Substrate effect on wear resistant transition metal nitride hard coatings: Microstructure and tribo-mechanical properties

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Substrate effect on wear resistant transition metal nitride hard coatings: microstructure and tribomechanical properties D. Dinesh Kumar, N. Kumar, S. Kalaiselvam, S. Dash, R. Jayavel

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S0272-8842(15)00804-4 http://dx.doi.org/10.1016/j.ceramint.2015.04.059 CERI10466

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21 February 2015 10 April 2015 11 April 2015

Cite this article as: D. Dinesh Kumar, N. Kumar, S. Kalaiselvam, S. Dash, R. Jayavel, Substrate effect on wear resistant transition metal nitride hard coatings: microstructure and tribo-mechanical properties, Ceramics International, http://dx.doi.org/10.1016/j. ceramint.2015.04.059 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Substrate effect on wear resistant transition metal nitride hard coatings: microstructure and tribo-mechanical properties D. Dinesh Kumar a, b, N. Kumar c,*, S. Kalaiselvam b, S. Dash c, R. Jayavel a a

Centre for Nanoscience and Technology, Anna University, Chennai, Tamil Nadu, India600025. b

Department of Applied Science and Technology, Anna University, Chennai, Tamil Nadu, India-600025.

c

Materials Science Group, Indira Gandhi centre for Atomic Research, Kalpakkam, Tamil Nadu, India-603102.

*

Corresponding Author: Tel./fax: +91 4427480081.

E-mail: [email protected] (N. Kumar) Abstract Four types of different hard transition metal nitride (TMN: ZrN, CrN, WN and TiN) coatings were deposited on Si (100) and 316LN stainless steel substrates using DC magnetron sputtering. A comprehensive study of microstructure and substrate dependent tribomechanical properties of TMN coatings was carried out. Higher hardness (H) and elastic modulus (E) were obtained for WN (H=40 GPa and E=440 GPa) and TiN (H=30 GPa and E=399 GPa) coatings. This is related to the formation of (100) and (111) preferred orientations in WN and TiN coatings, respectively. However, the less hardness and elastic modulus were obtained for ZrN and CrN coatings where (200) orientation is preferred. Remarkably, low friction coefficient (0.06 – 0.57) and higher wear resistance in the coatings deposited on steel substrates are directly associated with the higher resistance to plastic deformation (H3/E2) and the presence of intrinsic compressive stress. Three body wear modes enhanced the friction coefficient (0.15 – 0.62) and the wear rate in the coatings deposited on Si substrates. This is primarily associated with low fracture toughness of brittle single crystalline Si (100) substrates. Steel-on-steel contact was dominated in ZrN/steel sliding

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system. This occurs due to the severe adhesive wear mode of steel ball, whereas, the abrasive wear modes were attained for the CrN, WN and TiN coatings sliding against steel balls. Keywords: Microstructure, tribo-mechanical properties, friction, wear resistance

1. Introduction Low friction coefficient and high wear resistance are eventually required for the enhancing energy efficiency of moving mechanical components in semiconductor industries, cutting tools, spacecraft and biomedical implant industries [1-5]. It is known that due to the high friction and wear, the sustainability and work efficiency of the mechanical components reduces. Therefore, improving the hardness, elastic modulus and fracture toughness of materials are the key parameters for enhancing tribological properties of mechanical components [6]. The designing of nanocrystalline hard coating is an important way for improving the tribo-mechanical properties [7]. TMN coatings in binary phase are another important class of materials for several tribological applications owing to their excellent physicochemical, mechanical and tribological properties [8-12]. Among them, the tribomechanical properties of nanocrystalline TiN coatings deposited by several techniques have been extensively studied [13-16]. However, the reports on the microstructure, mechanical and tribological properties of other TMN coatings (CrN and ZrN) are limited [4,17,18]. The WN coating is another important refractory material with extreme hardness and barrier properties. This coating is largely studied for the microelectronic applications, but less attention is paid for tribological point of view [9,19,20]. In general, the tribo-mechanical properties are inherently related to the composition and microstructural characteristics, which can be typically controlled by the deposition conditions [21]. The PVD [22] techniques are useful to deposit the TMN coatings, amongst them, the magnetron sputtering is commonly applied due to its versatility, uniformity and its large industrial applications [23,24]. Unfortunately, there

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is still a lack of knowledge to understand the substrate dependent microstructure and tribomechanical characteristics of TMN coatings deposited by PVD techniques. The main aim of the present work was to deposition and characterization of ZrN, CrN, WN and TiN coatings on single crystalline Si (100) and polycrystalline 316 LN stainless steel substrates using reactive DC magnetron sputtering. Si substrate was chosen owing to its extensive applications in semiconductor industries; while, the stainless steel is applicable in cutting tool industries. The observed results reveal an interesting correlation with the composition and microstructure; and it provides a remarkable relationship between the microstructure and tribo-mechanical behavior of the coatings. The substrate dependent friction and wear behavior of these coatings were also elucidated and explained on the basis of coating microstructure.

2. Experimental details 2.1 Coatings deposition ZrN, CrN, WN and TiN coatings were deposited on Si (100) and 316LN SS substrates using DC magnetron sputtering technique (Hind High-Vacuum, Bangalore, India) at the substrate temperature (Ts) of 250°C. Commercially available polycrystalline AISI 316LN SS (15 × 15 × 1.0 mm) and Si (100) (10 × 10 mm × 500 µm) substrates were used for the deposition. The polishing and surface cleaning process for 316LN SS substrates were reported elsewhere [25]. The target-substrate distance was kept constant approximately 70 mm for all the samples. High pure (99.997%) Zr, Cr, W and Ti discs (50 mm diameter and 3 mm thickness) were used as target materials procured from Taewon Scientific Co. Ltd., Korea. Typical base pressure of the vacuum chamber was achieved around 9 × 10-6 m.bar. In order to remove the surface contamination, each target surface was pre-sputtered for 5 minutes. During deposition, ultra high pure (99.999%) Ar and reactive N2 gases were fed into the chamber and the total deposition pressure was maintained around 2×10-3 m.bar. For this, the Ar gas flow 3

rate was kept constant at 30 sccm; whereas, the N2 flow rate was varied for different layers (5 sccm for ZrN & TiN and 8 sccm for CrN & WN coatings). The nitrogen flow rates in TMN coatings were individually optimized by various characterization techniques, and the optimized flow rate (as mentioned above) was used to deposit TMN coatings. To improve the adhesion between the coating and substrate, the Ti metal interlayer with the thickness of 150 nm was deposited for 5 minutes at the sputtering power of 150W. For all the samples, the constant sputtering target power and deposition time were 150W and 60 minutes, respectively. The typical thickness of all the coatings was in the range of ~ 1.8 – 2.0 µm as observed by in-situ Quartz-crystal digital thickness monitor.

2.2 Characterization of the coatings The surface topographies of the coatings were analyzed using atomic force microscope (AFM, park system XE-70). Crystallography of the coatings were characterized by X-ray diffraction (XRD, Rigaku- Miniflex II) with a Cu Kα radiation (λ = 1.5406 Å), step size of 0.02º. The qualitative analyses of chemical behavior of the coatings were characterized by Xray photoelectron spectroscopy (XPS-SHIMADZU-ESCA3400), using Mg radiation with the acceleration voltage of 0 ~12 kV. Low ion energy sputtering was carried out for all the coatings to remove the surface contamination before the XPS analysis. The nanomechanical properties of the coatings were evaluated by Hysitron Nanoindentation technique (Triboindenter TI950, USA) using Berkovich diamond indenter with the tip curvature of 150 nm. The maximum load of 6 mN was applied for all the nanoindentation measurements. The loading-unloading rate was 1.5 mN/min. and the pause between loading-unloading was given for 10 seconds. Oliver-Pharr method was used to calculate the hardness and elastic modulus of the coatings [26]. The tribological behavior of the coatings and substrates were analyzed by linear reciprocating mode ball-on-disc tribometer (CSM Instruments, Switzerland). The 100Cr6 steel ball with the diameter of 6 mm was used as the sliding counter body. Load and 4

sliding speed were constant 1 N and 3 cm/s, respectively, for all the samples. These experiments were carried out in atmospheric condition at room temperature. The wear profile of tribological results was evaluated using Dektak 6M-stylus profilometer. The wear track morphology was observed by the optical microscope. Cross-section of the TiN coatings deposited on Si and steel substrates were examined using Field Emission Scanning Electron Microscope (FE-SEM, Carl Zeiss, SUPRA 55) at the operating voltage of 20 kV. For this, the coated steel substrate was processed by wire cutting method, and the cross-sectional area was polished.

3. Results and discussion 3.1 Atomic Force Microscope analysis The surface topographies of TMN coatings deposited on Si and steel substrates were analyzed by non-contact mode AFM. These are shown in Fig. 1 and 2, respectively. Three dimensional (3D) AFM topography show the uniform distribution of surface. The spherical columnar grains were found in all the coatings except WN deposited on both the substrates. The semi-crystalline (Fig. 1(c)) and amorphous (Fig. 2(c)) like surface features are observed in WN coating owing to the lower nucleation potential [27]. Change in topography is related to mixed phase formation as ascribed from the XPS and XRD results. These results will be discussed in later section. Nucleation rate drives to vary the grain size 70-80, 45-70, 20-30 and 30-50 nm for the ZrN, CrN, WN and TiN coatings deposited on Si substrates. In the case of coated steel substrates, the grain size generally reduces and exists in the range of 40-60, 30-80, 10-15 and 25-40 nm for the same series of coatings as mentioned above. It is reported that the crystallinity of the coating enhances, when the coatings deposited on single crystal Si than polycrystalline steel substrate [28]. Substrate dependent AFM analysis and its crystalline quality will be discussed in the next section. For the tribological point of view, it is indispensable to evaluate the surface root mean square (rms) roughness of the coatings. In 5

general, the surface roughness is determined by the height differences of the surface features in the particular area. The average rms of the ZrN, CrN, WN and TiN coatings deposited on Si substrates are 7.32, 5.19, 4.21 and 3.76 nm, respectively. However, these values decreased to 6.79, 5.04, 3.49 and 3.23 nm for the similar coatings deposited on steel substrates, and this is possibly related to the grain size effect.

3.2 X-ray Diffraction analysis The crystallography of TMN coatings deposited on both substrates including with uncoated Si and steel substrates are shown in Fig. 3. The results revealed that the coatings were exhibited in polycrystalline nature on both substrates. The diffraction pattern of ZrN, CrN and TiN coatings possess (111), (200), (220) and (222) orientation. In addition, the (311) plane was found in ZrN and TiN coatings, as shown in Fig. 3(a) and (b). The diffraction patterns are in agreement with the standard JCPDS data for the cubic phases of ZrN (File No: 35-0753), CrN (File No: 76-2494) and TiN (File No: 38-1420). The dominant orientations on both the substrates were observed along (200) planes for ZrN and CrN coatings; whereas, the (111) plane was preferred for TiN coatings (Fig. 3). The dominant orientation of (200) plane relies on the lowest energy state which was derived from the critical interaction between strain and surface energies [29]. In general, the mixed orientations of nuclei have been formed during the initial film growth even at lower Ts. Further increasing the film growth, the (200) oriented crystallites become dominant and it occupied more surface area due to the threshold thermal energy which promotes (200) orientations in ZrN and CrN coatings at 250 ºC. It has been reported that the crystalline ZrN and CrN phases were formed even at room temperature [30,31]. In the case of TiN coatings, the domination along (111) plane might be the possibility of obtaining minimum surface energy than that of (200) planes at lower Ts [32]. It is also well known that, the Ts play an inevitable role in determining the preferred orientations. In contrast, the mixed phases of α-W, β-W2N (JCPDS file No.: 25-1257) and δ6

WN (File No.: 25-1256) were exhibited in the WN coatings deposited on both substrates, as shown in Fig. 3(a) and (b). Amongst them, the (100) orientation of δ-WN peak is dominant, followed by the secondary characteristic of β-W2N (111) peak on both substrates. In addition, the metallic α-W crystallites were observed along (110) plane for the WN coatings deposited on Si substrates (Fig. 3(a)). The XRD results signify that the sub-stoichiometric WN coatings were formed at the N2 flow rate of 8 sccm. Insufficient activation energy of the ad-atoms to WN crystallization at 250 ºC resulting in the broader peaks [27]. Substrate peaks were also found in the XRD patterns for the coatings deposited on steel substrates (Fig. 3(b)). Lower intensity and broader XRD peaks for the coatings deposited on steel substrates were obtained as compared to the coated Si substrates. Hence, the variations in the calculated lattice parameters, strain and crystallite sizes for both substrates were attained, which are furnished in Table 1. The lattice parameters (a) of the coatings deposited on Si substrate reveals slightly higher than that of standard values. However, it was lower for the coatings deposited on steel substrates at similar deposition conditions. Hence, these results suggested that the presence of tensile stress in the coatings deposited on Si substrates; whereas, compressive stress were found in the coatings deposited on steel substrates. The calculated micro strain ( ) also reveals the similar results, and it clearly indicates substrate dependent crystallographic changes. Due to the presence of compressive stress, the crystallite size also decreased for TMN coated steel substrates (calculated using Debye-Scherer formula). Furthermore, the calculated crystallite sizes are significantly smaller than the grain size observed from the AFM results, which are caused by the coalescence of smaller adjacent crystals to form the larger grains [33].

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3.3 X-ray photoelectron spectroscopy analysis Core level XPS spectra were recorded after sputtering the coating surface to ascertain the chemical and phase compositions of TMN coatings (Fig. 4(a-d)). In all the coatings, well defined metal-nitride phases are observed. These are in good agreement with the XRD results (Fig. 3). Moreover, oxygen peak was also observed in all the samples. It is due to the strong affinities of the transition metal atoms with oxygen during the sputtering process leads to the formation of oxide and oxy-nitride phases. Figure 4(a) and (d) represents deconvoluted XPS spectra of Zr 3d and Ti 2p metal species with the reactive components of N 1s and O 1s peaks in the respective ZrN and TiN coatings. These coatings were deposited at the nitrogen flow rate of 5 sccm. Similarly, the deconvoluted spectra of CrN and WN coatings deposited at 8 sccm reveal metal species of Cr 2p and W 7f with the reactive species of N 1s and O 1s peaks. The results are shown in Fig. 4(b) and (c). These spectra clearly show Zr 3d, Cr 2p, W 4f and Ti 2p bands which exhibits doublet peaks due to spin-orbital coupling [34]. Figure 4(a) disclose the two dominant peaks in the Zr 3d spectra which are constituted by Zr 3d5/2 and Zr 3d3/2 components corresponding to 179.5 and 181.8 eV, which are in agreement with the ZrN phases [35]. In addition to this, weak binding energy is appeared at 183.9 eV and this could be arises due to the formation of ZrO2 or Zr-O-N phase. Formations of these phases relate to the large difference in electronegativity of Zr atoms with atmospheric oxygen and nitrogen used for reactive gas for deposition [30]. In the N 1s spectrum, the binding energies of 397.1 and 398.4 eV resembling to the formation of Zr-N and Zr-O-N phases, respectively. The XPS qualitative analysis of CrN coatings related to Cr 2p, N 1s and O 1s bands are presented in Fig. 4(b). The strong doublet of Cr 2p3/2 and Cr 2p1/2 peaks are located at 574.5 and 583.7 eV, revealing the formation of Cr-N phase [36,37]. The peak positions centered at 397.7 and 400 eV indicate formation of CrN and Cr-O-N phases, respectively. The Cr 2p spectra visualize that, there is no major evidence to development of pure metal-oxide phases in the CrN

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coating. In Fig. 4(b), the intensity counts of O 1s spectra are less compared to the other sample, which denotes the low affinity of oxygen atoms into the CrN lattice vacancy sites [35]. On the other hand, deconvoluted XPS peaks of W 4f, N 1s and O 1s are observed in WN coating (Fig. 4(c)). XPS results are in agreement with XRD indicating formation of mixed phases (Fig. 3). From the results, it can be identified that the W 4f7/2 and W 4f5/2 doublet peak positions at 31.8 and 33.8 eV are associated with W-N and W2N phases, respectively [20,38,39]. It has been reported that, the deconvoluted N 1s binding energies centered at 396.7 and 397.7 eV [40] are contributing to W-N and W2N species, respectively, which are typical metal-nitride phases. Moreover, higher binding energy state at 37.7 eV of W 4f5/2 peak reveals the WO3 phases [20]. Intensity of the O 1s band is predominantly higher than that of N 1s peak (Fig. 4(c)), which describes higher chemical affinity of W to oxygen atoms. Figure 4(d) shows the high resolution XPS spectra of deconvoluted Ti 2p, N 1s and O 1s peaks for the TiN coating. In the Ti 2p spectra, the spin orbit doublet Ti 2p3/2 and Ti 2p1/2 peaks are centered at the binding energies of 454.9 and 460.8 eV which is attributed to the TiN phases [41,42]. Additionally, the shakeup peaks of Ti 2p3/2 and Ti 2p1/2 at 457.1 and 462.8 eV are corresponding to the Ti-O-N and TiO2 phases. Such chemical interaction arises due to high chemical affinity of oxygen atoms into the TiN interstitial sites. The deconvoluted N 1s can be fitted with two Gaussian curves with the binding energies of 397.1 and 399.4 eV associated with Ti-N and Ti-O-N phases, respectively (Fig. 4(d)). The O 1s spectra with the binding energies at 530.9 and 533 eV represent the respective TiO2 and Ti-O-N phases [41]. However, no obvious oxide or oxy-nitride phases were observed from the XRD results. This depicts that the presence of oxide phase could be more significant on the surface as compared to the bulk of the coating. Since, the XPS is surface sensitive to chemical aspect and thus, the oxide peaks were appearing in all the coatings. 9

3.4 Nanomechanical properties Hardness (H) and elastic modulus (E) of coatings were measured by nanoindentation technique. To avoid the influence of substrate during nanoindentation measurements, less than 10% of coatings thickness was intended by applying the maximum load of 6 mN. Totally, 10 different places were intended in each sample to obtain the average H and E values. Figure 5(a) and (b) shows the load-displacement curves for the coatings deposited on Si and steel substrates, respectively. The amount of work done by the coatings can be determined from the area beneath the load-displacement curves. The ratio between plastic work after peak loading and the elastic wok after unloading describes the elastic recovery (%ER) of the materials during nanoindentation measurements [24,43]. A larger value of %ER represents the coating has higher stiffness, which indicating the high hardness. In the present work, the lower value of %ER was obtained for the CrN coating around 55%, and it was increased as in the order of 60%, 65% and 89% for the respective ZrN, TiN and WN coatings deposited on Si substrates. Similar trend were found in the coatings deposited on steel substrates. Hence, the H values were obtained as 26.11 ±2, 19.59 ±2, 40.18 ±2 and 30.61 ±2 GPa corresponds to the ZrN, CrN, WN and TiN coatings (Fig. 6). Similarly, the change in E values for the ZrN, CrN, WN and TiN coatings is 368.3 ±25, 236 ±20, 440.8 ±25 and 399 ±25 GPa, respectively. In the case of coatings deposited on steel substrates, slightly lower H (~ 1-2 GPa) and E (~ 30-40 GPa) values were obtained, as shown in Fig. 7. These H and E values were significantly higher than that of bulk hardness and modulus of the individual materials. Another important factor to determine the mechanical properties of the coating is resistance to plastic deformation (H3/E2). Figure 6(a) and (b) shows the H3/E2 values along with the H and E values corresponding to the coatings deposited on both the substrates. It was found that the higher H3/E2 values of 0.18 ±0.03 and 0.32 ± 0.02 GPa corresponds to the TiN

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and WN coatings deposited on Si substrate. These values were increased to 0.2 ±0.03 and 0.36 ±0.03 GPa, for the TiN and WN coatings deposited on steel substrates. Higher %ER of these coatings could be the reason to resist the plastic deformation and hence, the higher H3/E2 values were attained. Minimum H3/E2 values of ~ 0.13 GPa were found for the ZrN and CrN coatings which is directly related to the lower %ER (Fig. 5). In general, changes in H, E and H3/E2 values between the different TMN coatings could be explained by microstructure and chemical environment. It is shown that the coatings with (111) preferred orientation possess higher hardness. This is due to the fact that (111) orientation has a zero Schmid factor for all slip the planes [44]. In the present work, slightly larger grain size and (200) oriented texture facilitates higher plastic deformation in the ZrN and CrN coatings. This might be decreasing factor of H and E value. The TiN coatings on both substrates were exhibited improved mechanical properties than earlier reported value [16,24,44]. This is due to the formation of (111) texture (Fig. 2). Moreover, AFM results revealed the lower grain size (< 50 nm) in TiN coating on both the substrates. On the other hand, the enhanced hardness of WN coatings was obtained. This could be related to smaller grain/crystallite sizes (< 30 nm) as derived from AFM and XRD results. The nanoscale grains were effectively inhibited the crack propagation during indentation measurements owing to the existence of possible surface cracks and point defects at the grain boundaries, resulting in the higher hardness [45,46]. According to the Hall-Petch relation, the dislocation density is higher in the nanoscale grains (20-50 nm) at the grain boundaries, and thereby, it would enhance the mechanical properties of WN and TiN coatings. Besides, the strong chemical bonding between the W and Ti atoms with nitrogen is another reason to enhance the nanomechanical properties of WN and TiN coatings.

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3.5 Tribological properties Tribological behavior of TMN coatings deposited on Si and steel substrates was measured by ball-on-disc micro-tribometer. Investigating the substrate effect and tribological behaviors of uncoated Si and steel are also measured.

3.5.1 Friction behavior Figure 7(a) and (b) illustrates the friction coefficient of uncoated Si and steel substrates sliding against steel ball. The average friction coefficient value of 0.35 was observed for the Si whereas; it was higher 0.58 on the steel substrate. However, the standard deviation of this value is higher in Si due to the periodic instability in friction curve. The average values are found to be 0.62, 0.45, 0.3 and 0.15 for the ZrN, CrN, WN and TiN coatings deposited on Si substrates (Fig. 8). Average friction coefficient and standard deviation of these values are given in the inset of Fig. 8, which is higher in CrN and TiN coatings. Surprisingly, the drastic variation in friction coefficient was observed in the similar coatings deposited on steel substrates (Fig. 9). Average values of friction coefficient were accomplished as 0.57, 0.4 and 0.21 for the ZrN, CrN and WN coatings, respectively. Figure 9(d) depicts ultra-low friction coefficient 0.06 for the TiN coatings deposited on steel substrate. Such a low friction coefficient is attributed to (111) orientation of this coating [24]. However, standard deviation of TiN including with CrN coating is high and it follows the trend for Si substrate also. In dry sliding conditions, formation of surface oxides could also be the reason for obtaining the lowest friction coefficient in TiN coatings. However, the sharp discontinuity appears in the TiN coated Si substrate after 300 m (Fig. 8). This might be due to the periodic formation and delamination of thin lubricant oxide layers after the coating deformation [47,48]. The trend of friction curves of uncoated steel substrate (Fig. 7(b)) and ZrN coatings deposited on both substrates (Fig. 8(a) and 9 (a)) were quite similar. This could be due to the material transfer of steel ball onto the ZrN coating and hence, sliding occurs between steel-on-steel surfaces 12

[49,50]. Relatively, all the coatings deposited on steel substrate indicate lower and stable friction coefficient than that of the coatings deposited on Si substrate as illustrated in Fig. 10. Moreover, ZrN and CrN coatings illustrate the marginal difference in friction coefficient whereas, it was significant for the WN and TiN coatings. These results are well evident that the substantial variations in the frictional behavior of coatings on both substrates are not only directly related to the microstructure, but also the mechanical properties and microstructure of substrate.

3.5.2 Wear behavior Wear profile of the TMN coatings deposited on both the substrates is shown in Fig. 11 and 12. This shows 2D profile (width and depth) of wear mark which is directly subjected to loss of materials [48,51]. Generally wear loss is proportional to the friction coefficient. This proportionality breaks down when other physical-chemical factors dominates [52]. It is shown that the friction coefficient of TiN coatings deposited on both the substrate is less, but wear depth is higher than other coatings. Such behavior is unable to explain by hardness and elastic modulus of coatings, which is not in consistent with wear loss. There are other factors like chemical affinity of different TMN coatings with steel ball that could be the reason for high wear loss but less friction coefficient [53,54]. Understanding the wear mechanism, the wear scar morphology of the uncoated/bare substrates as well as the wear track morphology of coatings with corresponding ball scars were obtained by optical microscope. These results are shown in Fig. 13, 14 and 15. Figure 11 indicates the narrow wear width and higher wear depth for the various nitrides coated Si substrates as compared with the coated steel substrates (Fig. 12). Wear loss from the ZrN coatings deposited on both substrates is not observed, and this is confirmed by wear profile, which shows spikes in positive side only (Fig. 11(a) and 12(a)). The micrographs of uncoated steel (Fig. 13(b)) as well as the ZrN coatings on both substrates are well in agreement in terms of similar morphology of transferring material from 13

steel ball on the coating surface. Under dry sliding conditions, the oxidized ball material was effectively transferred on to the coatings due to the extreme adhesive wear mode of steel ball against ZrN coating. This confirmed that the sliding occurs between transferred oxidized steel and steel materials. The transferred materials are clearly shown in the wear morphology (Fig. 14(a) and 15(a)). Moreover, the contact stress is higher at the coating/ball interface, resulting wider wear track [47,48]. In the case of CrN, WN and TiN coatings, the abrasive wear mode was observed. This is illustrated in the wear profile and morphological analysis. Severely deformed wear profiles were obtained for the CrN coatings on both the substrates (Fig. 11(b) and 12(b)) indicating trapped wear debris/particles at the coating/ball interface during sliding. Wear track was contaminated by debris and delaminated particles as evident by wear track morphology, as shown in Fig. 14(b) and 15(b). This could be the reason for obtaining the wavered friction curves for CrN coatings. In contrast, polished wear profile and morphology were observed for the WN and TiN coatings on both the substrates; and these results are in agreement with fiction curves. Figures 14 and 15 shows polished wear surfaces, including micro-grooves and fine abrasive particles from ball and coatings. Local failure was observed on the TiN coating deposited on Si substrate, which is possibly due to the brittle fracture of Si substrate exhibited by dynamic contact stress. Moreover, the wear morphology of ball scars are the strongest evidence for obtaining wider deformation of ball materials when slid against the TMN coatings deposited on steel (Fig. 15(a-d)) than Si substrates (Fig. 14(a-d)). Therefore, the higher wear resistance was exhibited in the TMN coatings deposited on steel substrates.

3.5.3 Mechanism of substrate dependent tribological properties of the coatings. Based on the observations, tribological behavior of TMN coatings deposited on steel were superior to the coatings deposited on Si substrates. This could be explained by the two important factors; (i) coating microstructure and (ii) mechanical properties of substrates. The 14

nucleation and growth of grains in the various nitride coatings on Si substrates were predominant than the polycrystalline steel substrates. Moreover, tensile stress was obtained in the coated Si samples, thereby, it exhibits lower resistance to plastic deformation; resulting lower wear resistance. Besides, it is well known that the single crystalline Si is a highly brittle material. During stressed sliding contact, the plastic deformation can easily propagate into brittle materials along the crystalline orientation. Figure 16 obviously illustrated that the several brittle fractures and micro-pits were found in the uncoated Si substrate after it slid against steel ball. Several micrometers in fractured length were obtained across the sliding direction since, the tribological measurement was carried out against the substrate orientation (100) (Fig. 16). In the case of coated Si substrates, shear cracks were nucleated and developed from the substrate. During continuous sliding motions, this crack propagates into the coating due to the formation of plastic deformation around the wear track. Thus the hard wear particles of coating materials were detached (pits/debris) due to the higher shear cracking [3,55]. At the next stage, these hard wear particles were entrapped at the two body contact interface and provided third body abrasive wear mechanisms. These three-body wear produces deep scratches in the wear surfaces, as shown in the wear profile and morphology (Fig. 11 and 12) resulting higher friction. These substrate dependent tribological mechanisms of coated Si and steel substrates could be clearly explained by the cross-sectional FE-SEM images as shown in Fig. 17. Figure 17(a) illustrates scars, wear particles and sharp interface between the coatings and Si substrate. Moreover, the noticeable brittle fracture of Si substrate was obtained due to the continuous frictional sliding motions. The brittle fracture influencing nucleation and crack propagation into the coating materials eventually generating three-body wear mode. In this process, detached particles were deposited on the wear track as illustrated in Fig. 17(a) (ii). Higher friction and wear for the coatings deposited on Si substrates were initially enhanced by the tensile stress; and later on, three bodies wear mode involved by the

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brittle fracture of Si substrates. The wear depth of TiN coating was ~ 4.5 µm (Fig. 11(d)), which is more than two times higher than that of coating thickness. Typical thicknesses of all TMN coatings were around ~ 2 µm. The wear depths in all the other coatings on both substrates are quite less than the coating thickness. Unlike Si substrates, the steel substrates exhibit in polycrystalline as well as ductile in nature. It is well noticed that the polished surface was observed without the presence of major scars (Fig. 17(b)). Moreover, on the near-edge of the cross sectional surface, the coating/substrate interface could not be observed. It is resulting from the mixing of the coating and substrate materials because of the plastic deformation of steel substrate, as shown in Fig. 17(b). This could be anticipated from the highest contact stress and contact temperature at the center of wear track owing to the continuous sliding motions. Due to the ductility of steel substrate, the crack energy dissipation and contact stress relaxation took place towards the substrate, showing lower friction and higher wear resistance of coated steel as compared to the coated Si substrates. On the other hand, higher compressive stress and higher resistance to plastic deformation would be the factor for enhancing wear resistance [56]. Since, the presence of compressive stress in TMN coatings deposited on steel substrates is highly influenced in the inhibition of crack propagation during sliding motion [45,57].

4. Conclusion TMN coatings were deposited by reactive DC magnetron sputtering on Si and steel substrates at 250 ºC. Substrate dependent difference in microstructure and tribo-mechanical properties of TMN coatings was observed. XRD results showed predominant (111) orientation of TiN coatings whereas, (200) orientation was characteristic for ZrN and CrN coatings. Due to the insufficient activation energy, mixed phases were obtained with the orientation of δ-WN (100) plane. These results were well in agreement with the chemical analysis confirmed by XPS. AFM and XRD results signify less particle sizes of WN coating, resulting superior 16

hardness compare to ZrN and CrN coatings. Coatings deposited on steel substrates possess higher resistance to plastic deformation owing to the presence of compressive stress and finer microstructure. The abrasive wear mode is the most prominent mechanism involved in the coatings on both substrates except ZrN coatings. The steel-on-steel sliding appears because of the severe adhesive wear mode of steel ball against ZrN coatings, resulting in the similar friction behavior as steel ball sliding against uncoated steel sample. From the tribological results, it can be perceived that the wear mechanisms of the coatings are dependent on the nature of the substrate. Overall friction and wear resistance behavior of the coatings deposited on steel are superior to Si substrates. Friction coefficient in TiN coating deposited on steel is ~0.06. Such improved characteristics are originated from the finer microstructure, compressive stress, higher resistance to plastic deformation and ductile nature of steel substrate. On the contrary, third-body wear induces the higher friction and wear rate on the coated Si samples, which is typically originated from the brittle fracture of Si substrates under the sliding contact.

Acknowledgement One of the authors D. Dinesh Kumar gratefully acknowledges the Council for Scientific and Industrial Research (CSIR), Government of India for awarding Senior Research Fellowship (SRF), sanction No.: 09/468(460)/2012-EMR-I. Authors would like to thank the CSIR, Government of India for providing financial support to carry out this research work (CSIR project sanction order no. 03(1266)/12/EMR-II). Dr. M. Arivanandhan, Research Institute of Electronics, Shizuoka University, Japan is gratefully acknowledged for XPS characterizations. Ashok Bahuguna, MSG, IGCAR is acknowledged for wear track dimension measurement. Authors would like sincerely thank to Dr. Srinivasa Rao Bakshi and

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Mr. Karthi Selva, IIT Madras for Nanoindentation measurements. Authors are thankful to Dr. A.K. Tyagi and Mr. M.P. Janawadkar, MSG, IGACR for their support. References [1] A. Caron, D.V. Louzguine-Luzguin, R. Bennewitz, Structure vs chemistry: friction and wear of Pt-based metallic surfaces, ACS Appl. Mater. Interfaces 5 (2013) 1341−11347. [2] E.C. Samano, A. Clemente, J.A. Díaz, G. Soto, Mechanical properties optimization of tungsten nitride thin films grown by reactive sputtering and laser ablation, Vacuum 85 (2010) 69–77. [3] L. Vandoni, A.G Demir, B. Previtali, N. Lecis, D. Ugues, Wear behavior of fiber laser textured TiN coatings in a heavy loaded sliding regime, Materials 5 (2012) 2360–2382. [4] A.J. Nathanael, R. Yuvakkumar, S.I. Hong, T.H. Oh, Novel zirconium nitride and hydroxyapatite nanocomposite coating: detailed analysis and functional properties, ACS Appl. Mater. Interfaces 6 (2014) 9850−9857. [5] H.A. Ching, D. Choudhury, M. Nine, N.A. Abu Osman, Effects of surface coating on reducing friction and wear of orthopaedic implants, Sci. Technol. Adv. Mater. 15 (2014) 14402–14423. [6] I. Tzanakis, M. Hadfield, B. Thomas, S.M. Noya, I. Henshaw, S. Austen, Future perspectives on sustainable tribology, Renew. Sustain. Energy Rev. 16 (2012) 4126– 4140. [7] J. Zheng, J. Hao, X. Liu, Q. Gong, W. Liu, The plasma nitriding treatment of TiN/TiCN multilayer films, Appl. Surf. Sci. 268 (2013) 195– 203.

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nonequilibrium states and emergent behaviour, Sci. Technol. Adv. Mater. 13 (2012) 43001–43026. [55] A Matthews, S Franklin, K. Holmberg, Tribological coatings: contact mechanisms and selection, J. Phys. D: Appl. Phys. 40 (2007) 5463–5475. [56] L. Meshi, S. Samuha, S.R. Cohen, A. Laikhtman, A. Moshkovich, V. Perfilyev, I. Lapsker, L. Rapoport, Dislocation structure and hardness of surface layers under friction of copper in different lubricant conditions, Acta Mater. 59 (2011) 342–348. [57] K. Holmberg, A. Matthews, H. Ronkainen, Coatings tribology-contact mechanisms and surface design, Tribol. Int. 3(1–3) (1998) 107–120. List of figures caption: Figure 1. AFM topographies of (a) ZrN, (b) CrN, (c) WN and (d) TiN coatings deposited on Si substrate. Figure 2. AFM topographies of (a) ZrN, (b) CrN, (c) WN and (d) TiN coatings deposited on steel substrate. Figure 3. XRD pattern of CrN, TiN, ZrN and WN coatings deposited on (a) Si and (b) steel substrates. Figure 4. XPS spectra of (a) ZrN, (b) CrN, (c) WN and (d) TiN coatings deposited on Si substrate. Figure 4. XPS spectra of (a) ZrN, (b) CrN, (c) WN and (d) TiN coatings deposited on Si substrate. Figure 5. Load-displacement curves of CrN, ZrN, TiN and WN coatings deposited on (a) Si and (b) steel substrates. Figure 6. Hardness (H), elastic modulus (E) and H3/E2 value of ZrN, CrN, WN and TiN coatings deposited on (a) Si and (b) steel substrates. 25

Figure 7. Friction coefficient of uncoated (a) Si and (b) steel substrates sliding against steel ball. Figure 8. Friction coefficient of (a) ZrN (b) CrN (c) WN and (d) TiN caoatings deposited on Si substrates sliding against steel ball. Figure 9. Friction coefficient of (a) ZrN (b) CrN (c) WN and (d) TiN caoatings deposited on steel substrates sliding against steel ball. Figure 10. Comparision of mean friction coefficient of the TMN caoatings deposited on Si and steel substrates sliding against steel ball. Figure 11. Wear depth profile of (a) ZrN (b) CrN (c) WN and (d) TiN caoatings deposited on Si substrates. Figure 12. Wear depth profile of (a) ZrN (b) CrN (c) WN and (d) TiN caoatings deposited on steel substrates. Figure 13. Wear micrographs of uncoated substrates sliding against steel ball (a) uncoated Si (Scar on Si at (i) 5x & (ii) 50x and, scar on SS ball (iii) 5x & (iv) 50x magnifications) and (b) uncoated steel (Scar on steel substarte at (i) 50x and, scar on steel ball (ii) 5x & (iii) 50x magnifications). Figure 14. Wear micrographs of (a) ZrN (b) CrN (c) WN and (d) TiN coatings deposited on Si substrates sliding against steel ball. (Scar on coatings at (i) 5x & (ii) 50x and, scar on steel ball (iii) 5x & (iv) 50x magnifications). Figure 15. Wear micrographs of (a) ZrN (b) CrN (c) WN and (d) TiN coatings deposited on steel substrates sliding against steel ball. (Scar on coatings at (i) 5x & (ii) 50x and, scar on steel ball (iii) 5x & (iv) 50x magnifications). Figure 16. Optical micrographs of fractured Si substrates after sliding against steel ball (a) 50x and (b) 100x magnifications.

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Figure 17. FE-SEM cross-sectional wear micrographs of TiN coating deposited on (a) Si [(i) 500x and (ii) 15000x magnifications] and (b) steel substrates [(i) 500x and (ii) & (iii) 5000x magnifications]. List of figures

Figure 1. AFM topographies of (a) ZrN, (b) CrN, (c) WN and (d) TiN coatings deposited on Si substrate.

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Figure 2. AFM topographies of (a) ZrN, (b) CrN, (c) WN and (d) TiN coatings deposited on steel substrate.

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Figure 3. XRD pattern of uncoated substrates, CrN, TiN, ZrN and WN coatings deposited on (a) Si and (b) steel substrates.

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Figure 4. XPS spectra of (a) ZrN, (b) CrN, (c) WN and (d) TiN coatings deposited on Si substrate.

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Figure 5. Load-displacement curves of CrN, ZrN, TiN and WN coatings deposited on (a) Si and (b) steel substrates.

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Figure 6. Hardness (H), elastic modulus (E) and H3/E2 value of ZrN, CrN, WN and TiN coatings deposited on (a) Si and (b) steel substrates.

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Figure 7. Friction coefficient of uncoated (a) Si and (b) steel substrates sliding against steel ball.

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Figure 8. Friction coefficient of (a) ZrN (b) CrN (c) WN and (d) TiN caoatings deposited on Si substrates sliding against steel ball.

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Figure 9. Friction coefficient of (a) ZrN (b) CrN (c) WN and (d) TiN caoatings deposited on steel substrates sliding against steel ball.

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Figure 10. Comparision of mean friction coefficient of the TMN caoatings deposited on Si and steel substrates sliding against steel ball.

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Figure 11. Wear depth profile of (a) ZrN (b) CrN (c) WN and (d) TiN caoatings deposited on Si substrates.

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Figure 12. Wear depth profile of (a) ZrN (b) CrN (c) WN and (d) TiN caoatings deposited on steel substrates.

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Figure 13. Wear micrographs of uncoated substrates sliding against steel ball (a) uncoated Si (Scar on Si at (i) 5x & (ii) 50x and, scar on SS ball (iii) 5x & (iv) 50x magnifications) and (b) uncoated steel (Scar on steel substarte at (i) 50x and, scar on steel ball (ii) 5x & (iii) 50x magnifications).

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Figure 14. Wear micrographs of (a) ZrN (b) CrN (c) WN and (d) TiN coatings deposited on Si substrates sliding against steel ball. (Scar on coatings at (i) 5x & (ii) 50x and, scar on steel ball (iii) 5x & (iv) 50x magnifications).

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Figure 15. Wear micrographs of (a) ZrN (b) CrN (c) WN and (d) TiN coatings deposited on steel substrates sliding against steel ball. (Scar on coatings at (i) 5x & (ii) 50x and, scar on steel ball (iii) 5x & (iv) 50x magnifications).

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Figure 16. Optical micrographs of fractured Si substrates after sliding against steel ball (a) 50x and (b) 100x magnifications.

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Figure 17. FE-SEM cross-sectional wear micrographs of TiN coating deposited on (a) Si [(i) 500x and (ii) 15000x magnifications] and (b) steel substrates [(i) 500x and (ii) & (iii) 5000x magnifications].

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Table 1. Calculated lattice parameter (a) micro strain (ε) and crystallite size (d) for the ZrN, CrN, WN and TiN coatings on Si (100) and 316LN SS substrates

Lattice parameter, a

Micro strain, ε

Crystallite size (nm)

on silicon

on steel

on silicon

on steel

on silicon

on steel

ZrN

4.64

4.452

0.0137

0.027

18.84 ± 2

14.7 ± 2

CrN

4.152

4.02

0.028

0.03

16.64 ± 2

13.1 ± 2

WN

2.922

2.775

0.01

0.041

8.09 ± 2

6.4 ± 2

TiN

4.265

4.121

0.005

0.03

14.92 ± 2

12.8 ± 2

Coating

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