Influence of TiC/a-C phase on tribological properties of composite coatings

Introduction

Composite coatings consisting of crystallites of hard phases, such as carbides or nitrides impregnated in an amorphous matrix, provide significant improvement in hardness, toughness, wear resistance and lowering of friction.1^–4 Low friction is derived by creating a soft matrix of amorphous phase embedded with nanocrystallites. Nature of bonding, crystal size and amount of each phase determine the mechanical and tribological properties of the TiC/amorphous carbon (a-C) coatings. Thus, if the composite consists of crystals of hard phase embedded in a soft phase, it exhibits higher hardness and lower friction. 5 5,6 These soft phases usually have lower hardness with enhanced self-tribolubricant properties.5^–7 If this phase is hard, grain boundary sliding is blocked, which hinders the plasticity. Then, this phase is able to bear triboinduced stress to avoid deformation. Consequently, it may generate an extra amount of adhesive strength and cause higher friction where sliding of softer phase can be disrupted by hard crystallites. In a composite material, the content of the softer phase can cause lower wear resistance, yielding higher friction due to abrasion. Therefore, only an ideal combination of both these phases can improve the hardness with lower friction. 3 3,6

The motivation of the present work is to understand the tribological behaviour of TiC/a-C composites by establishing correlation among phase compositions, coating microstructure with tribological properties. For this purpose, TiC/a-C overlayer coatings on titanium substrate were prepared by gas phase carburisation in CH₄ atmosphere. The synthesis temperature was varied to change the relative extent of TiC with a-C in composite coatings.

Experimental

The synthesis of TiC overlayer coatings was carried out by controlled atmospheric exposure of titanium to CH₄ gas in a thermogravimetric analyser. The experiments were carried out at two processing temperatures, i.e. 850 and 950°C, with 4 sccm flow of CH₄ gas. Further experimental details can be found elsewhere.8 The crystal structure of the coatings was examined by glancing incident X-ray diffraction (GIXRD). The linear reciprocating mode of a ball on disc tribometer (CSM, Switzerland) was used to carry out tribological tests of these coatings. A spherical steel ball (100Cr6 SS) of Ø6 mm with surface roughness of 0·06 μm was used as a sliding body to measure the value of the coefficient of friction. Tribotest was performed at a normal load of 5 N with a sliding speed of 1·5 cm s⁻¹. Tests were performed at ambient (dry and unlubricated) condition having a relative humidity of 58%. Surface roughness and normalised wear rates of the coatings were evaluated from the cross-sectional profiles across the wear track after dimensional measurement by a surface profilometer (Dektak 6M). Microindentation test at 0·196 N normal loads was carried out to measure the hardness of the coatings. The SEM and atomic force microscopy (AFM) techniques were used to observe the coating's morphology and the surface topography of deformed wear tracks respectively. Micro-Raman (Horiba Jobin Yvon, HR 800, France) measurements were performed using a spectrometer equipped with an optical microscope to record the phase evolution of the coatings and on the surface of the wear tracks. For this purpose, laser radiation with wavelength of 514·5 nm and 5 mW power was used as an excitation source.

Results and discussion

The X-ray diffraction analysis presented in Fig. 1a and b clearly shows that the coatings grown at the temperatures of 850 and 950°C were predominantly TiC with some unreacted titanium, as seen by peaks matching with the JCPDF data card no. 44-1294. From the peak width of TiC (111), TiC (200) peaks, and the crystallite size of these coatings determined by the Scherrer formula was found to be 64 and 78 nm respectively, as shown in Fig. 1a and b. These peaks occur at diffraction angles of 42·33 and 36·19° respectively. The broadening of TiC (200) peak is found to be lower than TiC (111). The growth of the crystallite is found to change with the synthesis temperature. The intensity of TiC (200) shown in Fig. 1a is found to be significantly higher compared to TiC (200), as shown in Fig. 1b. The intensity of TiC (111) is found to be low in Fig. 1a. From the JCPDF data, the peak of the TiC (111) line is centred at a diffraction angle of 35·85°. The shift of TiC (111) to higher diffraction angles of 36·40 and 36·19°, as shown in Fig. 1a and b respectively, describes the presence of compressive residual stress, which restricts the expansion of crystallites at higher synthesis temperature. The TiC (200) peak is centred at a diffraction angle of 41·64°, where this is found to be shifted to the higher angle of 42·33°. Generally, all peaks corresponding to TiC and Ti are found to be shifted towards a higher diffraction angle. No peak corresponding to carbon phases is present in GIXRD due to its amorphous nature. The SEM images of these coatings are presented in Fig. 2, where a columnar cluster of grains is found to be organised. At lower synthesis temperature, this was found to be smaller in size, as shown in Fig. 2a compared to Fig. 2b. The thicknesses of the coatings were found to be 1·8 and 7·8 μm, synthesised at the temperatures of 850 and 950°C respectively. This could be explained based on the fact that the rate of decomposition of CH₄ is higher at higher temperature. The evolution of the TiC phase is governed by a reaction between metallic titanium and incoming carbon species arising from the dissociation of CH₄ on the surface. At lower processing temperature, the titanium content is higher than that of carbon obtained from the decomposition of CH₄, which results in the formation of a titanium rich coating. With increasing processing temperature, the ratio of Ti/a-C close to unity is achieved, providing favourable growth conditions for the TiC. 9 9,10 The surface roughness was found to increase with the coating thickness and found to be 73 and 132 nm at the synthesis temperatures of 850 and 950°C respectively. The increase in surface roughness is mainly related to the formation of larger clusters, as shown in Fig. 2b.

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Figure 1

Glancing incident X-ray diffraction of coatings synthesised at a 850°C and b 950°C

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Figure 2

Topography of coatings synthesised at a 850°C and b 950°C

From Fig. 3, it is clear that the value of the coefficient of friction of high temperature coating (curve b) is lower than the value for low temperature coating (curve a). A clear transition, indicated by (T), from lower to higher values of coefficient of friction is seen in both cases. This transition was found to directly depend on the thickness of the coatings. This transition is observed to ocuur at longer sliding distance (∼19 m) in high temperature coating, as shown in curve b, in comparison to low temperature coating when these occur at a sliding distance of ∼3·8 m (curve a). The transition to higher friction arises due to the sliding of ball against the titanium rich secondary sublayer of the coatings, whereas the upper layer is made by a TiC phase with a-C contents. This top layer breaks down and exposes the layers underneath during triboinduced sliding, causing the transition. The coefficient of friction presented in curve a in Fig. 3 is higher than the value obtained for titanium substrate, as shown in curve c in Fig. 3. The instability in curve c in Fig. 3 arises due to the adhesive nature of contact deformation. The coefficient of friction is also found to be influenced by the triboinduced deformation of wear surface that normally results in the formation of cracks and grooves that are clearly visible in the SEM images, as shown in Fig. 3. The magnitude of these cracks and the formation of grooves due to coating breakdown are found to be larger in high temperature coatings (Fig. 3, curve b), where the coefficient of friction was found to be lower. The formation of these cracks and grooves is mainly caused by rupturing of the coating surface attributed to fatigue fretting wear due to the generation of cyclic tribo stresses. In the SEM image corresponding to curve a, the scale of cracks on the wear surface is smaller, and grooves are absent, but the coefficient of friction is found to be higher. A partially distributed oxide scale is also visible in this image. This can be explained on the basis that failure of the coating occurred in the early stage of sliding and after the removal of coating contact sliding takes place with the titanium enriched substrate surface.

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Figure 3

Coefficient of friction of coatings synthesised at a 850°C and b 950°C and c titanium substrate

Atomic force microscopy imaging of coatings and wear surfaces is shown in Fig. 4. The image in Fig. 4a shows a cluster of grains arranged on the surface of the coating synthesised at 950°C. These clusters are found to be plastically deformed at microscopic scale due to the generation of cyclic stresses during triboinduced sliding. The deformation of these clusters is clearly noticed in Fig. 4b. In the inset of this image, a zoomed portion is shown, which clearly indicates plastic deformation of these clusters. Deformation of sharp elongated and spherical asperities of clusters during sliding smoothen the surface of the wear track. In this condition, contacts per unit area of the sliding surfaces increase, causing increasing adhesion between the surfaces. This causes dissipation of higher frictional energy with increasing wear distance. When the wear surface is found to be completely smooth, the coefficient of friction attains a maximum value and remains steady, as shown in Fig. 3a and b after ∼14 and 24 m of sliding distances respectively.

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Figure 4

Topography (AFM) of coating synthesised at a 950°C and b topography of wear surface after 40 m sliding distance

Raman spectroscopy measurements were performed to explain the contribution of structure and phase on the coefficient of friction. It is found that during the friction test, different tribo chemical reactions are induced between the counter probe and the surface of the wear track, producing significant modifications of the Raman spectra obtained from wear scars. This was compared with the spectra obtained from the surface of the coatings before tribo testing. These differences are useful to evaluate and quantify the coefficient of friction with respect to the phase and structure of the coatings. In this context, some peaks are detected in the range of 227–826 cm⁻¹ for low and high temperature coatings, as shown in Fig. 5a and b. These vibrational modes point to the formation of a hypostoichiometric TiC. It is established that a stoichiometric TiC has no Raman vibration modes, but the contribution of disorder due to carbon vacancies makes them active.10 The intensity of the TiC phase is found to be strongest at 227 cm⁻¹ in low temperature coating, as seen in Fig. 5a, whereas the vibrational modes at 445 and 609 cm⁻¹ shown in Fig. 5b are fairly strong in the high temperature coating. All these vibrational modes correspond to the formation of a substoichiometric TiC phase. The vibrational mode of the TiC phase is found to be very weak in the wear track formed in Fig. 5a compared to Fig. 5b. This region is related with wear loss of the TiC layer along with a-C during triboinduced sliding. Weak evolution of the metal oxide, such as α-Fe₂O₃, is also observed at a frequency band of 320 cm⁻¹ in the wear track presented in Fig. 5a. The origin of this peak can be attributed to the formation of α-Fe₂O₃ by tribo chemical reaction of the steel ball (100Cr6 steel) with the surrounding atmosphere under the influence of frictional heating, which generally results in a higher value of coefficient of friction. 6 6,11 The higher value is also attributed to the rise in adhesive strength due to the formation of interfacial bonding between oxides and metal surface. The vibrational modes from TiC at 445 and 609 cm⁻¹, observed in the wear track, as seen in Fig. 5b, are stronger than those seen in Fig. 5a. This phase is responsible for the low value of the coefficient of friction and the increased wear resistance, as shown in curve b in Fig. 3. The intensity of these peaks in the wear surface is found to be lower than the peaks observed at the coating surface. This is attributed to the removal of triboinduced TiC/a-C layer during sliding. It also shows that thin coatings have low load bearing capacity, and therefore, the substrate undergoes deformation after contact sliding. This generates stresses at the interface that weaken the coating adhesion and lead to wearing. Earlier transition from lower to higher friction regime is illustrated in curve a in Fig. 3 compared to curve b in Fig. 3. This has a direct link with the formation of a thin layer of TiC/a-C coatings on the substrate, which wears out early during tribo sliding. This explains the absence of significant peaks corresponding to TiC in the Raman spectra.

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Figure 5

Raman spectra of coatings synthesised at a 850°C and b 950°C on coating and wear surfaces

Along with the contribution of the TiC phase, the influence of a-C in the coating is also observed. It is known that Raman scattering is the inelastic scattering of photons by phonons due to the change of polarisability caused by the phonon mode. However, in a graphite-like structure, the bandgap lies in the visible range only within a small part of k space around the K point. All these bands have π character. In this case, photons resonantly excite states only at the k vector, where the bandgap is equal to the photon energy. This sets up a polarisation density wave of this k vector. Its intensity is strong because of the long range polarisability of π states. The vibrational G mode of graphite has E_2g symmetry. The eigenvector of this mode occupies in plane bond stretching motion of pairs of sp2 C atoms.12 This mode does not require the presence of sixfold ring, and so it occurs at all sp2 sites, not only those in rings. The D band is called a breathing mode of A_1g symmetry involving phonons near the K zone boundary. This mode is absent in perfect graphite and only becomes active in the presence of disorder. The D band vibrational mode is dispersive; it varies with the photon excitation energy, even when the G peak is not dispersive. In this regard, peak intensities of breathing D mode and stretching G band modes are obtained at 1348 and 1589 cm⁻¹ respectively. These correspond to ‘disordered’ and ‘graphitic’-like structures of the a-C network. On the coating surfaces, the intensity, bandwidth and peak position of these phases are found to be similar in both Fig. 5a and b. The appearance of a shoulder in the G band frequency at 1621 cm⁻¹ also seen in these figures is due to the D band mode, which arises due to the contribution from non-zero centre phonons.12 The intensity and broadening of these D and G band vibrational modes explain the nature of bonding and the degree of disorder present in graphite-like carbon atoms. The higher order of broadening of the G band mode with weak intensity relative to the D band vibrational mode describes the higher magnitude of disorder present in sp2 graphite-like structure.13 In spite of having diverse physical properties, the families of carbon material based on sp2 bonding network generally have low friction and wear resistance. Wear induced tribo chemical surface reaction can change the nature of bonding, which in turn can influence friction and wear. 14 14,15 The intensities of these bands (D and G modes) are found to be weak in the wear track, while these bands are found to be prominent on the surface of the coatings. In Fig. 5b, the intensities of these phases are observed to be stronger than the spectra observed in Fig. 5a on the wear surface. The presence of a a-C enriched layer can reduce the friction due to easy sliding, while after wear induced removal of this layer, the friction values increase as sliding occurs between the metallic titanium enriched surface and the steel ball. A negligible low value of coefficient of friction is exhibited due to the contribution of a-C with strong evolution of G band vibrational mode. The D band is found to dominate on the specimen surface synthesised at both the temperatures of 850 and 950°C. In this context, the intensity of the D band on the wear surface is significantly reduced as compared to G band vibrational mode, as shown in Fig. 5a and b. This could be explained by triboinduced sliding, which creates higher phase disordering in D bands. It is well established that the lamellar structure of the sp2 bonded carbon atoms has a lower friction due to the weak van der Waals forces exhibiting between the basal planes, which easily shear during triboinduced sliding. The contribution of sp2 phase with the G band vibrational mode acts as a tribolubricating layer that reduces the resistance to shearing. This phenomenon causes a lower value of the coefficient of friction. The contribution of TiC phase assumes significance as its embedment into amorphous matrix provides a strengthening mechanism that improves the wear resistance. It can be seen that although the magnitudes of D and G bands are similar under both synthesis temperatures, friction is found to be lower in Fig. 5b, where the TiC phase has a stronger signal than that in Fig. 5a observed in the wear track. Therefore, the effective tribolubricant properties of this composite are also determined by the contribution of TiC phase along with the presence of D and G band vibrational modes in sp2 sites.

The ideal combination of both D and G band vibrational modes can increase the hardness and wear resistance properties of composite coatings. In this regard, hardness was found to be 7·74 and 9·68 GPa with wear rates of 9·5×10⁻⁶ and 7·5×10⁻⁷ mm³ N⁻¹ m⁻¹, for coatings synthesised at 850 and 950°C respectively. The hardness and wear rates are measured to be 1·8 GPa and 5×10⁻⁵ mm³ N⁻¹ m⁻¹, respectively for the titanium substrate. These values, when compared with those obtained from composite coating, show a marked difference. The hardness of the substrate is significantly lower with higher wear rate. This relationship is seen in Fig. 6. The measured hardness has a direct significance with the volume fraction of hard TiC present in the a-C phase of the coatings. In addition, carbides have a high bond strength, which exceeds the strength of the host material. It acts as a hard binder phase in the composite. 1 1,5 The ratio of binder/carbide phase influences the resulting values of hardness and toughness of this composite. The increase in hardness is also associated with the lower crystallite size, as seen in Fig. 1b, which disrupts the dislocation motion during deformation.

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Figure 6

Hardness and wear rate of coatings synthesised at a 850°C and b 950°C and c titanium substrate

Conclusion

In the present study, composite coatings of TiC/a-C were synthesised using gas phase carburisation in CH₄ atmosphere on the titanium substrate. The thickness and morphology of the coating surface were found to depend on the processing temperature. Evolution of the TiC/a-C phase was found to be higher with higher processing temperature. The formation of a-C and TiC phases strongly influences the tribological properties of these composite coatings. Low friction was found to be driven by the formation of TiC and amorphous hexagonal sp2 graphite-like phase with the increase in intensity of the G band, a well known tribolubricant. The hardness and wear resistance properties are found to be enhanced by the combination of D band present in the TiC matrix. A clear transition from lower to higher friction was observed where wear loss and deformation of the tribolubricant phase cause contact with the sublayer of the titanium reach surface. In thicker coating, such low friction regime persisted over longer sliding distances.