deposition of the composites, constant powers of 450 W and 160 W were applied to the Cr and Ag targets, respectively, yielding deposition rates of 30 nm/min for CrN and 15 nm/min for Ag, as determined from cross-sectional micrographs of pure CrN and Ag layers. During deposition, the substrates were at a floating potential of −30 V and were continuously rotated about the polar axis with 50 rpm, in order to obtain optimal coating uniformity. The deposition temperature, including the contribution due to plasma heating, was measured using a pyrometer which was cross-calibrated with a thermocouple within the sample stage. The hardness H, modulus E, and resistance to plastic deformation in terms of the H3/E2 ratio [16–18], were measured on metallographically polished cross-sections (final polish with 0.03 µm γ-alumina slurry) using a Wilson–Tukon Microhardness Tester with a Knoop indenter. The elastic modulus was determined using the method developed by Marshall et al. [19], which has previously been applied to both hard [20,21] and soft [22] coating materials, and employs the expression: 0
0
b b b H ≈ = −α ; a E a0 a
ð1Þ
where b and b′ are the short diagonal of the Knoop indentation before and after elastic recovery, a and a′ are the long diagonal before and after elastic recovery, and α is an experimentally determined constant of 0.45. The change in a with elastic recovery is negligible, i.e. a = a′ [19]. Thus, E is directly determined from the difference in b′/a and b/a. Stated uncertainties for H and E correspond to±1 standard deviation. Room temperature tribological properties were measured by sliding 6 mm diameter 100Cr6 bearing steel and alumina balls with 7 and 15 GPa hardness, respectively, for 10,000 cycles in 35–45% relative humidity air, using a Nanovea Series ball-on-disk tribometer. Total sliding distances of 377 and 157 m were obtained for 12 mm and 5 mm diameter wear tracks against steel and alumina, respectively. A normal load of 15 N and sliding speed of 10 cm/s (159 rpm) was used against steel while a 5 N normal load and 5 cm/s (191 rpm) sliding
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speed was used against alumina. The lower load for the alumina counterface reduces the contact stress to a level comparable to that for the steel counterface, facilitating comparison between the two data sets. In addition, prior studies with alumina counterfaces also use a 5 N load [7,10]. These testing parameters, also listed in Table 1, give calculated values of the maximum Hertzian contact stress and initial Hertzian contact radius of 1.6 GPa and 0.07 mm for steel and 1.3 GPa and 0.04 mm for alumina. Friction coefficients were recorded continuously throughout testing. The wear scars were analyzed using a Nanovea ST400 non-contact optical profilometer. The total volume of material removed below the initial plane of the sample was determined using 3-D profilometer surface topography data from two segments of the wear track located 180° from each other. The two segments were 2.5 and 1 mm long for tests with steel and alumina, respectively, and exhibited wear volumes and wear track widths that agreed with each other to within ±5% for all coatings. The reported wear rates are based on the wear volumes obtained from the average of the two segments, extrapolated over the entire circular track. The surface composition and structure of the wear scars were determined by scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS), using a JEOL JSM-840A SEM with a Kevex Instruments model 2003 detector as well as a JEOL JSM-6330F FE-SEM. EDS linescans across the wear tracks were performed to evaluate microcompositional changes associated with tribological testing. These measurements were performed with an incident beam energy of 10 keV to increase surface sensitivity, which was estimated using [23]
1:67
R=
0:0276AE0 Z 0:89 ρ
;
ð2Þ
where R (μm) is the radius of a sphere centered on the specimen surface that represents the sampling volume, Z and A the average atomic number and weight (g/mol), E0 the incident beam energy in keV, and ρ the average density in g/cm3. The estimated maximum electron penetration depth for these coatings is 0.87 μm, with the major fraction calculated from simple geometrical considerations of the measured X-ray intensity originating from the top 300 nm. That is, these measurements provide a method to discern changes in the composition near the sample surfaces. Temperature dependent transient friction coefficients μt for CrN and CrN–Ag coatings were obtained from ball-on-disk tests against alumina with a continually increasing temperature of 10 °C/min from 25–400 °C and 5 °C/min from 400–700 °C. These transient temperature tests were completed with a wear track diameter of 5 mm, a constant sliding speed of 1 cm/s (38 rpm), and a constant load of 15 N to reduce wear track run-in time due to initial surface roughness effects. μt values for each 50 °C temperature interval were determined by averaging over ~ 190 and ~380 cycles for testing temperatures Tt below and above 400 °C, respectively.
Table 1 Tribological testing parameters for room temperature (RT) testing against 100Cr6 and Al2O3 counterfaces, and during a temperature ramp against Al2O3. Parameter
100Cr6, RT
Al2O3, RT
Al2O3, T ramp
Ball radius (mm) Ball hardness (GPa) Wear track radius (mm) Linear velocity (cm/s) Angular velocity (rad/s) Load (N) Hertzian pressure (GPa) Hertzian radius (mm)
6 7 6 10 16.7 15 1.6 0.07
6 15 2.5 5 20.0 5 1.3 0.04
6 15 2.5 1 4.0 15 3.9 0.06
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3. Results and discussion 3.1. Coating composition and structure Post-deposition EDS analyses and Rutherford back scattering (RBS) show that the layers have an average Ag concentration of 22 ± 0.8 at.%, which is in agreement with the expected composition from the deposition rate ratio. The CrN within the composite is stoichiometric, with a N:Cr ratio of 1.0, as determined by RBS and consistent with previous investigations using comparable deposition conditions [24–27]. All coatings are 5 μm thick and are fully dense (100 ± 3%), based on the mass-to-thickness ratio measured by a combination of SEM and RBS. The as-deposited root-mean-square surface roughness Rrms is 70 ± 10 nm for the pure CrN and 100 ± 20, 200 ± 20, and 1190 ± 70 nm for CrN–Ag composites with Ts = 500, 600, and 700 °C, respectively, as measured by profilometry. As previously reported [15], pure CrN exhibits a highly preferred 002 orientation with a columnar morphology, while the composites show a mixed grain orientation with Ag forming lamellae which disrupt the columnar morphology and are elongated parallel to the growing surface, as evident from the micrographs of polished crosssections shown in Fig. 1. These micrographs are obtained in the backscatter electron mode to show strong elemental contrast and are inverted so that the high-atomic mass Ag aggregates appear dark in the lighter CrN matrix. Fig. 1(a) is an image from a layer grown at Ts = 500 °C. It shows Ag aggregates with an average lateral width of 50 nm and an average height along the growth direction of 25 nm, yielding an estimated aggregate volume of ~5 × 104 nm3. We note that, despite the circular symmetric deposition conditions, not all Ag lamellae are necessarily equiaxed in the orthogonal in-plane directions. The micrograph in Fig. 1(b) has a 4× lower resolution and shows the Ag distribution for Ts = 600 °C. The aggregates are 6× wider and 4× higher, 300 and 100 nm, respectively, than for Ts = 500 °C. Raising the growth temperature further to Ts = 700 °C exacerbates Ag segregation and results in the formation of alternating Ag rich and Ag depleted zones within the CrN matrix. This is illustrated in Fig. 1 (c), which shows lamella that are 600 nm wide and 200 nm high on
Fig. 1. Cross-sectional backscatter electron micrographs from CrN–Ag coatings with 22 at.% Ag grown at (a) 500, (b) 600, and (c) 700 °C. The images are inverted so that the high-atomic-mass Ag rich regions appear as dark contrast.
average. The increase in Ag aggregate size and the decrease in homogeneity with Ts are attributed to the increase in surface adatom mobility that allows Ag to nucleate submicron precipitates on the growing layer surface. An additional effect of the increased Ag segregation at Ts ≥ 600 °C is the development of nodular surface defects and non-uniform Ag distributions which are most dominant for the highest growth temperature of 700 °C [15]. As discussed in the following, the dramatic increase in the Ag aggregate size by more than three orders of magnitudes, from 5 × 104 to 9 × 106 to 7 × 107 nm3 for Ts = 500 to 600 to 700 °C, respectively, strongly affects tribological properties at both room and elevated temperatures. 3.2. Tribological properties against 100Cr6 steel Fig. 2 is a plot of the friction coefficient μ, determined by room temperature ball-on-disk measurements, for CrN–Ag nanocomposite coatings deposited on stainless steel substrates at Ts = 500, 600, and 700 °C. The plot also includes the friction data for a pure CrN sample deposited at 600 °C. The pure CrN specimen displays an initial friction coefficient of ~0.7 which drops to an average of 0.64± 0.01 after a 1000 cycle running-in period, in close agreement with reported values of μ = 0.62 for CrN with a 100Cr6 steel counterface [28]. All CrN–Ag composite samples exhibit an initial room temperature friction coefficient of ~0.5. This is followed by an increase and subsequent decrease to μ = 0.58 ± 0.02 for Ts = 500 °C, a monotonic decrease to μ = 0.47 ± 0.02 for Ts = 600 °C, and an increase to μ = 0.85 ± 0.04 for Ts = 700 °C. These μ values correspond to the average over the final 5000 cycles, while the stated error bars are the standard deviation from the fluctuations over the same (5000–10,000) cycle interval. Optical profilometry surface images for samples tested against steel are shown in Fig. 3. The surface morphologies of the pure CrN and the Ts = 500 °C nanocomposite layers are comparable, showing evidence of grooves and intermittent pitting along with associated jagged surface asperities. They also show similar wear rates of 3.6 × 10− 6 and 3.8 × 10− 6 mm3/Nm, respectively, despite the greater H3/E2 ratio of 0.19 GPa for the pure CrN versus 0.09 GPa for the Ts = 500 °C composite, as given in Table 2. In contrast, the Ts = 600 °C coating exhibits a 16× lower wear, 2.3 × 10− 7 mm3/Nm, and only limited roughening which is likely due to material transfer from the steel counterface to the coating surface. This lower wear rate is observed despite the lower H3/E2 for the composite versus the pure CrN, indicating that the presence of Ag plays the primary role in determining friction coefficient and wear rate. The coating deposited at 700 °C shows the largest damage, the highest wear rate of 9.9 × 10− 6 mm3/Nm, despite a higher hardness of 24 GPa and
Fig. 2. Friction coefficient μ vs number of cycles during ball-on-disk testing against 100Cr6 steel for pure CrN deposited at Ts = 600 °C and for CrN–Ag composites with 22 at.% Ag and Ts = 500, 600 and 700 °C.
C.P. Mulligan et al. / Surface & Coatings Technology 204 (2010) 1388–1394
Fig. 3. 3-D surface profiles of a pure CrN coating and CrN–Ag composite coatings with Ts = 500, 600 and 700 °C, after testing against a 100Cr6 steel ball.
H3/E2 ratio of 0.19, and the surface exhibits a highly pitted, rough, and irregular surface. The surfaces roughness is attributed to the combination of material transfer from the steel counterface and the as-deposited roughness, which increases significantly from root-mean-square roughness values Rrms = 70 and 100 nm for the pure CrN and the Ts = 500 °C composite, to 200 and 1190 nm for the Ts = 600 and 700 °C composites, respectively. The measured wear rates, also presented in Table 2, should be treated as lower bound values, since they are determined from the wear volume below the initial plane of the sample surface and do not account for negative wear associated with material transfer from the counterface. In addition to the wearing of the coatings, circular wear scars develop on the steel ball counterfaces. The diameters of the wear scars on the steel correspond to the measured track width on the coating surfaces. They are 1.11 mm for pure CrN, and 1.02, 0.89, and 1.25 mm for the nanocomposites with Ts = 500, 600, and 700 °C, respectively. This corresponds to a 60% lower steel wear rate for the Ts = 600 °C composite vs. the pure CrN coating, based on spherical cap volume calculations as
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shown in Table 2. The reduced wear track widths and associated counterface wear scar diameters for Ts = 500 and 600 °C are concurrent with the lower friction coefficients μ = 0.58 and 0.47 in comparison to pure CrN, μCrN = 0.62, respectively, and are attributed to the presence of Ag as well as the lower coating hardness of 17 and 20 GPa, respectively, which are considerably lower than the value for pure CrN of 28 GPa. The counterface wear rate is particularly low for the Ts = 600 °C composite, which exhibits a lower friction than both the pure CrN and Ts = 500 °C composite. As the Ts = 600 °C composite is slightly harder than the Ts = 500 °C composite, this indicates that the friction coefficient is the primary factor in determining wear response. Evidence for the formation of an effective lubricious Ag layer on the Ts = 600 °C composite surface is shown in Fig. 4. The plot shows the Ag concentration, as determined from EDS linescans across the wear tracks. For Ts = 500 °C, the Ag concentration remains essentially constant across the track, fluctuating in the 18–25 at.% range with an average value of 22 at.%. This corresponds to the bulk coating composition and, in turn, indicates that CrN and Ag are worn uniformly. In contrast, the Ag concentration in the wear track of the Ts = 600 °C sample exhibits multiple peaks to 30–40 at.% Ag, resulting in an average concentration of 25 at.% which is higher than the 22 at.% measured outside of the wear track. This may be due the lamellar shaped Ag aggregates, which are well developed at Ts = 600 °C and smear along the surface, yielding a soft Ag rich solid lubricating interface which, in turn, results in the low μ = 0.47 ± 0.02. The sample with Ts = 700 °C exhibits a constant Ag intensity (not shown), comparable to the Ts = 500 °C result. All these line-scan results, and in particular the Ag increase for Ts = 600 °C, are confirmed by wide area EDS measurements in multiple locations within and outside the tracks. The wide area measurements also indicate that all samples exhibit comparable patches of Fe deposits within the wear track due to the counterface wear and material transfer. The effect of Ts on the room temperature friction coefficient and wear rate is attributed to the orientation and structure of the Ag within the nanocomposite, as discussed in the following. The Ts = 500 °C coating exhibits friction and wear that is comparable to pure CrN, indicating that the very fine grained Ag within the matrix fails to form a continuous lubricious interface, resulting in a negligible effect of the Ag and tribological properties that resemble those for pure CrN. In contrast, Ts = 600 °C leads to a nanocomposite with a 25% lower friction and 16× lower wear than pure CrN. This suggests that the distributed lamellar Ag inclusions in this coating are of suitable size and orientation relative to the surface to provide an effective solid lubricating interface. This lubrication is effective despite the 3× higher
Table 2 Microstructural, mechanical, and tribological properties at room temperature. Sample
Pure CrNa
Ts = 500 °C
Ts = 600 °C
Ts = 700 °C
Ag agg. size (w h, nm) Roughness, Rrms (nm) Hardness (GPa) Elastic modulus (GPa) H3/E2 (GPa)
– 70 ± 10 28 ± 2 340 ± 80 0.19
50 × 25 100 ± 20 17 ± 2 240 ± 40 0.09
300 × 100 200 ± 20 20 ± 3 240 ± 60 0.14
600 × 200 1190 ± 70 24 ± 4 270 ± 70 0.19
100Cr6 steel Friction coefficient—μ Coating wear (mm3/Nm) Ball wear (mm3/Nm)
0.64 3.6 × 10− 6 4.4 × 10− 6
0.58 3.8 × 10− 6 3.1 × 10− 6
0.47 2.3 × 10− 7 1.8 × 10− 6
0.85 9.9 × 10− 6 7.1 × 10− 6
Alumina Friction coefficient—μ Coating wear (mm3/Nm) Ball wear (mm3/Nm)
0.59 4.7 × 10− 6 –
0.50 7.0 × 10− 6 –
0.69 8.0 × 10− 6 –
0.8 ~ 10− 4 –
a
Pure CrN deposited with Ts = 600 °C.
Fig. 4. Plots of the Ag concentration versus position, obtained from EDS line scans across the wear tracks for Ts = 500 and 600 °C composites tested against steel.
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initial surface roughness of the composite coating. However, raising the deposition temperature to Ts = 700 °C leads to composites with significantly higher friction and wear than for pure CrN. We attribute this to the large segregated Ag grains that cause structural inhomogeneity and a dramatic 17× increase in surface roughness due to the development of nodular defects [15]. That is, the material is worn too easily, leaving behind a roughened surface and generating a significant amount of wear debris. 3.3. Tribological properties against alumina The plot in Fig. 5 shows the room temperature friction coefficient of samples tested against 6 mm diameter alumina balls. Pure CrN exhibits a μ that increases from 0.4 to 0.5 to reach an average value of 0.59±0.03. This is slightly lower than μ=0.64 for the softer steel counterface, and slightly higher than reported CrN values with alumina counterfaces of 0.37, 0.52, and 0.55 at normal loads of 5, 5, and 2 N for Refs. [7,10,28], respectively. The composite coatings exhibit room temperature friction coefficients that increase from 0.3 to a steady-state value of μ=0.50±0.03 for Ts =500 °C, rise rapidly from an initial μ=0.4 and then fluctuate around an average value of 0.69±0.06 for Ts =600 °C, and oscillate around μ~0.8 prior to coating failure after ~1000cycles for Ts =700 °C. Fig. 6 shows corresponding optical profilometry results. The wear tracks conform to the spherical shape of the alumina ball, indicating relatively little wear of the alumina counterface material. The measured wear rate is lowest for the pure CrN coating, 4.7×10− 6 mm3/Nm, and increases to 7.0×10− 6, 8.0×10− 6, and ~10-4 mm3/Nm for nanocomposites with Ts =500, 600, and 700 °C, respectively. The CrN–Ag coating grown at Ts =500 °C exhibits a slightly lower friction but a higher wear rate than pure CrN. We attribute this to the Ag inclusions which affect the mechanical behavior of the composite coating, with H and H3/E2 decreasing from 28 and 0.19 GPa for CrN to 17 and 0.09 GPa for CrN–Ag. This softening of the coating and relative reduction in resistance to plastic deformation increases wear and only slightly reduces friction without forming an effective and continuous lubricating interface against the hard alumina counterface. The Ts =600 °C coating also exhibits higher room temperature friction and wear against alumina than pure CrN. This is in stark contrast to the results with a steel counterface. The dramatic dependence on the counterface material is attributed primarily to the difference in counterface wear: For the steel counterface, the ball wears and conforms to the surface, forming a flat on flat type geometry with reduced contact stress due to load spread over a much wider area, which is effectively lubricated by the lamellar Ag aggregates that are expected to shear along the surface as discussed in the prior section. In contrast, the considerably harder alumina ball experi-
Fig. 6. 3-D surface profiles of a pure CrN coating and CrN–Ag composite coatings with Ts = 500, 600 and 700 °C, after testing against an alumina counterface.
ences negligible wear and does not conform to the coating surface, resulting in plowing and uniform wear of both CrN and Ag phases. We attribute the higher μ, in comparison to pure CrN, to higher roughness from localized nodular surface asperities [15] and the associated formation of brittle wear debris from deformation and fracture of these asperities. The argument of uniform wear of CrN and Ag is supported by EDS measurements (not shown), which indicate no significant change in Ag concentration in the wear track for sliding against alumina, which is opposite to the result observed for sliding against steel in Fig. 4.The rapid tribological failure for Ts =700 °C is attributed to fracture of the large, 1–5 μm high nodular defects along the rough surface which likely lead to brittle failure of the coating. The differences in behavior for the steel and alumina counterfaces against Ts = 600 and 700 °C composites, wherein the composites exhibit large nodular surface asperities with average asperity tip radii Rs of 3.2 and 2.3 μm and average asperity heights of 0.4 and 2.5 μm, as measured by profilometry, respectively, can be further explained using the plasticity index Ψ, which is a dimensionless quantity that describes deformation behavior independent of normal load and sliding speed. In contrast to H3/E2, in which higher numbers represent an increased ability of a contact to resist plastic deformation on the macroscale of the planar surface and ball counterface, the plasticity index describes the deformation behavior on the microscale of surface asperities in terms of the likelihood of a contact to experience plastic deformation, with higher numbers indicating a greater likelihood of plastic deformation. That is, for Ψ ≪ 1, deformation is expected to be entirely elastic, while Ψ ≫ 1 leads to plastic deformation. The plasticity index is obtained using [29]: Ψ=
Fig. 5. Friction coefficient μ vs number of cycles during ball-on-disk testing against alumina for pure CrN and for CrN–Ag composites with Ts = 500, 600 and 700 °C.
E* H
rffiffiffiffiffi σs ; Rs
ð3Þ
where σs is the composite standard deviation of asperity peak heights which we estimate with the root-mean-square surface roughness of
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the composite surfaces, 200 and 1190 nm for Ts = 600 and 700 °C, respectively. E⁎ is the contact modulus defined by the elastic moduli and Poisson's ratios of the two materials in contact according to
E* =
1−ν21 1−ν22 + E1 E2
!−1 ;
ð4Þ
and is determined using the elastic moduli of the composites given in Table 2, an estimated composite Poisson's ratio of 0.23 [30], and elastic moduli and Poisson's ratios of 210 GPa and 0.3 for steel and 375 GPa and 0.22 for alumina [31], respectively. We calculate Ψ = 4.3 and 1.5 for counterface and coating, respectively, for the case of a steel counterface against the Ts = 600 °C composite, and obtain Ψ = 13.1 and 3.8 for the much rougher Ts = 700 °C composite. For both growth temperatures, all Ψ are above unity and the Ψ values for the steel counterfaces are considerably higher than for the CrN–Ag surfaces, indicating that plastic deformation of the softer steel counterfaces dominate in relieving the contact stresses at asperity tips. This is in agreement with the observed wear of the steel counterface material. For the alumina counterface, the situation is complicated by the closer relative hardness of coating and counterface, resulting in comparable plasticity index values of 2.6:1.9 and 7.9:5.0 for alumina:CrN–Ag at Ts = 600 and 700 °C, respectively. For the alumina counterface, the high plasticity index for Ts = 700 °C is in agreement with the observed severe wear, while the lower but comparable numbers between alumina and CrN–Ag for Ts = 600 °C suggest a likely simultaneous onset of plastic deformation for both coating and counterface. Due to the brittle behavior of ceramics such as CrN, at ambient temperature, onset of yielding would likely lead to fracture of large surface asperities and the generation of wear debris leading to the higher friction and wear observed for Ts = 600 °C and most extensively for Ts = 700 °C, which exhibits sharper asperity radii and larger asperity heights. 3.4. Tribological response as a function of temperature In order to determine temperatures at which the self-lubricating behavior of these films is most prevalent, the transient tribological response as a function of temperature was analyzed. Fig. 7 is a plot of the average transient friction coefficient μt during a continuous temperature ramp from Tt = 25–700 °C. The μt for pure CrN decreases from 0.65 to 0.26 for Tt = 100 to 300 °C, remains relatively constant at
Fig. 7. Transient friction coefficients μt measured against alumina during a continuous temperature ramp from Tt = 25–700 °C, for pure CrN and for CrN–Ag composites deposited at Ts = 500, 600 and 700 °C.
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μ t = 0.25 for 300 < T t < 500 °C, and increases to μ t = 0.55 for Tt ≥ 600 °C. The CrN–Ag coating with Ts = 500 °C exhibits a nearly identical friction as pure CrN for Tt ≤ 300 °C. However, μt continues to drop for Tt > 300 °C to reach a minimum of 0.05 at Tt = 500 °C, followed by a rise to 0.33 for Tt ≥ 600 °C. The Ts = 600 °C coating shows a qualitatively similar curve as for Ts = 500 °C with, however, a consistently higher μt that decreases from 0.64 at Tt = 100 °C to a minimum of 0.34 at Tt = 500 °C and reaches 0.58 for Tt ≥ 600 °C. The Ts = 700 °C coating exhibits rapid failure after ~ 1000 cycles and at Tt = 200–250 °C, as expected from room temperature testing. At low temperatures (Tt < 300 °C), μt decreases with increasing Tt for all samples. This is attributed to a drying of the environment along with an effective softening of the coatings due to an increasing dislocation mobility in CrN and Ag as well as a reduction in shear strength of Ag, and a reduction in the hardness of alumina which decreases from 15 GPa at room temperature to 8.5 GPa at 500 °C [32]. The continued drop in μt at Tt > 300 °C for Ag containing coatings indicates the presence of an increasingly effective lubricious Ag surface layer. The formation of this Ag surface layer with increasing Tt > 300 °C is attributed to a temperature activated solid lubricant transport to the surface, which we have previously demonstrated to occur at elevated temperatures [9]. The very low transient friction (μt = 0.05) for Ts = 500 °C at Tt = 500 °C is comparable to reported values for steady-state μ < 0.05 for ultra-thin (<0.2 μm) Ag films tested at room temperature in vacuum [33] and μ < 0.1 for thin (~2 μm) Ag films tested at 570 °C in an inert atmosphere [14]. Similar to the room temperature results against alumina, the friction for Ts = 600 °C is higher than for Ts = 500 °C. This is attributed to the inability of the Ag to effectively lubricate the 3× rougher Ts = 600 °C surface against the non-conforming alumina ball. The increase in μt above Tt = 500 °C for all samples is likely due to oxidative degradation similar to what has been previously reported for CrN [10,28]. These results suggest, consistent with the room temperature data, that the as-deposited Ag aggregate size and, in turn, the lubricant transport and the surface morphology, strongly affect the high-temperature tribological performance of CrN–Ag nanocomposite coatings. 4. Conclusions Tribological characterization of reactive magnetron sputtered CrN and CrN–Ag nanocomposite coatings against both 100Cr6 steel and alumina indicate that the nanocomposite structure at fixed composition strongly affects friction and wear. For the steel counterface, pure CrN and composites grown at 500 °C show comparable room temperature friction and wear, while Ts = 600 °C composites exhibit a 25% reduced friction and a 16 fold reduced wear rate in comparison to pure CrN. This is attributed to the formation of a lubricious Ag surface layer, as evidenced by EDS analyses, which is facilitated by the orientation and structure of the fine Ag lamellae within the CrN matrix for Ts = 600 °C. The wear increases considerably for the Ts = 700 °C, which is attributed to rough surfaces exacerbated by the formation of large nodules. For tribological tests against alumina, μ is found to increase with surface roughness with the lowest value observed for the CrN–Ag composite deposited at Ts = 500 °C, which exhibits a comparable roughness as the pure CrN coating. Wear rates for the CrN–Ag composites are a factor of 1.5–2 higher than for pure CrN, for both Ts = 500 and 600 °C. This higher wear rate is attributed to a reduction in overall hardness and H3/E2 ratio due to the addition of Ag as well as the inability of the Ag to form an effective lubricating interface against the harder alumina counterface, and for Ts = 600 °C also the higher initial roughness. The most promising benefits for CrN–Ag composite coatings are expected at elevated temperatures. Tribological testing of the coatings over a continuous temperature range of Tt = 25–700 °C indicates significant friction reductions for Ag containing composites with the minimum transient friction coefficient of 0.05 occurring at a transient test temperature of 500 °C for the finest grained composite. The
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Acknowledgments
[12] [13] [14] [15] [16] [17] [18]
This research was supported by the United States Army TEX3 Program, through the Armament Research, Development and Engineering Center (ARDEC), and by the National Science Foundation under grant no. CMMI-0653843.
[19] [20] [21] [22] [23]
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reduction in friction for Tt > 300 °C is attributed to a temperature activated lubricant transport of Ag to the surface. The friction increases above Tt = 500 °C, which is attributed to oxidative degradation.
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