Influence of Ti addition on the structure and properties

SCT-18565; No of Pages 9 Surface & Coatings Technology xxx (2013) xxx–xxx

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Influence of Ti addition on the structure and properties of low-friction W–S–C coatings Jill Sundberg a,⁎, Harald Nyberg b, Erik Särhammar c, Fredrik Gustavsson b, Tomas Kubart c, Tomas Nyberg c, Staffan Jacobson b, Ulf Jansson a a b c

Department of Chemistry, Ångström Laboratory, Uppsala University, PO Box 538, SE-751 21 Uppsala, Sweden Tribomaterials group, Department of Engineering Sciences, Uppsala University, PO Box 534, SE-751 21 Uppsala, Sweden Solid State Electronics, Department of Engineering Sciences, Uppsala University, PO Box 534, SE-751 21 Uppsala, Sweden

a r t i c l e

i n f o

Article history: Received 30 November 2012 Accepted in revised form 22 May 2013 Available online xxxx Keywords: WS2 Sputtering Tribology Solid lubrication XPS

a b s t r a c t Transition metal dichalcogenides, such as WS2 and MoS2, are known for their layered structure and lubricating properties. When deposited as thin coatings, however, their use as solid lubricants is limited by their low hardness and load-bearing capacity. The addition of another element, such as carbon, can improve the mechanical properties, although the hardness of for example W–S–C coatings is still rather low. In this work, Ti has been added to W–S–C coatings in order to further increase the hardness by carbide formation. W–S–C and W–S–C–Ti coatings were deposited by non-reactive magnetron sputtering, and characterized with regard to composition, structure and mechanical and tribological properties. It was found that the addition of Ti leads to the formation of a new carbide phase, and a significant increase in hardness for coatings with moderate carbon contents. The friction properties of W–S–C–Ti coatings were found to be comparable to that of W–S–C coatings, with friction coefficients down to μ ≈ 0.02 and similar wear rates against steel in a dry atmosphere. Formation of WS2 in the wear track of W–S–C–Ti was confirmed by transmission electron microscopy. It has thus been shown that the addition of Ti to W–S–C coatings can increase the hardness, while still maintaining WS2 lubrication. © 2013 Elsevier B.V. All rights reserved.

1. Introduction Transition metal dichalcogenides, TMDs, are materials well-known for their solid lubrication properties. They have a general formula of MX2, with the metal M being Mo or W, and the chalcogen X being S or Se. Their crystal structure is highly anisotropic, and can be described as layers of metal atoms in-between layers of chalcogen atoms, creating sandwich-like units where each metal atom is coordinated by a trigonal prism of chalcogen atoms. The bonding within each unit, i.e. the M\X bond, is covalent, while the different units are held together by considerably weaker van der Waals forces [1]. Thus, the layered structure makes these materials easily sheared, leading to a beneficial tribological behaviour with low friction coefficients [2]. However, the TMDs in themselves exhibit very low hardness and low load-bearing capacity. Porous TMD coatings are prone to oxidation and easily react with oxygen or moisture, forming MO3 which does not exhibit the same lowfriction properties [3–6]. One way to improve the mechanical properties

⁎ Corresponding author. Tel.: +46 18 471 3775; fax: +46 18 51 35 48. E-mail addresses: [email protected] (J. Sundberg), [email protected] (H. Nyberg), [email protected] (E. Särhammar), [email protected] (F. Gustavsson), [email protected] (T. Kubart), [email protected] (T. Nyberg), [email protected] (S. Jacobson), [email protected] (U. Jansson).

and density is the addition of a third element — either a metal, such as Ti [7–11] or Cr [12,13], or a second main-group element, such as nitrogen [14–16] or carbon [14–20]. This generally leads to denser coatings with a higher hardness. The resulting material is usually described as a composite of TMD crystals in an amorphous matrix. Promising results have been presented for, e.g., the W−S−C system, where coatings become denser and harder with the addition of carbon, while still showing beneficial tribological properties [17,21]. The W−S−C coatings consist of nanocrystalline WS2, and possibly also some carbide grains, in an amorphous carbon matrix. Voevodin et al. suggested a “chameleon” behaviour, where WS2 is responsible for low friction in dry atmospheres, while the carbon matrix provides low friction in humid environments [17,22,23]. However, the results presented by Polcar et al. indicate that only the WS2 phase provides lubrication [24]. In either case, W−S−C coatings have better tribological properties in dry than in humid environments, but the addition of carbon can improve the mechanical properties and the performance in humid conditions as compared to undoped WS2 coatings. Even though W− S − C coatings have good mechanical properties compared to pure WS2 coatings, their hardness is still rather low. One possibility of increasing the hardness, and thereby possibly also the wear rate, is to incorporate a hard phase, such as a carbide. It is possible to form tungsten carbide, but crystalline carbides are not formed to any large extent when combining WS2 with carbon. One route could

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Please cite this article as: J. Sundberg, et al., Surf. Coat. Technol. (2013), http://dx.doi.org/10.1016/j.surfcoat.2013.05.032

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J. Sundberg et al. / Surface & Coatings Technology xxx (2013) xxx–xxx

therefore be to add a fourth element that is a strong carbide-former — such as titanium. The incorporation of an additional carbide phase could increase the hardness, and possibly the wear resistance, of the coating while still maintaining low friction provided by the WS2 phase. Polcar et al. alloyed W−S−C coatings with Cr [25], which is also a carbide-forming metal. The W–S–C–Cr coatings were amorphous, without any crystalline chromium carbide, and the hardness was only slightly higher than for W–S–C. However, sputtered Cr–C coatings are often found to contain amorphous rather than crystalline carbides [26–28]. The aim of this work is to study the effect of adding titanium to W–S–C coatings. Titanium forms a cubic (fcc) carbide, and is more prone to form crystalline carbides than either Cr or W. Its addition should therefore increase the amount of crystalline carbide and thus the hardness of a W–S–C coating. In this study, W–S–C and W–S–C–Ti coatings have been deposited and characterized in detail with regard to their structure, phase composition and chemical bonding, in order to relate this to their mechanical properties and tribological behaviour. 2. Experimental details The coatings were deposited by non-reactive pulsed DC magnetron sputtering, using a von Ardenne system fitted with two magnetrons and a substrate heater. A sintered WS2 target (99.98%) and a carbon target (99.995%) were used, both with a diameter of 100 mm. For the deposition of W–S–C–Ti, a titanium foil (ASTM grade 1) was mounted onto the WS2 target so that it covered part of the race track. The carbon content of the coatings was varied by changing the power applied to the carbon target. Carbon target powers of 0, 350, 750 and 1000 W were used. For all depositions, the power applied to the WS2(Ti) target was 450 W. The substrates used were single-crystalline silicon wafers of (100) orientation for characterization purposes, and powder metallurgical high-speed steel (ASP 2023) for tribological testing. The high-speed steel substrates were hardened to 9 GPa and polished to a surface roughness of about 10 nm Ra. Prior to deposition, the substrates were cleaned by argon ion etching in the deposition chamber for 20 min, at a substrate bias voltage of − 600 V. The deposition was performed at a discharge pressure of 4 mTorr. The sample stage was rotated during the deposition, and no bias was applied to the substrates, which were kept at floating potential. Due to the different deposition rates of the different coatings, the deposition time was adjusted within the range 75 to 110 min, in order to obtain coatings with a thickness of approximately 1 μm. Each type of coating was deposited both with and without heating of the substrates. When heating was applied, the substrates reached a temperature of approximately 300 °C. The combination of four different carbon target powers, two WS2 target configurations (with and without added titanium) and two deposition temperatures resulted in a total of 16 coatings with different compositions being deposited. The coatings will hereafter be referred to by their constituting elements, with a number indicating the carbon target power used in the deposition (when applicable), and by the designations “HT” or “RT”, for coatings deposited with or without heating of the substrates, respectively. As an example, the designation W–S–C350–Ti HT describes a coating containing titanium, deposited using a carbon target power of 350 W and substrate heating. Energy-dispersive X-ray spectroscopy (EDS), performed on a Zeiss LEO 440 scanning electron microscope with a LaB6 crystal as the electron source, was used to determine the elemental composition of the coatings. Furthermore, elemental composition as well as chemical bonding was studied by X-ray photoelectron spectroscopy (XPS), using a PHI Quantum 2000 instrument with monochromatic Al Kα radiation. Prior to analysis, the coatings were mildly cleaned by sputtering with 200 eV Ar+ ions, rastered over an area of 1 mm2. The resulting spectra were curve fitted using Casa XPS software, using

a Shirley background and Voigt profiles. It should be noted that the surface sensitivity of XPS made sputter cleaning of the surfaces necessary, even though preferential sputtering of sulphur from WS2 and MoS2 is a well-known problem [29,30]. Even when using low-energy (200 eV) ions, the ion dose needed to remove the oxidized surface might have affected the composition of the measured sample volume, and thus also the chemical shifts measured by XPS. Raman spectroscopy was performed using a Renishaw microRaman system 2000 with a 514 nm laser at a power of 10 mW. The structure of the coatings was analysed by grazing incidence XRD, at an incidence angle of 1°, in parallel beam geometry on a Siemens D5000 instrument with Cu Kα radiation. The coatings were imaged by scanning electron microscopy (SEM) in top-view and as fractured cross-sections, using a Zeiss 1550 instrument with a field-emission gun as the electron source. Selected coatings were also studied by transmission electron microscopy (TEM) on a FEI Tecnai F30 ST, for which cross-section samples were made using a focused ion beam (FIB) in a FEI Strata D235 instrument. The hardness of the coatings was determined by nanoindentation, using an Ultra Nano Hardness tester from CSM Instruments, with a Berkovich diamond tip. The indentations were made using a maximum load of 900 μN, in order to obtain a penetration depth of less than 10% of the total coating thickness in all cases. For each sample, 20 indentations were made and the hardness and Young's modulus was calculated using the Oliver–Pharr method [31]. The tribological properties of the coatings were studied using a ball-on-disc setup with rotating geometry, the rotational radius being 2.5 mm and the sliding speed 100 mm/s. The counter surface was ball bearing steel balls with a diameter of 6 mm, subjected to a normal load of 5 N. The test duration was 10,000 revolutions (corresponding to a sliding distance of almost 160 m) and the measurements were performed both in dry air and under ambient conditions — the relative humidity being b5% and approximately 50%, respectively. After testing, the wear tracks and scars were analysed by TEM and Raman spectroscopy as described above. Wear of the coatings was estimated by measuring the volume loss in sections of the wear tracks using white light interference profilometry (WYKO NT 1100), while wear of the counter surfaces was determined by measuring the wear scar diameters of the worn balls using light optical microscopy (Nikon Measurescope). 3. Results and discussion 3.1. Chemical and structural characterization 3.1.1. Elemental composition The chemical composition of the coatings is presented in Fig. 1, along with the denominations for each sample. The compositions were determined by EDS, with the carbon contents corrected for each sample using XPS data. As expected, the increase in the power applied to the carbon target during deposition leads to a progressive increase in the carbon content. The carbon content is up to 40 at.% for the W–S–C coatings, and up to 30 at.% for the W–S–C–Ti coatings. The titanium content of the W–S–C–Ti coatings was between 35 and 50 at.%. The W–S–C coatings have S/W ratios of 0.7–0.9, i.e. significantly lower than 2. This sulphur substoichiometry of sputtered WS2 or MoS2 coatings, also when sputtered from stoichiometric targets, is wellknown and usually attributed to the preferential resputtering of sulphur [14,32–34]. The WS2 phase is thus more correctly referred to as WSx, where usually x is below 2. The W–S–C–Ti coatings, on the other hand, have S/W ratios above 1.2, i.e. higher than for W–S–C coatings but still substoichiometric. However, this ratio does not necessarily reflect the stoichiometry of the WSx phase, since W and S could also be bonded to the other elements present. All coatings also contained oxygen; 2–3 at.% for the W–S–C coatings, and up to 9 at.% for the W–S–C–Ti coatings, which is not included in the presented elemental compositions. The higher oxygen content in the W–S–C–Ti



Fig. 1. Elemental composition of all coatings of the four different series, determined by EDS, with the carbon contents corrected using XPS data for each sample.

coating is reasonable as titanium has a strong affinity for oxygen. In all series, the W–S–C0–(Ti) coatings contain more oxygen than those with carbon, which is not surprising as they are either more porous (in the W–S case) or contain titanium not bonded to carbon (in the W–S–Ti case). 3.1.2. Structure and morphology The diffraction pattern for the W–S–C0 HT coating features the typical hexagonal WS2 peaks, as shown in Fig. 2. The most prominent peak around 2θ = 35° is clearly asymmetric, which is due to the turbostratic stacking of the WS2 basal planes. The turbostratic stacking results in a sharp peak at approximately the position for the (100) reflection, with a tail towards higher angles representing other reflections of the (10l) family with l = 1, 2, 3… [14,15,19,33,35,36]. Another slightly asymmetric peak is seen at the position of the (110) reflection. X-ray diffraction measured by a θ–2θ scan (not shown) also features mainly the (10 L) reflection, and only a low-intensity (002) reflection, indicating that very few of the WSx crystals are oriented with basal planes parallel

a

3

to the surface. The pattern is typical for sputtered TMD coatings, which usually have a type I structure with the basal planes perpendicular to the substrate surface, as opposed to the type II structure where the basal planes are horizontally aligned and parallel to the surface [37–39]. The sample has a porous and columnar morphology (see Fig. 3a), with thin platelets growing perpendicular to the surface in a manner characteristic for this type of coatings. The pure W–S coating deposited without heating, W–S RT, has a dense morphology and a diffraction pattern with broader peaks, indicating smaller crystallites than in the high-temperature coating. The diffraction patterns become less distinct with the addition of carbon, and coatings with high carbon contents are virtually X-ray amorphous. The as-deposited W–S–C1000 HT coating has some short-range ordering in TEM (see Fig. 4), but no crystal grains. The change is accompanied by a shift towards a dense and featureless morphology, as seen in the SEM cross-sections shown in Fig. 3. Coatings deposited with substrate heating are more crystalline than their non-heated counterparts, as can be expected since the atoms should have a larger surface mobility. Thus, the WS2 phase in the W–S–C coatings becomes more dispersed with the addition of carbon, which is in agreement with previous studies of W–S–C and other doped TMD systems [15,16,19,40,41]. Furthermore, a slight shift towards lower diffraction angles is observed with increasing carbon content, see Fig. 2. The lattice parameter of WSx is thus increased with the addition of carbon, which can be related to changes in stoichiometry, a fact that was discussed but not observed by Nossa et al. [19]. The W–S–C0–Ti coatings show a broad peak situated at a higher angle than WS2, which is consistent with the results of Scharf et al. regarding W–S–Ti coatings with high titanium contents [10]. Also the morphology of W–S–C–Ti coatings with low carbon contents, e.g. W–S–C350–Ti HT in Fig. 3e, is dense and featureless. The diffraction pattern from XRD measurements (see Fig. 2) shows only broad, undefined peaks, and TEM micrographs and selected area electron diffraction (SAED) show only short range ordering. However, for W–S–C–Ti coatings with higher carbon contents, a new set of reflections appears in the X-ray diffraction as well as in SAED, indicating the presence of a crystalline phase with the cubic NaCl structure. The cell parameter is 4.50 Å, which is significantly larger than the cell parameter of 4.33 Å for common TiC. The morphology also changes towards a more columnar appearance (see Fig. 3f), although it is different from the lamellae of crystalline WSx. In TEM, crystal planes are clearly seen (Fig. 4b), and the SAED pattern indicates a polycrystalline material. Given its occurrence for high contents of

b W−S−C1000 HT

Intensity (a.u.)

W−S−C750 HT

W−S−C1000−Ti HT

W−S−C350 HT W−S−C750−Ti HT W−S−C350−Ti HT W−S−C0 HT 30

40

50

60

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70

80

W−S−C0−Ti HT 30

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Fig. 2. XRD patterns for a) W–S–(C) and b) W–S–(C)–Ti coatings deposited using substrate heating. The coatings deposited without heating show the same trend, but are comparatively less crystalline. For the W–S–C coatings, the WSx peaks (▼) become broader with increasing carbon content. For W–S–C–Ti coatings, reflections from a crystalline cubic TiCxSy phase (♦) appear for higher carbon contents.


4


Fig. 3. SEM cross-sections showing the development of morphology with the addition of carbon for coatings without (a–c) and with (d–f) titanium: a) W–S–C0 HT, b) W–S–C350 HT, c) W–S–C1000 HT, d) W–S–C0–Ti HT, e) W–S–C350–Ti HT, and f) W–S–C1000–Ti HT.

titanium and carbon, it is reasonable to expect these two elements to be present in the crystalline phase. However, the larger cell parameter could be due to impurities or doping with other elements. Titanium carbide is known for being substoichiometric in carbon, since the homogeneity range for TiCx ranges from x = 0.47 to x = 0.97 [42]. Furthermore, sputtering is a non-equilibrium technique, which makes possible the formation of metastable phases. The phase formed in this case is a sulphur-doped carbide, which could be denoted as TiCxSy and is described in more detail elsewhere [43]. 3.1.3. Chemical bonding The Raman spectrum for the pure W–S–C0 HT coating (see Fig. 5a) features two prominent peaks at 350 and 420 cm−1, belonging to the WSx phase [44], along with some other peaks that could be due to surface oxidation [45]. When a small amount of carbon is added, i.e. for the W–S–C350 HT coating, the spectrum shows broader and less intense WSx peaks as a result of the loss of crystallinity for the WSx phase. For higher carbon contents, and for the RT coatings, the

peaks are vaguely discernible, which is in agreement with the observation from XRD, that the WSx phase is poorly crystalline. However, as carbon is added to the coatings, two broad peaks gradually appear at about 1390 and 1570 cm−1. These peaks correspond to the D and G bands of carbon [21,46–49], respectively, and thus originate from the carbon matrix phase. The Raman spectra for the W–S–C–Ti coatings are similar to those of W–S–C coatings, with broad carbon peaks for the higher carbon contents. The low-wavenumber region does not feature any clearly distinguishable peaks. However, the titanium carbide phase is not expected to be Raman active. XPS analysis of W–S–C coatings, as seen in Fig. 6, showed that the peaks in the W4f region consist of contributions from two different chemical states. The peaks can be fitted using two sets of W4f doublets, each with a W5p3/2 line as it has a similar binding energy (BE) and thus has to be included in the fit. One contribution has the W4f7/2 line situated at a binding energy of about 32.0 eV, and is ascribed to WSx. Binding energies of 32.2 to 33.2 are reported for WS2 [50], but the slightly lower values in this case could be due to sulphur deficiency. To fit the

Fig. 4. HRTEM of as-deposited coatings: a) W–S–C1000 HT and b) W–S–C1000–Ti HT, the latter featuring crystalline grains. SAED is shown as insets (note that the bright dots in b) originate from the silicon substrate, due to the SAD aperture being significantly larger than the area shown in the HRTEM images).



a

W−S−C1000 HT

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Intensity (a.u.)

W−S−C750−Ti HT W−S−C750 HT

W−S−C350−Ti HT W−S−C350 HT W−S−C0−Ti HT W−S−C0 HT 400

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1200

1600

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Fig. 5. Raman spectra for a) W–S–C coatings and b) W–S–C–Ti coatings, all deposited with substrate heating.

measured spectra, another peak doublet at a larger binding energy is needed. This is found around 31.5 eV, and can be attributed to W–C, for which values of 31.5 to 32.2 eV are reported in the literature [50]. Before sputtering, all W4f spectra show mostly W\O bonds, as the surface oxide masks the unaffected coating. After sputter cleaning, however, the contribution of W\O bonds is negligible except in the case of the high-temperature-deposited W–S–C0 coating, which is highly porous and thus oxidized throughout the whole depth of the coating. Our observations agree with those of Nossa et al., who found W–S, W–C and some W\O bonds in W–S–C coatings [19]. However, the low-BE contribution attributed to W–C has a larger relative area than would be expected from the composition of the as-deposited coatings. Furthermore, a contribution at a similar binding energy is observed also for W–S coatings without carbon. The binding energy for metallic W is reported as 31.4 eV, and could thus also give a low-BE contribution, for example as an effect of preferential sputtering of sulphur. Sputter damage is a well-known problem in WS2-based coatings [25,26], and

S2p

C1s

Intensity (a.u.)

W−S−C750 HT

W4f

the preferential sputtering of sulphur should change the immediate surroundings of the tungsten atoms, thereby reducing their binding energy and leading to an overrepresentation of low-BE bonds in the spectra. Rumaner et al. studied the preferential sputtering from WS2 single crystal surfaces, and observed the appearance of a contribution ascribed to metallic W as an effect of sputter damage [29]. Due to their similar binding energies, it is not possible to separate the contributions of W–C and of metallic W from sputter damage. Quantifications could thus overestimate the amount of W\C bonds. The S2p peak (see Fig. 6) was, for all W–S–C coatings, situated around a bond energy of 161.8 to 162.0 eV, which is slightly lower than the literature values of 162.1 to 163.0 eV [50,51] but similar to the results of Nossa et al. [19]. The slight shift can be attributed to sulphur deficiency in the WSx phase. The sulphur atoms are thus mainly bonded to tungsten. In most cases, however, it is necessary to incorporate a second doublet pair at a slightly higher bond energy of 162.8–163.2 eV in order to obtain a good fit. This peak has been


Ti2p

39

37

35

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31

165 164 163 162 161

287 286 285 284 283 282

466

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458

Binding energy (eV) Fig. 6. XPS spectra for the W–S–C750 HT (top) and W–S–C750–Ti HT (bottom) coatings.


454

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suggested to originate from elemental sulphur [52], but is more likely due to the presence of sulphur–carbon bonds [19], which have a literature value of 163.6 eV [50] and should exist in the interface between the WSx phase and the matrix. The C1s region in the spectra measured for W–S–C coatings is best fitted using three different contributions, as shown in Fig. 6. There is one contribution at around 284.0 to 284.3 eV, which is typical for carbon bonded to carbon [50], and should originate from the amorphous matrix [19]. A second peak is found at 283.2 eV, which is slightly lower than reported values of 283.5 for WC [50,53] but still indicates the presence of W\C bonds [19]. The relative contribution of amorphous carbon successively increases with the total carbon content of the coating, in accordance with the results of Nossa et al. [19]. The third contribution needed to fit the C1s region is a small peak around 286 eV, which could originate from C\S bonds [50], contributions from which were also used in the fit for the S2p peak. In the current study, as well as in literature [19,54], W\C bonds were observed in coatings where no tungsten carbide crystallites were detected by XRD or electron diffraction. However, several studies show that amorphous tungsten carbides are formed when amorphous carbon coatings are doped with moderate amounts of tungsten. Monteiro et al. deposited amorphous carbon coatings doped with tungsten by pulsed vacuum arc deposition, and found that coatings with up to 20 at.% of W contained amorphous WC [55]. For W-doped carbon coatings deposited by magnetron co-sputtering, Silva et al. observed that the coatings were amorphous up to 23 at.% of W for un-biased conditions [56]. In previous studies of the W–S–C system, the nature of the tungsten carbide phase is not entirely clear. For coatings with 40–50 at.% of C, Voevodin et al. claimed the presence of WC nanocrystallites based on TEM images, although no such phase was detected by XRD [17]. Nossa et al., Polcar et al., and Evaristo et al., detected W\C bonds but did not clearly observe crystalline tungsten carbide except possibly for coatings with high carbon contents of more than 70 at.% [19,40,57]. We suggest that the matrix phase in W–S–C coatings consists not only of amorphous carbon, but of amorphous carbon and amorphous tungsten carbide, in a manner similar to W-doped carbon coatings [55,56,58]. The formation of W\C bonds leads to less tungsten being available for W\S bonds, and could therefore increase the x of the WSx phase, which is consistent with the observed increase in lattice parameter seen in Fig. 2 [19]. For W–S–C–Ti coatings with no or low carbon content (i.e. W–S–C0–Ti and W–S–C350–Ti coatings), the W4f spectra are similar to those of the W–S–C coatings. For coatings with higher carbon contents, however, the resulting peak is situated at a slightly higher binding energy of 32.2 to 32.3 eV and can be fitted using only small W–C contributions, as seen in Fig. 6. This could indicate that there is less sputter damage of the WSx phase in these coatings, and/or that they contain less tungsten carbide — possibly since the carbon is instead bonded to titanium. Furthermore, to obtain a good fit for the W–S–C–Ti coatings, a third set of peaks is needed at a higher bond energy around 34 eV, which could be from WO3 [50] — possibly due to increased porosity with the evolution of a columnar microstructure. The S2p peak of the W–S–C–Ti coatings has a different shape than that of the W–S–C coatings (see Fig. 6). Three sets of peaks are needed to obtain a good fit, representing S\W bonds at 161.6–161.9 eV, S\C bonds at 162.8–163.2 eV, and a third state of sulphur at an intermediate binding energy of 162.4 to 162.5 eV. As previously mentioned, the titanium carbide phase is doped with sulphur, and the intermediate state is thus attributed to sulphur in the titanium carbide (TiCxSy). Also for the C1s region, the addition of Ti to W–S–C leads to the appearance of an additional chemical state. The new contribution has a binding energy of 282.6 to 282.7 eV, which is lower than for tungsten carbide, but higher than the 281.6 eV reported for TiC [50,53]. However, as previously mentioned, the titanium carbide phase contains sulphur, and a chemical environment different from that in pure TiC should therefore be expected. As in W–S–C, peaks

representing C\W, C\C and C\S bonds are observed. The fraction of amorphous carbon is smaller in W–S–C–Ti than in W–S–C, due to the formation of the titanium carbide phase. The Ti2p peaks in spectra for the W–S–C–Ti coatings have, as their main contribution, a doublet pair with the 2p3/2 peak at a binding energy of 454.6 to 455.1 eV, as seen in Fig. 6. This position matches TiC, which should have a binding energy around 454.9 eV according to literature [59]. As previously mentioned, the carbide phase is doped with sulphur atoms, which could explain some variations in the bond energies. The exchange of carbon in TiC for sulphur should increase the electronegativity of the environment around titanium, which can be related to the increase in intensity on the high-BE side of the carbide peaks. For the W–S–C–Ti coatings with higher carbon contents, i.e. the W–S–C750–Ti and W–S–C1000–Ti coatings, there is also a contribution with the 2p3/2 peak situated at 458.7– 458.8 eV, closely matching the literature values for TiO2. Since these coatings have a somewhat porous microstructure (as seen in Fig. 3), it is not surprising that they are more oxidized than the dense W–S–C350–Ti coatings, particularly as titanium has a high affinity for oxygen. The Ti2p peaks for the coatings without carbon feature several oxidized contributions, but are of limited interest in this study. To summarize the conclusions drawn from XPS spectra, the W–S–C coatings are seen to consist of WSx in a matrix of amorphous carbon and amorphous tungsten carbide. Since part of the tungsten is bonded to carbon instead of sulphur, the actual x for the WSx phase is significantly higher than the apparent one calculated from the total elemental composition (i.e., x = 0.7–0.9). It may still be below 2, although the wellknown problem of sputter damage makes quantifications unreliable. The addition of Ti leads to the formation of Ti\C bonds, and an increase in the amount of carbon that is present as carbide. The TiC phase also contains sulphur, which is reflected in an additional chemical state for sulphur as compared to W–S–C coatings. In the case of W–S–C–Cr coatings studied by Polcar et al., the addition of Cr was not seen to alter the chemical bonding of any other element, although the Cr 2p peak matched the position for metallic Cr as well as chromium carbide [25]. The coatings were amorphous and thus did not contain any crystalline chromium carbide, but possibly amorphous chromium carbide. This is consistent with the fact that, unlike for Ti, it is not uncommon for sputtered chromium carbides to be amorphous [26–28].

3.2. Mechanical properties The mechanical properties of the coatings were evaluated by depth-sensing nanoindentation using a Berkovich tip. The pure W–S–C0 HT coating proved to have a very low hardness and Young's modulus, which is expected considering its high porosity. The maximum load was decreased to 200 μN in this case because of the low hardness, as opposed to 900 μN for all other measurements. The less crystalline and denser W–S–C0 RT coating proved to have a higher hardness (4 GPa) and Young's modulus (70 GPa). As seen in Fig. 7, the addition of carbon increased the hardness of the coatings, which is in agreement with literature results [14,41]. The trend for the Young's modulus closely follows the trend of the hardness. The doping of W–S coatings with carbon or other elements generally leads to densification, which should account for at least part of the increase in hardness. Formation of harder phases such as tungsten carbide can also add to the hardness. In the case of W–S–C coatings, the hardness increases with increasing carbon content up to a maximum, after which further addition reduces the hardness. The W–S–C coatings in the current work contain up to about 40 at.% of carbon, and the hardness is seen to level out or possibly decrease over these levels, which agrees with the hardness maximum around 40 at.% of carbon reported by Polcar et al. [18,21]. The addition of moderate amounts of carbon improves the density of the coatings and facilitates the formation of some W\C bonds. At higher carbon contents, however, the morphology



W−S−C RT W−S−C HT W−S−C−Ti RT W−S−C−Ti HT

15

showed reasonably low initial friction, and that post-test examination of the wear scars of failed coatings by light optical microscopy revealed that all failures were due to coating delamination. This indicates that the failures are not necessarily due to poor intrinsic frictional behaviour of the coatings, but rather to poor adhesion of the coatings to the substrate. The results presented are therefore mainly based on coatings that exhibited sufficient adhesion, and thus could be evaluated and compared.

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1000

Carbon target power (W) Fig. 7. Hardness of W–S–C and W–S–C–Ti coatings as a function of the carbon target power. The lines are guides for the eye.

does not change further, and the additional carbon likely adds to the matrix phase rather than forming bonds with tungsten. Although the W–S–C coatings are significantly harder than pure W–S coatings, they still only reach hardness values up to about 10 GPa. The purpose of adding titanium was to enable the formation of a titanium carbide phase, thereby increasing the hardness further. As seen in Fig. 7, the W–S–C–Ti coatings are indeed in several cases significantly harder than the corresponding W–S–C coatings. Also for coatings without carbon (i.e. with a carbon target power of 0 W), the hardness is increased by the addition of titanium. This effect is typical for doping TMD coatings with a third element, as the microstructure and morphology of the coating change [7,10]. For low-carbon coatings, the hardness is significantly increased by the addition of titanium, especially for the HT coatings where it increases from 4 to 18 GPa under otherwise identical deposition conditions. Since in this case both coatings have a similar dense and featureless appearance, the hardness increase is attributed to the titanium carbide phase in W–S–C–Ti. In Fig. 7, it is seen that titanium addition to coatings with high carbon contents only slightly increases the hardness, or does not increase it at all. In these cases, the W–S–C–Ti coatings contain the titanium carbide phase, but at the same time they are more porous than their W–S–C counterparts, and the combined effects result in only a small net effect on the measured hardness. It should be mentioned that, like for the W–S–C coatings, the trends for the Young's modulus follow the trends in hardness closely, and the maximum value of 220 GPa is reached by the W–S–C350–Ti HT sample with a hardness of 18 GPa. When doping W–S–C coatings with chromium, Polcar et al. did not observe a significant change in hardness: the W–S–C–Cr coatings had a hardness of 6–7 GPa, which is close to that of the comparable W–S–C coating at 5 GPa [25]. The addition of Cr lead to the coatings becoming amorphous, and no crystalline chromium carbides were formed. However, chromium is not a strong carbide-former as titanium, and non-reactive sputtering usually favours the formation of amorphous chromium carbide, which gives lower hardness values than coatings containing nanocrystallites [27,28]. Its addition should therefore not be expected to increase the hardness.

3.3.2. Tribological tests in dry air All coatings were tested in unlubricated sliding against a ball bearing steel ball, in an atmosphere of dry air (relative humidity b 5%). Fig. 8 shows the mean coefficient of friction of coatings that showed no obvious signs of failure during testing, whereas Fig. 9 shows the corresponding wear rates. Each point in Fig. 8 corresponds to the mean value of sliding friction during the entire test. In applicable cases, the presented value is the average for both performed tests. All W–S–C coatings that could be evaluated showed low and stable friction (μ b 0.05). Failure occurred only for coatings deposited at the two highest carbon target powers (750 and 1000 W), although the W–S–C750 RT coating showed low and stable friction in one of the two performed tests. For the W–S–C–Ti coatings, all tests started out at a high friction level, typically exceeding 0.7 during the first tens of revolutions. The coatings without carbon (i.e. the W–S–C0–Ti coatings) failed almost instantly, without ever reaching a low friction state. In contrast, the carbon containing coatings showed almost no tendency to flaking, and survived all of the tests. The W–S–C–Ti coatings deposited at room temperature also showed reasonably low friction. In contrast, the W–S–C–Ti HT coatings showed a large variation of the friction with the carbon content. The W–S–C350–Ti HT coating showed friction similar to that of the corresponding W–S–C coating (μ ≈ 0.02), whereas the coatings with higher carbon contents (750 and 1000 W) showed dramatically higher friction levels (0.18 and 0.46, respectively). A detailed study of the running-in process of selected W–S–C–Ti coatings, compared to the corresponding W–S–C coatings, is presented in another study by the current authors [60]. The specific wear rates of the coatings tested in dry air are shown in Fig. 9. In general, the wear rate followed the friction, so that low friction was accompanied by a low wear rate. This is perhaps most clearly seen for the W–S–C–Ti coatings deposited with substrate heating, where both friction and wear rate increased strongly with the carbon target power. A clear exception from the correlation between friction and wear was the W–S–C0 HT coating, which had a more than five times higher wear rate than the W–S–C350-HT

0.5

Mean coefficient of friction

Hardness (GPa)

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3.3. Tribological properties 3.3.1. Adhesion of coatings All coatings were tested using a pin-on-disc setup. For a number of coatings, evaluation of the tribological properties of the coating material was not possible due to failure of the coatings. It should be noted that most coatings that failed at early stages of testing

0 0

350

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Carbon target power (W) Fig. 8. Mean coefficient of friction in dry air for all successfully tested coatings.




104

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Fig. 9. Specific wear rate in dry air of all successfully tested coatings, corresponding to the friction results shown in Fig. 8. Please note the logarithmic vertical scale.

Fig. 11. Specific wear rate in humid air of all successfully tested coatings, corresponding to the friction results shown in Fig. 10.

coating, and almost thirty times higher than that of the W–S–C0 RT coating, despite the three coatings showing almost identical frictional behaviour. Considering the very low hardness and large porosity of the W–S–C0 HT coating, as compared to the other coatings, it is not surprising that it is severely worn. Similar behaviour of pure, porous W–S coatings has been observed previously [14]. The wear of the counter surfaces (balls) was also studied, but is not presented separately. The results were very similar to those of the coating wear, as large wear of the coating was associated with large wear of the ball. The only exception from this was the W–S–C0 HT coating, discussed above. Despite showing the highest coating wear of all tested coatings, this coating caused the least wear of the counter surface, which again should be related to the extreme softness of this coating.

the performance of the W–S–C coatings with high carbon contents due to problems with adhesion, although they sometimes showed low initial friction. The addition of Ti to the coatings had no obvious positive effect on the performance in humid air. Both friction and wear were higher for the W–S–C–Ti coatings, compared to corresponding W–S–C coatings (in the cases were this comparison was possible). As in dry air, the W–S–C–Ti coatings did however suffer from failure less frequently, except for the W–S–C0–Ti coatings. The wear of the W–S–C–Ti was mainly of continuous nature, rather than the flaking observed for the failing W–S–C coatings. The addition of titanium might thus have improved the adhesion to the substrate.

3.3.3. Tribological tests in humid air Figs. 10 and 11 show the mean coefficient of friction and specific wear rate, respectively, from the tests in humid air (relative humidity approximately 50%). As expected, the friction levels were generally higher than in dry air, the only exception being the W–S–C1000–Ti HT coating, which improved slightly from the very high level seen in dry air. The lowest friction, approximately 0.1, was observed for the W–S–C750 RT coating. As in the dry case, it is difficult to evaluate

Mean coefficient of friction

0.5

3.3.4. Tribofilm characterization The W–S–C–Ti coating with a low carbon content deposited using substrate heating (W–S–C–Ti 350 HT) had the highest hardness of all coatings in this study and showed good tribological properties with μ b 0.02 in the dry-atmosphere test. Therefore, it was chosen for post-test analysis by TEM. A cross-section sample was prepared by FIB from the wear track, perpendicular to the sliding direction, after 5000 cycles, and imaged by TEM (Fig. 12). The cross-section showed a tribofilm with a thickness up to about 40 nm, consisting mainly of amorphous material, which contains oxygen and thus consists of


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0

0

350

750

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Carbon target power (W) Fig. 10. Mean coefficient of friction in humid air for all successfully tested coatings.

Fig. 12. TEM cross-section perpendicular to the sliding direction from the wear track of a W–S–C350–Ti HT coating. Horizontally aligned WS2 planes have formed on the top surface, as well as within the amorphous tribofilm.



oxide of titanium and/or tungsten. In the case of W–S–C–Cr coatings, Polcar et al. observed a layer consisting mostly of chromium oxide underneath the outermost layer of WS2 [25]. In some areas, the tribofilm is thin and amorphous — likely areas which have not been points of contact. In other parts, characteristic WS2 planes have formed and horizontally aligned on the top surface as well as within the amorphous material. These planes have a plane distance matching that of WS2. Thus, it is confirmed that crystalline WS2 is generated and aligned towards the ideal angle for low-friction sliding in the contact, even though no crystalline WSx is detected in the as-deposited coatings. Similar observations have been published for other TMD-based materials [25,57,61,62]. 4. Conclusions W–S–C and W–S–C–Ti coatings with varying carbon contents were deposited by non-reactive DC magnetron sputtering from a carbon and a WS2(Ti) target, with and without heating of the substrates. The W–S–C coatings were found to consist of WSx in a matrix of amorphous carbon and amorphous tungsten carbide, with the crystallinity of the WSx phase decreasing with increasing carbon content so that the coatings were virtually X-ray amorphous for higher carbon contents. The W–S–C–Ti coatings were amorphous for low carbon contents, but the chemical bonding indicated the same WSx and matrix phases as in W–S–C, along with a titanium-based carbide (TiCxSy). For higher carbon contents, the W–S–C–Ti coatings featured a columnar morphology and a crystalline carbide phase, which is interpreted as TiC with an expanded lattice due to sulphur doping-TiCxSy. Carbon addition to pure W–S increases the hardness by change in morphology and the formation of some W\C bonds. However, the W–S–C–Ti coatings are harder than their W–S–C counterparts for moderate carbon contents, a fact attributed to the formation of more carbide while still maintaining the dense and amorphous morphology. In the tribological tests, the addition of Ti was seen to improve the durability of the coatings. Failure was less frequent for the coatings containing both Ti and C than for those without Ti. The addition of Ti generally leads to increased friction, although the best Ti-containing coating did reach the same low level as the best W–S–C coatings. TEM studies in the wear track of a selected W–S–C–Ti sample showed tribo-induced formation WS2 planes aligned along the sliding direction. It can thus be concluded that the hardness of W–S–C coatings can be increased by the addition of Ti, and that horizontally aligned WS2 can be generated in the tribological contact also for W–S–C–Ti films. Acknowledgements The support from the Swedish Foundation for Strategic Research via the programme Technical advancement through controlled tribofilms is gratefully acknowledged. References [1] P.D. Fleischauer, Thin Solid Films 154 (1987) 309. [2] S.R. Cohen, L. Rapoport, E.A. Ponomarev, H. Cohen, T. Tsirlina, R. Tenne, C. Levy-Clement, Thin Solid Films 324 (1998) 190. [3] T.B. Stewart, P.D. Fleischauer, Inorg. Chem. 21 (1982) 2426. [4] P.D. Fleischauer, J.R. Lince, Tribol. Int. 32 (1999) 627. [5] S.V. Prasad, J.S. Zabinski, N.T. McDevitt, Tribol. Trans. 38 (1995) 57. [6] V. Buck, Wear 114 (1987) 263. [7] V. Rigato, G. Maggioni, D. Boscarino, L. Sangaletti, L. Depero, V.C. Fox, D. Teer, C. Santini, Surf. Coat. Technol. 116–119 (1999) 176.

9

[8] V. Rigato, G. Maggioni, A. Patelli, D. Boscarino, N.M. Renevier, D.G. Teer, Surf. Coat. Technol. 131 (2000) 206. [9] M. Evaristo, A. Nossa, A. Cavaleiro, Mater. Sci. Forum 514–516 (2006) 687. [10] T.W. Scharf, A. Rajendran, R. Banerjee, F. Sequeda, Thin Solid Films 517 (2009) 5666. [11] D.G. Teer, Wear 251 (2001) 1068. [12] B. Deepthi, H.C. Barshilia, K.S. Rajam, M.S. Konchady, D.M. Pai, J. Sankar, A.V. Kvit, Surf. Coat. Technol. 205 (2010) 565. [13] B. Deepthi, H.C. Barshilia, K.S. Rajam, M.S. Konchady, D.M. Pai, J. Sankar, Surf. Coat. Technol. 205 (2010) 1937. [14] A. Nossa, A. Cavaleiro, Surf. Coat. Technol. 142 (2001) 984. [15] A. Nossa, A. Cavaleiro, Surf. Coat. Technol. 163 (2003) 552. [16] A. Nossa, A. Cavaleiro, N.J.M. Carvalho, B.J. Kooi, J.T.M. De Hosson, Thin Solid Films 484 (2005) 389. [17] A.A. Voevodin, J.P. O'Neill, J.S. Zabinski, Tribol. Lett. 6 (1999) 75. [18] T. Polcar, M. Evaristo, A. Cavaleiro, Vacuum 81 (2007) 1439. [19] A. Nossa, A. Cavaleiro, J. Mater. Res. 19 (2004) 2356. [20] M. Roy, T. Koch, A. Pauschitz, Appl. Surf. Sci. 256 (2010) 6850. [21] T. Polcar, M. Evaristo, A. Cavaleiro, Plasma Processes Polym. 4 (2007) 541. [22] A.A. Voevodin, J.S. Zabinski, Thin Solid Films 370 (2000) 223. [23] J.H. Wu, D.A. Rigney, M.L. Falk, J.H. Sanders, A.A. Voevodin, J.S. Zabinski, Surf. Coat. Technol. 188–89 (2004) 605. [24] T. Polcar, M. Evaristo, A. Cavaleiro, Wear 266 (2009) 388. [25] T. Polcar, F. Gustavsson, T. Thersleff, S. Jacobson, A. Cavaleiro, Faraday Discuss. 156 (2012) 383. [26] G. Gassner, J. Patscheider, P.H. Mayrhofer, E. Hegedus, L. Toth, I. Kovacs, B. Pecz, V. Srot, C. Scheu, C. Mitterer, J. Vac. Sci. Technol. B 24 (2006) 1837. [27] C. Ziebert, J. Ye, M. Stüber, S. Ulrich, M. Edinger, I. Barzen, Surf. Coat. Technol. 205 (2011) 4844. [28] M. Andersson, J. Högström, S. Urbonaite, A. Furlan, L. Nyholm, U. Jansson, Vacuum 86 (2012) 1408. [29] L.E. Rumaner, T. Tazawa, F.S. Ohuchi, J. Vac. Sci. Technol. A 12 (1994) 2451. [30] J.C. Bernede, Appl. Surf. Sci. 171 (2001) 15. [31] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [32] J. Moser, F. Levy, F. Bussy, J. Vac. Sci. Technol. A 12 (1994) 494. [33] M. Regula, C. Ballif, J.H. Moser, F. Levy, Thin Solid Films 280 (1996) 67. [34] G. Weise, N. Mattern, H. Hermann, A. Teresiak, I. Bacher, W. Bruckner, H.D. Bauer, H. Vinzelberg, G. Reiss, U. Kreissig, M. Mader, P. Markschlager, Thin Solid Films 298 (1997) 98. [35] J. Moser, F. Levy, Thin Solid Films 240 (1994) 56. [36] I. Martin-Litas, P. Vinatier, A. Levasseur, J.C. Dupin, Thin Solid Films 416 (2002) 1. [37] R. Bichsel, F. Levy, Thin Solid Films 116 (1984) 367. [38] P.A. Bertrand, J. Mater. Res. 4 (1989) 180. [39] M. Genut, L. Margulis, G. Hodes, R. Tenne, Thin Solid Films 217 (1992) 91. [40] M. Evaristo, A. Nossa, A. Cavaleiro, Surf. Coat. Technol. 200 (2005) 1076. [41] T. Koch, M. Evaristo, A. Pauschitz, M. Roy, A. Cavaleiro, Thin Solid Films 518 (2009) 185. [42] J.L. Murray, Phase Diagrams of Binary Titanium Alloys, ASM International, 1989. [43] J. Sundberg, H. Nyberg, E. Särhammar, K. Kádas, L. Wang, O. Eriksson, T. Nyberg, S. Jacobson, U. Jansson, Mater. Res. Lett. (2013), http://dx.doi.org/10.1080/21663831. 2013.802262. [44] T. Sekine, T. Nakashizu, K. Toyoda, K. Uchinokura, E. Matsuura, Solid State Commun. 35 (1980) 371. [45] M.F. Daniel, B. Desbat, J.C. Lassegues, B. Gerand, M. Figlarz, J. Solid State Chem. 67 (1987) 235. [46] J. Schwan, S. Ulrich, V. Batori, H. Ehrhardt, S.R.P. Silva, J. Appl. Phys. 80 (1996) 440. [47] A.C. Ferrari, J. Robertson, Phys. Rev. B 61 (2000) 14095. [48] A.C. Ferrari, Diamond Relat. Mater. 11 (2002) 1053. [49] S. Prawer, K.W. Nugent, Y. Lifshitz, G.D. Lempert, E. Grossman, J. Kulik, I. Avigal, R. Kalish, Diamond Relat. Mater. 5 (1996) 433. [50] J. Chastain, R.C. King, in, Physical Electronics, Inc., Eden Prairie, MN, 1995. [51] K.T. Ng, D.M. Hercules, J. Phys. Chem. 80 (1976) 2094. [52] J.S. Zabinski, M.S. Donley, S.V. Prasad, J. Mater. Sci. 29 (1994) 4834. [53] L. Ramqvist, K. Hamrin, G. Johansson, A. Fahlman, C. Nordling, J. Phys. Chem. Solids 30 (1969) 1835. [54] J.V. Pimentel, T. Polcar, M. Evaristo, A. Cavaleiro, Tribol. Int. 47 (2012) 188. [55] O.R. Monteiro, M.-P. Delplancke-Ogletree, I.G. Brown, Thin Solid Films 342 (1999) 100. [56] C.W.E. Silva, J.R.T. Branco, A. Cavaleiro, Thin Solid Films 515 (2006) 1063. [57] T. Polcar, M. Evaristo, A. Cavaleiro, Plasma Processes Polym. 6 (2009) 417. [58] B. Trindade, M.T. Vieira, E. Bauer-Grosse, Acta Mater. 46 (1998) 1731. [59] T. Zehnder, J. Patscheider, Surf. Coat. Technol. 133–134 (2000) 138. [60] H. Nyberg, J. Sundberg, E. Särhammar, F. Gustavsson, T. Kubart, T. Nyberg, U. Jansson, S. Jacobson, Wear (2013), http://dx.doi.org/10.1016/j.wear.2013.01.065. [61] T. Polcar, A. Cavaleiro, Surf. Coat. Technol. 206 (2011) 686. [62] J.J. Hu, R. Wheeler, J.S. Zabinski, P.A. Shade, A. Shiveley, A.A. Voevodin, Tribol. Lett. 32 (2008) 49.
