- Email: [email protected]
DLC coatings
Surface and Coatings Technology 162 (2002) 42–48
Magnetron sputtering of nanocomposite (Ti,Cr)CNyDLC coatings Sam Zhanga,*, Yongqing Fua, Hejun Dua, X.T. Zengb, Y.C. Liub a
School of Mechanical and Production Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798, Singapore b Singapore Institute of Manufacturing Technology, 71 Nanyang Drive, Singapore 638075, Singapore Received 12 February 2002; accepted in revised form 19 August 2002
Abstract Superhard nanocrystalline (Ti, Cr)CNyDLC coatings were prepared through co-sputtering of Ti, Cr and graphite targets in an argonynitrogen atmosphere. Results from both transmission electron microscopy (TEM) and grazing incident X-ray diffraction (GIXRD) indicated that the grain size of (TiCr)Cx Ny crystals was approximately 10–20 nm. X-ray photoelectron spectroscopic studies confirmed that an increase in the sputtering power at the Ti target not only increased the Ti composition in the film but also brought about an increase in sp3 bonding in DLC matrix, in agreement with the raising hardness with Ti sputtering power. Film hardness and elastic modulus were measured with a nano-indenter, and film hardness reached 40 GPa. Tribological behaviors of the films were evaluated using a ball-on-disk tribometer, and the films demonstrated properties of low-friction and good wear resistance. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Nanocomposite coating; Diamond-like carbon; (Ti,Cr)N; Sputtering; Tribology; Wear; Coefficient of friction
1. Introduction Nanostructured coatings have recently attracted increasing interest because of the possibilities of synthesizing materials with unique physical–chemical properties, e.g. superior hardness, toughness, chemical stability, low friction and wear-resistance w1–4x. Highly sophisticated surface related properties, such as superplasticity, optical, magnetic, electronic and catalytic properties can be obtained by advanced nanostructured coatings, making them attractive for industrial applications in highspeed machining, tooling, biomedical, automotive, optical applications and magnetic storage devices w5– 10x. There are many types of design models for nanostructured coatings, such as nanocomposite coatings, nano-scale multilayer coating, superlattice coating, nanoscale-graded coatings, etc. Among them, superhard nanocomposite coatings attracted more interest due to the endless possibilities of the synthesizing materials of unique properties w11–21x. Designing of nanocomposite *Corresponding author. Tel.: q65-790-4400; fax: q65-791-1859. E-mail address: [email protected] (S. Zhang).
coating needs consideration of many factors, for examples, the interface volume, grain size, single layer thickness, surface and interfacial energy, texture, epitaxial stress and strain, etc., all of which depend significantly on the materials selection, deposition methods and parameters w22–27x. Grain boundary hardening is one of the possibilities to control the microstructure in order to increase coating hardness. With the decrease in crystal size, the hardness of materials increases according to the ‘Hall–Petch’ relationship, especially for crystal size down to tens of nanometer. However, a new deformation mechanism, called grain boundary sliding, replaces the dislocation activity that dominates deformation process in conventional materials. Softening caused by the grain boundary sliding is mainly attributed to the large amount of defects in the grain boundaries, which allows fast diffusion of atoms and vacancies with the applying of stress. A further increase of the strength and hardness with decreasing crystallite size can be achieved only if grain boundary sliding is blocked by appropriate coating design and materials selection.
0257-8972/02/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 2 . 0 0 5 6 1 - 3
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
43
Fig. 1. AFM morphology of the film deposited at Ti target power density of 4.5 Wycm2 and nitrogen gas flow ratio of (a) 33% (b) 100%.
There are many design ideas for nanocomposite coatings, and embedding nanocrystalline phases in amorphous phase matrix is quite often applied which can be fulfilled easily by physical vapor deposition (PVD) and chemical vapor deposition (CVD) methods. Diamondlike carbon (DLC), amorphous carbon nitride and other hard amorphous materials (with high hardness and elastic modulus) have been recognized as the primary candidates for the amorphous matrix while nano-sized refractory nitrides, such as TiN, Si3N4, AlN, BN, etc., could be used as strengthening phases w28x. For this coating design, the size and the distribution of these nanocrystals in the amorphous structures need to be optimized to obtain high hardness. These nanocrystalline grains should have random orientation (high angle grain
boundaries) to minimize grain incoherence strain, and high toughness could be achieved by strain release via nanocrystals sliding in the DLC matrix. Termination of nanocracks by deflection at grain boundaries and by energy loss within the amorphous matrix can dramatically improve the toughness. Carbon is one of the most important and widely used materials in many areas and diamond-like carbon (DLC) has been recognized as one of the primary candidates used for wear resistant and solid lubricating coating. TiN (or TiC) and CrN coatings prepared using sputtering have been applied in industry for wear protective applications. Following the above design idea, in this study, we combine these two types of coatings to synthesize a nanocomposite structure. In film design, DLC acts as
44
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
Fig. 2. GIXD analysis of the films deposited under 300 W Ti target power with AryN2s2:1.
hard, tough and lubricating matrix, while nano-particles act as reinforcing crystallites to improve hardness and mechanical properties. 2. Experimental Nanocomposite (TiCr)CxNy yDLC coatings were prepared through co-sputtering of graphite, Ti and Cr targets
in argonynitrogen atmosphere in a magnetron sputtering system. Silicon wafers and WC discs were used as the substrates. WC discs with a diameter of 50 mm and a thickness of 5 mm were mechanically ground and polished, then ultrasonically rinsed in acetone. The substrate holder rotated during the deposition for uniformity. In this study, nitrogen gas flow ratio wN2 y(N2q Ar)x was controlled at 33 and 100%, with Ti target power density of 4.5 and 6.5 Wycm2 (d.c.). The power density of chromium and graphite target were controlled at 4.5 Wycm2 (d.c.) and 6.5 Wycm2 (r.f.), respectively. Deposition was performed at a low temperature of 150 8C for 2 h. The substrate-to-target distance was 100 mm. Surface and cross-section morphologies of the coatings were investigated using an atomic force microscopy (AFM) and a JEOL scanning electron microscope. The crystalline structure was obtained by grazing incidence X-ray diffraction (GIXD) method. The nanocomposite coating was also deposited on potassium bromide (KBr) pellets for 20 min, and then the pellets were dissolved in water to float off the film for TEM study. The thin film thus obtained was examined using a JEOL 200 kV TEM. X-ray photoelectron spectroscopy (XPS) analysis was performed on film surface using a Kratos AXIS spectrometer with monochromatic Al Ka (1486.6 eV)
Fig. 3. TEM photos of deposited nanocomposite thin films.
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
45
Fig. 4. Load-displacement curve during load and unloading process for nano-indentation.
X-ray radiation. Film hardness and elastic recovery were evaluated using a nano-indentation tester with a penetrating depth of 100 nm. Load bearing capacity of the nanocomposite coatings on WC substrate was assessed using a scratch tester. A diamond stylus was driven across the coating at a constant speed of 1 mmys and a continuously increased load rate of 1.5 Nys. Tribological behaviors of nanocomposite coatings were evaluated using a ball-on-disk tribometer under dry sliding conditions at room temperature. Alumina balls (with a diameter of 9.5 mm and a surface roughness better than Ras0.05 mm) were used as the counterface materials. The normal loads were 10, 20 and 50 g. All the tests were run in laboratory air (25 8C and relatively humidity of 65"3%) with a sliding distance of 300 m and a sliding speed of 0.2 mys. Coefficient of friction was recorded during each test. Wear volumes were calculated from the wear tracks measured with a laser profilometer.
Fig. 5. XPS depth profile of the film deposited at Ti power density of (a) 4.5 Wycm2 and (b) 6.5 Wycm2.
density of 6.5 Wycm2 with nitrogen to argon gas flow ratio of 1:3. There are some broad peaks of (TiCr)CN crystalline phases. As a first degree approximation, the average crystallite size can be estimated by the Debye– Scherrer formula (shown in Eq. (1)) w30x:
3. Results and discussions Cross-section of the coating is dense and featureless with a thickness of approximately 1 mm. Fig. 1a shows AFM morphologies of the coating deposited at a power density of 4.5 Wycm2 on the Ti target with nitrogen to argon gas flow ratio of 1:3. The surface features a wavy morphology or hilly humps of approximately 30–40 nm in width and 6–8 nm in height. In the case of deposition under a pure nitrogen gas, the coating roughness increases significantly (see Fig. 1b). This is probably owing to the rapid formation and growth of nitrides, and the formation of polymeric C–N based phase w29x. Fig. 2 shows the GIXRD profile of the crystalline structures for the film deposited at Ti target power
Table 1 Film bonding structure distribution (%) of C 1s and N 1s peaks Bonding structure
C 1s 284.7 285.7 286.6 288.3
eV(C–C bonding) eV(sp2 bonding) eV(sp3 bonding) eV (C–O bonding)
N 1s 398.3 eV(N–C sp3 bonding) 399.1 eV(N–C sp2 bonding) 401 eV (N–O bonding)
%, at Ti power density of 4.5 Wycm2
at Ti power density of 6.5 Wycm2
45.53 36.46 7.57 10.44
57.17 16.04 15.99 10.79
42.30 40.43 17.27
61.64 25.44 12.92
46
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
Fig. 6. Typical XPS high-resolution carbon 1s and nitrogen 1s core level spectra.
Fig. 7. Scratch profile of the film deposited at Ti target power density of 6.5 Wycm2 and nitrogen flow ratio of 33%. The insets show SEM morphology of the scratch track at different loading stage.
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
47
(1)
DsKlybcosu
Fig. 8. Wear curves for the film deposited at Ti target power density of 6.5 Wycm2 and nitrogen gas flow ratio of 33% at increasing normal load: (a) 10 g, (b) 20 g, (c) 50 g.
Where K is a constant (Ks0.91), D is the mean crystallite dimension normal to diffracting planes, l is the X-ray wavelength (ls0.15406 nm for Cu target), b in rad is the peak width at half-maximum peak height and u is the Bragg angle. The calculated average grain size of (Ti,Cr)CN crystals is approximately 20 nm. For all the films in this study, the calculated grain size of (Ti,Cr)CN crystals are within 10–20 nm. High magnification bright-field TEM photo of the films is shown in Fig. 3a. In an amorphous DLC matrix, many tiny crystals can be observed with dimensions approximately 8–12 nm. These nanocrystalline crystals have different orientations and lattice spacing. It should be pointed out that these crystals form only after 20min deposition, while those in the coating deposited for 2 h may be larger than those shown in Fig. 3a due to grain growth, as is illustrated in the AFM morphology in Fig. 1. Fig. 3b shows the dark-field TEM photos, clearly revealing these nano-size crystals. Fig. 3c shows the diffraction ring patterns, indicating these nanocrystalline phases of TiCrCN. Fig. 4 is the nanoindentation load-displacement profile of the film deposited at Ti target power density of 6.5 Wycm2. The calculation of hardness and elastic modulus is performed using Oliver and Pharr method w31x. The measured film hardness is approximately 40 GPa. The Young’s modulus of the film is approximately 300 GPa. With the decrease of Ti power density to 4.5 W/cm2, the film hardness decreases to approximately 25 GPa. Fig. 5 gives the atomic composition of the films obtained from XPS analysis. XPS results revealed that with the increase of Ti target power density 4.5–6.5 Wy cm2, Ti in the film increases from 12.3 to 22 at.%, while the contents of other elements only change slightly. Therefore, one possible reason for the hardness increasing with Ti target power could come from the increase in the amount of nano-sized (TiCr)CN phases in the film. With the increase in the nitrogen flow ratio from 33 to 100%, film hardness decreases dramatically from 40 to 18 GPa. The high content of polymeric C– N based phase formed under high content of nitrogen gas could be detrimental for the film hardness. Fig. 6 shows typical high-resolution carbon 1s and nitrogen 1s core level spectra. The peaks were fitted using the binding energy interval values from the literature w32,33x. A Gaussian decomposition of C 1s spectra
Table 2 Average coefficient of friction for specimens prepared under different Ti contents and normal loads 10 N
2
4.5 Wycm 6.5 Wycm2
20 N
50 N
COF
Wear volume
COF
Wear volume
COF
Wear volume
0.03 0.02
4.06 not appreciable
0.02 0.03
0.61 not appreciable
0.01 0.03
2.02 not appreciable
48
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
gives rise to the following peaks: 284.7 eV (C–C bonding), 285.7 eV (sp2 bonding), 286.6 eV (sp3 bonding) and 288.3 eV (C–O bonding). Similar analysis on N 1s spectra yields 398.3 eV (N–C sp3 bonding), 399.1 eV (N–C sp2 bonding) and 401 eV (N–O bonding). The concentrations of the different bonding structures of C 1s and N 1s peaks are calculated and the results are given in Table 1. With the increase of Ti target power density, the concentration of sp3 bonding structure increased for both C 1s and N 1s peaks, and at the same time, sp2 decreased, which supports the significant increase in film hardness as discussed earlier. Scratch test was used to evaluate the dynamic load bearing capacity and friction properties of the coatingsubstrate system. The ‘lower critical load’ or the normal load at which the first damage or a sharp increase in friction coefficient is observed is widely used as a measure of the load bearing capacity. Fig. 7 is such a scratch profile which clearly shows the coating fails at a normal load of approximately 60 N, indicating a good adhesion and load bearing capacity of nanocomposite coatings on WC–Co substrate. The coefficient of friction remains at a very low value of 0.05, indicating good friction properties. SEM micrographs for the different stages during scratch testing are also shown as inset in Fig. 7. The critical load slightly decreased to 53 N at Ti target power density 4.5 Wycm2. With the increase of Ti target power, there is no significant change in coefficient of friction. The coefficient of friction remains at an extremely low value of 0.01–0.03 within the 300-m wear distance under different normal loads (Fig. 8). With the increase of normal load, the coefficient of friction decreases slightly, which can probably be explained by the easier graphitization of wear debris and generating of lubricating surface layer with the increase of normal load. Table 2 lists the average coefficient of friction and wear volume for specimens prepared under different Ti target powers and normal loads. The long-term coefficient of friction remains a low and stable value (less than 0.03). With the increase of Ti target power density from 4.5 to 6 Wycm2, the wear volume decreases so much that the surface profilometry does not reveal noticeable wear track. 4. Conclusions Superhard nanocrystalline (Ti, Cr)CNyDLC coatings were prepared through co-sputtering of Ti, Cr and graphite targets in an argonynitrogen atmosphere. Results from both TEM and GIXRD indicated that the grain size of the (TiCr)CxNy crystals were approximately 10–20 nm. XPS results confirmed that an increase in the sputtering power at the Ti target not only increased the Ti composition in the film but also brought about increase in sp3 bonding in the DLC matrix, in agreement
with the raising hardness with Ti sputtering power. Film hardness and elastic modulus were measured with a nano-indenter. Under the experiment conditions, the film hardness fell within 18–40 GPa. The tribological behaviors of the films were evaluated using a ball-on-disk tribometer, and the films demonstrated properties of low-friction and good wear resistance. References w1x P. Holubar, M. Jilek, M. Sima, Surf. Coat. Technol. 133–134 (2000) 145. w2x R.A. Andrievski, Mater. Trans. 42 (2001) 1471. w3x J. Musil, Surf. Coat. Technol. 125 (2000) 322. w4x S. Veprek, A.S. Argon, J Vac. Sci. Technol. B 20 (2002) 650. w5x S. Veprek, S. Reiprich, Thin Solid Films 268 (1995) 64. w6x V. Provenzano, R.L. Holtz, Mater. Sci. Eng. A 204 (1995) 125. w7x F. Mazaleyrat, L.K. Varga, J. Magn. Magnetic Mater. 215–216 (2000) 253. w8x B. Cantor, C.M. Allen, Scr. Mater. 44 (2001) 2055. w9x H. Zeng, M.L. Yan, Y. Liu, J. Appl. Phys. 89 (2001) 810. w10x J.L. Solis, S. Saukko, L. Kish, Thin Solid Films 391 (2001) 255. w11x S. Veprek, A.S. Argon, Surf. Coat. Technol. 146 (2001) 175. w12x S. Veprek, Thin Solid Films 317 (1998) 449. w13x J. Musil, F. Kunc, H. Zeman, Surf. Coat. Technol. 154 (2002) 304. w14x J. Musil, J. Vlcek, Surf. Coat. Technol. 142 (2001) 557. w15x M.A. Baker, S. Klose, C. Rebholz, Surf. Coat. Technol. 151 (2002) 338. w16x E. Ribeiro, A. Malczyk, S. Carvalho, Surf. Coat. Technol. 151 (2002) 515. w17x D.V. Shtansky, K. Kaneko, Y. Ikuhara, Surf. Coat. Technol. 148 (2001) 206. w18x L. Rebouta, C.J. Tavares, Surf. Coat. Technol. 133–134 (2000) 234. w19x S. Carvalho, L. Rebouta, Thin Solid Films 398–399 (2001) 391. w20x M. Misina, J. Musil, Surf. Coat. Technol. 110 (1998) 168. w21x J. Musil, H. Polakova, Surf. Coat. Technol. 127 (2000) 99. w22x J. Musil, H. Hruby, Thin Solid Films 36 (2000) 104. w23x J. Musil, P. Karvankova, J. Kasl, Surf. Coat. Technol. 139 (2001) 101. w24x A.A. Voevodin, J.P. O’Neill, J.S. Zabinski, Surf. Coat.Technol. 116–119 (1999) 36. w25x S. Veprek, A. Niederhofer, K. Moto, T. Bolom, H.D. Mannling, P. Nesladek, G. Dollinger, A. Bergmaier, Surf. Coat.Technol. 133–134 (2000) 152. w26x F. Vaz, L. Rebouta, P. Goudeau, Surf. Coat.Technol. 146 (2001) 274. w27x J. Patscheider, T. Zehnder, M. Diserens, Surf. Coat. Technol. 146 (2001) 201. w28x H. Holleck, in: A. Kumar, Y.W. Chung, J.J. Moore, J.E. Smugeresky (Eds.), Surface Engineering: Science and Technology I, The Minerals, Metals and Materials Society, 1999, pp. 207–231. w29x A. Stanishevsky, L. Khriachtchev, I. Akula, Diamond Relat. Mater. 7 (1998) 1190–1195. w30x M.V. Zdujic, O.B. Milosevic, L.C. Karanovic, Mater Lett. 13 (1992) 125–128. w31x T. Berlind, N. Hellgren, Surf. Coat. Technol. 141 (2001) 145. w32x E.G. Wang, Prog. Mater. Sci. 41 (1997) 241–298. w33x K. Sharma, J. Narayan, Int. Mater. Rev. 42 (1997) 137–192.
Magnetron sputtering of nanocomposite (Ti,Cr)CNyDLC coatings Sam Zhanga,*, Yongqing Fua, Hejun Dua, X.T. Zengb, Y.C. Liub a
School of Mechanical and Production Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798, Singapore b Singapore Institute of Manufacturing Technology, 71 Nanyang Drive, Singapore 638075, Singapore Received 12 February 2002; accepted in revised form 19 August 2002
Abstract Superhard nanocrystalline (Ti, Cr)CNyDLC coatings were prepared through co-sputtering of Ti, Cr and graphite targets in an argonynitrogen atmosphere. Results from both transmission electron microscopy (TEM) and grazing incident X-ray diffraction (GIXRD) indicated that the grain size of (TiCr)Cx Ny crystals was approximately 10–20 nm. X-ray photoelectron spectroscopic studies confirmed that an increase in the sputtering power at the Ti target not only increased the Ti composition in the film but also brought about an increase in sp3 bonding in DLC matrix, in agreement with the raising hardness with Ti sputtering power. Film hardness and elastic modulus were measured with a nano-indenter, and film hardness reached 40 GPa. Tribological behaviors of the films were evaluated using a ball-on-disk tribometer, and the films demonstrated properties of low-friction and good wear resistance. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Nanocomposite coating; Diamond-like carbon; (Ti,Cr)N; Sputtering; Tribology; Wear; Coefficient of friction
1. Introduction Nanostructured coatings have recently attracted increasing interest because of the possibilities of synthesizing materials with unique physical–chemical properties, e.g. superior hardness, toughness, chemical stability, low friction and wear-resistance w1–4x. Highly sophisticated surface related properties, such as superplasticity, optical, magnetic, electronic and catalytic properties can be obtained by advanced nanostructured coatings, making them attractive for industrial applications in highspeed machining, tooling, biomedical, automotive, optical applications and magnetic storage devices w5– 10x. There are many types of design models for nanostructured coatings, such as nanocomposite coatings, nano-scale multilayer coating, superlattice coating, nanoscale-graded coatings, etc. Among them, superhard nanocomposite coatings attracted more interest due to the endless possibilities of the synthesizing materials of unique properties w11–21x. Designing of nanocomposite *Corresponding author. Tel.: q65-790-4400; fax: q65-791-1859. E-mail address: [email protected] (S. Zhang).
coating needs consideration of many factors, for examples, the interface volume, grain size, single layer thickness, surface and interfacial energy, texture, epitaxial stress and strain, etc., all of which depend significantly on the materials selection, deposition methods and parameters w22–27x. Grain boundary hardening is one of the possibilities to control the microstructure in order to increase coating hardness. With the decrease in crystal size, the hardness of materials increases according to the ‘Hall–Petch’ relationship, especially for crystal size down to tens of nanometer. However, a new deformation mechanism, called grain boundary sliding, replaces the dislocation activity that dominates deformation process in conventional materials. Softening caused by the grain boundary sliding is mainly attributed to the large amount of defects in the grain boundaries, which allows fast diffusion of atoms and vacancies with the applying of stress. A further increase of the strength and hardness with decreasing crystallite size can be achieved only if grain boundary sliding is blocked by appropriate coating design and materials selection.
0257-8972/02/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 2 . 0 0 5 6 1 - 3
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
43
Fig. 1. AFM morphology of the film deposited at Ti target power density of 4.5 Wycm2 and nitrogen gas flow ratio of (a) 33% (b) 100%.
There are many design ideas for nanocomposite coatings, and embedding nanocrystalline phases in amorphous phase matrix is quite often applied which can be fulfilled easily by physical vapor deposition (PVD) and chemical vapor deposition (CVD) methods. Diamondlike carbon (DLC), amorphous carbon nitride and other hard amorphous materials (with high hardness and elastic modulus) have been recognized as the primary candidates for the amorphous matrix while nano-sized refractory nitrides, such as TiN, Si3N4, AlN, BN, etc., could be used as strengthening phases w28x. For this coating design, the size and the distribution of these nanocrystals in the amorphous structures need to be optimized to obtain high hardness. These nanocrystalline grains should have random orientation (high angle grain
boundaries) to minimize grain incoherence strain, and high toughness could be achieved by strain release via nanocrystals sliding in the DLC matrix. Termination of nanocracks by deflection at grain boundaries and by energy loss within the amorphous matrix can dramatically improve the toughness. Carbon is one of the most important and widely used materials in many areas and diamond-like carbon (DLC) has been recognized as one of the primary candidates used for wear resistant and solid lubricating coating. TiN (or TiC) and CrN coatings prepared using sputtering have been applied in industry for wear protective applications. Following the above design idea, in this study, we combine these two types of coatings to synthesize a nanocomposite structure. In film design, DLC acts as
44
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
Fig. 2. GIXD analysis of the films deposited under 300 W Ti target power with AryN2s2:1.
hard, tough and lubricating matrix, while nano-particles act as reinforcing crystallites to improve hardness and mechanical properties. 2. Experimental Nanocomposite (TiCr)CxNy yDLC coatings were prepared through co-sputtering of graphite, Ti and Cr targets
in argonynitrogen atmosphere in a magnetron sputtering system. Silicon wafers and WC discs were used as the substrates. WC discs with a diameter of 50 mm and a thickness of 5 mm were mechanically ground and polished, then ultrasonically rinsed in acetone. The substrate holder rotated during the deposition for uniformity. In this study, nitrogen gas flow ratio wN2 y(N2q Ar)x was controlled at 33 and 100%, with Ti target power density of 4.5 and 6.5 Wycm2 (d.c.). The power density of chromium and graphite target were controlled at 4.5 Wycm2 (d.c.) and 6.5 Wycm2 (r.f.), respectively. Deposition was performed at a low temperature of 150 8C for 2 h. The substrate-to-target distance was 100 mm. Surface and cross-section morphologies of the coatings were investigated using an atomic force microscopy (AFM) and a JEOL scanning electron microscope. The crystalline structure was obtained by grazing incidence X-ray diffraction (GIXD) method. The nanocomposite coating was also deposited on potassium bromide (KBr) pellets for 20 min, and then the pellets were dissolved in water to float off the film for TEM study. The thin film thus obtained was examined using a JEOL 200 kV TEM. X-ray photoelectron spectroscopy (XPS) analysis was performed on film surface using a Kratos AXIS spectrometer with monochromatic Al Ka (1486.6 eV)
Fig. 3. TEM photos of deposited nanocomposite thin films.
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
45
Fig. 4. Load-displacement curve during load and unloading process for nano-indentation.
X-ray radiation. Film hardness and elastic recovery were evaluated using a nano-indentation tester with a penetrating depth of 100 nm. Load bearing capacity of the nanocomposite coatings on WC substrate was assessed using a scratch tester. A diamond stylus was driven across the coating at a constant speed of 1 mmys and a continuously increased load rate of 1.5 Nys. Tribological behaviors of nanocomposite coatings were evaluated using a ball-on-disk tribometer under dry sliding conditions at room temperature. Alumina balls (with a diameter of 9.5 mm and a surface roughness better than Ras0.05 mm) were used as the counterface materials. The normal loads were 10, 20 and 50 g. All the tests were run in laboratory air (25 8C and relatively humidity of 65"3%) with a sliding distance of 300 m and a sliding speed of 0.2 mys. Coefficient of friction was recorded during each test. Wear volumes were calculated from the wear tracks measured with a laser profilometer.
Fig. 5. XPS depth profile of the film deposited at Ti power density of (a) 4.5 Wycm2 and (b) 6.5 Wycm2.
density of 6.5 Wycm2 with nitrogen to argon gas flow ratio of 1:3. There are some broad peaks of (TiCr)CN crystalline phases. As a first degree approximation, the average crystallite size can be estimated by the Debye– Scherrer formula (shown in Eq. (1)) w30x:
3. Results and discussions Cross-section of the coating is dense and featureless with a thickness of approximately 1 mm. Fig. 1a shows AFM morphologies of the coating deposited at a power density of 4.5 Wycm2 on the Ti target with nitrogen to argon gas flow ratio of 1:3. The surface features a wavy morphology or hilly humps of approximately 30–40 nm in width and 6–8 nm in height. In the case of deposition under a pure nitrogen gas, the coating roughness increases significantly (see Fig. 1b). This is probably owing to the rapid formation and growth of nitrides, and the formation of polymeric C–N based phase w29x. Fig. 2 shows the GIXRD profile of the crystalline structures for the film deposited at Ti target power
Table 1 Film bonding structure distribution (%) of C 1s and N 1s peaks Bonding structure
C 1s 284.7 285.7 286.6 288.3
eV(C–C bonding) eV(sp2 bonding) eV(sp3 bonding) eV (C–O bonding)
N 1s 398.3 eV(N–C sp3 bonding) 399.1 eV(N–C sp2 bonding) 401 eV (N–O bonding)
%, at Ti power density of 4.5 Wycm2
at Ti power density of 6.5 Wycm2
45.53 36.46 7.57 10.44
57.17 16.04 15.99 10.79
42.30 40.43 17.27
61.64 25.44 12.92
46
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
Fig. 6. Typical XPS high-resolution carbon 1s and nitrogen 1s core level spectra.
Fig. 7. Scratch profile of the film deposited at Ti target power density of 6.5 Wycm2 and nitrogen flow ratio of 33%. The insets show SEM morphology of the scratch track at different loading stage.
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
47
(1)
DsKlybcosu
Fig. 8. Wear curves for the film deposited at Ti target power density of 6.5 Wycm2 and nitrogen gas flow ratio of 33% at increasing normal load: (a) 10 g, (b) 20 g, (c) 50 g.
Where K is a constant (Ks0.91), D is the mean crystallite dimension normal to diffracting planes, l is the X-ray wavelength (ls0.15406 nm for Cu target), b in rad is the peak width at half-maximum peak height and u is the Bragg angle. The calculated average grain size of (Ti,Cr)CN crystals is approximately 20 nm. For all the films in this study, the calculated grain size of (Ti,Cr)CN crystals are within 10–20 nm. High magnification bright-field TEM photo of the films is shown in Fig. 3a. In an amorphous DLC matrix, many tiny crystals can be observed with dimensions approximately 8–12 nm. These nanocrystalline crystals have different orientations and lattice spacing. It should be pointed out that these crystals form only after 20min deposition, while those in the coating deposited for 2 h may be larger than those shown in Fig. 3a due to grain growth, as is illustrated in the AFM morphology in Fig. 1. Fig. 3b shows the dark-field TEM photos, clearly revealing these nano-size crystals. Fig. 3c shows the diffraction ring patterns, indicating these nanocrystalline phases of TiCrCN. Fig. 4 is the nanoindentation load-displacement profile of the film deposited at Ti target power density of 6.5 Wycm2. The calculation of hardness and elastic modulus is performed using Oliver and Pharr method w31x. The measured film hardness is approximately 40 GPa. The Young’s modulus of the film is approximately 300 GPa. With the decrease of Ti power density to 4.5 W/cm2, the film hardness decreases to approximately 25 GPa. Fig. 5 gives the atomic composition of the films obtained from XPS analysis. XPS results revealed that with the increase of Ti target power density 4.5–6.5 Wy cm2, Ti in the film increases from 12.3 to 22 at.%, while the contents of other elements only change slightly. Therefore, one possible reason for the hardness increasing with Ti target power could come from the increase in the amount of nano-sized (TiCr)CN phases in the film. With the increase in the nitrogen flow ratio from 33 to 100%, film hardness decreases dramatically from 40 to 18 GPa. The high content of polymeric C– N based phase formed under high content of nitrogen gas could be detrimental for the film hardness. Fig. 6 shows typical high-resolution carbon 1s and nitrogen 1s core level spectra. The peaks were fitted using the binding energy interval values from the literature w32,33x. A Gaussian decomposition of C 1s spectra
Table 2 Average coefficient of friction for specimens prepared under different Ti contents and normal loads 10 N
2
4.5 Wycm 6.5 Wycm2
20 N
50 N
COF
Wear volume
COF
Wear volume
COF
Wear volume
0.03 0.02
4.06 not appreciable
0.02 0.03
0.61 not appreciable
0.01 0.03
2.02 not appreciable
48
S. Zhang et al. / Surface and Coatings Technology 162 (2002) 42–48
gives rise to the following peaks: 284.7 eV (C–C bonding), 285.7 eV (sp2 bonding), 286.6 eV (sp3 bonding) and 288.3 eV (C–O bonding). Similar analysis on N 1s spectra yields 398.3 eV (N–C sp3 bonding), 399.1 eV (N–C sp2 bonding) and 401 eV (N–O bonding). The concentrations of the different bonding structures of C 1s and N 1s peaks are calculated and the results are given in Table 1. With the increase of Ti target power density, the concentration of sp3 bonding structure increased for both C 1s and N 1s peaks, and at the same time, sp2 decreased, which supports the significant increase in film hardness as discussed earlier. Scratch test was used to evaluate the dynamic load bearing capacity and friction properties of the coatingsubstrate system. The ‘lower critical load’ or the normal load at which the first damage or a sharp increase in friction coefficient is observed is widely used as a measure of the load bearing capacity. Fig. 7 is such a scratch profile which clearly shows the coating fails at a normal load of approximately 60 N, indicating a good adhesion and load bearing capacity of nanocomposite coatings on WC–Co substrate. The coefficient of friction remains at a very low value of 0.05, indicating good friction properties. SEM micrographs for the different stages during scratch testing are also shown as inset in Fig. 7. The critical load slightly decreased to 53 N at Ti target power density 4.5 Wycm2. With the increase of Ti target power, there is no significant change in coefficient of friction. The coefficient of friction remains at an extremely low value of 0.01–0.03 within the 300-m wear distance under different normal loads (Fig. 8). With the increase of normal load, the coefficient of friction decreases slightly, which can probably be explained by the easier graphitization of wear debris and generating of lubricating surface layer with the increase of normal load. Table 2 lists the average coefficient of friction and wear volume for specimens prepared under different Ti target powers and normal loads. The long-term coefficient of friction remains a low and stable value (less than 0.03). With the increase of Ti target power density from 4.5 to 6 Wycm2, the wear volume decreases so much that the surface profilometry does not reveal noticeable wear track. 4. Conclusions Superhard nanocrystalline (Ti, Cr)CNyDLC coatings were prepared through co-sputtering of Ti, Cr and graphite targets in an argonynitrogen atmosphere. Results from both TEM and GIXRD indicated that the grain size of the (TiCr)CxNy crystals were approximately 10–20 nm. XPS results confirmed that an increase in the sputtering power at the Ti target not only increased the Ti composition in the film but also brought about increase in sp3 bonding in the DLC matrix, in agreement
with the raising hardness with Ti sputtering power. Film hardness and elastic modulus were measured with a nano-indenter. Under the experiment conditions, the film hardness fell within 18–40 GPa. The tribological behaviors of the films were evaluated using a ball-on-disk tribometer, and the films demonstrated properties of low-friction and good wear resistance. References w1x P. Holubar, M. Jilek, M. Sima, Surf. Coat. Technol. 133–134 (2000) 145. w2x R.A. Andrievski, Mater. Trans. 42 (2001) 1471. w3x J. Musil, Surf. Coat. Technol. 125 (2000) 322. w4x S. Veprek, A.S. Argon, J Vac. Sci. Technol. B 20 (2002) 650. w5x S. Veprek, S. Reiprich, Thin Solid Films 268 (1995) 64. w6x V. Provenzano, R.L. Holtz, Mater. Sci. Eng. A 204 (1995) 125. w7x F. Mazaleyrat, L.K. Varga, J. Magn. Magnetic Mater. 215–216 (2000) 253. w8x B. Cantor, C.M. Allen, Scr. Mater. 44 (2001) 2055. w9x H. Zeng, M.L. Yan, Y. Liu, J. Appl. Phys. 89 (2001) 810. w10x J.L. Solis, S. Saukko, L. Kish, Thin Solid Films 391 (2001) 255. w11x S. Veprek, A.S. Argon, Surf. Coat. Technol. 146 (2001) 175. w12x S. Veprek, Thin Solid Films 317 (1998) 449. w13x J. Musil, F. Kunc, H. Zeman, Surf. Coat. Technol. 154 (2002) 304. w14x J. Musil, J. Vlcek, Surf. Coat. Technol. 142 (2001) 557. w15x M.A. Baker, S. Klose, C. Rebholz, Surf. Coat. Technol. 151 (2002) 338. w16x E. Ribeiro, A. Malczyk, S. Carvalho, Surf. Coat. Technol. 151 (2002) 515. w17x D.V. Shtansky, K. Kaneko, Y. Ikuhara, Surf. Coat. Technol. 148 (2001) 206. w18x L. Rebouta, C.J. Tavares, Surf. Coat. Technol. 133–134 (2000) 234. w19x S. Carvalho, L. Rebouta, Thin Solid Films 398–399 (2001) 391. w20x M. Misina, J. Musil, Surf. Coat. Technol. 110 (1998) 168. w21x J. Musil, H. Polakova, Surf. Coat. Technol. 127 (2000) 99. w22x J. Musil, H. Hruby, Thin Solid Films 36 (2000) 104. w23x J. Musil, P. Karvankova, J. Kasl, Surf. Coat. Technol. 139 (2001) 101. w24x A.A. Voevodin, J.P. O’Neill, J.S. Zabinski, Surf. Coat.Technol. 116–119 (1999) 36. w25x S. Veprek, A. Niederhofer, K. Moto, T. Bolom, H.D. Mannling, P. Nesladek, G. Dollinger, A. Bergmaier, Surf. Coat.Technol. 133–134 (2000) 152. w26x F. Vaz, L. Rebouta, P. Goudeau, Surf. Coat.Technol. 146 (2001) 274. w27x J. Patscheider, T. Zehnder, M. Diserens, Surf. Coat. Technol. 146 (2001) 201. w28x H. Holleck, in: A. Kumar, Y.W. Chung, J.J. Moore, J.E. Smugeresky (Eds.), Surface Engineering: Science and Technology I, The Minerals, Metals and Materials Society, 1999, pp. 207–231. w29x A. Stanishevsky, L. Khriachtchev, I. Akula, Diamond Relat. Mater. 7 (1998) 1190–1195. w30x M.V. Zdujic, O.B. Milosevic, L.C. Karanovic, Mater Lett. 13 (1992) 125–128. w31x T. Berlind, N. Hellgren, Surf. Coat. Technol. 141 (2001) 145. w32x E.G. Wang, Prog. Mater. Sci. 41 (1997) 241–298. w33x K. Sharma, J. Narayan, Int. Mater. Rev. 42 (1997) 137–192.