Mo-S-Ti-C Nanocomposite Films for Solid-State Lubrication - ACS

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Mo-S-Ti-C Nanocomposite Films for Solid-State Lubrication Zewen Duan, Xiaoyu Zhao, Zhenggang Nai, Li Qiao, Jiao Xu, Peng Wang, and Weimin Liu ACS Appl. Nano Mater., Just Accepted Manuscript • DOI: 10.1021/acsanm.8b02184 • Publication Date (Web): 18 Jan 2019 Downloaded from http://pubs.acs.org on January 18, 2019

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ACS Applied Nano Materials

Mo-S-Ti-C Nanocomposite Films for Solid-State Lubrication Zewen Duana,b, Xiaoyu Zhaoa,b, Zhenggang Naic, Li Qiaoa,b, Jiao Xu*a, Peng Wang*a, Weimin Liua a

State Key Laboratory of Solid Lubrication, Lanzhou Institude of Chemical Physics, Chinese

Academy of Sciences, Lanzhou 730000, China b Center

of Materials Science and Optoelectronics Engineering, University of Chinese Academy of

Science, Beijing 100049, China c School

of Stomatology, Lanzhou University, Lanzhou 730000, PR China

Postal address: No. 18, Tianshui middle road, Lanzhou 730000, China

*E-mail: [email protected] ORCID iDs: 0000-0002-8128-5372

Abstract A Mo-S-Ti-C nanocomposite film with both high-crystalline MoS2 in (002) basal planes and high hardness was successfully synthesized in the condition of pure MoS2 films growing with edge-terminated surface. The evolution of their structure and the transformation of crystal orientation during film growth were investigated by X-ray diffraction (XRD) patterns and focused ion beam combing with the transmission electron microscopy (FIB&TEM). TEM revealed that the MoS2 edge sites usually considered as catalytic materials in catalysis field can provide active sites for titanium and carbon atoms aggregating in deposition. This aggregation of dopants not only largely passivates the activated MoS2 edge sites and induce the initial (100) textured MoS2 lamellar gradually transforming into (002) orientation, but also effectively reduce the damage the dopants do the highly ordered MoS2 crystallites. As a result, a Mo-S-Ti-C composite film with well-crystallized MoS2 in (002) basal planes was formed. The nanoindented experiments reveal that the hardness of this nanocomposite film reaches up to 6.81 GPa. Benefiting from the enhanced hardness and well-crystallized MoS2 in (002) basal planes, the corresponding pin-on-disk tribotests indicate that this nanocomposite film exhibits low friction coefficient (about 0.02) and high wear resistance whether in vacuum or ambient air. Keywords：

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Edge

terminated

MoS2

film;

structural

evolution;

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growth

mechanism;

high-crystalline; mechanical properties; tribological properties. 1. Introduction The edges of layered materials, such as MoS2 and MoSe2, have aroused great interest in recent years because of its excellent properties and wide applications in some nanodevices and functional nanomaterials. Particularly, one most interest is to prepare highly efficient catalyst because of their highly intrinsic surface energy (25000 mJ/m-2) [1]. Typically, the efficient hydrogen evolution reaction (HER) correlates directly with the density of the exposed edge sites of MoS2 films [2], electrochemical ammonia preparation via nitrogen reduction reaction on

MoS2

catalysts [3] and the productive photocatalytic H2 productions on UiO-66/CdS hybrids modified by the catalysts[4]. In contrast, compared with the diverse applications of edge sites of MoS2 basal planes in catalytic reactions, less attentions have been made to utilize the edges of these layered nanomaterials in the fields of tribology because of its full of dangling bonds and sensitivity active to oxidation [5, 6, 7]. In comparison with edge terminated MoS2 basal planes (Type Ⅰ, (100) texture), another type of surface sites, terrace terminated MoS2 basal planes (Type Ⅱ, (002) texture, shown in Fig. 1) with minimal dangling bonds, covalent bonding in S-Mo-S basal planes and well arranged basal planes parallel to the substrate exhibits denser structure and less-sensitive oxygen properties [8-11]. Therefore, insuring pure MoS2 films growing with Type Ⅱ texture was widely considered as an essential condition to fabricate MoS2-based composite films with excellent lubricant properties in previous all studies. For example, some self-assembled or manual nanomultilayer films, such as Mo-S-C, Mo-S-Ti, Wo-S-C and Mo-S-C-N films [12-16] were synthesised in the condition of pure MoS2 films growing with Type Ⅱ texture. Some amorphous composite films, such as Ti [17-22] or Cr [23, 24], or some nonmetal elements, such as nitrogen [25-27] or carbon [22, 23, 24, 27] doped MoS2 films, were also fabricated in the condition of pure MoS2 films growing with Type Ⅱ texture.


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Unfortunately, there are two great challenges for fabricating MoS2-based composite films with preferential Type Ⅱ basal planes. First, it’ hard to fabricate pure MoS2 films with dominant Type Ⅱ texture by magnetron sputtered-systems, since most deposition processes will lead to the formation of Type Ⅰ films [9, 10, 11]. Buck once reported that the partial pressure of H2O (PH2O) plays a crucial role in determining the structure MoS2 films. The Type Ⅰ films were obtained at high PH2O, and the Type Ⅱ films were obtained at low PH2O [28]. While at low working pressure, the S/Mo ratio of this film is less than 1 because of the fierce bardment of some neutral argon atoms reflected on the targets [29], which results in difficulty in the formation of the tribofilm. Muratore et al. [30] reported that deposition processes with ion energy ＞ 175 eV greatly promote the growth of MoS2 crystallites in (002) orientation. Therefore, preparing pure Type Ⅱ films as the requirement of some special equipments and fussy deposition process is becoming a great challenge for researchers. Highly ordered intrinsic lamellar structure of MoS2 phases ensures low friction performance of deposited film, and the amorphous layers provide hardness phase for supporting the upper soft lubricant phase, both of which play key roles in determining the wear live of deposited film. Therefore, another major obstacle to lubrication by MoS2 is that even with a small amount of elements into Type Ⅱ MoS2 films, the highly ordered intrinsic lamellar of MoS2 was broken inevitably. Namely, the enhanced hardness by doping elements is achieved with a sacrifice on the reducing highly ordered intrinsic lamellar structure of MoS2. As reported by Xu, with the carbon content, the (002) peaks of MoS2 crystallites become broadened and the intensity was reduced clearly because of the serious loss of highly ordering of MoS2 crystallites [12-16]. Multilayer film, as reported in previous studies [12-15], the MoS2 (002) basal planes in nanomultilayered MoS2-riched layers were largely broken and become disordered even in the low dopant content, a phenomenon also observed in Mo-S-Ti composite films [31, 32], which creats a great challenge for forming a self-adaptive tribofilm in sliding.



Fig. 1 Type Ⅰ and Type Ⅱ orientations All MoS2-basied composite films reported before were fabricated by doping dopant into the pure MoS2 films with terrace terminated surface. Here, a Mo-S-C-Ti composite film with both highly MoS2 crystallites (002) basal planes and high hardness was successfully synthesized in the condition of pure MoS2 films growing with edge-terminated surface. We surprisingly find that the edges of MoS2 films can be maximumly utilized in tribology, and it is to say that compared with the terraced-terminated surface, the edge-terminated surface can provide active sites for carbon and titanium metal aggregating. This aggregation largely passivates the active sites and effectively inhibits the (100) textured MoS2 film growth. Meanwhile, it greatly reduces damage the dopants do MoS2 crystallites in (002) basal planes. MoS2 based films cross sections were investigated by using of high resolution transmission electron microscopy (HRTEM), and the chemical compositions and microstructure of the composite films were studied by using of Energy dispersive spectroscopy (EDS), X-ray diffraction spectroscopy (XRD) and X-ray photoelectron spectroscopy (XPS). The tribological and mechanical properties of MoS2-based films containing various dopants have also been studied. Particularly, based on the analysis of HRTEM images and XRD spectrum of MoS2-based composite films, we carefully explained the MoS2 structural transition (from (100) to (002) texture) in deposition and concluded that it’s reasonable to prepare MoS2-based films with excellent tribological properties in the condition of pure MoS2 film growing with (100) basal planes, which may provide a novel and easier method for developing of MoS2-basied composite materials with highly MoS2 crystallites in (002) basal planes.


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2. Experimental methods 2.1 Preparation of Mo-S-C, Mo-S-Ti and Mo-S-Ti-C composite films. The MoS2 based composite films were sputtered on monocrystalline silicon (100) substrates by using of a r.f. (ratio frequency) magnetron sputtering system composed of four controllable magnetron sputtering sources. In this experiment, two of sputtering sources hold titanium and carbon targets respectively, and the other two MoS2 targets. Prior to deposition, all substrates for films growth dipped in acetone for 10 minutes and followed by ultrasonically cleaning in alcohol for 20 minutes, and after that, they were placed at the substrate holder which is about 60 mm away from the sputtering targets. The sputtering chamber vacuum was continually decreased to 1.0×10-3 Pa, and after that, the substrates were etched by Ar+ ions about 10 minutes to eliminate oxides adsorbing on the substrates surface at a substrate bias voltage of -500 V. During deposition, the gas flow was fixed at 40 sccm, the working gas pressure 0.65 Pa, and sample stage rotational speed 9 rev./min. The r.f. (ratio frequency) sputtering power applied to the MoS2 target was adjusted to a constant of 250 W, and the r.f. sputtering power applied to carbon target was increased from 200 W to 400 W to prepare Mo-S-C composite films with various amounts of carbon. To fabricate Mo-S-C-Ti composite films, the d.c. (direct current) sputtering power applied to titanium target was maintained as a constant at 0.11 A, and the r.f. sputtering power applied to carbon target was increased from 50 W to 200 W. The total deposition time for each batch was controlled within 120 min. Besides, prior to deposition, a Ti interlayer in thickness about 100 nm was first deposited by d.c. sputtering power (0.5 A) to enhance the adhesion between films and substrates. 2.2 Characterization of composite film structure Firstly, the high resolution transmission electron microscopy (HAADF-HRTEM, TECNAI G2 S-TWIN F20, FEI, USA; accelerating voltage 200 kV) were used to clearly observe MoS2 crystals nanostructure and their evolutions in composite films growth. Besides, Energy dispersive spectroscopy (EDS, JSM-5601LV, JEOL, Japan), was applied to further investigate the chemical compositions. To observe the crystallographic phase changes of MoS2 based composite films, the XRD experiments



were conducted by using of grazing incidence X-ray diffraction (GIXRD, Rigau RINT2400) equipping with Cu Kα radiation λ = 1.54056, together with the acquired diffraction patterns from 10° to 90°. Finally, X-ray photoelectron spectroscopy (XPS) equipping with Al Kα irradiation was carried out to measure the chemical bonds changes of these sputtered films. 2.3. Characterizations and mechanical and tribological properties A pin-on-disk tribotester equipping with a vacuum chamber was used to investigate the tribological behaviors of all composite films. A GCr15 steel ball with 3 mm in diameter was applied as a counterpart, together with a normal load within 3 N (Hertz contact pressure about 1.5 GPa) applying in all friction experiments. Besides, a high-pressure vacuum for the HV (high vacuum) friction testing was controlled as 9.0×10-3 Pa, and the relative humidity in tribometer after pumping was about 25%. The disk rotational speed for all tribotests was controlled as 1000 rev/min (sliding velocity of 0.52 m/s) together with a wear track in radius of 3 mm. A Nano Indenter DCM nano-mechnical system ((MTS, America) with Berkovich diamond indenter was used to investigate the nanoindented hardness and Young’s modulus of sputtered films, and besides, the maximum indentation depth was set to be less than 10% of bulk film thickness to avoid any effects induced by substrate deformation [27]. Finally, five repeated trials were performed at each sample different regions for average evaluation. 3. Results and Discussion 3.1 Chemical and structural characterization. The corresponding chemical compositions and deposition conditions of MoS2 based composite films analyzed by EDS are presented in Table 1. It can be seen that, the S/Mo ratios of all MoS2 based composite films are low than 2. As reported by previous studies [12, 13], the sub-stoicheiometry of S and Mo in MoS2 based composite films was mainly due to the preferential reaction of sulfur atoms with H2 and O2 in residual atmosphere, together with the preferential sputtering effects of sulfur caused by the collision response of argon atoms reflected on the targets. Table 1. The deposition conditions and chemical composition of MoS2 based


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composite films Process PAr (Pa) T(℃) 150 150 150 150 150 150 150 150

0.65 0.65 0.65 0.65 0.65 0.65 0.65 0.65

Target power (W/A) C target MoS2 Ti target (W) target (W) (A) 250 200 250 300 250 400 250 250 0.11 50 250 0.11 100 250 0.11 200 250 0.11

Chemical composition (at.%) C Ti S/Mo 2 19 30 41 9 13 18

4.1 4.4 4.1 3.8

1.87 1.78 1.81 1.71 1.83 1.79 1.74 1.81

Fig.2 (a) XRD patterns of Mo-S-C, Mo-S-Ti films and Mo-S-Ti-C composite films sputtered at various graphite target power (50 W～200 W). (b) The intensity ratio of (100) and (002) peaks over the summation intensity of each spectrum for MoS2, Mo-S-C, Mo-S-Ti and Mo-S-Ti-C films, I(100)/I(002)+I(100) and I(002)/I(002)+I(100). Fig. 2 shows the XRD patterns of all sputtered MoS2 based composite films. It can be seen that pure MoS2 films exhibited one dominant peak located at 33.5° together with two weak peaks located at 13.7 and 58, which corresponds to the (100), (002) and (110) crystal orientation, respectively. Except for these dominant peaks mentioned above, some weak peaks locating at 36.1°, 39.3°, and 61.3°, which corresponding to the (102), (103) and (008) orientation of MoS2 crystals, also appeared on the XRD patterns. While, with addition of dopants into pure MoS2 films, a change of preferred orientation from (100) to (002) MoS2 basal planes was clearly observed. At the low graphite sputtering power of 200 W, the weak (002) peaks



disappeared and only leaving the dominant (100) peaks in Mo-S-C composite film, which provides a hint of internal structure of the columnar grains in Mo-S-C film. When graphite sputtering power was increased from 200 W to 400 W, the intensity of (002) peaks of the Mo-S-C composite films were enhanced gradually. In comparison with carbon atoms, with a small amount of titanium (4 at.%) into sputtered MoS2 films, (100) peaks were exhaustively inhibited, only leaving of dominant (002) peak in Mo-S-Ti film. Furthermore, when doped carbon atoms into this Mo-S-Ti film with dominant (002) texture, the intensity of (002) peak is continually intensified and increases to the maximum at graphite power of 200 W, indicating an intensified reorientation of MoS2 crystals in (002) texture in simultaneous co-sputtering of carbon and titanium atoms. Unlike the previous results of reduced and broadened (002) peak in Mo-S-M composite films (M: C, Ti, Au, Ag, Cr et al.), the dominant (002) peak in Mo-S-Ti film was further intensified and became sharpen after doping carbon atoms, indicating that co-sputtering of carbon and titanium atoms did not cause severe damage to the highly ordered structures of MoS2 crystallites, a result never being found in previous studies. XPS was used to investigate the chemical bonding in Mo-S-C and Mo-S-C-Ti composite films and the results was treated with the XPSPEAK41 software using a Shirley background and Guassian-Lorentian function in ratio of 80%:20%. The C 1s, Mo 3d, and S 2p spectrum lines of Mo-S-C composite films deposited at different sputtering power applied on graphitic targets are shown in Fig. 3, where the XPS results of Mo-S-C composite films are quoted from our previous work [33]. Besides, the C 1s, Mo 3d, S 2p and Ti 2p spectrum of Mo-S-C-Ti film deposited at Titanium power 0.11 A and graphite power 200 W are also presented. In Fig. 3(a1), the C 1s spectrum are deconvoluted into two peaks, of which C=C (284.6 eV) and C-C (288.4 eV) bonds respectively. The S 2p spectrum, as shown in Fig. 3(a2), also exhibits a doublet binding energy, 163.6 and 162.4 eV respectively, which corresponds to the S 2p1/2 and S 2p3/2 spectral lines of S2- in MoS2 crystals, together with the peaks locating at 161.6 and 162.8 eV corresponding to the S 2p3/2 and S 2p1/2 in MoS2-x, which fully


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explained the sub-stoicheiometry (＜ 2) of S to Mo [34]. The Mo 3d peaks locating at the binding energy of 231.9 and 228.8 eV, are consistent with the spectrum of Mo 3d3/2 and Mo 3d5/2 in MoS2 (Mo4+) respectively. Because of the substoichiometric ratio of S to Mo in MoS2-x, the Mo 3d spectrum was deconvoluted into peaks at 231.2 and 228.1 eV, corresponding to the Mo 3d3/2 and Mo 3d5/2 in MoS2-x respectively, which agrees well with the results of S 2p spectrum above. Besides, it can be seen from Mo 3d (see Fig. 3), a small hump appeared at around 226.1 eV, which corresponds to the S 2s peak in the Mo 3d spectrum [34]. With increasing power applied on graphite target power from 200 W to 400 W, no visibly change is found on the C 1s, S 2p and Mo 2d spectrum, and no formation of CS2 and C-Mo chemical bonding in composite films, as shown in Fig. 3(a2, a3). Furthermore, when doped low level of titanium (4 at.%) into Mo-S-C film, an obvious change was found for all spectrums, as shown in Fig. 3(b1-b3). In Fig. 3 (b1), there are visible decoupling peaks in C 1s spectrum and an shifting peak locating at 283.9 eV corresponding to graphene peaks shifting down from 284.5 to 283.9 eV, which can be explained by that of the intercalation of oxygen under graphene leading to its decouping and the shifting peak at 283.9 eV consistenting with the parts of free-standing graphene [35]. Besides, as shown in Fig. 3(b1), there is a tiny shoulder peak locating at approximately 281.4 eV, which represents the Ti-C bond formation in deposition. In Fig. 3 (b2), the deconvoluted S 2p spectrum at the binding energy of 162.8 and 161.6eV corresponding to S 2p1/2 and S 2p3/2 in MoS2-x disappeared, only leaving S 2p1/2 and S2p3/2 spectrum lines of S2- in MoS2, which indicates that the higher ratio of S to Mo this Mo-S-C-Ti film exhibits in comparison with the ratio in Mo-S-C composite film. The Ti 2p spectra of Mo-S-C-Ti films composed of peaks around 460.8, 461.5, 459.8, 456.9, 455.1 and 454.8 eV was separated, as shown in Fig. 3(d4). According to the studies reported before [27, 34], (Ti 2p1/2) 461.5 eV and (Ti 2p3/2) 455.15 eV are mainly ascribed to oxide states of titanium metal, which suggested that titanium atoms at the surface have been oxided largely. Besides, the other doublet peaks located at Ti (2p3/2) 454.8 eV and Ti (2p1/2) 460.8 eV binding energies are attributed to Ti-C bonds [34, 36], which agrees well with the Ti-C peaks (284.1 eV) in C 1s spectrum.



Fig. 3 XPS spectrum of Mo-S-C and Mo-S-C-Ti composite films (Mo 3d, S 2p, C 1s and Ti 2p): (a1-a3) [33] Mo-S-C composite films deposited at different graphite power (200 W～400 W). (b) Mo-S-C-Ti film deposited at graphite power of 200 W as well as Titanium power of 0.11 A. 3.2 Nanostructural MoS2 crystals evolution during film growth 3.2.1 Mo-S-C and Mo-S-Ti films growth


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Fig.4 (a1) The cross-sectional HRTEM observations of Mo-S-C (graphite sputtering power 200 W), (b1) Mo-S-Ti (Titanium sputtering power 0.11 A), and (c1) Mo-S-Ti-C (graphite sputtering power 200 W and Titanium sputtering power 0.11 A respectively) composite films, with insets showing the cross-section images of whole



films. (a2), (b2), (c2) Their corresponding high-magnification TEM images of Mo-S-C, Mo-S-Ti and Mo-S-Ti-C films. As exhibited in XRD patterns above (Fig. 2(a)), with a small amount of carbon atoms (19 at.%) doping into sputtered MoS2 films within dominant (100) texture, the initial (002) basal planes in MoS2 films disappeared and were completely transformed into (100) texture. In contrast, compared with the carbon atoms, with smaller addition of titanium atoms (4 at.%) into pure MoS2 film, the MoS2 basal planes in (100) orientation disappeared completely, only leaving the (002) textured MoS2 in Mo-S-Ti composite film. To investigate the transformation of MoS2 crystals orientation in Mo-S-Ti and Mo-S-C composite films growth, FIB&SEM system was used to prepare samples for cross-section HRTEM observations. The cross-sectional TEM images of Mo-S-C composite film sputtered at graphite power 200 W are presented in Fig.4 (a1, a2). It can be seen that a complete nanocylindrical structure composed of MoS2 enrichment phases (coded as-α) and carbon rich phases (coded as-β) arranging alternately in vertical patterns was formed. These MoS2 and carbon-rich phases, wider than 20 nm, growed more than hundreds of nanometer in a direction perpendicular to the films surface and without any stagnation during film growth, which is similar to previous work [33]. The nanocylindrical MoS2 lamellars, as reflected by the inset of Fig. 4((a2)), being not inhibited by the carbon atoms, are well crystallized with (100) texture, which is in good agreement with the dominant and sharp (100) peak shown in Mo-S-C film XRD patterns. Similarly, the SAED pattern exhibits obvious diffraction rings of deposited MoS2 crystals dominant in (100) orientation but the weak rings of (002) orientation, shown in Fig.5 (a1). Beisdes, it can be seen that the carbon enrichment phases (inset of Fig. 4(a2)) are mostly composed of amorphous carbon matrix together with the highly distorted MoS2 nanoclusters, a nanostructure similar to the ‘capping layers’ in Mo-S-C nanoperiod multilayers [22, 23]. The cross-sectional TEM images of Mo-S-Ti composite film deposited at titanium power of 0.11 A are shown in Fig. 4 (b1, b2). It can be seen that in comparison with the carbon atoms, titanium atoms exhibit stronger ability in the inhibition of the MoS2


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phases within dominant (100) texture. Namely, with a small addition of titanium atoms to the MoS2 film, the (100) textured MoS2 basal planes disappeared and were completely transformed into (002) basal planes. Besides, it’s noted that titanium atoms in Mo-S-Ti film not aggregated forming rich phase but dispersed homogenously in MoS2 basal planes. Therefore, these MoS2 crystals within dominant (002) basal planes lost the long-range ordering of MoS2 crystallites and exhibited short-range ordered nanostructure, which can fully explain the dominant but broadened (002) peaks shown in Mo-S-Ti film XRD patterns. In contrast with the SAED characterization of Mo-S-C composite film, the SAED patterns from Mo-S-Ti film also show an inverse trend: the clear ring of (002) orientation, together with weak ring of (100) orientation, shown in Fig.5 (a2).

Fig.5 (a1) The corresponding diffraction patterns collected from Mo-S-C film (a1), Mo-S-Ti film (a2) and Mo-S-Ti-C film (a3) respectively. 3.2.2 Mo-S-Ti-C films growth Further observations of the Mo-S-Ti-C composite film XRD patterns above (Fig. 2(a)), with increasing addition of carbon atoms into this Mo-S-Ti film, the diffraction



angle of (002) peaks were continually intensified and exhibited a sharper tendency with the carbon content, and the (002) peak intensity in Mo-S-Ti-C film increases to the maximum at graphite target power of 200 W. This indicates that the disordered MoS2 crystallites in (002) basal planes in Mo-S-Ti composite film were well-crystallized after doping of carbon atoms. To investigate the transformation of MoS2 crystals orientation in Mo-S-Ti-C film growth, FIB&SEM system was again carried out to fabricate samples for closer observations, and the cross-sectional HRTEM images of the Mo-S-Ti-C composite film deposited at graphite power (200 W) and titanium power (0.11 A) are shown in Fig. 4(c1, c2). As reflected by the insets of Fig. 4(a1, b1, c1), the Mo-S-Ti-C composite film exhibited lowest thickness in comparison with the Mo-S-C and Mo-S-Ti composite films (It can be attributed to the film densification by co-doping of carbon and titanium atoms). Besides, it’s surprisingly find that compared with the Mo-S-Ti composite film with loss of long-range of MoS2 crystallites, the MoS2 crystals in Mo-S-Ti-C film exhibit an ordering tendency, and these highly-ordered MoS2 crystals in (002) orientation stacked approximately in thickness about 20 nm, as shown in Fig. 6(c2). Further observations of Mo-S-Ti-C film, as demonstrated in high magnification images (Fig. 6(c3)), found that this Mo-S-Ti-C composite film is mostly composed of alternating structures with α, β and γ phases. Unlike the previous MoS2-based composite films [22-25, 37], where carbon and titanium atoms usually disperse homogenously in MoS2 crystals basal planes, in this case, carbon and titanium atoms exhibit a tendency to agglomerate around (100) textured MoS2 crystallites, forming (100) textured MoS2 based Mo-S-Ti-C compositional mixed phases (denoted as α phase). With film growth, the MoS2 crystals in (100) basal planes began to deflect, forming β phases. Finally, based on the β phases, the MoS2 crystals in (100) basal planes were completely transformed to (002) basal planes, forming nanocrystals composed of finely-ordered MoS2 molecular layers (denoted as γ phase), which are responsible for the dominant and sharp (002) peak shown in Mo-S-Ti-C film XRD patterns. In comparison with the SAED patterns from Mo-S-Ti film, the diffraction rings of (002) basal planes in Mo-S-Ti-C composite film become more clear, supporting the XRD


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observations. According to the observations of nanostructural evolution involved in Mo-S-Ti-C film, it can be concluded that a small amount of MoS2 reactive edge sites ((100) basal planes) can efficiently provide active sites for carbon and titanium atoms aggregating during film growth, and these aggregation largely reduce the homogeneous dispersion of dopants in MoS2 basal planes, preventing the damage the dopants do MoS2 crystallites and protecting the long range MoS2 crystallinity in (002) basal planes, which may ensure an easy formation of a protective film in sliding and a long wear live.

Fig.6 The high magnification images of Mo-S-Ti-C composite film 3.3 Growth mechanisms of nanostructured MoS2 based composite films 3.3.1 The enhanced (100) texture in Mo-S-C film & the inhibited (100) texture in Mo-S-Ti film With small addition of carbon (19 at.%) atoms into the sputtered MoS2 in (100) planes, the (100) peak intensity was not inhibited but enhanced, a phenomenon having been observed in our previous work [33]. It was explained detailedly as follows [33]: when pure MoS2 crystallites are deposited with edge-terminated growing surface, the edge sites of the crystallites always expose to the film surface and exhibit high surface energy. The incident carbon atoms with weak interaction with edge sites may have enough activation energy to diffuse along the MoS2 edged terminated surface to lamellar gaps, as shown in Fig. 7 (a1), forming nanostructure composed of highly



ordered MoS2 rich phase encapsulated by carbon matrix (Fig. 4 (a1)). Because of MoS2 crystallites activate edge-sites not being covered by carbon atoms still exposing to the surface in film growth, accompanying with a tendency of minimizing MoS2 active energy in film growth, the incidient MoS2 moleculars would spontaneously bind themselves to activate region, which promoted MoS2 basal planes perpendicular to the film surface extending up to tens of nanometers, together with the enhanced intensity of (100) peak in Mo-S-C film. While, in comparison with the carbon atoms, the titanium atoms exhibit larger atomic radii [35], almost twice than that of carbon atoms, which are expected to have slower motion and longer retention on the MoS2 edge sites. Besides, as J. Moser reported before [9-11], the root cause of the inhibition of (100) texture was the elimination of MoS2 reactive edge sites during film growth. Based on the analysis above, here, we attributed the inhibition of (100) texture in Mo-S-Ti film growth to the efficient coverage of titanium atoms to the MoS2 edge sites. This evolution process was explained detailedly as follows. When MoS2 crystallites grow with edge terminated surface, unlike the carbon atoms behavior mentioned above, a portion of inciding titanium ions with slower motion having no enough time to diffuse must reside onto these active sites. As a result, because of the residing of titanium atoms on edge sites, the MoS2 active sites were largely covered, and the high surface energy were passivated effectively (Fig. 7(b1, b2)). With the elimination of the active edge sites in film growth, the sputtered MoS2 in (100) basal planes were gradually transformed into (002) basal planes. Once MoS2 crystallites transforming into (002) orientation, the titanium serving as solution, would disperse homogenously in MoS2 basal planes [31, 39], which results in the loss of long-range ordering of MoS2 crystallites and broadened (002) peaks.


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Fig.7 Proposed schematics of the growth mechanism of MoS2-based composite films: (a1, a2) nanocylindrical Mo-S-C film (b1, b2) Mo-S-Ti film and (c1, c2) Mo-S-Ti-C nanocomposites. 3.3.2 Further enhanced (002) texture in Mo-S-Ti-C films Compared with the Mo-S-Ti composite film, the Mo-S-Ti-C nanocomposite film exhibits sharper (002) peaks (Fig. 2 (a)) and enhanced MoS2 crystallinity. This indicates that the titanium atoms homogeneously dispersing in Mo-S-Ti composite film may accumulate after doping of carbon atoms, which reduced destroy the dopants to the MoS2 crystallites. Besides, according to the XPS spectrum of



Mo-S-Ti-C composite film, there are some Ti-C bonds forming in deposition. Based on the analysis above, this enhanced (002) texture in Mo-S-Ti-C film can be explained as follows. When MoS2 crystallites grow with (100) texture, in comparison with the carbon atoms, the titanium atoms with large atomic radii, as precursor, may preferentially reside on the MoS2 reactive sites to minimize the system energy. Subsequently, the inciding carbon atoms, having negligible interaction with the exposed MoS2 edge sites, may be catched by titanium atoms as they traveled along the MoS2 edge sites and formed a spot of TiC (Fig. 7(c1, c2). As a result, the carbon and titanium clusters with partial TiC are formed along the edges of MoS2 crystallites (Fig.6, α phase), which efficiently covers the active sites and passivates the high surface energy of the MoS2 edge sites. On the passivated interface, with loss of the active sites, the MoS2 crystals in (100) basal planes were transformed into (002) basal planes (Fig. 6, γ phase). Because of the carbon and titanium atoms aggregating along the MoS2 edge sites, the (002) textured MoS2 nanocrystals maintained their ordered MoS2 molecular layers and exhibited enhanced crystallinity. Compared with the pure sputtered MoS2 dominant in (100) basal planes, the pure sputtered MoS2 films growing with (002) texture [12-16, 34, 38] were lack of active edge sites for dopants aggregating, therefore, the carbon and titanium atoms have no locations to accumulate but to distribute dispersedly into the MoS2 basal planes, which greatly destroyed the MoS2 crystal structure and lead to the amorphous characterization. Based on the analysis above, it can be concluded that it’s exactly the existence of these edge sites in film, which provides active sites for carbon and titanium atoms preferentially aggregating on the edge sites of (100) textured MoS2 crystallites. This aggregation not only passivates the MoS2 edge sites, but also greatly reduces the damage the dopants do the (002) textured MoS2 crystallites. 3.4 Nano-indentation measurements and tribological tests The mechanical properties of films are evaluated by Nano Indenter DCM nano-mechnical system ((MTS, America) equipping with Berkovich diamond indenter, and the hardness and Young’s modulus of pure MoS2 film and MoS2 based composite films, including Mo-S-C, Mo-S-Ti and Mo-S-Ti-C, are present in Fig.8. As


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shown in Fig. 8(a), the pure MoS2 films exhibit a low hardness and Young’s modulus, which are attributed to its high porosity. With addition of the carbon (deposited at the graphite sputtering power of 200 W), the hardness of this Mo-S-C composite film with dominant (100) basal planes was enhanced to 1.96 GPa because of the existence of nanocylindrical amorphous carbon matrix. For Mo-S-Ti composite film (Ti, 4 at.%), it can be noted that the film hardness is significantly increased to 7.53 GPa, which is much higher than the hardness of pure MoS2 and Mo-S-C composite film due to the existence of titanium dopant together with its denser (002) basal planes [22, 23]. However, when doped carbon into this Mo-S-Ti film, it’ surprisingly finds that this Mo-S-Ti-C film exhibits a decreasing trend in hardness in comparison with Mo-S-Ti film, from 7.53 to 6.81 GPa. It can be explained by the fact that due to enhanced MoS2 crystallinity caused by the aggregation of the dopants, the densification in microscale and amorphization of this Mo-S-Ti-C film were slightly reduced in comparison with Mo-S-Ti film, as reflected by the XRD results (Fig. 2(a)) and HRTEM images (Fig. 4(b2)) above. This Mo-S-Ti-C composite film not only maintains its well crystallized MoS2 basal planes, but also exhibits the enhanced mechanical properties, a phenomenon never reported in all previous studies [22, 23, 25, 39]. Finally, it can be seen that, the Young’s modulus and elastic recovery ratio (delastic/dmax) of all composite films mentioned above, as reflected by the insets of Fig. 8(a), followed a trend similar to the hardness change. The tribological properties of pure MoS2, Mo-S-C, Mo-S-Ti and Mo-S-Ti-C composite films were tested by using a ball-on-disk tribometer. All tests were conducted under the same loading conditions and the friction curves acquired from vacuum and humid air are shown in Fig. 8(b) and (c). Particularly, all sputtered films wear lives were defined as the friction coefficient exceeded 0.5. As shown in Fig. 8(b), in vacuum, except for pure MoS2 films, MoS2 based composite films, including Mo-S-C, Mo-S-Ti and Mo-S-Ti-C films, exhibited a low friction coefficient (varing from 0.03 to 0.05) during the whole friction tests. Particularly, even this Mo-S-C composite film exhibits MoS2 phase dominant in (100) basal planes, as shown in Fig. 2, it also displays excellent lubricant properties in vacuum due to their structure



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transformation (from (100) to (002) texture) in sliding [9-11]. In comparison with the Mo-S-C composite film, the Mo-S-Ti and Mo-S-Ti-C composite films exhibited lower friction coefficient in vacuum, which can be attributed to the preferential orientation of MoS2 in (002) basal planes. In ambient air, it can be found that pure MoS2 film, Mo-S-C and Mo-S-Ti films all exhibit high friction coefficient and short wear live (less than 2 × 104). Surprisingly, this Mo-S-Ti-C composite film still remains a low friction coefficient ranging from 0.02 to 0.03 during the 105 revolutions in ambient air, which is far below than the friction coefficient (0.1～0.2) the previous MoS2-based composite films exhibited [22, 23, 24].

Fig.8 (a) The hardness and Young’s modulus of pure MoS2, Mo-S-C, Mo-S-Ti and Mo-S-Ti-C

composite

films,

with

the

insets

showing

the

corresponding

load-displacement curves of composite films. The friction coefficient curves of MoS2, Mo-S-C, Mo-S-Ti and Mo-S-Ti-C composite films in vacuum (b) and humid air (c). A Micro XAM 3D non-contact surface profilometer was used to calculate the wear rates of mentioned tribotests in vacuum and ambient air above, and the results are shown in Fig. 9. The acquired wear time of all tribo-tests was controlled as 20


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minutes. It can be seen that for nanocylindrical Mo-S-C and Mo-S-Ti films, they all exhibit low friction coefficient in vacuum, 0.05 for Mo-S-C film and 0.02 for Mo-S-Ti film, shown in Fig. 9(a). Nevertheless, the wear rates for them are up to 29.2×10-16 m3/N·m and 57.8×10-16 m3/N·m respectively, which means that the low friction coefficient is achieved with high sacrifice on the wear rates of film [8, 40]. Compared with the tribotests in vacuum, the wear tracks created in ambient air are much wider and the corresponding wear rates increased significantly, up to 33.7×10-16 m3/N·m and 95.4×10-16 m3/N·m respectively. Obviously different to the wear tracks and wear rates of Mo-S-C and Mo-S-Ti films, shown in Fig. 9(a1-b2), the wear tracks which formed on Mo-S-Ti-C films surface become much narrowed and the wear rates are obviously smaller than that of Mo-S-C and Mo-S-Ti films, and the corresponding wear rates are 2.86×10-16 m3/N·m in vacuum and 1.56×10-16 m3/N·m under ambient air atmosphere respectively, which means the best wear resistance of this Mo-S-Ti-C film exhibits among these films. Such an excellent lubricant properties of Mo-S-Ti-C composite films can be attributed to the improved hardness after multiple doping as well as the unbroken MoS2 crystallites. Benefiting from these highly ordered MoS2 crystallites in Mo-S-Ti-C composite films, a protective tribofilm can quickly form through self-organization under the induction of a rotating shear stress, and the instinct lubricant properties of MoS2 phase was fully achieved whether in vacuum or ambient air sliding condition, which results in the low friction coefficient and high antiwear properties.



Fig.9 3D non-contact surface mapping of the wear tracks of Mo-S-C, Mo-S-Ti and Mo-S-Ti-C composite films after sliding for 2×104 revolutions in vacuum (a1, b1, c1) and ambient air (a2, b2, c2) and wear rate of MoS2 based films (d). 4. Conclusions In this present work, a Mo-S-C-Ti composite film with low friction coefficient both in vacuum and ambient air was successfully synthesized in the condition of pure MoS2 films growing with edge-terminated surface. HRTEM images show that the edge-terminated MoS2 crystallites can efficiently provide active sites for carbon and titanium aggregating in film growth. This aggregation largely passivates the active sites and inhibit (100) textured MoS2 film growth, and similarly, it greatly reduces destruction the dopants do MoS2 crystallinity. Particularly, based on the analysis of HRTEM images, XRD spectrum and molecular dynamics simulation of MoS2-based films, the MoS2 structural transition (from (100) to (002) texture) in deposition was


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explained detailedly. Consequently, it’s reasonable to prepare MoS2-based films with excellent tribological properties in the condition of pure MoS2 growing with (100) basal planes, which may provide a novel and easier method for developing of MoS2-basied composite materials with well crystallized MoS2 (002) basal planes. Acknowledgements This research was supported by the National Natural Science Foundation of China (No. 51705503, 11875305) and by the National Magnetic Confinement Fusion Energy Research Project with project numbers of 2017YFE0302500 and 2015GB121004 for financial support. We thank for L. Qiao and Q. L. Chai for invaluable help on the film deposition and insights on the composite films evolution mechanism.

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Mo-S-Ti-C Nanocomposite Films for Solid-State Lubrication - ACS

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