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a-C:H nanocomposite coatings deposited by unbalanced magnetron sputtering
Surface & Coatings Technology 206 (2012) 3299–3308
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Microstructure and tribology of TiC(Ag)/a-C:H nanocomposite coatings deposited by unbalanced magnetron sputtering Yunfeng Wang a, Jun Wang a, Guangan Zhang b, Liping Wang b,⁎, Pengxun Yan a a b
School of Physics Science & Technology, Lanzhou University, Lanzhou 73000, PR China State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, PR China
a r t i c l e
i n f o
Article history: Received 11 July 2011 Accepted in revised form 18 January 2012 Available online 27 January 2012 Keywords: Nanocomposite TiC(Ag)/a-C:H High vacuum Superlow friction
a b s t r a c t TiC(Ag)/a-C:H nanocomposite coatings with various Ag concentrations were fabricated on Si p(100) substrates. The composition and structure of as-deposited nanocomposite coatings were systemically investigated, and the friction and wear behaviors were also evaluated under the ambient, high temperature and high vacuum, respectively. Results show that the TiC nanocrystallites were formed in the amorphous hydrogenated carbon matrix near the substrate. The co-dopant Ag possessed nanocrystalline structure in the asfabricated coatings whilst it formed Ag clusters (10–50 nm) on the surface. Furthermore, the introduction of Ag caused a significant reduction in the residual compressive stress without considerable decrease of the hardness and improved the adhesive strength of nanocomposite coatings. Tested as-deposited and after annealed at 500 °C coatings, the TiC(Ag)/a-C:H coatings showed a reduction of friction coefficients and wear rates with increment of Ag concentration. Under high vacuum condition, the TiC(Ag)/a-C:H coatings presented superlow friction behavior where the friction coefficient was reduced from 0.01 to 0.005 and lifetime increased from 0 to 1500 cycles. The significant improvement in tribological properties was mainly attributed to the low shear strength of Ag clusters on the surface as well as Ag diffusion to surface and wear track of coatings. The superior friction and wear behaviors of TiC(Ag)/a-C:H coatings make them good candidates as solid lubrication materials in space and aircraft applications. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Space and aircraft systems have many various moveable devices. An urgent challenge is to develop solid lubrication coatings with low friction coefficient and long lifetime in broad temperature range of operation, such as −100 to +100 °C for space-borne deices and −40 to +300 °C for aircraft components under high vacuum [1]. Diamond-like carbon (DLC) coatings as a kind of solid lubrication coatings with both low friction coefficient and high wear resistance, have been intensively studied in theoretical and application field [2]. Although the friction behavior of DLC films provides the opportunity for the practical application in space and aircraft systems, there are also some problems. DLC can maintain low friction in highvacuum environment if the H content is above 50%, however, DLC graphitization in friction contacts and the associated increase the friction coefficient in high-vacuum environment [3]. However, the superlow friction behavior of a-C:H provides the opportunity for the practical application in high vacuum condition [4,5]. At present, achieving long lifetime for superlow friction seems to be the most prominent issue. A possible method to improve the lifetime is
⁎ Corresponding author. Tel.: + 86 931 4968080; fax: + 86 931 4968163. E-mail address: [email protected] (L. Wang). 0257-8972/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2012.01.036
through compositional modification of a-C:H, such as preparation of a-C:H by incorporating certain elements into their amorphous microstructures [6–9]. However, this method also reduces the harness and wear resistance. Recently, significant progress in the design and production of nanocomposite coatings has resulted in advanced mechanical characteristic and low friction coefficient [10–13]. In this work, the selected Ti doping formed hard TiC phase can restore hardness of coatings [14,15]. The soft and ductile Ag is a potential lubricant material in vacuum and high temperature applications [16–18]. For Ag-incorporated DLC films, studies have been reported on the electrochemical, biological, and antibacterial activities and tribological properties [19,20], however tribological behavior under high vacuum and different conditions correlated with corresponding microstructures and mechanical properties have seldom been reported. Thus, the duplex doped TiC(Ag)/a-C:H nanocomposite coatings were deposited on Si p(100) substrates by magnetron sputtering Ti and Ag targets in argon (Ar) and methane (CH4) gas mixture atmosphere. Compared with simplex Ti-doped TiC/a-C:H coating. 2. Experimental TiC(Ag)/a-C:H nanocomposited coatings with thickness about 1.5 μm were deposited on Si p(100) substrates by a magnetron sputter deposition system. The Si p(100) substrates were first
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ultrasonically cleaned, then etched by Ar + plasma bombardment for 30 min at a pulsed bias voltage of −1000 V to remove residual impurities and native oxide on the substrates. Prior to coating deposition, a Ti interlayer of 150 nm thickness was deposited under a bias voltage of −500 V to enhance the adhesion between the final coating and substrate. Deposition was carried out in an argon (Ar) and methane (CH4) mixture, with background pressure pumped to lower than 1 × 10 − 3 Pa. Twinborn targets focusing the substrate were applied a mid-frequency AC power with the current of 2.0 A at a bias voltage of −1000 V with the duty factor of 20%. The deposition time was 2 h. The target is a rectangle shape and consists of a stack of Ag and Ti bars. The total number of the bars was kept at 30 and the ratio of the number of Ag bars to that of Ti ones, Ag/Ti, which was equivalent to the target area ratio, was set to 0/30, 3/27, 6/24, and 9/21. The temperature during deposition was in the range 180–200 °C monitored by a thermocouple inserted into the substrate holder. In our experiments, the substrate was not specially heated, however, it could be heated due to the magnetron discharge and ion bombardment. After substrate cleaning for about 30 min and the transition layer deposition for about 40 min, the substrate temperature had reached about 110 °C. When twinborn targets focusing the substrate were applied a mid-frequency AC power, the substrate temperature reached and maintained a final temperature about 180–200 °C. Typical deposition parameters are listed in Table 1. The chemical composition and bond state were characterized by a Perkin-Elmer PHI-5702 multi-functional photoelectron spectroscopy (XPS) with Al Kα (1476.6 eV) X-ray radiation to investigate the compositions and chemical states of carbon in the coatings. The XPS spectra were collected with a constant analyzer energy mode, at a chamber pressure of 10 − 8 Pa and pass energy of 29.4 eV, with 0.125 eV/step. An Ar + ion beam was used to sputter cleaning about 30 s to remove contaminants on the surface of the samples before obtaining the Ti2p, C1s, Ag3d and O1s spectra. High-resolution transmission electron microscopy (HRTEM) was performed using the JEM2010 TEM operated at 200 kV. The nanohardness was measured by a Nanotest600 nanoindenter apparatus (MicroMaterials Ltd.) using a Berkovich diamond tip and the maximum indentation depth being kept at 150 nm (about 10% of the total coating thickness) to minimize the substrate contribution. The internal stress was obtained from change in the radius of curvature of the Si wafer measured before and after the coating deposition by the observation of Newton's rings using an optical interferometer system, and then the internal stress was calculated by the Stoney's equation [21] σ¼
Es t2s 1 1 − 6ð1−vs Þ tc R2 R1
where Es/(1 − vs) is the substrate biaxial modulus; ts and tc are wafer and coating thickness; and R1 and R2 are the radius of curvature of Si wafer before and after deposition of coatings, respectively. The scratch test was used to evaluate the adhesion strength on UMT2MT test rig, using a scratch mode. Thus the scratch tests were carried out using a diamond indenter of 0.4 mm in radius by continuously increasing the normal load by 100 N/min. The load at which the friction force assumed a sharp increase was defined as the critical load (LC) of the coating and used as a quasi-quantitative criterion to evaluate the
adhesion strength of the coatings on the substrate. The wear surfaces were characterized by JSM-5600LV scanning electron microscope (SEM). The surface morphology and fractured cross sections of the coatings were examined using a JSM-6701F cold field Scanning electron microscope (FESEM). In order to investigate the thermal stability of as-deposited coatings, annealing experiments were carried out in a furnace in Ar atmosphere at pressure of 50 Pa. The coatings were annealed at fixed temperature 500 °C for 1 h, respectively, and cooled down to room temperature. After annealed, the tribology performances of coatings were characterized so as to evaluate evolution of the coatings caused by annealing. CCr15 steel balls were used for all tests (approximate 1.623 GPa Hertzian contact pressure). Besides, friction experiments in high vacuum were performed at room temperature using a ballon-flat tribometer in a vacuum chamber. Before the test, the chamber was evacuated with a base pressure of 3.0 × 10 − 4 Pa. The sample rotated at a constant speed of ~ 500 rpm against the CCr15 steel balls (3.175 mm in diameter) under the normal load of 1 N, which resulted in a Hertzian contact pressure of 950 MPa. 3. Results and discussion 3.1. Chemical composition and microstructure Table 2 lists the composition of as-deposited coatings which are arranged by Ag content determined by EDS. The range of Ag concentration of as-deposited coatings, 0 to 1.62 at.%, is correlated with the deposition parameters, while the Ti concentration in the coatings is not remarkably changed. The XPS technique is used to monitor the composition of the TiC(Ag)/a-C:H nanocomposite coatings due to its chemical sensitivity so as to explore the compounds formed in the coatings. The spectra are illustrated in Fig. 1. As shown in Fig. 1(a) and (c), the XPS spectra shown the inexistence of TiC and/or Ti at the surface of as-deposited coatings, probably due to the significant poisoning of Ti bars. The binding energy of C1s in Fig. 1(a) of all as-deposited coatings shifted to low binding energy indicating that more sp 2-C has formed with increasing Ag concentration in the coatings. The increasing of sp 2-C content was attributed to the Ag species had formed nanoclusters whilst these nanoclusters could absorb the compressive stress from the a-C:H matrix and reduced the carbon densification, which have been described by other researchers [22]. The binding energy of the Ag3d5/2 peak in Fig. 1(b) was centered at 368.4 eV for all asdeposited coatings, which suggested the existence of Ag phase in the coatings. The intensity of Ag peak increased as a function of Ag concentration indicating the increase of volumetric fraction of Ag phase. On the other hand, combining the binding energy value of the Ag3d5/2 peak, it further indicated that the Ag species were distributed in carbon network in the form of metallic phase. This provided evidence of the existence of Ag nanoclusters in the a-C:H matrix. The binding energy of the O1s peak in Fig. 1(d) was centered at 531.6 eV for all as-deposited coatings. The O1s peak of TiO2 was reported at 529.7 eV [23] binding energy, while the carbonyl and C\O were assigned to 531.5 eV and 533.2 eV [24]. In accordance to the observations of all these authors the observed peak components should not include TiO2. The effect of addition of Ag to TiC/a-C:H on the microstructure of as-deposited coatings can be shown by the X-ray diffraction patterns
Table 1 Summary of the coatings deposition conditions. Item
Parameter
Ar gas flow rate CH4 gas flow rate Deposition pressure Mid-frequency AC power Applied bias
120 sccm 40 sccm 1 Pa 2.0 A − 1000 V
Table 2 Composition of as-deposited coatings. Sample
0/30
3/27
6/24
9/21
Ag/Ti [Ti] (at.%) [Ag] (at.%)
0/30 2.81 0
3/27 2.68 0.56
6/24 2.69 1.13
9/21 2.74 1.62
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Fig. 1. XPS spectra of as-deposited coatings with different Ag concentrations: (a) C1s, (b) Ag3d, (c) Ti2p and (d) O1s.
in Fig. 2. The coating with 0 at.% Ag content exhibits diffraction peaks at (111) TiC and (220) TiC [25,26]. The coating with Ag doping exhibits two crystalline phases in the XRD patterns, one is TiC with a predominant (111) orientation and the other is Ag with a predominant (111) orientation. It is suggested that the nucleation of Ag with (111) orientation disrupts TiC growth along the (220) axis. In order to prove the existence of Ag nanoparticles in the asfabricated coatings, the HRTEM analysis was performed. HRTEM images of TiC(Ag)/a-C:H coatings with different Ag concentration are presented in Fig. 3. As shown in Fig. 3(a), HRTEM image of coating without Ag incorporation presented a fine nanocrystalline structure. Both XRD and HRTEM analyses point to the fact that the nanocrystalline grains are the TiC phase. Fig. 3(b) shows that TiC(Ag)/a-C:H coating with 0.56 at.% Ag formed Ag amorphous phase (dark spots) with a diameter of approximately 2–5 nm and dispersed uniformly within the a-C:H matrix (gray areas). The Ag amorphous phases were apparent as dark spots due to their higher atomic mass and density. As shown in Fig. 3(c), the microstructure of the coating changed
Fig. 2. XRD pattern of TiC(Ag)/a-C:H coatings with different Ag concentration.
completely with 1.13 at.% Ag incorporation compared to that of the coating with 0.56 at.% Ag incorporation. It clearly exhibited the presence of the coarsening of remaining Ag in the a-C:H matrix, and the Ag nanocrystalline particles of 5 nm diameter were evenly distributed in the a-C:H matrix. For the coating with 1.62 at.% Ag incorporation (as shown in Fig. 3(d)), the microstructure changed slightly compared to that of the coating with 1.13 at.% Ag incorporation. The only change is the size of Ag nanocrystalline particles, from about 5 nm to 10 nm. Furthermore, the crystallographic planes with interplanar distances of approximately 0.235 nm are observed in these regions, which corresponds to the (111) plane of face-centered Ag. Other planes in the TiC and Ag crystal systems also contributed to the SAED pattern of coating, indexing in Fig. 4. The strongest diffraction rings from the TiC phase were (111) and (220) planes, while the dominant diffraction ring from the Ag phase was (111) plane orientations. As shown by above results, the HRTEM information agreed nearly with those from XRD. The increasing grain size with Ag content may be explained as the dispersion of the formed Ag nanocrystallites in the amorphous carbon matrix (DLC) which was favorable to the constant renucleation of Ag nanoparticles [27]. Besides, such microstructure with large fraction of grain boundary could provide ductility by activating grain boundary slip and crack termination by nanocrack splitting, which would provide a unique combination of high hardness and toughness for as-deposited coatings [28]. Typical surface morphology images of as-deposited TiC(Ag)/a-C:H coatings are shown in Fig. 5. The difference in brightness in the FESEM image allows us to distinguish between the substances of the nucleated precipitates (Ag) and the a-C:H matrix. The FESEM images also show that Ag nanoparticles had a tendency to form clusters on or near the surface during the deposition process. Fig. 5(a) and (b) shows the morphology of the coatings containing 0 at.% Ag and 0.56 at.% Ag, respectively. No Ag clusters were found on the surface. As the Ag content in these coatings was increased, the increment in
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Fig. 3. HRTEM micrographs of as-deposited coatings with different concentration: (a) 0 at.% Ag, (b) 0.56 at.% Ag, (c) 1.13 at.% Ag and (d) 1.62 at.% Ag.
the number and the size of the lighter colored clusters (Ag clusters) on the surface was more notable, and that these clusters were still submicron in size ranging from 10 nm to 50 nm, as shown in Fig. 5(c) and (d). This process more likely occurred in the higher Ag concentration coatings, due to the Ag diffused to the surface and nucleated a particle at the surface. For the most Ag content (1.62 at.%), the Ag clusters (the white points) with the sizes ranging from 10 nm to 50 nm were evenly distributed through the entire coating surface, as presented in Fig. 5(d). FESEM micrographs presented in Fig. 6 show the fractured crosssections of the nanocomposite coatings. Both the interlayer and the transition layer were characterized by slight columnar and granular structures, which could contribute to the large numbers of
nanocrystalline TiC and Ag particles embedded in the layer. During the deposition process, the flow rate of Ar gas kept at 120 sccm, and the transition layer deposition was carried out by transforming the methane flow rate from 0 to 40 sccm at a 2-min interval increments. Changes in the flow rate of methane gas lead to a variation of the carbon content in the nanocomposite coatings. Both the interlayer and the transition layer were characterized by slight columnar and granular structures, which could contribute to the large numbers of nanocrystalline TiC and Ag particles embedded in the layer. Moreover, a transition from slight columnar and granular to glassy microstructure was also observed to respond to an increase in carbon content and the Ti concentration continuously decreased. Carbonaceous compound species from the plasma condensed on the surface of Ti bars. With
Fig. 4. SAED pattern indexing for TiC(Ag)/a-C:H nanocomposite coating with 1.13 at.% Ag, (b) 1.62 at.% Ag and (c) 0 at.% Ag.
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Fig. 5. SEM surface images of as-deposited coatings with different concentration: (a) 0 at.% Ag, (b) 0.56 at.% Ag, (c) 1.13 at.% Ag and (d) 1.62 at.% Ag.
relatively high CH4 flow rate (high CH4 concentrations in the atmosphere), carbonaceous compound species could be sputtered from the surface of Ti bars and the amount of Ti sputtered from the bars continuously decreased. In the case of low C amount, there was an ensemble of island-shape structure of Ag-rich layer which caused by Ag nanoparticles self-assembled into arrays attributed to the high surface energy of Ag relative to C observed in [29] during Ag deposition.
Subsequent carbon layer deposited hinders Ag atom diffusion and fixed the Ag-rich layer appearance. To summarize the results of the structural investigations, the following microstructure evolution in TiC(Ag)/a-C:H nanocomposite coatings with increasing Ag content was proposed: when the Ag content below 1 at.%, the coatings consisted of mostly amorphous structure; as Ag content increased, the coatings consisted of nanocrystalline Ag particles in
Fig. 6. FESEM micrographs of the fractured cross-sections of the nanocomposite coatings with different Ag concentration (a) 0 at.% Ag, (b) 0.56 at.% Ag, (c) 1.13 at.% Ag and (d) 1.62 at.%.
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Fig. 7. (a) Typical curve of nano-indentation load vs displacement into the surface of coating with 1.13 at.% Ag; (b) the mechanical properties of TiC(Ag)/a-C:H with different Ag concentration.
an amorphous a-C:H matrix and Ag clusters formed on the surface of the coatings. 3.2. Mechanical properties Coatings with different microstructures are expected to have different mechanical properties. The hardness was obtained by nanoindentation. Typical curve of 1.13 at.% Ag incorporation is shown in Fig. 7(a). Fig. 7(b) exhibits the effects of the Ag concentration on the hardness and the residual compressive stress of as-deposited coatings. As seen in Fig. 7(b), Ag doping leads to a monotonic decrease in hardness (H) with increment of Ag content. Besides, the incorporation of Ag also resulted in the reduction of the residual compressive stress. Results show that the incorporation of 0.56 at.% Ag caused a slight reduction of compressive stress. Increasing the Ag concentration further from 0.56 at.% to 1.13 at.% caused significant decrease of compressive stress from 2.7 to 1.5 GPa. Further incorporation of Ag up to 1.62 at.% results in a gradual decrease of the residual compress stress down to 1.4 GPa. Incorporation of Ag will induce a gradual reduction of the residual compressive stress concurrent with a similar reduction of hardness. This drastic decrease of the magnitude of residual compressive stress, however, was mainly due to formation of a metallic Ag phase (as shown in Fig. 3(c)) acting as efficient buffer sites to absorb the stress in the a-C:H matrix. The process of doping Ag nanoparticles into the amorphous coatings likely lead to two competing and contradicting effects on the hardness: (1) the addition of Ag nanoparticles deteriorated the mechanical properties which contributed to softening; and (2) the addition of Ag nanoparticles played a role as strengthening dopants to increase the coating hardness. This behavior has also been reported in the past in other systems including TiC/Ag, TiN/Ag and Mo2N/Ag [30–33]. According to the literature, the embedded nanoparticles of
silver in an amorphous matrix could form a heterogeneous structure and decrease the plastic flow. The final condition of the coating hardness depended on the compromise of these two effects. Suitable amount of Ag nanoparticles doped in the amorphous carbon matrix can promote the formation of Ag nanoparticles of favorable sizes, which maintained the hardness and reduced the intrinsic stress of a coating. The adhesive behaviors of the as-deposited TiC(Ag)/a-C:H coatings with 0 and 1.13 at.% Ag incorporation were evaluated by the scratch tests. It could be found that the 1.13 at.% Ag incorporation increased the critical load to 27 ± 1 N compared to the critical load of 18 ± 2 N of 0 at.% Ag incorporation. It is notable that the TiC(Ag)/a-C:H coatings with 1.13 at.% Ag incorporation showed a significant reduction in the residual compressive stress without a considerable decrease of the hardness, as well as improvement in adhesive strength, indicating that TiC(Ag)/a-C:H coatings with 1.13 at.% Ag incorporation is promising as solid lubrication coating from the perspective of tribological properties.
3.3. Tribological performances Fig. 8 shows the average friction coefficients of TiC(Ag)/a-C:H nanocomposite coatings were sensitive to Ag content at asdeposited and after annealed at 500 °C conditions, where increased Ag concentration generally resulted in a reduction of the friction coefficient. The lowest friction coefficient was obtained for the coating with 1.13 at.% Ag which showed an average value of 0.07 at ambient condition and 0.04 after annealed at 500 °C condition. While the higher friction coefficient obtained for the coating with 1.62 at.% Ag may be due to the higher Ag content on the surface of the coating, which can be seen in Fig. 5(d). Therefore, a proper Ag concentration in the coatings is required to facilitate lubrication.
Fig. 8. Friction coefficients of TiC(Ag)/a-C:H nanocomposite coatings with different Ag concentration: (a) as-deposited and (b) after annealed at 500 °C condition.
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The wear rates were calculated after scanning the wear track with various Ag concentrations, as shown in Fig. 9. As a whole, the wear rates of TiC(Ag)/a-C:H nanocomposite coatings both as-deposited and after annealed at 500 °C condition decreased with the increment of Ag concentration. The wear rates remarkably decreased for the coatings contained Ag from 0 at.% to 1.13 at.%, and then slightly fluctuated with the further increase of the Ag concentration to 1.62 at.%. Generally, the wear rates of carbon-based coatings after annealed at 500 °C condition should always be higher than that tested at ambient condition. However, the wear rates of as-deposited nanocomposite coatings after annealed at high temperature (500 °C) were always lower than those at ambient condition with Ag incorporation. Therefore, it can be concluded that a fraction of Ag in the nanocomposite coatings diffused to the surface of the coatings as an effective solid lubricant. The results show that the wear rate of TiC(Ag)/a-C:H nanocomposite coating with 1.13 at.% Ag incorporation (1.18 × 10− 16 m 3/N·m) was less than 15% of coating without Ag incorporation at high temperature condition. Fig. 10 shows the FESEM surface images of coating 1.13 at.% Ag concentration (a) as-deposited coating and (b) coating after annealed at 500 °C. The Ag clusters on the surface are significantly larger and denser than the initial Ag clusters on the surface of asdeposited coating, which indicates that Ag in the nanocomposite coatings diffused to the surface so that it facilitates to grow and coalesce of Ag particles. The larger and denser Ag clusters on the surface of the coating after heating could lubricate the coating surface to reduce friction coefficients and wear rates. The variation of friction behaviors tested at high vacuum (HV) condition of the nanocomposite coatings with different Ag content are shown in Fig. 11. As reported [34], the friction results under HV condition exhibit two different behaviors: after a running-in period, the friction coefficients stabilizes at very low values less than 0.01 (superlow friction coefficients) for the coatings with only Ag clusters on the surface of the coatings, while below this threshold the friction coefficient is in the milli-range during the first several hundred cycles and then increases rapidly up to 0.3–0.4. The experiments were then rapidly stopped. For each deposition process, the two different behaviors can be discriminated by the different Ag concentration of asdeposited coatings. Fig. 11 shows the evolution of the average friction coefficient in HV condition versus the number of sliding cycles for all four coatings. All the coatings had an initial running-in period, where the friction coefficient decreased from about 0.20 to b0.01 at steady state. The friction curve of 1.13 at.% Ag incorporation was characterized by a running-in period of around 700 cycles where the friction coefficient decreased slowly to 0.01 followed by a typical value between 0.002 and 0.01, the ultralow friction regime persists for about 1500 cycles. In the case of coating with 1.62 at.% Ag incorporation, the friction coefficient starts at 0.2 and diminishes rapidly, reaching the milli-range after only 200 cycles, nevertheless, the ultralow
Fig. 9. Specific wear rates of TiC(Ag)/a-C:H nanocomposite coatings with different Ag concentration: as-deposited and after annealed at 500 °C condition.
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Fig. 10. FESEM surface images of coating 1.13 at.% Ag concentration (a) as-deposited coating and (b) coating after annealed at 500 °C.
friction regime persists for about 500 cycles. The friction curve was disappearance between 300 and 700 cycles. The disappearance due to the machine was capable of reliable sensing loads that are 1 × 10 − 3 N. We used 1 N force in a sliding experiment, then being able to sense 0.001 N lateral force, which can translate into friction coefficients as low as 0.001. If the friction coefficient values were lower than 0.001 to be measured by the friction force sensor, it looks like the measurement ‘disappears’ between 300 and 700 cycles. For the coatings with the Ag concentration of 0 at.% and 0.56 at.%, the running-in time is only about 40 cycles, while the ultralow friction regime persists for just about 25 cycles. The sudden increase in friction coefficients above 0.1 for all as-deposited coatings indicates the coatings failure. Although the Ag concentration was different for each sample, the friction coefficient in the initial stage was at around 0.20 which should be due to the hydrogen atoms in a-C:H matrix. However, superlow friction for without Ag-doped coatings could not be kept for long periods of time at HV condition. The hydrogen atoms at the surface possibly disappeared during the friction tested. However, for Ag-doped coatings, the Ag clusters on the surface of coatings could be beneficial for tribological applications because the Ag clusters exhibit a low shear strength which could reduce the friction force on hard surface and improve the lifetime during sliding wear at HV condition. It was also anticipated that the Ag in the coatings should be drawn to diffuse both vertically and horizontally through the a-C:H matrix to the worn area so as to improve anti-wear properties. SEM images of wear tracks and scars tested in HV condition for the coatings with 0 and 1.13 at.% Ag content are estimated, as shown in Fig. 12. Many folds could be observed on the worn surface of coating with 1.13 at.% Ag content, as shown in Fig. 12(c). The existence of folds suggested that the worn surface of the coating was easily plastically deformed. However, no such folds existed on the worn surface of the coating without Ag incorporation, but the worn surface was characteristic with some small pits, as shown in Fig. 12(a). The occurrence of pits suggested that the worn surface of the coating without Ag incorporation was more brittle. The wear fragments for both coatings were also different. Many granular wear fragments were scattered on the worn surface of the coating without Ag incorporation, while many flattened wear fragments were observed on the worn surface of the coating with 1.13 at.% Ag content. Fig. 12(b) and (d) presents the SEM images of worn surface on mating balls, respectively. The transferred materials were negligible on the contact surface, while the wear fragments were clearly seen in the direction of sliding. A dark gray to black material accumulated mainly outside the initial contact area for the worn surface of the ball
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Fig. 11. Friction coefficients of TiC(Ag)/a-C:H nanocomposite coatings with different Ag concentration (a) 0 at.% Ag, (b) 0.56 at.% Ag, (c) 1.13 at.% Ag and (d) 1.62 at.% Ag tested at high vacuum.
mating for the sample without Ag incorporation, while this was not observed from the worn surface of ball mating for the coating with 1.13 at.% Ag content. No visible transfer film could be found on the worn surface for both coating mating balls, which agreed with the report by J. Fontaine et al. [35]: after the drastic friction increase, no transfer film can be found on the ball.
Fig. 13 shows the back scattered electron (BSE) images of wear scar on the surface of the coating with 1.13 at.% Ag content. The elemental distribution analysis of wear scar on the surface included three kinds of elements which are the Ag, C, Ti and O, respectively, and the relative content of each element could be speculated by the difference of color. As shown in Fig. 13, it could be proved directly
Fig. 12. SEM images of wear tracks and scars with different Ag concentration (a), (b) 0 at.% Ag and (c), (d) 1.13 at.% Ag tested at high vacuum.
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Fig. 13. Elemental distribution maps of wear scar of the coating with 1.13 at.% Ag tested at high vacuum : (a) Ag-La, (b) C-Ka, (C) Ti-Ka, (d) O-Ka and (e) back scattered electron (BSE) images of wear scar.
that the Ag and C contents was increased while Ti content was decreased in wear scar. We could say that these tribological performances should be attributed to the lubrication of the a-C:H coupled with effect of Ag clusters on the surface as well as the diffusion of the Ag in the coatings to the worn area so as to improve the antiwear properties. The friction evolution of TiC(Ag)/a-C:H coatings has demonstrated that a friction coefficient of 0.005 and a lifetime of 1500 cycles can be achieved in HV condition. These tribological performances should be attributed to the lubrication of the a-C:H coupled with effect of Ag clusters on the surface as well as the diffusion of the Ag in the coatings to the worn area so as to improve the anti-wear properties. The process is illustrated in the schematic picture shown in Fig. 14, and it indicates that the TiC(Ag)/a-C:H coating can achieve an excellent self-lubricating behavior during the sliding friction. The lifetime of the coatings with 0 at.% and 0.56 at.% Ag contents was short compared to the coatings with higher Ag concentration (the coatings with 1.13 at.% and 1.62 at.% Ag contents), which is mainly attributed that there was no Ag clusters on the surface of the coatings with 1.13 at.% and 1.62 at.% Ag contents.
4. Conclusions Nanocomposite TiC(Ag)/a-C:H coatings were successfully fabricated by mid-frequency dual-magnetron sputtering on silicon wafer substrates. Ag nanocrystallites were observed in the coatings with the Ag concentration of 1.13 and 1.62 at.%, and Ag clusters (10–50 nm) also formed on the surface. The friction coefficient and wear rates were sensitive to Ag content. In atmosphere, the friction coefficient and the wear rates was reduced as increase of Ag concentration. A properly Ag concentration in the coatings was required to facilitate lubrication. In high vacuum condition, at least 1.13 at.% Ag was needed to form a Ag lubrication mechanism, which reduced the friction coefficient from the 0.01 to 0.005, and increased the lifetime from 0 to 1500 cycles. The low friction coefficient and long lifetime of the TiC(Ag)/a-C:H coatings were attributed to the presence of low shear Ag clusters on the surface and the diffusion of the Ag in the coatings to the worn area so as to improve the anti-wear properties. Acknowledgments This work was carried out with financial support from the National Natural Science Foundation of China (No.11172300 and No. 11004203). References [1] [2] [3] [4] [5] [6] [7] [8]
Fig. 14. Schematic diagram of tribofilm induced by sliding friction for TiC(Ag)/a-C:H nanocomposite coatings tested at high vacuum.
[9] [10] [11]
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Surface & Coatings Technology journal homepage: www.elsevier.com/locate/surfcoat
Microstructure and tribology of TiC(Ag)/a-C:H nanocomposite coatings deposited by unbalanced magnetron sputtering Yunfeng Wang a, Jun Wang a, Guangan Zhang b, Liping Wang b,⁎, Pengxun Yan a a b
School of Physics Science & Technology, Lanzhou University, Lanzhou 73000, PR China State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, PR China
a r t i c l e
i n f o
Article history: Received 11 July 2011 Accepted in revised form 18 January 2012 Available online 27 January 2012 Keywords: Nanocomposite TiC(Ag)/a-C:H High vacuum Superlow friction
a b s t r a c t TiC(Ag)/a-C:H nanocomposite coatings with various Ag concentrations were fabricated on Si p(100) substrates. The composition and structure of as-deposited nanocomposite coatings were systemically investigated, and the friction and wear behaviors were also evaluated under the ambient, high temperature and high vacuum, respectively. Results show that the TiC nanocrystallites were formed in the amorphous hydrogenated carbon matrix near the substrate. The co-dopant Ag possessed nanocrystalline structure in the asfabricated coatings whilst it formed Ag clusters (10–50 nm) on the surface. Furthermore, the introduction of Ag caused a significant reduction in the residual compressive stress without considerable decrease of the hardness and improved the adhesive strength of nanocomposite coatings. Tested as-deposited and after annealed at 500 °C coatings, the TiC(Ag)/a-C:H coatings showed a reduction of friction coefficients and wear rates with increment of Ag concentration. Under high vacuum condition, the TiC(Ag)/a-C:H coatings presented superlow friction behavior where the friction coefficient was reduced from 0.01 to 0.005 and lifetime increased from 0 to 1500 cycles. The significant improvement in tribological properties was mainly attributed to the low shear strength of Ag clusters on the surface as well as Ag diffusion to surface and wear track of coatings. The superior friction and wear behaviors of TiC(Ag)/a-C:H coatings make them good candidates as solid lubrication materials in space and aircraft applications. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Space and aircraft systems have many various moveable devices. An urgent challenge is to develop solid lubrication coatings with low friction coefficient and long lifetime in broad temperature range of operation, such as −100 to +100 °C for space-borne deices and −40 to +300 °C for aircraft components under high vacuum [1]. Diamond-like carbon (DLC) coatings as a kind of solid lubrication coatings with both low friction coefficient and high wear resistance, have been intensively studied in theoretical and application field [2]. Although the friction behavior of DLC films provides the opportunity for the practical application in space and aircraft systems, there are also some problems. DLC can maintain low friction in highvacuum environment if the H content is above 50%, however, DLC graphitization in friction contacts and the associated increase the friction coefficient in high-vacuum environment [3]. However, the superlow friction behavior of a-C:H provides the opportunity for the practical application in high vacuum condition [4,5]. At present, achieving long lifetime for superlow friction seems to be the most prominent issue. A possible method to improve the lifetime is
⁎ Corresponding author. Tel.: + 86 931 4968080; fax: + 86 931 4968163. E-mail address: [email protected] (L. Wang). 0257-8972/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2012.01.036
through compositional modification of a-C:H, such as preparation of a-C:H by incorporating certain elements into their amorphous microstructures [6–9]. However, this method also reduces the harness and wear resistance. Recently, significant progress in the design and production of nanocomposite coatings has resulted in advanced mechanical characteristic and low friction coefficient [10–13]. In this work, the selected Ti doping formed hard TiC phase can restore hardness of coatings [14,15]. The soft and ductile Ag is a potential lubricant material in vacuum and high temperature applications [16–18]. For Ag-incorporated DLC films, studies have been reported on the electrochemical, biological, and antibacterial activities and tribological properties [19,20], however tribological behavior under high vacuum and different conditions correlated with corresponding microstructures and mechanical properties have seldom been reported. Thus, the duplex doped TiC(Ag)/a-C:H nanocomposite coatings were deposited on Si p(100) substrates by magnetron sputtering Ti and Ag targets in argon (Ar) and methane (CH4) gas mixture atmosphere. Compared with simplex Ti-doped TiC/a-C:H coating. 2. Experimental TiC(Ag)/a-C:H nanocomposited coatings with thickness about 1.5 μm were deposited on Si p(100) substrates by a magnetron sputter deposition system. The Si p(100) substrates were first
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ultrasonically cleaned, then etched by Ar + plasma bombardment for 30 min at a pulsed bias voltage of −1000 V to remove residual impurities and native oxide on the substrates. Prior to coating deposition, a Ti interlayer of 150 nm thickness was deposited under a bias voltage of −500 V to enhance the adhesion between the final coating and substrate. Deposition was carried out in an argon (Ar) and methane (CH4) mixture, with background pressure pumped to lower than 1 × 10 − 3 Pa. Twinborn targets focusing the substrate were applied a mid-frequency AC power with the current of 2.0 A at a bias voltage of −1000 V with the duty factor of 20%. The deposition time was 2 h. The target is a rectangle shape and consists of a stack of Ag and Ti bars. The total number of the bars was kept at 30 and the ratio of the number of Ag bars to that of Ti ones, Ag/Ti, which was equivalent to the target area ratio, was set to 0/30, 3/27, 6/24, and 9/21. The temperature during deposition was in the range 180–200 °C monitored by a thermocouple inserted into the substrate holder. In our experiments, the substrate was not specially heated, however, it could be heated due to the magnetron discharge and ion bombardment. After substrate cleaning for about 30 min and the transition layer deposition for about 40 min, the substrate temperature had reached about 110 °C. When twinborn targets focusing the substrate were applied a mid-frequency AC power, the substrate temperature reached and maintained a final temperature about 180–200 °C. Typical deposition parameters are listed in Table 1. The chemical composition and bond state were characterized by a Perkin-Elmer PHI-5702 multi-functional photoelectron spectroscopy (XPS) with Al Kα (1476.6 eV) X-ray radiation to investigate the compositions and chemical states of carbon in the coatings. The XPS spectra were collected with a constant analyzer energy mode, at a chamber pressure of 10 − 8 Pa and pass energy of 29.4 eV, with 0.125 eV/step. An Ar + ion beam was used to sputter cleaning about 30 s to remove contaminants on the surface of the samples before obtaining the Ti2p, C1s, Ag3d and O1s spectra. High-resolution transmission electron microscopy (HRTEM) was performed using the JEM2010 TEM operated at 200 kV. The nanohardness was measured by a Nanotest600 nanoindenter apparatus (MicroMaterials Ltd.) using a Berkovich diamond tip and the maximum indentation depth being kept at 150 nm (about 10% of the total coating thickness) to minimize the substrate contribution. The internal stress was obtained from change in the radius of curvature of the Si wafer measured before and after the coating deposition by the observation of Newton's rings using an optical interferometer system, and then the internal stress was calculated by the Stoney's equation [21] σ¼
Es t2s 1 1 − 6ð1−vs Þ tc R2 R1
where Es/(1 − vs) is the substrate biaxial modulus; ts and tc are wafer and coating thickness; and R1 and R2 are the radius of curvature of Si wafer before and after deposition of coatings, respectively. The scratch test was used to evaluate the adhesion strength on UMT2MT test rig, using a scratch mode. Thus the scratch tests were carried out using a diamond indenter of 0.4 mm in radius by continuously increasing the normal load by 100 N/min. The load at which the friction force assumed a sharp increase was defined as the critical load (LC) of the coating and used as a quasi-quantitative criterion to evaluate the
adhesion strength of the coatings on the substrate. The wear surfaces were characterized by JSM-5600LV scanning electron microscope (SEM). The surface morphology and fractured cross sections of the coatings were examined using a JSM-6701F cold field Scanning electron microscope (FESEM). In order to investigate the thermal stability of as-deposited coatings, annealing experiments were carried out in a furnace in Ar atmosphere at pressure of 50 Pa. The coatings were annealed at fixed temperature 500 °C for 1 h, respectively, and cooled down to room temperature. After annealed, the tribology performances of coatings were characterized so as to evaluate evolution of the coatings caused by annealing. CCr15 steel balls were used for all tests (approximate 1.623 GPa Hertzian contact pressure). Besides, friction experiments in high vacuum were performed at room temperature using a ballon-flat tribometer in a vacuum chamber. Before the test, the chamber was evacuated with a base pressure of 3.0 × 10 − 4 Pa. The sample rotated at a constant speed of ~ 500 rpm against the CCr15 steel balls (3.175 mm in diameter) under the normal load of 1 N, which resulted in a Hertzian contact pressure of 950 MPa. 3. Results and discussion 3.1. Chemical composition and microstructure Table 2 lists the composition of as-deposited coatings which are arranged by Ag content determined by EDS. The range of Ag concentration of as-deposited coatings, 0 to 1.62 at.%, is correlated with the deposition parameters, while the Ti concentration in the coatings is not remarkably changed. The XPS technique is used to monitor the composition of the TiC(Ag)/a-C:H nanocomposite coatings due to its chemical sensitivity so as to explore the compounds formed in the coatings. The spectra are illustrated in Fig. 1. As shown in Fig. 1(a) and (c), the XPS spectra shown the inexistence of TiC and/or Ti at the surface of as-deposited coatings, probably due to the significant poisoning of Ti bars. The binding energy of C1s in Fig. 1(a) of all as-deposited coatings shifted to low binding energy indicating that more sp 2-C has formed with increasing Ag concentration in the coatings. The increasing of sp 2-C content was attributed to the Ag species had formed nanoclusters whilst these nanoclusters could absorb the compressive stress from the a-C:H matrix and reduced the carbon densification, which have been described by other researchers [22]. The binding energy of the Ag3d5/2 peak in Fig. 1(b) was centered at 368.4 eV for all asdeposited coatings, which suggested the existence of Ag phase in the coatings. The intensity of Ag peak increased as a function of Ag concentration indicating the increase of volumetric fraction of Ag phase. On the other hand, combining the binding energy value of the Ag3d5/2 peak, it further indicated that the Ag species were distributed in carbon network in the form of metallic phase. This provided evidence of the existence of Ag nanoclusters in the a-C:H matrix. The binding energy of the O1s peak in Fig. 1(d) was centered at 531.6 eV for all as-deposited coatings. The O1s peak of TiO2 was reported at 529.7 eV [23] binding energy, while the carbonyl and C\O were assigned to 531.5 eV and 533.2 eV [24]. In accordance to the observations of all these authors the observed peak components should not include TiO2. The effect of addition of Ag to TiC/a-C:H on the microstructure of as-deposited coatings can be shown by the X-ray diffraction patterns
Table 1 Summary of the coatings deposition conditions. Item
Parameter
Ar gas flow rate CH4 gas flow rate Deposition pressure Mid-frequency AC power Applied bias
120 sccm 40 sccm 1 Pa 2.0 A − 1000 V
Table 2 Composition of as-deposited coatings. Sample
0/30
3/27
6/24
9/21
Ag/Ti [Ti] (at.%) [Ag] (at.%)
0/30 2.81 0
3/27 2.68 0.56
6/24 2.69 1.13
9/21 2.74 1.62
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Fig. 1. XPS spectra of as-deposited coatings with different Ag concentrations: (a) C1s, (b) Ag3d, (c) Ti2p and (d) O1s.
in Fig. 2. The coating with 0 at.% Ag content exhibits diffraction peaks at (111) TiC and (220) TiC [25,26]. The coating with Ag doping exhibits two crystalline phases in the XRD patterns, one is TiC with a predominant (111) orientation and the other is Ag with a predominant (111) orientation. It is suggested that the nucleation of Ag with (111) orientation disrupts TiC growth along the (220) axis. In order to prove the existence of Ag nanoparticles in the asfabricated coatings, the HRTEM analysis was performed. HRTEM images of TiC(Ag)/a-C:H coatings with different Ag concentration are presented in Fig. 3. As shown in Fig. 3(a), HRTEM image of coating without Ag incorporation presented a fine nanocrystalline structure. Both XRD and HRTEM analyses point to the fact that the nanocrystalline grains are the TiC phase. Fig. 3(b) shows that TiC(Ag)/a-C:H coating with 0.56 at.% Ag formed Ag amorphous phase (dark spots) with a diameter of approximately 2–5 nm and dispersed uniformly within the a-C:H matrix (gray areas). The Ag amorphous phases were apparent as dark spots due to their higher atomic mass and density. As shown in Fig. 3(c), the microstructure of the coating changed
Fig. 2. XRD pattern of TiC(Ag)/a-C:H coatings with different Ag concentration.
completely with 1.13 at.% Ag incorporation compared to that of the coating with 0.56 at.% Ag incorporation. It clearly exhibited the presence of the coarsening of remaining Ag in the a-C:H matrix, and the Ag nanocrystalline particles of 5 nm diameter were evenly distributed in the a-C:H matrix. For the coating with 1.62 at.% Ag incorporation (as shown in Fig. 3(d)), the microstructure changed slightly compared to that of the coating with 1.13 at.% Ag incorporation. The only change is the size of Ag nanocrystalline particles, from about 5 nm to 10 nm. Furthermore, the crystallographic planes with interplanar distances of approximately 0.235 nm are observed in these regions, which corresponds to the (111) plane of face-centered Ag. Other planes in the TiC and Ag crystal systems also contributed to the SAED pattern of coating, indexing in Fig. 4. The strongest diffraction rings from the TiC phase were (111) and (220) planes, while the dominant diffraction ring from the Ag phase was (111) plane orientations. As shown by above results, the HRTEM information agreed nearly with those from XRD. The increasing grain size with Ag content may be explained as the dispersion of the formed Ag nanocrystallites in the amorphous carbon matrix (DLC) which was favorable to the constant renucleation of Ag nanoparticles [27]. Besides, such microstructure with large fraction of grain boundary could provide ductility by activating grain boundary slip and crack termination by nanocrack splitting, which would provide a unique combination of high hardness and toughness for as-deposited coatings [28]. Typical surface morphology images of as-deposited TiC(Ag)/a-C:H coatings are shown in Fig. 5. The difference in brightness in the FESEM image allows us to distinguish between the substances of the nucleated precipitates (Ag) and the a-C:H matrix. The FESEM images also show that Ag nanoparticles had a tendency to form clusters on or near the surface during the deposition process. Fig. 5(a) and (b) shows the morphology of the coatings containing 0 at.% Ag and 0.56 at.% Ag, respectively. No Ag clusters were found on the surface. As the Ag content in these coatings was increased, the increment in
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Fig. 3. HRTEM micrographs of as-deposited coatings with different concentration: (a) 0 at.% Ag, (b) 0.56 at.% Ag, (c) 1.13 at.% Ag and (d) 1.62 at.% Ag.
the number and the size of the lighter colored clusters (Ag clusters) on the surface was more notable, and that these clusters were still submicron in size ranging from 10 nm to 50 nm, as shown in Fig. 5(c) and (d). This process more likely occurred in the higher Ag concentration coatings, due to the Ag diffused to the surface and nucleated a particle at the surface. For the most Ag content (1.62 at.%), the Ag clusters (the white points) with the sizes ranging from 10 nm to 50 nm were evenly distributed through the entire coating surface, as presented in Fig. 5(d). FESEM micrographs presented in Fig. 6 show the fractured crosssections of the nanocomposite coatings. Both the interlayer and the transition layer were characterized by slight columnar and granular structures, which could contribute to the large numbers of
nanocrystalline TiC and Ag particles embedded in the layer. During the deposition process, the flow rate of Ar gas kept at 120 sccm, and the transition layer deposition was carried out by transforming the methane flow rate from 0 to 40 sccm at a 2-min interval increments. Changes in the flow rate of methane gas lead to a variation of the carbon content in the nanocomposite coatings. Both the interlayer and the transition layer were characterized by slight columnar and granular structures, which could contribute to the large numbers of nanocrystalline TiC and Ag particles embedded in the layer. Moreover, a transition from slight columnar and granular to glassy microstructure was also observed to respond to an increase in carbon content and the Ti concentration continuously decreased. Carbonaceous compound species from the plasma condensed on the surface of Ti bars. With
Fig. 4. SAED pattern indexing for TiC(Ag)/a-C:H nanocomposite coating with 1.13 at.% Ag, (b) 1.62 at.% Ag and (c) 0 at.% Ag.
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Fig. 5. SEM surface images of as-deposited coatings with different concentration: (a) 0 at.% Ag, (b) 0.56 at.% Ag, (c) 1.13 at.% Ag and (d) 1.62 at.% Ag.
relatively high CH4 flow rate (high CH4 concentrations in the atmosphere), carbonaceous compound species could be sputtered from the surface of Ti bars and the amount of Ti sputtered from the bars continuously decreased. In the case of low C amount, there was an ensemble of island-shape structure of Ag-rich layer which caused by Ag nanoparticles self-assembled into arrays attributed to the high surface energy of Ag relative to C observed in [29] during Ag deposition.
Subsequent carbon layer deposited hinders Ag atom diffusion and fixed the Ag-rich layer appearance. To summarize the results of the structural investigations, the following microstructure evolution in TiC(Ag)/a-C:H nanocomposite coatings with increasing Ag content was proposed: when the Ag content below 1 at.%, the coatings consisted of mostly amorphous structure; as Ag content increased, the coatings consisted of nanocrystalline Ag particles in
Fig. 6. FESEM micrographs of the fractured cross-sections of the nanocomposite coatings with different Ag concentration (a) 0 at.% Ag, (b) 0.56 at.% Ag, (c) 1.13 at.% Ag and (d) 1.62 at.%.
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Fig. 7. (a) Typical curve of nano-indentation load vs displacement into the surface of coating with 1.13 at.% Ag; (b) the mechanical properties of TiC(Ag)/a-C:H with different Ag concentration.
an amorphous a-C:H matrix and Ag clusters formed on the surface of the coatings. 3.2. Mechanical properties Coatings with different microstructures are expected to have different mechanical properties. The hardness was obtained by nanoindentation. Typical curve of 1.13 at.% Ag incorporation is shown in Fig. 7(a). Fig. 7(b) exhibits the effects of the Ag concentration on the hardness and the residual compressive stress of as-deposited coatings. As seen in Fig. 7(b), Ag doping leads to a monotonic decrease in hardness (H) with increment of Ag content. Besides, the incorporation of Ag also resulted in the reduction of the residual compressive stress. Results show that the incorporation of 0.56 at.% Ag caused a slight reduction of compressive stress. Increasing the Ag concentration further from 0.56 at.% to 1.13 at.% caused significant decrease of compressive stress from 2.7 to 1.5 GPa. Further incorporation of Ag up to 1.62 at.% results in a gradual decrease of the residual compress stress down to 1.4 GPa. Incorporation of Ag will induce a gradual reduction of the residual compressive stress concurrent with a similar reduction of hardness. This drastic decrease of the magnitude of residual compressive stress, however, was mainly due to formation of a metallic Ag phase (as shown in Fig. 3(c)) acting as efficient buffer sites to absorb the stress in the a-C:H matrix. The process of doping Ag nanoparticles into the amorphous coatings likely lead to two competing and contradicting effects on the hardness: (1) the addition of Ag nanoparticles deteriorated the mechanical properties which contributed to softening; and (2) the addition of Ag nanoparticles played a role as strengthening dopants to increase the coating hardness. This behavior has also been reported in the past in other systems including TiC/Ag, TiN/Ag and Mo2N/Ag [30–33]. According to the literature, the embedded nanoparticles of
silver in an amorphous matrix could form a heterogeneous structure and decrease the plastic flow. The final condition of the coating hardness depended on the compromise of these two effects. Suitable amount of Ag nanoparticles doped in the amorphous carbon matrix can promote the formation of Ag nanoparticles of favorable sizes, which maintained the hardness and reduced the intrinsic stress of a coating. The adhesive behaviors of the as-deposited TiC(Ag)/a-C:H coatings with 0 and 1.13 at.% Ag incorporation were evaluated by the scratch tests. It could be found that the 1.13 at.% Ag incorporation increased the critical load to 27 ± 1 N compared to the critical load of 18 ± 2 N of 0 at.% Ag incorporation. It is notable that the TiC(Ag)/a-C:H coatings with 1.13 at.% Ag incorporation showed a significant reduction in the residual compressive stress without a considerable decrease of the hardness, as well as improvement in adhesive strength, indicating that TiC(Ag)/a-C:H coatings with 1.13 at.% Ag incorporation is promising as solid lubrication coating from the perspective of tribological properties.
3.3. Tribological performances Fig. 8 shows the average friction coefficients of TiC(Ag)/a-C:H nanocomposite coatings were sensitive to Ag content at asdeposited and after annealed at 500 °C conditions, where increased Ag concentration generally resulted in a reduction of the friction coefficient. The lowest friction coefficient was obtained for the coating with 1.13 at.% Ag which showed an average value of 0.07 at ambient condition and 0.04 after annealed at 500 °C condition. While the higher friction coefficient obtained for the coating with 1.62 at.% Ag may be due to the higher Ag content on the surface of the coating, which can be seen in Fig. 5(d). Therefore, a proper Ag concentration in the coatings is required to facilitate lubrication.
Fig. 8. Friction coefficients of TiC(Ag)/a-C:H nanocomposite coatings with different Ag concentration: (a) as-deposited and (b) after annealed at 500 °C condition.
Y. Wang et al. / Surface & Coatings Technology 206 (2012) 3299–3308
The wear rates were calculated after scanning the wear track with various Ag concentrations, as shown in Fig. 9. As a whole, the wear rates of TiC(Ag)/a-C:H nanocomposite coatings both as-deposited and after annealed at 500 °C condition decreased with the increment of Ag concentration. The wear rates remarkably decreased for the coatings contained Ag from 0 at.% to 1.13 at.%, and then slightly fluctuated with the further increase of the Ag concentration to 1.62 at.%. Generally, the wear rates of carbon-based coatings after annealed at 500 °C condition should always be higher than that tested at ambient condition. However, the wear rates of as-deposited nanocomposite coatings after annealed at high temperature (500 °C) were always lower than those at ambient condition with Ag incorporation. Therefore, it can be concluded that a fraction of Ag in the nanocomposite coatings diffused to the surface of the coatings as an effective solid lubricant. The results show that the wear rate of TiC(Ag)/a-C:H nanocomposite coating with 1.13 at.% Ag incorporation (1.18 × 10− 16 m 3/N·m) was less than 15% of coating without Ag incorporation at high temperature condition. Fig. 10 shows the FESEM surface images of coating 1.13 at.% Ag concentration (a) as-deposited coating and (b) coating after annealed at 500 °C. The Ag clusters on the surface are significantly larger and denser than the initial Ag clusters on the surface of asdeposited coating, which indicates that Ag in the nanocomposite coatings diffused to the surface so that it facilitates to grow and coalesce of Ag particles. The larger and denser Ag clusters on the surface of the coating after heating could lubricate the coating surface to reduce friction coefficients and wear rates. The variation of friction behaviors tested at high vacuum (HV) condition of the nanocomposite coatings with different Ag content are shown in Fig. 11. As reported [34], the friction results under HV condition exhibit two different behaviors: after a running-in period, the friction coefficients stabilizes at very low values less than 0.01 (superlow friction coefficients) for the coatings with only Ag clusters on the surface of the coatings, while below this threshold the friction coefficient is in the milli-range during the first several hundred cycles and then increases rapidly up to 0.3–0.4. The experiments were then rapidly stopped. For each deposition process, the two different behaviors can be discriminated by the different Ag concentration of asdeposited coatings. Fig. 11 shows the evolution of the average friction coefficient in HV condition versus the number of sliding cycles for all four coatings. All the coatings had an initial running-in period, where the friction coefficient decreased from about 0.20 to b0.01 at steady state. The friction curve of 1.13 at.% Ag incorporation was characterized by a running-in period of around 700 cycles where the friction coefficient decreased slowly to 0.01 followed by a typical value between 0.002 and 0.01, the ultralow friction regime persists for about 1500 cycles. In the case of coating with 1.62 at.% Ag incorporation, the friction coefficient starts at 0.2 and diminishes rapidly, reaching the milli-range after only 200 cycles, nevertheless, the ultralow
Fig. 9. Specific wear rates of TiC(Ag)/a-C:H nanocomposite coatings with different Ag concentration: as-deposited and after annealed at 500 °C condition.
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Fig. 10. FESEM surface images of coating 1.13 at.% Ag concentration (a) as-deposited coating and (b) coating after annealed at 500 °C.
friction regime persists for about 500 cycles. The friction curve was disappearance between 300 and 700 cycles. The disappearance due to the machine was capable of reliable sensing loads that are 1 × 10 − 3 N. We used 1 N force in a sliding experiment, then being able to sense 0.001 N lateral force, which can translate into friction coefficients as low as 0.001. If the friction coefficient values were lower than 0.001 to be measured by the friction force sensor, it looks like the measurement ‘disappears’ between 300 and 700 cycles. For the coatings with the Ag concentration of 0 at.% and 0.56 at.%, the running-in time is only about 40 cycles, while the ultralow friction regime persists for just about 25 cycles. The sudden increase in friction coefficients above 0.1 for all as-deposited coatings indicates the coatings failure. Although the Ag concentration was different for each sample, the friction coefficient in the initial stage was at around 0.20 which should be due to the hydrogen atoms in a-C:H matrix. However, superlow friction for without Ag-doped coatings could not be kept for long periods of time at HV condition. The hydrogen atoms at the surface possibly disappeared during the friction tested. However, for Ag-doped coatings, the Ag clusters on the surface of coatings could be beneficial for tribological applications because the Ag clusters exhibit a low shear strength which could reduce the friction force on hard surface and improve the lifetime during sliding wear at HV condition. It was also anticipated that the Ag in the coatings should be drawn to diffuse both vertically and horizontally through the a-C:H matrix to the worn area so as to improve anti-wear properties. SEM images of wear tracks and scars tested in HV condition for the coatings with 0 and 1.13 at.% Ag content are estimated, as shown in Fig. 12. Many folds could be observed on the worn surface of coating with 1.13 at.% Ag content, as shown in Fig. 12(c). The existence of folds suggested that the worn surface of the coating was easily plastically deformed. However, no such folds existed on the worn surface of the coating without Ag incorporation, but the worn surface was characteristic with some small pits, as shown in Fig. 12(a). The occurrence of pits suggested that the worn surface of the coating without Ag incorporation was more brittle. The wear fragments for both coatings were also different. Many granular wear fragments were scattered on the worn surface of the coating without Ag incorporation, while many flattened wear fragments were observed on the worn surface of the coating with 1.13 at.% Ag content. Fig. 12(b) and (d) presents the SEM images of worn surface on mating balls, respectively. The transferred materials were negligible on the contact surface, while the wear fragments were clearly seen in the direction of sliding. A dark gray to black material accumulated mainly outside the initial contact area for the worn surface of the ball
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Fig. 11. Friction coefficients of TiC(Ag)/a-C:H nanocomposite coatings with different Ag concentration (a) 0 at.% Ag, (b) 0.56 at.% Ag, (c) 1.13 at.% Ag and (d) 1.62 at.% Ag tested at high vacuum.
mating for the sample without Ag incorporation, while this was not observed from the worn surface of ball mating for the coating with 1.13 at.% Ag content. No visible transfer film could be found on the worn surface for both coating mating balls, which agreed with the report by J. Fontaine et al. [35]: after the drastic friction increase, no transfer film can be found on the ball.
Fig. 13 shows the back scattered electron (BSE) images of wear scar on the surface of the coating with 1.13 at.% Ag content. The elemental distribution analysis of wear scar on the surface included three kinds of elements which are the Ag, C, Ti and O, respectively, and the relative content of each element could be speculated by the difference of color. As shown in Fig. 13, it could be proved directly
Fig. 12. SEM images of wear tracks and scars with different Ag concentration (a), (b) 0 at.% Ag and (c), (d) 1.13 at.% Ag tested at high vacuum.
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Fig. 13. Elemental distribution maps of wear scar of the coating with 1.13 at.% Ag tested at high vacuum : (a) Ag-La, (b) C-Ka, (C) Ti-Ka, (d) O-Ka and (e) back scattered electron (BSE) images of wear scar.
that the Ag and C contents was increased while Ti content was decreased in wear scar. We could say that these tribological performances should be attributed to the lubrication of the a-C:H coupled with effect of Ag clusters on the surface as well as the diffusion of the Ag in the coatings to the worn area so as to improve the antiwear properties. The friction evolution of TiC(Ag)/a-C:H coatings has demonstrated that a friction coefficient of 0.005 and a lifetime of 1500 cycles can be achieved in HV condition. These tribological performances should be attributed to the lubrication of the a-C:H coupled with effect of Ag clusters on the surface as well as the diffusion of the Ag in the coatings to the worn area so as to improve the anti-wear properties. The process is illustrated in the schematic picture shown in Fig. 14, and it indicates that the TiC(Ag)/a-C:H coating can achieve an excellent self-lubricating behavior during the sliding friction. The lifetime of the coatings with 0 at.% and 0.56 at.% Ag contents was short compared to the coatings with higher Ag concentration (the coatings with 1.13 at.% and 1.62 at.% Ag contents), which is mainly attributed that there was no Ag clusters on the surface of the coatings with 1.13 at.% and 1.62 at.% Ag contents.
4. Conclusions Nanocomposite TiC(Ag)/a-C:H coatings were successfully fabricated by mid-frequency dual-magnetron sputtering on silicon wafer substrates. Ag nanocrystallites were observed in the coatings with the Ag concentration of 1.13 and 1.62 at.%, and Ag clusters (10–50 nm) also formed on the surface. The friction coefficient and wear rates were sensitive to Ag content. In atmosphere, the friction coefficient and the wear rates was reduced as increase of Ag concentration. A properly Ag concentration in the coatings was required to facilitate lubrication. In high vacuum condition, at least 1.13 at.% Ag was needed to form a Ag lubrication mechanism, which reduced the friction coefficient from the 0.01 to 0.005, and increased the lifetime from 0 to 1500 cycles. The low friction coefficient and long lifetime of the TiC(Ag)/a-C:H coatings were attributed to the presence of low shear Ag clusters on the surface and the diffusion of the Ag in the coatings to the worn area so as to improve the anti-wear properties. Acknowledgments This work was carried out with financial support from the National Natural Science Foundation of China (No.11172300 and No. 11004203). References [1] [2] [3] [4] [5] [6] [7] [8]
Fig. 14. Schematic diagram of tribofilm induced by sliding friction for TiC(Ag)/a-C:H nanocomposite coatings tested at high vacuum.
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