Thick Low-Friction nc-MeC/a-C Nanocomposite Coatings on Ti-6Al-4V Alloy: Microstructure and Tribological Properties in Sliding Contact with a Ball

Coating thickness measured from TEM micrographs was set as 1.8 µm (Figure 1). It was found that the nc-WC/a-C coating has a columnar morphology, characteristic for the zone “T” in Thornton’s structure-zone model (SZM).[19,20] The column diameter was estimated to be in the range of 10 to 200 nm. It is known from literature that the columnar structure is often observed in the magnetron sputtered coatings.[21,22]

The different zones (1 to 3 and the T one) in the SZM model are classified depending on process parameters (e.g., substrate temperature, gas pressure in the recipient, or ions’ energy reaching the substrate surface). The substrate temperature is of great importance due to the fact that, if the homologous temperature T _s/T _m (where T _s is the substrate temperature during coating deposition and T _m means the melting point temperature of the coating material) is lower than 0.3, the surface diffusion during deposition is negligible and the coating consists of columns of a 10 to 20 nm diameter, approximately, separated by voids of a few nanometers in diameter. The columns would have numerous defects and can be amorphous.[19,20] In the present work, the homologous temperature during magnetron sputtering was approximately 0.1, however the column diameter was distinctly greater than that expected from the SZM model.

The columnar structure of the similar coatings (TiC/a-C:H) was explained as a consequence of shadowing effects that occur during growth, originating from the TiC intermediate layer.[23] It was reported in the literature that the boundary of the columns was rich in carbon.[23,24]

The electron diffraction patterns (SAEDs) taken from the coating were diffused and difficult to unequivocally interpret. Therefore, identification of the phases present in the coating was performed with use of the XRD, GIXRD as well as HRTEM techniques. The XRD and GIXRD patterns (Figure 2) show that different tungsten carbides are present in the coating: W₂C (trigonal primitive; tp), WC_0.98 (face-centered cubic; fcc), and W₃C (cubic primitive; cp). In addition, the presence of a β-WC_1−x (fcc) carbide cannot be excluded. The size and spatial distribution of the carbide nanocrystallites in the coating were investigated with use of the HRTEM technique. It was found that the coating was composed of nanocrystallites of a size 2 to 5 nm embedded in an amorphous carbon matrix (Figures 3(a), (b)). The calculated diffraction patterns were obtained with use of the FFT. Analysis of the FFT patterns for particular nanocrystals confirmed the presence of different tungsten carbides in the coating, most frequently of W₂C (tp) and, sporadically, of WC_0.98 (fcc) as well as W₃C (cp) ones. As an example, a FFT pattern from a W₂C (tp) nanocrystal oriented along [234] the direction marked as A is presented in Figure 3(b). It was observed that the nanocrystallites were non-uniformly distributed in the coating. The neighboring nanocrystallites were separated by an amorphous carbon matrix at a distance not exceeding 12 nm.

A 160-nm-thick graded interlayer was present between the nc-WC/a-C coating and the substrate material. Three sub-layers, 70, 30, and 60 nm thick (marked in Figure 4 as A1-A3, respectively) might be distinguished. Analysis of the SAED pattern marked as 2 in Figure 1 revealed that the main crystalline component in this interlayer was W _α (body-centered cubic; bcc). However, some rings indicating WC_0.98 (fcc) and W₂C (tp) were also found. The other component of these three sub-layers (visible in A1 to A3) was evidently an amorphous carbon. The STEM-EDS line profile microanalysis of the chemical composition confirmed a gradient concentration of W atoms in the interlayer, which decreased in the direction from the substrate to the coating surface (Figure 4). Contrary to that, the fraction of the amorphous carbon in the interlayer increased together with the distance from the substrate. In the substrate itself directly under the interlayer, a zone consisting of the Ti α (O) phase was demonstrated (Figure 1) with grains of the size 0.2 to 1 µm. This zone was produced during a preliminary diffusion hardening of the near-to-surface zone of the specimens from the Ti-6Al-4V alloy at 1173 K (899.9 °C) in a glow discharge plasma excited in the atmosphere Ar+O₂ before coating deposition.

Figure 5 shows a microstructure of the nc-TiC/a-C coating formed on an oxygen-hardened Ti-6Al-4V alloy, as well as the SAED patterns taken from the areas marked in the micrograph as 1 to 3 with their identification. SEM and TEM investigation of the nc-TiC/a-C coating cross-section showed that the coating was dense and of a uniform thickness of 3.3 µm (Figure 5). Microstructural analysis revealed that the nc-TiC/a-C coating was composed of TiC _x nanocrystallites embedded in an amorphous carbon matrix. Similar to the previously described nc-WC/a-C one, this coating also has a columnar morphology. A columnar morphology of the magnetron sputtered nc-TiC/a-C coating was also observed by Shaha.[25] With increasing intensity of the Ar ion impingement, brought about by increasing the frequency of pulsed DC sputtering, the coating microstructure changed from columnar to column-free.[26]

Due to the high volume fraction of an amorphous carbon, the rings present in the SAED patterns taken from the coating are diffused and, therefore, difficult to unequivocally identify. However, a series of rings corresponding to plausibly different interplanar distances of TiC and TiC_0.62 phases (both of them face-centered cubic; fcc) are visible in the electron diffraction pattern intensity profile of the coating (Figure 5, diffraction pattern no. 1, top right corner). Similar findings were observed in our earlier work on a thin (~500 nm) nc-TiC/a-C coating.[3]

The phase composition of the coating was also investigated using XRD, GIXRD, and HRTEM. The diffraction peaks on XRD and GIXRD spectra correspond to TiC (fcc) and TiC_0.62 (fcc) phases (Figure 6).

As shown in Figure 5, a 200-nm-thick graded interlayer formed between the nanocomposite coating and the substrate. This interlayer was produced in order to increase adhesion of the coating to the underlying bulk material. Two phases—Ti α (hcp) and TiC_0.62 (fcc) were identified in the interlayer by the SAED method (Figure 5, electron diffraction no. 2). The size of the nanocrystallites of both phases in the interlayer evaluated from the dark-field TEM images was estimated in the range 10 to 70 nm. In contrast, it was found from the HRTEM images of the top nc-TiC/a-C coating (Figure 7(a)) that the size of the nanocrystallites embedded in an amorphous matrix was in the range 1 to 3 nm. With use of the FFT (Figure 7(a)), the nanocrystallites embedded in the carbon matrix of the coating were identified as TiC_0.62 and TiC phases. The distance between the nanocrystallites varied from 2 to 15 nm. On the other hand, the Ti α and TiC_0.62 nanocrystallites of a size up to 25 nm were found at the interface between the nc-TiC/a-C coating and the interlayer (Figure 7(b)).

Quantitative metallography was carried out to quantify coating microstructure. The following parameters were determined: particles’ equivalent circle diameter (ECD), mean diameter (d), particle projection area A _A (particle part of the projection area), and the volume fraction V _V. In order to gain representative results, 77 and 90 nanoparticles revealed by HRTEM were analyzed for the nc-TiC/a-C and nc-WC/a-C coatings, respectively. The volume fraction of the nanoparticles was estimated according to a procedure described in detail by Wosik et al.[27] Table I contains the statistical characteristics of the nanocrystalline particles embedded in amorphous carbon matrices of both the nc-WC/a-C and nc-TiC/a-C nanocomposite coatings. Distribution of the nanocrystallites’ size (ECD) for the nc-WC/a-C and nc-TiC/a-C coatings is given in Figures 8(a) and (b), respectively.

The tribological properties of the coatings were investigated in sliding contact with the ball in two modes—at a constant load in rotation and at a progressive load in translational motion. The effect of the load on the friction coefficient and the load-bearing capacity in contact load conditions of a coating/substrate system were studied during sliding of an Al₂O₃ ball of 6 mm diameter under a load in the range of 0 to 30 N. Microscopic observation of the wear tracks in translational motion revealed neither cracks nor delamination of the coatings, even at 30 N. Independently, a stable and uniform friction without cracking or delamination was confirmed by the low and invariable value of acoustic emission (Figure 9). Pressing of the rigid ceramic ball into the coating under the load in the range 0 to 25 N caused an elastic deformation only of the coating/substrate system, as evidenced by the lack of a groove in the wear track. The residual depth R _d measured after unloading, corresponding to the load of 25 N, reached an average value of 0.03 µm for the nc-WC/a-C coating and 0.04 µm for the nc-TiC/a-C one. After exceeding the value of 25 N, the residual depth R _d reached at most 0.04 and 0.12 µm for the coatings with WC and TiC nanocrystallites, respectively (Figure 9).

The analysis of stresses and strains in the contact load conditions of the coating/substrate systems is often carried out with the application of Hertz’s theory, which is extended for such systems[28] or forms the basis of FEM modeling.[29] According to the basic assumption of Hertz’s theory, the calculations must be carried out in the elastic range. For this reason, in the contact of the ceramic ball with the coating under consideration, at first a mean value of the contact pressure was determined and then it was compared with the yield stress for the systems.

In this paper, the mean value of the contact pressure (p _m) at the given load (assuming the stationary contact of the ball with the plane at the given moment) was designated in accordance with Hertz’s theory as[30]:andwhere p _m is mean contact pressure, a is contact radius, R is radius of the ball, F is load, E* is composite (reduced) elastic modulus of the bodies in contact 1/E* = (1 − ν ²)/E + (1 − ν ₁ ² )/E ₁, where E, E ₁, and ν, ν ₁ are elastic modulus and Poisson’s ratio of the plane and the ball, respectively.

The mean value of the contact pressure determined for the loads from 1 to 30 N changed in the range of 0.46 to 1.43 GPa for the system with the nc-WC/a-C coating and 0.37 to 1.14 GPa for the nc-TiC/a-C one. Plastic deformation can occur when the contact pressure is greater than the yield stress Y in accordance with the relationship, p _m > 1.1Y.[30‐33] If p _m is in the range 1.1Y < p _m < 3Y the yield stress can be dependent on the hardness (H), and for greater loads the H/Y ratio receives a relatively constant value equal to 3.

The analysis of the deformation of the coating/substrate system in contact with the Al₂O₃ ball of a 6 mm diameter shows that the maximum load of 30 N used in the tests does not cause plastic deformation, while the mean contact pressure p _m corresponding to the applied load is less than Y = 4.7 GPa for nc-WC/a-C and Y = 3.4 GPa for nc-TiC/a-C coatings. This state of deformation will also be retained when the mean contact pressure is twofold greater, which corresponds to the load distributed only on half of the maximum contact area of the ball (πa ² /2) sliding on a flat surface.

The COF (µ) in sliding contact is generally composed of two parts: an adhesive component, which is dependent on the shear strength of the interface (S), and the deformation component, also called the plowing term:where F _F is friction force, A is area of the contact, S is shear strength of the interface, and F _p is plowing term, F is applied normal load to the ball.

The shear strength S at a high pressure depends on the material parameters S ₀ and α, which represent material properties controlling friction and can be described approximately as, S = S ₀ + αp _m.[34] The friction coefficient can be reduced when the area of contact, the shear strength, and the plowing contribution are minimized.

The values of these components depend on many factors, mostly on the load-carrying ability of the tribological contact. This is of particular importance during friction of the coating/substrate systems when exceeding the critical load at a contact point is conducive to cracking of the coating and to subsequent delamination of the coating. Below the elastic limit and, for a relatively thin coating, the plowing component can be omitted, and the friction coefficient is determined by the shear strength of the interface and the area of contact.

The friction tests showed a decrease in the COF during an increase in the load, particularly for the loads greater than 1 N (Figure 9). The change of COF in the initial period below 1 N was the result of contact formation and the resistance to motion depended on the real surface roughness, which was not perfectly smooth. Hence, the results are shown only for the range of 1 to 30 N. The nc-TiC/a-C coating has a lower friction in comparison with the nc-WC/a-C one, which is due to the higher volume fraction of a lubricating amorphous carbon phase in the nc-TiC/a-C coating than in the coating with the WC nanocrystallites. Due to the lower hardness and stiffness of the nc-TiC/a-C coating, it will have a lower shear strength too.

The decrease in the COF with the load increase can be explained by the Hertzian contact model, which has been previously used for analyzing friction of other solid lubricant coatings.[34,35] The COF for a smooth and rigid ball in sliding contact with a flat element with a perfectly smooth surface in the range of elastic deformation can be calculated with the formula:where S ₀ is interfacial shear strength between the ball surface and the underlying wear track, R is radius of the ball, E* is composite elastic modulus of the bodies in contact, α represents the pressure dependence of the shear strength.

The change in the value of the friction coefficient is distinct from 0.13 to 0.09, but its range is small, due to the specific nature of co-operation in such a test, where the ball under progressive loading is in contact with the coating only once on the path friction (Figure 10). Over the load of 25 N the friction coefficient increased to a very low extent. Hard carbide particles strengthen the structure of the composite but may also disturb the friction process and, in extreme cases, may cause a big stress concentration. Tested coatings belong to the group of so-called solid lubricants and their low-friction properties are time-dependent, as confirmed in the friction tests under a constant load during a long time (Figure 11).

Comparison of the tribological properties of the investigated coatings possessing different thicknesses is possible because they are relatively thick and deposited on a substrate with a hardness and elastic modulus comparable to the coating. Hence, an oxygen-hardened alloy provides strong support for the coatings and creates a uniform coating/substrate system.

The friction processes for both coatings in dry sliding contact with the alumina ball at a constant load were stable and ran in a similar way. The greatest changes in friction force were observed in the initial period of the sliding friction, which was related to the matching of the contact surfaces. The average value of the initial COF was equal to 0.07 for the nc-WC/a-C and 0.06 for the nc-TiC/a-C-coated alloy (Figure 11). Then the friction force slightly decreased with time until a constant value was reached. The COF values in the steady-state period were approximately 0.05 for the nc-WC/a-C coating and slightly less (0.04) for the nc-TiC/a-C one. This low resistance to motion, as well as its decrease as a function of time, was due to the formation of a carbonaceous third body material during friction. As a result, a self-lubricating film was formed on the counter-face (i.e., on the ceramic Al₂O₃ ball) in the contact zone, which enabled the easy shear of the sliding couple and protected the surfaces from destruction (Figure 12). This film is assumed to be composed mainly of graphite-like carbon atom clusters.[4,5,7] The rate of the self-lubricating film formation was greater for the nc-TiC/a-C coating as a result of a greater volume fraction of the amorphous a-C carbon matrix in the coating. Due to that, as well as to a lower hardness of the coating, the contact area during dry friction of this coating underwent faster and easier matching. The lubricating film separated the surface of the ceramic ball from that of the nanocomposite coating and effectively decreased the resistance to motion. It is worth mentioning that, due to the low friction as well as to a low load applied to the ball which did not bring about the plastic deformation of the coating, the deformation component of the friction force was very small. The calculated value of the Hertzian contact pressure at the beginning of each sliding test was in the range of 0.5 to 0.62 GPa.

The wear mechanism of the nanocomposite coatings with an amorphous carbon a-C matrix is frequently explained in the literature in terms of the graphitization process.[4,35‐37] This process occurs due to transformation of the sp³-bonded carbon to the sp²-bonded one. The latter is characteristic of a graphite-like hexagonal structure with superior sliding properties, however, of a smaller strength. The process of sliding friction provides the energy required for the transformation of the diamond-like to graphite-like carbon under thermal and strain effects from the repeated friction cycles.[36]

In sliding contact with the Al₂O₃ ball, the wear of both composite coatings was mainly of a micro-abrasive and adhesive nature with no significant wear of the ball surface. Adhesive wear occurred after the formation of the self-lubricating film and may be due to the transfer of this film to the counter-face and back to the wear track, which was also observed by other researchers.[36] The occurrence of hard particles in the contact zone disturbed the steady-state sliding process with lubrication by the graphite-like film and intensified the abrasion. The different mixtures of the carbides W₂C, WC_0.98, and W₃C in the nc-WC/a-C coating and TiC_0.62 and TiC ones in the nc-TiC/a-C coating could bring about changes in the sliding condition due to different crystallographic structures of different phases. As a result, distinct scratches along the sliding direction and the wear debris at the edges of grooves, characteristic of abrasive wear, were observed in the wear tracks. The cross-sectional profiles of the grooves formed in the coatings by the ball are shown in Figure 13. The absence of pile-ups at the grooves’ edges gives evidence of elastic deformations only in the contact zone. Suitable load-bearing capacity of the coating/substrate system was provided by a particular duplex treatment which ensured an increase of the coating lifetime on the hardened substrate in comparison with the one on a baseline substrate and protected against severe wear.[2,38] The profile of the wear groove in the nc-TiC/a-C coating was more irregular than that in the nc-WC/a-C one. The greater roughness of the worn surface was due to a greater number of scratches formed as a result of the micro-abrasion of the softer coating by the hard carbides. The wear resistance of the a-C matrix without reinforcing particles was significantly lower in comparison with the composite coatings with WC and TiC nanocrystallites. It is worth noticing that the wear of the a-C coating during friction was caused by their abrasion and delamination as shown in the previous studies.

During the scratch test, the first cohesive cracks occurred when the load L _C1 achieved the value of 5 and 7.5 N for the nc-WC/a-C and nc-TiC/a-C coatings, respectively, whereas the adhesive cracks were observed at the critical load L _C2 equal to 12 or 13 N for the nc-WC/a-C or nc-TiC/a-C coatings, respectively (Table II). Even though the resistance to scratching was rather low, the coatings exhibited a good wear resistance during sliding friction. The wear rate (W _V) of the nc-WC/a-C coating was equal to 0.08 × 10⁻⁶ mm³ N⁻¹ m⁻¹, whereas for the nc-TiC/a-C one it was almost four times higher (0.28 × 10⁻⁶ mm³ N⁻¹ m⁻¹), even though the COF values of both coatings were close to each other. This behavior can be explained in part by the lower hardness of the nc-TiC/a-C coating, as well as by the lower volume fraction of the titanium carbides in this coating (merely 3 pct), which did not strengthen their structure enough. In contrast to that, the wear resistance of the coating with tungsten carbides nanocrystallites was much greater, plausibly due to its greater elasticity (Figure 9). The elastic strain to failure, which is related to the ratio of hardness (H) and elastic modulus (E), has been shown by a number of authors to be a more suitable parameter for predicting a coating’s resistance to wear than its hardness alone.[39‐41] The high values of the H/E ratio, above 0.1, result in high wear resistance as reported by Pei et al.[41] Additionally, the H ³/E ² ratio should be a strong indicator of a coating’s resistance to plastic deformation in this type of contact loading. The value of the H ³/E ² ratio for the nc-WC/a-C coating is close to 0.1 and is higher than nc-TiC/a-C one, what results in the higher resistance to plastic deformation of the nc-WC/a-C coating as well as its higher wear resistance. The ratio (R _d/P _d) of the residual depth R _d to the penetration depth P _d obtained from the scratch test was 0.16 for the specimen with the nc-WC/a-C coating and was approximately twice smaller (0.29) than that for the specimen with the nc-TiC/a-C one (Figure 14). This result was due to a higher volume fraction of tungsten carbide nanocrystallites, which increased the load-bearing capacity and rigidity of the nc-WC/a-C coating. Additionally, the crack propagation in this coating is more difficult, what is reflected in the reduced plastic deformation during the hardness measurement.