Towards high adherent and tough a-c coatings

Transcription

1 Thin Solid Films 482 (2005) Towards high adherent and tough a-c coatings Sam Zhang a, *, Xuan Lam Bui a, X.T. Zeng b, Xiaomin Li c a School of Mechanical and Production Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore , Singapore b Singapore Institute of Manufacturing Technology, 71 Nanyang Drive, Singapore , Singapore c State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai , P.R. China Available online 30 December 2004 Abstract How to increase the adhesion and toughness of hard or superhard amorphous carbon (a-c) coatings deposited on engineering substrates (steels, cemented carbide, etc.) is the subject of recent intensive study. Substrate bias grading and incorporation of metals in the coating are effective means to achieve high adhesion and high toughness yet maintain adequate hardness. This paper summarizes the newest development on this subject (bias grading and metal doping) and presents findings in adhesion and tribological studies. Four different types of magnetron-sputtered carbon-based coatings have been deposited on stainless steel substrates: (1) pure a-c deposited under constant substrate bias voltage of 140 V, (2) a-c coating deposited with substrate bias voltage graded stepwise from 20 to 150V, (3) nanocomposite coating obtained by co-sputtering of Ti and graphite targets and (4) nanocomposite coating obtained by co-sputtering of Ti, Al and graphite targets. The structure, adhesion, residual stress and tribological properties of these coatings are presented. The high adhesion strength and good tribological properties found in these coatings are highly desirable in engineering applications. D 2004 Elsevier B.V. All rights reserved. Keywords: Nanocomposite; Adhesion strength; Residual stress; Tribology 1. Introduction Amorphous carbon (a-c, also called diamond-like carbon or DLC) is classified as hydrogenated amorphous carbon (a- C:H) and nonhydrogenated amorphous carbon or also called hydrogen-free amorphous carbon (a-c). Comparing to a- C:H, in which hydrocarbon gases were employed as the source of carbon, a-c exhibits more beneficial properties such as higher hardness and elastic modulus, lower friction in humid environment, better thermal stability, etc. [1]. The interest in commercializing carbon-based technologies for engineering applications has been growing with time. However, the deposition of high-quality thick a-c coatings is still difficult owing to the excessive residual stresses developed during deposition. Under certain residual stress and coating-to-substrate bonding situation, there is a maximum coating thickness that can ensure a good coating * Corresponding author. Tel.: ; fax: address: msyzhang@ntu.edu.sg (S. Zhang). service without adhesion failure [2]. Usually, residual stresses can be divided into three parts: the growth-induced stress (intrinsic stress) as a result of ion bombardment during the growth of the coating, thermal stress due to mismatch of coefficient of thermal expansion between the coating and the substrate and other stress (extrinsic stress) as a result of post-deposition bcontaminationq of the coating such as gas and moisture absorption, etc. mostly seen in porous coatings [3,4]. For pure and dense a-c coatings, which are mostly deposited at low temperatures, the thermal and post-deposition stresses should be negligible; thus, the main part of the residual stress is the growth-induced stress. Depending on the structure, the magnitude of residual stress in a-c ranges from a few GPa to 10 GPa [5 7] that will greatly limit the maximum thickness of adherent coatings. For instance, Hou et al. reported that hydrogen-free a-c films deposited by pulsed laser deposition peel off as the coating thickness exceeds 200 nm [8]. Hydrogen-free a-c films deposited by filtered cathodic vacuum arc delaminate from the substrate at a thickness of about 180 nm [9] /$ - see front matter D 2004 Elsevier B.V. All rights reserved. doi: /j.tsf

2 S. Zhang et al. / Thin Solid Films 482 (2005) Sputtered a-c (with unbalanced configuration) deposited at a bias voltage of 50 V could not adhere to Si substrate if the coating thickness exceeds 1 Am [3]. It should be noted that the performance of a coating is judged not only by its wear resistance but also by its durability. The wear resistance is related to hardness and coefficient of friction whereas the durability requires high adhesion and toughness. Obtaining extreme high hardness alone is not difficult, but producing a combination of hardness, adhesion and toughness is very challenging. To achieve low residual stress and good adhesion, using a bond layer is a common practice. Other methods used include: (1) precoating cleaning, such as plasma etching of substrate surface [10] or water peening [11]; (2) composition grading [12,13]; (3) structural grading (multilayer coating) [14]; (4) annealing [15] and doping with metallic or nonmetallic elements [16,17]. Annealing treatment may affect the microstructure of the substrate, which limits the choice of substrate. Recently, we proposed substrate bias-graded deposition [18], whereby during the deposition process, the bias voltage was applied to the substrate in a bgradedq manner: gradually increase as the deposition progressed and coating thickness increased, therefore creating a graded sp 3 /sp 2 bonding throughout the thickness of the coating. The result was a coating of high adhesion with a graded hardness: the lowest at the coating substrate interface, higher as the coating grows, and the highest at the coating surface. This effectively improves the bonding adhesion of the coating on the substrate surface and at the same time provides high hardness on the coating surface. The incorporation of metals into hydrogen-free a-c matrix is another effective way in reduction of growth-induced stress through co-sputtering metal to act as a bstress-relaxant.q Among stress relaxing elements (Ti, Al, Si, etc.), Al is found to be one of the most effective elements in relieving stress [17,19]. However, the incorporation of Al results in reduction of hardness due to reduced sp 3 hybridization: about 60% of the coating hardness is lost when it is doped with only 10 at.% of aluminum [19]. To restore hardness, we have co-sputtered Ti, Al and C to produce randomly orientated nanocrystalline TiC grains embedded in Al-containing a-c matrix and form a nanocomposite coating of nc-tic/a-c(al) [20]. This coating has low residual stress (thus can be made thick), high hardness and toughness. The present paper summarizes the parametric studies of the substrate bias-grading and bdopingq (co-sputtering) of Al and Ti in a-c coatings. The results of adhesion and tribological studies are also presented with a comparison of four different types of magnetron-sputtered carbon-based coatings on stainless steel substrate: (1) pure a-c deposited under constant substrate bias voltage of 140 V, (2) a-c coating deposited with substrate bias voltage graded stepwise from 20 to 150 V, (3) nanocomposite coating obtained by co-sputtering of Ti and graphite targets and (4) nanocomposite coating obtained by co-sputtering of Ti, Al and graphite targets. 2. Experimental 2.1. Deposition of coatings Pure a-c and nanocomposite coatings were deposited on 440C stainless steel discs (with diameter of 55 mm and thickness of 5.5 mm) polished to surface roughness of R a =60 nm and Si wafers (100 mm diameter, 450 Am thickness and 2.0 nm in R a ) by DC magnetron sputtering using an E303A system (Penta Vacuum-Singapore) [21]. The a-c coatings were deposited using graphite target (100 mm in diameter, % purity) at a power density of 10.5 W/cm 2. The substrate bias voltage was applied in two modes: bconstant biasq and bgraded biasq. In bconstant bias,q a bias voltage of 140 V was applied to the substrate during the whole deposition process for the whole thickness of the coating. In bbias-graded deposition,q the substrate bias voltage was increased stepwise from 20 to 150 V at a step size of 2 V for every 100 s. bdopingq of the a-c matrix was done through co-sputtering of aluminum (100 mm in diameter, % purity) and graphite targets. The nanocomposite coatings of nc-tic/a-c were deposited by sputtering of titanium (100 mm in diameter, % purity) and graphite targets. The nanocomposite coatings of nc-tic/a-c(al) were deposited by sputtering of the titanium, aluminum and graphite targets all at the same time. During deposition of the nanocomposite coatings, the substrate bias voltage was kept at 150 V, the substrate temperature at 150 8C, the power density of the graphite target at 10.5 W/cm 2 and that of metal targets was varied for composition variation (Ti target: W/cm 2 ; Al target: W/cm 2 ). The base pressure of the chamber was Pa and the process pressure was kept constant at 0.4 Pa at an Ar flow rate of 50 cm 3 /min for all coatings. Prior to deposition, the substrates were ultrasonically cleaned for 20 min in acetone followed by 10 min in ethanol. After loading, the substrates were heated to and maintained at 150 8C for 30 min before plasma cleaning for 30 min at bias voltage of 300 V to remove surface oxides and contaminants. ATi bond layer of 100 nm in thickness was deposited on the substrate before deposition of the coating Characterization The thickness of coatings was measured using profilometer (Dektak 3 SJ) through a sharp step created by masking. The surface morphology of coatings was investigated by atomic force microscopy (AFM) using SPM- 9500J2 (Shimadzu) system. The surface roughness of coatings was calculated using the software combining with AFM over an area of 22 Am. The structure of the coatings was investigated with a Renishaw Raman spectroscope at the 633-nm line excited with a He Ne laser. The composition was determined by X-ray photoelectron spectroscopy (XPS) using a Kratos AXIS spectrometer with monochromatic AlK a ( ev) radiation. The hardness was determined using a nanoindenter (XP from MST, USA)

3 140 S. Zhang et al. / Thin Solid Films 482 (2005) Table 1 Specification of Shell Helix 15W-50 oil Kinematic viscosity (cst) Viscosity 40 8C 100 8C index Density (15 8C) (kg/m 3 ) Flash point a (8C) Pour point b (8C) a Flash point: The lowest temperature at which the oil gives off enough flammable vapor to ignite and produce a flame when an ignition source is present. b Pour point: The lowest temperature at which the oil is observed to flow. Table 2 Composition and properties of the coatings Coating Composition (at.%) R a C Ti Al (nm) Hardness (GPa) Residual stresses (GPa) I D /I G a-c(al) a-c ( 140 V bias) a-c (bias-graded) nc-tic/a-c nc-tic/a-c(al) with a Berkovich diamond indenter. The indentation depth was set not to exceed 10% of the coating thickness to avoid possible interference from the substrate. The adhesion strength was studied using the scanning microscratch tester (Shimadzu SST-101) where a diamond tip stylus of 15 Am in radius was dragged on the coating with a gradually increased load. The scanning amplitude was set at 50 Am at a speed of 10 Am/s. The lower critical load was used to indicate the adhesion strength. The magnitude of the residual stresses was calculated from the change in radius of curvature of Si wafers measured by a Tencor laser scanner before and after deposition using the Stoney s equation: E s r ¼ 61 ð m s Þ ts t c R 2 R 1 where E s /(1 m s ) is the substrate biaxial modulus (180.5 GPa for Si(100) wafer [22]); t s and t c are the substrate and coating thickness, respectively; R 1 and R 2 are the radius of curvature of Si wafer before and after deposition of the coating. Tribological tests were carried out using CSEM tribometer with ball on disc configuration under ambient (relatively humidity of 75% and room temperature of 22 8C) and oil lubrication conditions. The ambient test was conducted at a sliding speed of 20 cm/s at a load of 5 N for 1 km. Under oil lubrication, the sliding speed used was 5 cm/s at a load of 10 N for 1 km. Stainless steel (100Cr6) balls of diameter 6 mm were used as counterpart for both ambient and oil-lubricated tests. Table 1 lists the specification of the lubricant (Shell Helix 15W-50 engine oil). sp 3 /sp 2 fraction originated from the high-energy ion bombardment and back-sputtering, which hinders partially the formation of graphite structure [21]. Raman spectrum of a-c coating was deconvoluted into two peaks termed D band (at 1350 cm 1 ) and G band (at 1530 cm 1 ) with I D as the intensity of D band and I G as the intensity of G band [21]. Results from Raman spectrum indicated that an increase in bias voltage resulted in a decrease in I D /I G ratio (cf. Fig. 1). Though the I D /I G ratio is not a direct measurement of sp 3 or sp 2 bonding fraction, the I D /I G ratio is inversely proportional to sp 3 /sp 2 fraction [23]. Thus a decrease in I D /I G ratio signals an increase in sp 3 hybridization. The minimum I D /I G ratio (about 1.0) corresponding to a maximum sp 3 /sp 2 fraction was observed for the coating deposited on substrate biased at 150 V. Therefore, it is expected that in the bias-graded a-c coating, the sp 3 /sp 2 fraction decreases gradually from the surface to the coating substrate interface. Similarly, the hardness also decreases from the top of the surface to the interface. At indentation depth of 110 nm, the hardness value of the bias-graded a-c coating of 1.5 Am in thickness is 25.1 GPa. This value is inserted in Table 2 for easy reference Effect of co-sputtering of Al and Ti with C Fig. 2 shows C 1s and Al 2p XPS peaks of a-c, a-c(al), nc-tic/a-c and nc-tic/a-c(al) coatings. C 1s peak at ev is from amorphous carbon matrix, its chemical shift at 3. Results and discussion The thickness of the coatings in this study ranges from 1 to 1.6 Am. The chemical composition, surface roughness, hardness, residual stresses and I D /I G ratio from Raman studies are tabulated in Table Effect of the substrate bias voltage on the bonding structure of a-c coatings The a-c coating deposited on substrate biased at 140 V has the highest hardness of 28.1 GPa. This is a result of high Fig. 1. I D /I G ratio obtained from Raman spectra as a function of substrate bias voltage.

4 S. Zhang et al. / Thin Solid Films 482 (2005) Fig. 2. C 1s and Al 2p XPS spectra of a-c, a-c(al), nc-tic/a-c and nc-tic/ a-c(al) coatings ev is from the bonding with Ti in TiC [24]. Elemental Al 2p peak is at 74.2 ev [17]. Aluminum carbide has a binding energy at ev and aluminum oxycarbide has a binding energy at ev [25], both are not seen in the a- C(Al) coating. The binding energy of aluminum carbide and aluminum oxycarbide are within the spread of the C(1s) peak for TiC. Comparing the two XPS spectra of nc-tic/a-c and nc-tic/a-c(al), the peak shapes did not have noticeable difference. With those observations, it is believed that aluminum does not form bonds with carbon and it exists as elemental aluminum. As thus, we used the notation ba- C(Al)Q for Al-doped a-c. Co-sputtering of Ti and C targets facilitated formation of TiC (as is evident from Fig. 2). The size of the TiC crystals was estimated, using Scherrer formula from XRD spectra of the coating, to be in the range of 6 16 nm. The result was also confirmed by TEM microscopy [20] as 4 14 nm. The close proximity of the result from XRD and that from TEM indicates that the strain broadening in XRD peaks is not happening owing to the low residual stress (1.2 GPa) in the coating. Since the TiC crystals were nanosized, they were denoted as nc-tic. As such, the coating produced was termed bnc-tic/a-cq nanocomposite coating. As Al and Ti were co-sputtered with carbon at the same time, the above two effects were additive: formation of a- C(Al) matrix and nc-tic took place simultaneously to give rise to an bnc-tic/a-c(al)q nanocomposite coating. The size of the nc-tic was determined as 4 7 nm by XRD and 3 7 nm by TEM [20]. the coating (or increase in sp 2 bonding). Comparing with Ti, the addition of aluminum caused more reduction in the amount of sp 3 bonding since the aluminum added went into the a-c matrix that disturbed the carbon structure whereas adding Ti formed nc-tic. Compared with the amorphous carbon (cf., Table 2), the addition of 29 at.% of Ti led to an increase in I D /I G by about 120%, but about 19 at.% of Al going into the carbon matrix resulted in an increase in I D /I G of almost 210%. The addition of 13 at.% of Al and 40 at.% of Ti at the same time was responsible for an increase in I D /I G of almost 230% increase. That reveals a large increase in sp 2 bonding, which helped the residual stress relaxation and promoted toughness at the expense of hardness. However, the hardness of coatings was still relatively high (27.4 GPa for nc-tic/a-c coating and 19.6 GPa for nc-tic/a-c(al)). This hardness did not come from the high sp 3 /sp 2 fraction, but from the formation of hard nc-tic. In the meantime, the residual stress in the nc-tic/a-c(al) coating dropped down to a very low level of 0.4 GPa. That is a combined consequence of (1) low sp 3 /sp 2 bonding fraction (2) incorporation of soft metal Al in the a-c matrix Adhesion strength The adhesion strength is given in Fig. 4 in terms of lower critical load obtained from the scratch tests. The a-c coating deposited on substrate biased at 140 V exhibited the lowest adhesion (187 mn). For bias-graded a-c coating, the adhesion strength was more than doubled. Since the surface roughness of the a-c coating deposited at a constant bias voltage of 140 V and bias-graded a-c was at the same low level of ~3.5 nm in R a on Si wafer, the considerable increase in adhesion obtained in the bias-graded a-c coating was attributed to the combination of low residual stress and high toughness. As studied in Ref. [18,21], coatings deposited under lower bias voltage exhibited higher toughness and lower residual stress. As such, in bias-graded coating, it is 3.3. sp 3 /sp 2 fraction and the coating hardness Raman spectra of the bias-graded a-c, a-c(al), nc-tic/ a-c and nc-tic/a-c(al) coatings are shown in Fig. 3. Detailed analysis of the spectra was given in in Ref. [20]. Addition of Ti and/or Al caused a decrease in sp 3 bonding in Fig. 3. Raman spectra of bias-graded a-c, a-c(al), nc-tic/a-c and nc-tic/ a-c(al) coatings.

5 142 S. Zhang et al. / Thin Solid Films 482 (2005) Fig. 4. Scratch adhesion strength (lower critical load) of coatings: (a) a-c deposited under constant bias of 140 V, (b) bias-graded a-c, (c) nc-tic/ a-c and (d) nc-tic/a-c(al). natural that the substrate coating interface becomes more ductile, which results in higher adhesion strength. The nc-tic/a-c coating possessed an adhesion strength of 359 mn (Fig. 4c), slightly lower than that of the biasgraded a-c coating (381 mn, Fig. 4b) even though its residual stress was lower (1.2 GPa compared to 1.5 GPa, cf., Table 2). The lower residual stress should result in higher adhesion. The slight discrepancy is understood because firstly the residual stress was obtained from the change in the curvature of a whole Si wafer of 100 mm in diameter, i.e., an average of the whole coating regardless of possible variation due to structural grading in case of bias-graded deposition. In bias-graded a-c coating, sp 3 /sp 2 fraction increases from the substrate coating interface towards the outer surface of the coating, the local residual stress at the interface should be a lot lower than that close to the surface (where sp 3 /sp 2 ratio is the highest). Furthermore, the roughness, R a, of nc-tic/a-c was 7.3 nm on Si wafer (more than double that of the bias-graded coating) that results in higher friction thus higher shear stress at the contact area, giving rise to slightly lower critical load. In the case of nc-tic/a-c(al) nanocomposite coating, a very high adhesion of 697 mn was obtained. Evidently, the extremely low residual stress of 0.4 GPa played an important role. Another important contribution should be attributed to the extreme toughness obtained from the coating structure. The propagation of the microcracks generated in the scratch process will be hindered at the boundaries between the matrix and the grains. Meanwhile, more crack propagation energy will be relaxed in the tough a-c(al) matrix. The smoothness of the surface (5.5 nm in R a on Si wafer) also contributed towards the high critical load. As clearly seen from the optical micrograph of the scratch track, the a-c coating deposited at a constant substrate bias voltage of 140 V delamined in a brittle manner (Fig. 5a) as the load reached 187 mn, where the lower critical load and the higher critical load were not distinguishable. For the bias-graded a-c coating and the nanocomposite coatings (nc-tic/a-c and nc-tic/a-c(al)), however, damages inflicted onto the coating were not continuous (sporadic) with increasing load (Fig. 5b d). The fracture surface of nc-tic/a-c(al) coating appeared very bplasticq (Fig. 5d): the cracks formed but could not Fig. 5. Optical image of scratch tracks on coatings: (a) a-c deposited under constant bias of 140 V, (b) bias-graded a-c, (c) nc-tic/a-c and (d) nc-tic/a-c(al).

6 S. Zhang et al. / Thin Solid Films 482 (2005) inflict spallation instead, the tip was seen plough into the coating. As the tip ploughed deeper, the scanning amplitude decreased because the force exerted on the tip to vibrate in the transverse direction was not enough to overcome the resistance created by the material piled up Tribology The coefficient of friction vs. sliding distance under ambient and oil-lubrication conditions is plotted in Fig. 6. For clarity of the plot, results from two coatings are demonstrated: bias-graded a-c and nc-tic/a-c(al). Fig. 7 summarizes the coefficients of friction (steady-state value at the end of the test) obtained for all the coatings under ambient condition. These values are tremendously lower than the most popular ceramic coatings. For instance, for TiN the coefficient of friction is in the range of [26]. During the wear test of a-c coating, a graphite-rich layer forms between the two sliding surfaces [21,27]. In a humid environment, the graphite layer absorbs moisture and acts as a lubricant that contributes to reduce the friction coefficient. The same is true in composite coatings based on a-c [28]. Comparing a-c coatings (Fig. 7a and b) with nanocomposite coatings (Fig. 7c and d), the nanocomposite coatings had appreciably higher coefficients of friction because the embedded nc-tic caused rougher surface. Also, there is less carbon in nanocomposite coatings thus less graphite forms between the surface of coating and the wearing counterpart. The coefficients of friction of the four coatings and an uncoated steel substrate tested under oil lubrication are given in Fig. 8. The coefficients of friction for the coatings fluctuated less compared to that in the dry tests (also see Fig. 6a vs. c and b vs. d). Under oil lubrication, oil prevents the formation of the tribolayer and governs the friction behavior in the contact. In our test, because the oil is not pressurized into the contact, the oil film between the two Fig. 7. Coefficient of friction under ambient condition for (a) a-c deposited under constant bias of 140 V, (b) bias-graded a-c, (c) nc-tic/a-c and (d) nc-tic/a-c(al). surfaces is not thick enough to separate the coating surface and the counterpart; therefore contacts between the asperities of the two surfaces occur. All other conditions being the same, the surface morphology or roughness will dominate the friction process: rougher surface results in higher coefficient of friction. The a-c deposited at 140 V (a) and that under bias-graded condition (b) have about the same surface roughness (3.4 and 3.5 nm), the coefficient of friction is almost the same: ~0.14 under ambient condition and under oil lubrication. That was expected because the surface morphology of these two coatings was almost the same as a result of similar deposition condition at the top surface (bias voltage of 140 to 150 V). The difference in the coefficient of friction of nc-tic/a-c (Fig. 8c) and that of nc-tic/a-c(al) (Fig. 8d) reflects two effects: (1) roughness and (2) metal incorporation. Both aspects work towards resulting lower coefficient of friction for nc-tic/a-c(al): smoother surface morphology and more metal incorporation (Ti and Al). Pure a-c is a chemically inert material: the bond between lubricating oil and the a-c is not as strong as that between oil and metal-contained a-c. With stronger bonding between Fig. 6. Coefficient of friction vs. sliding distance of bias-graded a-c and nc- TiC/a-C(Al) coatings in ambient and oil lubrication conditions. Fig. 8. The coefficient of friction under oil lubrication for (a) a-c deposited under constant bias of 140 V, (b) bias-graded a-c, (c) nc-tic/a-c, (d) nc- TiC/a-C(Al) and (e) uncoated stainless steel substrate.

7 144 S. Zhang et al. / Thin Solid Films 482 (2005) oil and the coating, the lubrication effect will be more prominent leading towards lower coefficient of friction. This is seen in Fig. 8 by comparing (c) and (d) with (a) and (b). More investigation is under way. 4. Conclusion In quest of bsuperhardq and yet tough ceramic coatings, a-c-based coatings prove to be good candidates. bbiasgradedq deposition creates a graded bonding structure such that the sp 2 hybridization become more towards the substrate coating interface to provide better toughness and adhesion while sp 3 bonding gets more towards the surface of the coating to render higher hardness for tribological performance. Co-sputtering of Al with graphite embeds elemental Al into a-c to form a-c(al) coating which greatly reduces growth-induced stresses acquired during deposition, at the expense of hardness because of the reduction in the amount of sp 3 bonding. Co-sputtering of Ti and Al with graphite produces nc-tic to embed in a-c(al) matrix. The formation of nc-tic helps partially restore the hardness lost due to bdopingq of Al. Comparing with a-c coating deposited under constant substrate bias, the bias-graded deposition results in less friction and improves adhesion of amorphous carbon by twofolds. Embedding of nanocrystalline TiC in amorphous carbon increases adhesion by two to four times with a slight increase in coefficient of friction. The incorporation of Al in the amorphous carbon matrix doubles the adhesion strength and reduces the coefficient of friction by 20%. Under ambient condition, the nanocomposite coatings experience higher friction than a-c coatings while under oil lubrication the nanocomposite coatings have less friction. For nc-tic/a-c(al) coating, a low coefficient of friction of only 0.04 is achieved. Acknowledgment This work was supported by Nanyang Technological University Research Grant RG12/02. References [1] Y. Lifshitz, Diamond Relat. Mater. 8 (1999) [2] J. Musil, Proceeding of Symposium 1992 of Research Center for Ultra High Energy Density Heat Source, Osaka University, Japan, 27 March 1992, p. 23. [3] E. Mounier, Y. Pauleau, Diamond Relat. Mater. 6 (1997) [4] S. Zhang, D. Sun, Y.Q. Fu, H.J. Du, Q. Zhang, Diamond Relat. Mater. 13 (2004) [5] D.R. McKenzie, Y. Yin, N.A. Marks, C.A. Davis, B.A. Pailthorpe, G.A.J. Amaratunga, V.S. Veerasamy, Diamond Relat. Mater. 3 (1994) 353. [6] S. Zhang, H. Xie, X. Zeng, P. Hing, Surf. Coat. Technol. 122 (1999) 219. [7] D. Sheeja, B.K. Tay, S.P. Lau, X. Shi, Wear 249 (2001) 433. [8] Q. Hou, J. Gao, Mod. Phys. Lett. B11 (1997) 757. [9] B.K. Tay, D. Sheeja, L.J. Yu, Diamond Relat. Mater. 12 (2003) 185. [10] S. Zhang, H. Xie, Surf. Coat. Technol. 113 (1999) 120. [11] H.K. Tonshoff, A. Mohlfeld, C. Gey, G. Winkler, Surf. Coat. Technol (1999) 440. [12] M. Stuber, S. Ulrich, H. Leiste, A. Kratzsch, H. Holleck, Surf. Coat. Technol (1999) 591. [13] A.A. Voevodin, M.A. Capano, S.J.P. Laube, M.S. Donley, J.S. Zabinski, Thin Solid Films 298 (1997) 107. [14] A.A. Voevodin, C. Rebholz, J.M. Schneider, P. Stevenson, A. Matthews, Surf. Coat. Technol. 73 (1995) 185. [15] A.C. Ferrari, B. Kleinsorge, N.A. Morrison, A. Hart, V. Stolojan, J. Robertson, J. Appl. Phys. 85 (10) (1999) [16] A. Grill, Wear 168 (1993) 143. [17] B.K. Tay, P. Zhang, Thin Solid Films (2002) 177. [18] S. Zhang, X.L. Bui, Y.Q. Fu, D.L. Butler, H.J. Du, Diamond Relat. Mater. 13 (2004) 867. [19] B.K. Tay, Y.H. Cheng, X.Z. Ding, S.P. Lau, X. Shi, G.F. You, D. Sheeja, Diamond Relat. Mater. 10 (2001) [20] S. Zhang, X.L. Bui, Y.Q. Fu, Thin Solid Films 467 (2004) 261. [21] S. Zhang, X.L. Bui, Y.Q. Fu, Surf. Coat. Technol. 167 (2003) 137. [22] W.A. Brantley, J. Appl. Phys. 44 (1973) 534. [23] S. Zhang, X.T. Zeng, H. Xie, P. Hing, Surf. Coat. Technol. 123 (2000) 256. [24] A.A. Voevodin, J.S. Zabinski, Mater. Sci. 33 (1998) 319. [25] B. Mauyrama, F.S. Ohuchi, L. Rabenberg, Mater. Sci. Lett. 9 (1990) 864. [26] K. Holmberg, A. Matthews, Tribology Series, vol. 28, Elsevier, [27] A.A. Voevodin, A.W. Phelps, J.S. Zabinski, M.S. Donley, Diamond Relat. Mater. 5 (1996) [28] A.A. Voevodin, J.P. O Neill, J.S. Zabinski, Thin Solid Films 342 (1999) 194.