Nanostructured WS2–Ni composite films for improved oxidation, resistance and tribological performance

Applied Surface Science 288 (2014) 15–25

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Nanostructured WS2 –Ni composite films for improved oxidation, resistance and tribological performance Shusheng Xu a,b , Xiaoming Gao a , Ming Hu a , Jiayi Sun a , Dong Jiang a , Feng Zhou a , Weimin Liu a,∗ , Lijun Weng a,∗∗ a b

State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, PR China University of Chinese Academy of Sciences, Beijing 100049, PR China

a r t i c l e

i n f o

Article history: Received 29 April 2013 Received in revised form 5 September 2013 Accepted 5 September 2013 Available online 13 September 2013 Keywords: WS2 film Ni dopant Microstructure Oxidation resistance Tribology

a b s t r a c t WS2 films were prepared by radio frequency sputtering. In order to improve its mechanical properties and oxidation resistance, Ni was used as the dopant and the effect of Ni content on the microstructure, anti-oxidation capability, mechanical and tribological properties of composite films were studied by energy dispersive X-ray spectroscopy (EDS), X-ray photoelectron spectroscopy (XPS), Raman spectroscopy, grazing incidence X-ray diffraction (GIXRD), high resolution transmission electron microscope (HRTEM), field emission scanning electron microscopy (FESEM), nano-indentation tester, scratch tester and ball-on-disk tribometer. WS2 existed in nanocrystalline 2H–WS2 structure and Ni an amorphous phase. Increasing the Ni content resulted in a microstructural change from columnar platelet structure of pure WS2 film to a fiber-like structure of the composite film at low Ni content (5.0 at%), and to a featureless structure at high Ni content (>10 at%). Meanwhile, the films became more and more compact and showed improved anti-oxidation capability. The films represented an increase in hardness with Ni content ranging from 0 to 10.3 at% due to the microstructure being densified, but exhibited high brittleness as the Ni content higher than 10 at%. The composite film at low Ni content of about 5.0 at% showed much better wear resistance than pure WS2 film, but became brittle and had poor wear resistances at high Ni content of above 10.3 at%. The WS2 –5.0 at% Ni composite film exhibited the longest wear life of 5.8 × 105 cycles about sevenfold better than that of pure WS2 film in humid air. The wear mechanism was discussed in terms of the anti-oxidation capability of the films, morphology of the wear track and formation of transfer film. © 2013 Elsevier B.V. All rights reserved.

1. Introduction Transition metal disulfides (MS2 , M is molybdenum or tungsten) have been used as solid lubricants in various space motion mechanisms. The low friction coefficients of MS2 are mainly attributed to the layered structure in which an M layer is sandwiched between two hexagonally packed S layers. Chemical bonding within the inplane atoms is covalent, while weak Van der Waals force links the layers together, where lateral shear of layers occurs easily [1–3]. The major application of MS2 films is for spacecraft, but the films sometimes have to be subjected to terrestrial tests or storage before the spacecraft launch [4,5]. Therefore, MS2 films could be oxidized due to the presence of dangling or unsaturated bonds at edge planes in the layered structures, leading to deterioration of their tribological properties in terrestrial tests [6–9].

∗ Corresponding author. Tel.: +86 931 4968166; fax: +86 931 8277088. ∗∗ Corresponding author. Tel.: +86 931 4968003; fax: +86 931 8277088. E-mail addresses: [email protected] (W. Liu), [email protected] (L. Weng). 0169-4332/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2013.09.024

It was found that an unexpected lubricating effect could be obtained, when more than two solid lubricants were combined utilization, overcoming the disadvantage of single lubricant [10–13]. There have been many works on improving the tribological properties of MS2 film in air by means of co-sputtering MS2 and oxides (Sb2 O3 [4,14,15], PbO [16], etc.), metals (Pb [17], Au [4,11], Al [18], Ti [12,19], Cr [20,21], etc.) as well as ceramic [22]. However, the action mechanisms of the dopant did not reach consensus on the improvement of tribological properties for the MS2 films. Teer [23] speculated that Ti atoms were situated between the neighboring S planes in the MoS2 matrix. The resulted distortion was responsible for the increase in hardness and the excellent humidity resistance for the MoS2 –Ti composite film. Scharf et al. [4] demonstrated that the Au dopant was homogeneous distributed in the MoS2 film in the form of small island. Meanwhile, coarsening of Au particles induced by friction in the wear subsurface restricted deep oxidation of the MoS2 –Au composite film in humid air, and then the tribological performance could be significantly improved. In general, the incorporation of second phase into the MS2 matrix film might not reduce the high friction coefficient in humid air, but could

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Table 1 Film deposition parameters and chemical composition. Run no.

WN1 WN2 WN3 WN4

Retention time (s) WS2 target

Ni target

6 6 5 4

0 1.5 1.5 1.5

Alternate cycles

Ni (at%)

W (at%)

S (at%)

O (at%)

S/W ratio

120 130 150 165

0 5.0 10.3 16.5

36.3 33.3 31.0 28.5

56.9 57.2 53.9 50.8

6.8 4.5 4.8 4.2

1.57 1.72 1.74 1.78

improve the wear resistance due to (i) the significant increase in the film hardness with more than one order of magnitude higher than the pure MS2 films (MS2 –C, MS2 –Ti), (ii) the increase of the adhesive strength, or/and (iii) the preferential oxidation of the dopant compared to the MS2 film (MS2 –Ti, MS2 –Al) [18,19,23–25]. The formation of outermost surface in the wear track and the transfer films on the counterpart ball consisted of a well-ordered thin layer of WS2 with basal planes parallel to substrate was attributed to the excellent lubricanting property of the compact films. Meanwhile, the compact subsurface of the films without significant cracks or pores could provide high load-bearing capability [24,26,27]. WS2 is one of transition metal disulfides with lamellar structure, and previous investigations reveal that WS2 has an improved oxidation resistance and thermally stability withstanding temperature of about 100 ◦ C higher than MoS2 [28–30]. Nevertheless, the studies on the structure and tribological properties of the WS2 composite films are much less as compared with MoS2 . Due to the lattice parameters and thermal expansivity of metallic Ni being similar to the Fe substrate [31,32], the Ni dopant in WS2 films might help in forming beneficial eutectic layer near the film-substrate interface and so make for improving its adhesive strength and tribological properties [33]. So far, systematic studies on the variation of microstructure and tribological performance of WS2 –Ni composite films have not been reported. It was expected to modify the microstructure and improve the tribological performance in air by incorporating Ni in WS2 matrix film. In the present paper, the WS2 based composite films with varying amounts of Ni dopant were deposited by using radio frequency (RF) sputtering method to investigate the effect of Ni contents on the structure and mechanical/tribological properties of the films. 2. Experimental detail 2.1. Film preparation WS2 –Ni composite films were co-deposited using WS2 (80 mm in diameter, 99.9% purity) and Ni (80 mm in diameter, 99.9% purity) targets by a RF sputtering system under Ar pressure of 2.0 Pa. The center distance of two targets was 200 mm. The commercial n-type Si (100) wafer substrates were selected for film composition and structure analysis and AISI 440C steel substrates (25 mm × 25 mm × 4 mm) were applied for tribological properties test. The AISI 440C steel substrates were polished to surface roughness (Ra ) lower than 0.03 ␮m and then ultrasonically cleaned with alcohol. The vertical distance between substrates and targets was 60 mm. Before deposition, the vacuum chamber was evacuated to a background vacuum below 1.0 × 10−3 Pa, and then the substrates were Ar ion etched for 15 min at a DC bias of −600 V. Afterwards, WS2 –Ni composite films were co-deposited at powers of 320 and 300 W for WS2 and Ni targets, respectively. During the film deposition, two targets have been sputtered continuously and the substrates were alternately located under WS2 and Ni targets by revolving the sample holder. The amount of Ni dopant in the composite films was changed by adjusting the retention time of substrates under WS2 and Ni targets. This process was controlled

by a programmed motor. For comparison, all the film thicknesses were controlled to about 2 ␮m. Detailed deposition parameters are listed in Table 1. 2.2. Composition, structure and properties characterization Crystal structure of the films was analyzed by grazing incidence X-ray diffraction (GIXRD, Philips, X’ Pert Pro) with Cu K␣ radiation ( = 1.540598 nm, 40 kV, 30 mA) in grazing mode (1◦ ) in the scanning range of 2 from 5◦ to 75◦ . The morphology and elemental composition were characterized by a field emission scanning electron microscopy (FESEM, JSM-6701F) equipped with an energy dispersive X-ray spectroscopy (EDS). The chemical composition was analyzed by Raman spectroscopy and X-ray photoelectron spectroscopy (XPS). The Raman signals were detected by a Jobin Yvon Horiba HR800 with spectrometer with an excitation wavelength of 532 nm at a resolution of 1 cm−1 . The XPS spectra were obtained using X-ray photoelectron spectroscopy equipped with an Ar ion sputtering gun and a monochromatic Al K␣ radiation at pass energy of 29.4 eV under a background vacuum of 2.0 × 10−7 Pa at room temperature. The XPS spectra were referenced with respect to C 1s line at 284.8 eV. The microstructure of films was observed by a high resolution transmission electron microscope (HRTEM, Tecnai G2 F20S-Twin) operating at 200 kV. The thin foils from film’s cross-section were prepared by mechanical polishing and then Ar ion-milling (Gatan 691) at a small angle with respect to the milled surface. To inspect the anti-oxidation capability, the films were firstly stored in humid air for 48 h at 50 ± 5% RH and 25 ± 2 ◦ C and then the chemical composition of the films was analyzed by XPS. Film hardness was evaluated using nano-indenter (TI950, Hysitron TriboIndenter, USA) with a Berkowich diamond tip. In order to exclude the influence of substrates, the indentation depth was controlled at 200 nm, about 10% of the film thickness. At least five replicate indentations were done for each sample. The indentation hysteresis curves were analyzed following the procedure proposed by Oliver and Pharr [34]. The film toughness was investigated by scratch tester (Kaihua MFT-4000) with a conical diamond tip of 0.2 mm radius and 120◦ taper angle. Adhesive strengths were obtained by averaging the results of the three different scratches for each sample. The minimum load at which the first flaking of the film inside the scratch track occurs was termed the adhesion strength. Scratch test was driven across the films deposited on the steel substrate at a continuously increasing loading rate of 1 N s−1 and the nominal maximum load was 50 N. The adhesion strength was determined by the friction force and acoustic emission (AE). Friction and wear performances of the films were evaluated using a ball-on-disk tribometer. The disk was the coated steel substrate. AISI 440C steel ball (HRC ∼ 60, Ra ∼ 0.10 ␮m) of 8 mm in diameter was used as counterpart, which was cleaned with alcohol before each test. The test conditions: normal load of 5 N, rotational speed of 1000 r min−1 , room temperature (20 ± 2 ◦ C) and (∼10%, 30% and 75%) relative humidity. After the friction test, the morphology and composition of wear track and wear scar were analyzed by FESEM, MicroXAM 3D non-contact surface mapping

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Fig. 1. GIXRD spectra of pure WS2 and WS2 –Ni composite films.

profiler (AD Corporation, Massachusetts, USA) and Raman spectroscopy. 3. Results and discussion 3.1. Structure Fig. 1 exhibits the GIXRD spectra of pure WS2 and WS2 –Ni composite films with different Ni content of 5.0, 10.3 and 16.5 at%, respectively. It can be seen that all the films mainly showed hcp-WS2 (1 0 1), (0 0 2), (1 0 3) and (1 1 2) peaks besides another belonging to the AISI 440C steel substrate, while no Ni signal is observed. Referred to the powdered WS2 [JCPDS Card No. 08-0237], it is found that for the pure WS2 film, the (1 0 1) and (1 1 2) peaks shift to high diffraction angles while the (0 0 2) peak shift to a low one. This phenomenon was generally observed from sputtered MoS2 film [10,36] and accounted for the presence of O atoms which was substituted for S atoms in the MoS2 lattice during film deposition. Meanwhile, the residual stress of the film could sometimes induce the shift of the diffraction peaks. For the WS2 –Ni composite films, the shift of (1 0 1), (1 1 2) and (0 0 2) peaks is smaller than that of pure WS2 film, suggesting that doping Ni is of benefit for decreasing the oxygen contents and changing the film stress. It can be confirmed by the EDS results (Table 1) that for the WS2 –Ni composite films the S/W ratios are at range of 1.72–1.78, and all the S/W ratios are higher than that of pure WS2 film (about 1.57). Furthermore, as the Ni content increasing from 0 to 16.5 at%, the intensity of both (1 0 1) and (0 0 2) peaks, especially (1 0 1) peak, decreases accompanied by broadening and strengthening of the WS2 (1 0 3) peaks. This suggests that Ni dopant restricted the growth of crystalline WS2 and induced the turbostractic stacking of (1 0 3) crystalline plane, similar to Cr, Ti, C and N doped WS2 films [19,21,37]. To further analyze the film structure, cross-sectional HRTEM lattice imaging was performed for the composite films on Si substrate, and the typical results of WS2 –5.0 at% Ni and WS2 –10.3 at% Ni films are shown in Fig. 2. Both the composite films exhibits a composite structure of crystalline and amorphous phase, but the amorphous phase is more remarkable at higher Ni content. The WS2 —5.0 at% Ni film shows a change in the preferred orientation of crystallite growth along the film thickness direction. The crystallites mainly exhibited a basal plane oriented growth in the substrate-near part [Fig. 2(a)], above which edge plane oriented growth was predominate [Fig. 2(b)]. However, the WS2 –10.3 at% Ni film is composed of the crystallites with random orientations, together with amorphous phase either in the substrate-near or substrate-far parts

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[Fig. 2(c) and (d)]. The HRTEM results verified that the crystallinity and edge plane oriented growth of the films were restrained significantly as the Ni content increasing. Fig. 2(e) is the selected area electron diffraction (SAED) pattern obtained from the corresponding zone in substrate-far part of the WS2 —10.3 at% Ni composite film. Clearly, the (0 0 2), (1 0 3) and (1 1 2) rings are seen to coexist with an extremely strong (1 0 1) ring. The identification of the 2H–WS2 phase and a mixture of different oriented crystallites were allowable, which were well consistent with the GIXRD results. Numerous studies [4,12,20] have been carried out to improve the tribological properties of MS2 based film by doping metal, such as Ti, Au and Cr, etc., but to our knowledge, the location and state of metal dopant was not fully acknowledged. Teer et al. [23] speculated that Ti replaced Mo in the MoS2 matrix, or intercalated its compounds around the MoS2 lattice, based on the investigation of the structure and tribological properties of MoS2 –Ti composite films. Other studies [12,38] also suggested that the metal dopants resided at the grain boundaries, such as Au, Fe and Ni doped MoS2 films. As for the chemical state of metal dopants in composite films, metal and oxidized states were also observed from different reports [18,35,38]. In this study, the XRD and HRTEM results mightily implied that the Ni was present as amorphous phase in the composite films, but could not verify its chemical state (Ni or Ni oxides). To solve this problem, the composite films were analyzed by Raman spectroscopy and XPS, and the typical results are shown in Fig. 3(a) and (b), respectively. Raman spectrum of WS2 —10.3 at% Ni composite film mainly represents two characteristic peaks at 352 and 418 cm−1 corresponding to WS2 , besides the peak positions at 698 and 809 cm−1 ascribed to WO3 [19]. No signal of Ni oxides can be detectable [Fig. 3(a)], which indicates that Ni dopant in the composite film was not present as oxidation state. From the XPS Ni 2p spectrum [Fig. 3(b)], it can be seen that the Ni 2p3/2 and 2p1/2 peaks centered at binding energy of 853.6 and 870.9 eV, respectively, corresponding to elemental Ni [39,40], but no NiO signal is detectable. So, it can be confirmed by the Raman and XPS results that the Ni dopant is in the presence of elemental state, not its oxides. The FESEM micrographs of pure WS2 and WS2 –Ni composite films are presented in Fig. 4. The pure WS2 film in Fig. 4(a) exhibits a typical morphology of sputtered pure MoS2 or WS2 films [38,41], characterized by a duplex layer microstructure consisting of a dense and coherent layer (approximately 100 nm) where the crystallites normally exhibited basal plane oriented growth, and above a loose columnar platelet layer where the crystallites predominantly had edge plane oriented growth. Moreover, the film shows a coarse needle-like surface morphology. The WS2 —5.0 at% Ni film shows a fiber-like structure along with a dentrite-like surface [Fig. 4(b)]. As the Ni content was higher (10.3 and 16.5 at%), the films were further densified and exhibited a featureless cross-sectional microstructure, as well as denser dentrite-like surface morphology [Fig. 4(c) and (d)]. The above results revealed that the structure of WS2 films was significantly influenced by Ni dopant. The change of film microstructure could be explained by a structure zone model (SZM, Thornton model) [19], developed to explain the microstructural variation of nanocomposite films by Thornton et al. [42]. As shown in Fig. 4(a), the pure WS2 film represents a typical zone 2 morphology exhibiting columnar platelet microstructure with significant porosity, mainly caused by sufficient surface diffusion of adatoms during film growth. Nevertheless, XRD, HRTEM, XPS and Raman results revealed that the composite films were composed of crystalline WS2 and amorphous Ni, and the Ni amorphous phase should reside at the boundaries of the crystalline WS2 . In this case, the Ni dopant could hinder the grain growth and stimulate a renucleation of the grains [19], and this effect became remarkable with the further increase of Ni content. Thus, the composite films

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Fig. 2. Cross-sectional HRTEM micrographs of the composite films: (a) and (b) obtained from the substrate-near/far regions of the WS2 —5.0 at% Ni film, respectively, (c) and (d) obtained from the substrate-near/far regions of the WS2 —10.3 at% Ni film, and (e) was SAED patterns taken from the middle of the WS2 —10.3 at% Ni film films, respectively.

Fig. 3. Raman and XPS Ni 2p spectra obtained from the WS2 —10.3 at% Ni composite film.

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Fig. 4. Surface and cross-sectional FESEM micrographs of (a) pure WS2 and composite films containing various Ni contents of (b) 5.0 at%, (c) 10.3 at% and (d) 16.5 at%.

exhibited a fiber-like structure at lower Ni content of 5.0 at% [Fig. 4(b)] and a denser featureless structure at higher Ni content of 10.3 and 16.5 at% [Fig. 4(c) and (d)]. Moreover, the XRD and EDS revealed that the composite films showed relatively high S/W ratio as compared with the pure WS2 film. The presence of oxygen atoms and residual stress in the films were contributed to the shift of the diffraction peaks, such as (0 0 2), (1 0 1) and (1 1 2) peaks in those WS2 composite films. The change of residual stress of the films might originate from the addition of Ni dopant [43]. Previous studies revealed that for the sputtered MoS2 films, the ratio of S to Mo was generally lower than stoichiometric ratio of 2. The S loss was mainly attributed to two factors. One [44] was that residual oxygen in vacuum chamber was substituted for S atoms in the MoS2 lattice during film deposition. The other [45,46] was the preferential resputtering of S from growing film and so the films had lower S/W or S/Mo ratio. In general, the two factors were coexisted during the film growth. Based on the first factor, composite films with higher S/Mo ratio could be achieved by co-depositing with easily oxidized metals, due to the function of preferential oxidation of metals [21]. However, in the present study, the Ni oxides was absent, indicating that for the composite films the higher S/W ratio was probably attributed to the second factor. During the co-deposition process of composite films, the plasma density was increased due to the introduction of Ni in comparison to that of pure WS2 , resulting in short mean molecular free path of argon particles. Thus, the rescattered argon particles with lower energies were obtained and their preferential resputtering to the S atoms in growing film was declined. Consequently, the composite films achieved relatively high S/W ratios.

3.2. Properties 3.2.1. Anti-oxidation capability in humid air To check the effect of Ni dopant on anti-oxidation capability of WS2 based films, the pure WS2 and WS2 –Ni composite films were analyzed by XPS after stored in humid air (50 ± % RH, 48 h). The typical XPS spectra of W 4f and Ni 2p obtained from the pure WS2 and WS2 –Ni composite films are shown in Fig. 5. As reported previously, the W 4f7/2 and W 4f5/2 peaks for W4+ in WS2 are at binding energy of 32.7 and 34.9 eV, and at 36.0 and 38.2 eV for W6+ in WO3 [21,47]. The binding energies of Ni 2p3/2 for Ni and NiO are located at 853.6 and 856.2 eV [48], respectively. It can be observed that the pure WS2 film was composed of WO3 and WS2 either at the film top surface or inner layer [Fig. 5(a)], indicating that the film was heavily oxidized. For the WS2 –5.0 at% Ni composite film [Fig. 5(b) and (c)], at the top surface the chemical components were a mixture of NiO, WO3 and WS2 , but the relative intensity of peaks corresponding to the WO3 were lower than those of WS2 , and in the inner layer the W element was mainly present as WS2 and the Ni existed as metallic state. The WS2 –10.3 at% Ni composite film shows the similar compositional variation trend of tungsten-containing compounds both at the top surface and in inner layers after stored in humid air [Fig. 5(d)]. This indicated that after the storage test in humid air, the pure WS2 film was oxidized either at the surface or in inner layer, but for the composite film, only the surface layer was oxidized partially. That is to say, the WS2 based composite films had better anti-oxidation capability than the pure WS2 film. Previous studies [6,7] revealed that the oxidation of the sputtered WS2 or MoS2 films in the humid air was caused by the

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Fig. 5. XPS spectra of the pure WS2 and WS2 –Ni composite films: (a) W 4f of WS2 film, (b) W of WS2 —5.0 at% Ni composite film, and (c) W 4f and (d) Ni 2p of WS2 —10.3 at% Ni composite film. 4f.

existence of unsaturated or dangling bonds, mostly sourced at the edge planes in layer structure. Furthermore, both the films normally had a loose columnar platelet structure as shown in Fig. 4(a), and the porosity between the platelets was in favor of transporting O or H2 O from the film surface to inner layer, so they were heavily oxidized as exposed to humid air [5]. XRD, HRTEM and Raman results revealed that the WS2 –Ni composite films were mainly composed of crystalline WS2 and amorphous Ni, which was resided at the grain boundaries/dendrite surfaces and could cover the partial unsaturated bond sites at the edge planes. Thus, the anti-oxidation capability of the WS2 based films was improved. Moreover, FESEM analysis revealed that due to the Ni dopant, the microstructure of the composite films was densified significantly and the loose columnar platelet structure was restrained. This indicated that the transport of the O or H2 O from the film surface to inner became difficult when the composite films were exposed to the humid air. So the oxidation was confined to the surface layers.

3.2.2. Hardness and adhesive strength The hardness and adhesive strength of the pure WS2 and WS2 composite films are showed in Fig. 6. As the Ni content is changed from 5.0 to 10.3 at%, the hardness of the composite films increases from 0.8 to 2.2 GPa significantly, higher than that of the pure WS2

film, about 0.5 GPa. However, there was no substantial difference in the hardness for composite films with further increase of Ni content from 10.3 to 16.5 at%. For the pure WS2 film, the low hardness was attributed to its loose structure with coarse columnar platelets.

Fig. 6. Microhardness and adhesive strength measured for the pure WS2 and composite films containing different Ni concentration.

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Fig. 7. Typical fracture morphologies after scratch test for (a) pure WS2 films, and composite films deposited at Ni contents of (b) 5.0 and (c) 10.3 at%.

Fig. 8. Average friction coefficient (a) and wear rate (b) of pure WS2 and composite films with different Ni content as a function of the relative air humidity.

The densification of the composite films and solid solution hardening effect, i.e. the distortion of the matrix lattice caused by the incorporation of doping metal atoms, were both considered to be responsible for the hardness enhancement of the MS2 –metal composite films [49]. The scratch results shows that, in comparison to the adhesive strength of 18 N of pure WS2 films, the WS2 –Ni composite film has an improved adhesion strength of 22.5 N with the suitable addition of Ni content (about 5.0 at%). However, too much dopant led to sharp decrease in the adhesive strength of the WS2 –Ni composite film. Generally, scratch testing is also employed to evaluate the toughness of the films [50]. Typical fracture morphologies of the initial stage of scratch for the composite films are shown in Fig. 7. It can be found that the scratch failures are disseminated along the

Fig. 9. Friction curves of pure WS2 and various WS2 –Ni composite films in humid air.

scratching direction and different failure mode occurs for various WS2 –Ni composite films. For the pure WS2 film [Fig. 7(a)], wear debris distribute in succession at the edges of the scratch track, originating from the easy plastic fracture of the coarse columnar platelets with significant pores [41]. Fig. 7(b) concerns the dense composite film with 5.0 at% Ni, which presents big plastic fracture debris distributed discontinuously at the edges of the scratch channel. The dense film with suitable Ni dopant delayed the film flaking from the substrate and so high adhesion strength was achieved. Besides, there is also small debris lying around scratch channel due to film’s brittle fracture to a lesser extent. However, for the composite film with high Ni content (above 10.3 at%), plentiful of brittle flaking occurs around the scratch channel even at the initial running of the conical diamond tip [Fig. 7(c)], indicating that the film had high brittleness, poor adhesion strength and low film toughness. The above results revealed that a suitable Ni dopant in the WS2 based film induced film brittleness to a lesser extent, but was able to improve adhesion strength effectively, which was in favor of improvement of film’s tribological properties. 3.2.3. Friction and wear Fig. 8 shows the mean friction coefficient and wear rate of the films after sliding of 3.0 × 104 cycles, as the humidity was varied from 10% to 70% RH. The results shows that for all the films, the mean friction coefficient is increased with the relative air humidity, from 0.04 in 10% RH condition, to 0.05–0.06 in 30% RH condition, and to 0.1–0.15 in 70% RH condition. Moreover, the WS2 —5.0 at% Ni composite film is less sensitive to humid air than both the pure WS2 and composite film with above 10.3 at% Ni [Fig. 8(a)]. The wear rates of the films also exhibits the same tendency with the change of test conditions. Meanwhile, the wear trend of pure WS2 > WS2 —16.5 at% Ni > WS2 —10.3 at% Ni > WS2 —5.0 at% Ni films is found in all the relative air humidity conditions [Fig. 8(b)]. For evaluating the wear life of the films in 30% RH condition, failure criterion was defined as friction coefficient () = 0.20. Fig. 9 shows the friction coefficient vs. sliding cycles for the pure WS2

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Fig. 10. SEM micrographs of wear track of (a) pure WS2 , and WS2 –Ni composite films with content of (b) 5.0, (c) 10.3 and (d) 16.5 at% after test in 30% RH condition.

and WS2 –Ni composite films against AISI 440C steel balls measured with ball-on-disk tester. For the pure WS2 film, the wear life is about 8.4 × 104 cycles. All the composite films show better wear resistance than the pure WS2 film, especially the WS2 –5.0 at% Ni composite film exhibiting the long wear life of 5.8 × 105 cycles about sevenfold better than pure WS2 film. Friction and wear tests revealed that the tribological performance of sputtered WS2 film could be improved significantly by doping Ni, but the Ni content must be carefully controlled to obtain the long wear life.

To elucidate the wear mechanisms of the films, detailed analysis of the wear morphologies of the worn surface and wear scan on the counterpart ball, and the chemical composition of relative transfer film were performed after sliding 3.0 × 104 cycles in 30% RH condition. Figs. 10 and 11 show the images of wear track and the corresponding 3D non-contact surface mapping of the pure WS2 and WS2 composite films. The pure WS2 film has undergone adhesive wear as well as severe plastic deformation which can be proved by the adhesive patches, and exhibits the maximum track width with bulky fragmentation wear debris accumulated along

Fig. 11. 3D non-contact surface mappings of the wear track of (a) pure WS2 , and WS2 –Ni composite films with Ni content of (b) 5.0, (c) 10.3 and (d) 16.5 at% after test in 30% RH condition.

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Fig. 12. SEM micrographs of wear scar of the AISI 440 C counterpart ball of (a) pure WS2 and WS2 –Ni composite films with Ni content of (b) 5.0, (c) 10.3 and (d) 16.5 at% after test in 30% RH condition, and (f) EDS spectrum of transfer film on the counterpart corresponding to WS2 —5.0 at% Ni composite film.

the track [10(a) and 11(a)]. The fragmentation originated from the columnar platelet fracture of the loose films and the remaining thin effective lubricating layer was contributed to the limited wear life [41]. The grooves of the wear surface [10(c) and (d) and 11(c) and (d)] indicated that the composite films with above 10.3 at% Ni had undergone severe abrasive wear and slight adhesive wear. Besides, brittle wear debris can be also observed around the wear track for the WS2 —16.5 at% Ni composite film, in good accordance with the scratch results. On the contrary, the WS2 —5.0 at% Ni composite film exhibits adhesive as well as abrasive wear to some extent [10(b) and 11(b)]. It can be also found that this kind of film have smaller wear track depth than the others well corresponding to the better wear resistance. It has been believed that transfer film is one of the important factors that influence the friction and wear properties of the lubricant films. Fig. 12 gives the SEM images of wear scar on the counterpart ball and the EDS spectrum of the transfer film for WS2 —5.0 at% Ni composite film. For all the films, a great deal of the wear debris can be observed from counterpart surfaces [Fig. 12(a), (c) and (d)]. Furthermore, the wear debris are more considerable and powdered

for the composite films at higher Ni contents (above 10.3 at%) as well as the pure WS2 film. In comparison, the effective transfer film against the WS2 —5.0 at% Ni composite film is large and compact, and surrounded by only small amount of wear debris [Fig. 12(b)]. The EDS result shows that the transfer film for the composite films with 5.0 at% Ni contains a certain amount of Ni dopant, which may be contributed to form the denser transfer film. Nevertheless, high brittleness went against the formation of effective transfer film and induced to high wear rate as the film doped with too much Ni. Raman results reveals that for the composite films the transfer film is a mixture of WS2 and partial WO3 , and mostly WO3 for the pure WS2 film, typical Raman results shown in Fig. 13. Polcar et al. [24,27,51] demonstrated that for the WS2 /MoSe2 composite films, transfer film adhered on the counterpart ball surface consisted of two separated layers. The upper layer was thin WS2 or MoSe2 film with basal planes oriented parallel to the surface that could play a dominant role on low friction coefficient in the sliding process. The lower layer was a mixture phase of random orientation WS2 (or MoSe2 ), amorphous WO3 (or MoO3 ) or/and oxide of metals. In humid air, WS2 could be oxidized to

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S. Xu et al. / Applied Surface Science 288 (2014) 15–25

Acknowledgments The authors gratefully acknowledge the National Key Basic Research Program of China (973) (grant no. 2013CB632300) for financial support.

References

Fig. 13. Raman spectra of the transfer film on counterpart surfaces for (a) pure WS2 and (b) WS2 —5.0 at% Ni composite films after wear tests.

non-lubricious WO3 because of the presence of unsaturated or dangling bonds at edge planes in the layered structure, and so its tribological properties was deteriorated. In present study, the better wear resistances of the composite films were mainly attributed to the formation of effective transfer film and its lubrication function, which resulted from the improvement of anti-oxidation properties confirmed by XPS analysis (Fig. 5). For the pure WS2 film, the transfer film was quickly oxidized in humid air under tribo-chemical interaction and the production was mainly WO3 (Fig. 13). The nonlubricious WO3 could influence the further formation of effective transfer film, which caused high wear. For the composite films, due to the improved anti-oxidation properties, the transfer film could sustain lubricating action for longer duration and the further formation of tenacious and adherent transfer films was less influenced by the WO3 , and hence the wear resistance was improved significantly.

4. Conclusion The WS2 –Ni composite films comprising of crystalline WS2 and amorphous Ni were co-deposited by RF-sputtering technique. Ni doping resulted in a microstructure change from columnar platelet structure of pure WS2 film to fiber-like structure at lower Ni content (5.0 at%), and to featureless structure at higher Ni content (>10 at%), and the films became more and more compact. As a result, the composite films exhibited better anti-oxidation capability in humid air as compared with the pure WS2 film. Meanwhile, the films exhibited a significant improvement in hardness (from 0.5 to 2.2 GPa) and the hardness achieved the maximum values of 2.2 GPa as the Ni content above 10.3 at%, but excessive Ni induced high brittleness and the films flaking easily from the substrate even at low scratch contact load. A small amount of Ni dopant could improve the wear resistance of WS2 based film and the film with Ni content of 5.0 at% exhibited best tribological properties and had about sevenfold better wear life than the pure WS2 film in humid air environment. At high contents (>10 at%), Ni dopant made the films brittleness and wear resistances became worse.

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