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Microstructure and mechanical properties of WC–C nanocomposite films
Diamond and Related Materials 11 (2002) 1747–1752
Microstructure and mechanical properties of WC–C nanocomposite films Se Jun Parka,b, Kwang-Ryeol Leea,*, Dae-Hong Kob, Kwang Yong Euna a
Future Technology Research Division, Korea Institute of Science and Technology, P.O. Box 131, Cheongryang, Seoul, 130-650, South Korea b Department of Ceramic Engineering, Yonsei University, Seoul, 120-701, South Korea Received 5 February 2002; received in revised form 8 May 2002; accepted 13 May 2002
Abstract WC–C nanocomposite film was prepared by using a hybrid deposition system of r.f.-PACVD and DC magnetron sputtering. W concentration in the film was varied from 5.2 to 42 at.% by changing the CH4 fraction of the mixture sputtering gas of Ar and CH4. Hardness, residual compressive stress and electrical resistivity were characterized as a function of W concentration. Raman spectroscopy, XRD and high resolution TEM were employed to analyze the structural change in the film for various W concentrations. In the present W concentration range, the film was composed of nano-sized WC particles of diameter less than 5 nm and hydrogenated amorphous carbon matrix. Content of the WC particles increased with increasing W concentration. However, the mechanical properties of the film increased only when the W concentration was higher than 13 at.%. Structural analysis and electrical conductance measurements evidently showed that the increase in hardness and residual stress occurred as the WC particles were in contact with each other in the amorphous carbon matrix. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: WC–C nanocomposite films; Hybrid r.f.-PACVD and magnetron sputtering; Microstructure evolution; Mechanical and electrical properties
1. Introduction Although diamond-like carbon (DLC) film has many attractive properties of high hardness, optical transparency, chemical inertness, low friction and high wear resistance, a limited number of applications was implemented due to high residual compressive stress up to 10 GPa and weak film–substrate bonding especially with ferrous substrates w1x. Furthermore, environmental dependence of the tribological properties also limits the applications of pure DLC films w2x. Various third elements including Si, N, B and transition metals were considered to be an alloying element to remedy these disadvantages without losing favorable properties. Silicon, N and Ti were the most common elements used for third element addition w3–9x. It is well known that Si addition to the DLC films can improve their friction *Corresponding author. Tel.: q82-2-958-5494; fax: q82-2-9585509. E-mail address: [email protected] (K.-R. Lee).
behaviors against steel in ambient environment, while losing their wear resistance w5,10x. Among the transition metals, Ti addition effects on the structure and properties of DLC films were investigated in a wide range of deposition conditions w8,9,11,12x. It was observed that in most cases of transition metal incorporation, its carbide particles of various sizes were distributed in the amorphous carbon matrix w11,12x. Recently, Meng et al. reported the formation of nanocomposite films composed of TiC particles in hydrogenated carbon matrix when Ti concentration in the film was higher than 5.5 at.% w12x. Nanostructured films such as nanocomposite or nano-scale multilayer have drawn much attention owing to the possibility of excellent mechanical and thermal properties, which cannot be expected from those of each phase w13,14x. Completely new properties were reported in the nanocrystalline materials with a grain size of approximately 10 nm or less. In this range of grain size, the number of atoms in the grain is comparable to that in the boundary region. Hence, the properties of these materials are determined
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Fig. 1. Schematic of hybrid deposition system used in the present work.
mainly by the behaviors of the grain boundaries. Superhard properties of nanocomposite materials were discussed in terms of prohibited dislocation formation w15x or nanocrystalline subatomic structures w16x. In the case of nano-scale multilayers, the interlayer interface could play a similar role to that of the grain boundary in the nanocomposite materials. Investigations on the relationship between the microstructure and the physical properties would be thus a prerequisite for developing nanostructured thin films. In the present work, hydrogenated WC–C nanocomposite films were deposited by using a hybrid deposition system. We focused on the relationship between their mechanical properties including residual stress and the microstructure of the film in a wide range of W concentration. In the present work, we could identify an important structural factor to determine the mechanical properties of the nanocomposite film. It was observed that the hardness and the residual stress of the film significantly increased when the physical contact occurred between the nano-sized WC particles. On the other hand, when the particles were isolated from each other, the mechanical property of the film was mainly dependent on that of the carbon matrix. 2. Experimental A hybrid deposition system which combined DC magnetron sputtering and a 13.56-MHz r.f.-PACVD was used to deposit the WC–C nanocomposite films. Fig. 1 shows the schematic of the deposition system used in the present work. The substrates were placed on the
water cooled cathode to which an r.f. power was delivered via impedance matching network. During the deposition, the temperature of the sample was estimated to be less than 100 8C. The film was deposited by sputtering a high purity (99.95%) W target using a mixture of Ar and CH4 as the sputtering gas. The fraction of CH4 in the sputtering gas was varied from 33 to 58% to change the W concentration in the deposited film. A constant power density of 4 Wycm2 was used for the sputtering; 13.56 MHz of r.f. power was applied to the substrate during film deposition, which generated a self bias voltage of y150 V. The base pressure of the reaction chamber was less than 10y3 Pa. The pressure during deposition was kept at 1.33 Pa. A Si (100) wafer was used for the substrate. Thin (100"10 mm thick) Si strips of size 5=50 mm2 were also used for residual stress measurement. Before deposition, substrates were cleaned by an r.f. glow discharge of Ar at a bias voltage of y400 V at a pressure of 0.5 Pa for 15 min. Deposition rate increased from 8.5 to 12.6 nmymin as the CH4 fraction in the sputtering gas increased. Films of thickness 270"30 nm were deposited in the present work by adjusting the deposition time. The concentration of W in the film was measured by Rutherford backscattering spectrometry using a collimated 4He2q beam of 2.0 MeV. Because CH4 was used for carbon deposition, hydrogen should be incorporated in the film. In the present work, however, we reported the fraction of W in W and C without considering hydrogen. Film thickness was measured by an Alphastep profilometer using a step formed by masking the substrate during the deposition. All of the film–substrate composites were convex, showing that the residual stress is compressive. The curvature of the strip sample was measured by a laser reflection method. The residual stress of the film was then calculated from the equilibrium equation for a bending plate w17x. The hardness of the film was measured by nanoindentation in a continuous stiffness measurement (CSM) mode. Because of the substrate effect, the measured hardness varied with penetration depth. Hardness values at a penetration depth of 50 nm were reported in the present paper. Although these values could be different from the genuine hardness of the film, they could characterize the hardness if the values were compared between films of identical thickness deposited on the same substrate. The structure of the film was analyzed by employing transmission electron microscopy (TEM) and X-ray diffraction (XRD). XRD was measured by using Cu Ka X-ray radiation. The samples were tilted by 108 to exclude Si diffraction peaks. The TEM plan-view sample was prepared by a typical procedure composed of grinding, dimpling and ion milling. A cold stage was used during the ion milling process to suppress any change in microstructure. The sample was observed under a TEM
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Fig. 2. W concentration in the film for various CH4 fractions in the sputtering gas.
Fig. 3. Dependence of the residual compressive stress of the film on W concentration.
(Hitachi H-9000 NAR) at the electron acceleration voltage of 300 KV. Macro Raman spectra were obtained using Ar ion laser at a resolution of 1 cmy1. The electrical resistivity of the film was measured by a fourpoint probe system.
a-C:H film (2.4 GPa). This result agrees with the previous observation that the carbon films deposited by sputtering have a graphitic bond structure due to the low energy of the sputtered atoms w18,19x. The residual stress varied in two different ways depending on the W concentration. When the W concentration was smaller than 13 at.%, the residual stress was similar to that of the sputtered pure carbon film. However, beyond 13
3. Results and discussion Fig. 2 summarized the W concentration of the film for various fractions of CH4 in the sputtering gas. W concentration of the film deposited without r.f. bias was also presented for comparison. W concentration was linearly proportional to the Ar fraction in the sputtering gas. As the Ar fraction increased from 0.4 to 0.67, W concentration of the film increased from 5.2 to 42 at.%. This result shows that the W concentration in the film can be systematically controlled by changing the CH4 fraction in the sputtering gas. Composition of the film was independent of the applied bias voltage. Even if the substrate bias can change the energy of the ionized sputtered atoms, its effect on the chemical composition would be negligible in the present experimental condition because of the low ionization ratio of the sputtered atom. Fig. 3 shows the dependence of residual stress on the W concentration. Two values of the residual stress were presented for pure carbon film: one for the hydrogenated amorphous carbon (a-C:H) film deposited by CH4 glow discharge (solid triangle); and the other for the films deposited by sputtering from a highly poisoned W target with carbon (solid circle). The sputtered carbon film had much smaller residual stress (0.4 GPa) than that of
Fig. 4. Dependence of the measured hardness on W concentration.
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Fig. 5. XRD spectra of deposited film for various W concentrations.
at.% of W, the residual stress monotonically increased with increasing W concentration. At a W concentration of 42 at.%, the residual stress was 2.7 GPa. The hardness of the film was summarized in Fig. 4. As in the residual
stress, the hardness of pure carbon film deposited by sputtering was much lower than that of a-C:H film. In the range of W concentration from 0 to 9.5 at.%, the hardness slightly decreased from 10 to 8.4 GPa. However, when W concentration was larger than 10 at.%, the hardness of the film increased from 8.4 to 18.3 GPa as the W concentration increased to 42 at.%. The most interesting behavior observed in the present work would be that the hardness and residual stress of the film significantly increased only when the W concentration was larger than approximately 10 at.%. The XRD spectrum and TEM microstructure of the film showed that the film is composed of nano-sized tungsten carbide particles and an amorphous matrix. Fig. 5 shows the XRD spectra for various W concentrations. The intensity was normalized to the film thickness and shifted upward for ease of comparison. The number on the spectrum is the W concentration of the film in atomic percent. No significant diffraction peak was observed when the W concentration was 5.2 at.%. However, the diffraction peak at a diffraction angle of approximately 37.58 started to appear at a W concentration of 6.9 at.%. The diffraction peak corresponds to the (111) plane of NaCl type b-WC1yx phase w20x. (TEM electron diffraction of Fig. 6 also confirmed the formation of a NaCl type b carbide phase.) The shift of the diffraction peak from 378 might occur due to the residual
Fig. 6. TEM plan view and electron diffraction pattern of the films. (a) 6.9, (b) 14.5 and (c) 26 at.% W concentration.
S.J. Park et al. / Diamond and Related Materials 11 (2002) 1747–1752
stress of the nano-sized carbide particles. The intensity of the diffraction peak increased with W concentration, showing that the W in the film preferred the formation of the carbide phase. The position of the peak also moved to 378. The full width at half maximum (FWHM) of Lorentzian fitting decreased from 9.538 to 6.058 as the W concentration increased from 6.9 to 42 at.%. The value of FWHM was intimately related with the crystallinity of the carbide phase. This spectral change thus showed that the crystallinity of the carbide phase also increased with increasing W concentration. Fig. 6 shows the TEM microstructures and corresponding electron diffraction patterns of the deposited films. As can be seen in Fig. 6a, small crystalline carbide particles of diameter less than 5 nm were uniformly distributed in the amorphous matrix when the W concentration was 6.9 at.%. TEM electron diffraction exhibited only a diffuse ring of WC1yx (111) plane, which reflects the poor crystallinity of the carbide phase. (Because the beam size for the electron diffraction was much larger than the carbide particles, a diffraction ring was observed rather than the diffraction spots.) However, in the sample of W concentration equal to 14.5 at.% (Fig. 6b), the diffraction ring of WC1yx (111) became sharper and the WC1yx (220) diffraction ring could be observed. (A very sharp ring of diameter smaller than that of WC1yx (111) is presumably due to the deep focusing of aperture.) This change shows that the crystal size and crystallinity of the WC1yx particles increased with W concentration. Fig. 6c shows that further increasing W concentration to 26 at.% resulted in highly crystalline WC phase formation as can be judged from two clear diffraction rings. Because of the overlap of particle images in the projected TEM microstructure, the change in particle size could not be clearly observed. However, the sharper diffraction rings show that the increased number of carbide particles of high crystallinity exists in the films of higher W concentration. Figs. 5 and 6 show that the films were nanocomposite of WC1yx particles embedded in an amorphous carbon matrix. The number density and crystallinity of the carbide particles increased with W concentration. The change in residual stress and hardness observed in the present work could be understood in terms of the interlinks between the carbide particles. When the W concentration was low, the carbide particles might be isolated in the amorphous carbon matrix. In this case, the mechanical properties of the film could not be related with those of the carbide particles, but determined by those of the amorphous carbon matrix. This explanation agrees with the experimental observation: when the W concentration was less than approximately 10 at.%, the mechanical properties of the film were essentially the same as those of the sputtered carbon film (see Figs. 3 and 4). However, at higher W
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Fig. 7. Resistivity of the film for various W concentrations.
concentrations where physical contact between the carbide particles could occur, the mechanical properties of the film were affected by those of the carbide particles. As the content of the carbide particles increased, a larger number of carbide particles could contact each other resulting in higher hardness and residual stress. The microstructure of Fig. 6c shows the increased possibility of contact between the carbide particles at higher W concentration. The electrical resistivity of the film definitely showed that physical contact between the carbide particles occurred when the W concentration was larger than 13 at.%. Fig. 7 shows that electrical resistivity varied with W concentration. The resistivity gradually decreased from 7.9 to 2.31 Vcm with increasing W concentration from 5.2 to 9.5 at.%. The gradual change in resistivity could be due to changes in the matrix phase, as will be discussed in the Raman spectrum analysis. However, the resistivity suddenly decreased by approximately two orders of magnitude at a W concentration of approximately 13 at.%. In the range of W concentration from 14.5 to 42 at.%, a gradual decrease in the resistivity was again observed. The sudden decrease in resistivity could be possible only when the efficient conducting paths resulted from the contact of the WC1yx particles. Fig. 8 shows the G-peak position of carbon Raman spectra for various W concentrations. It is empirically known that the G peak position of Raman spectra shifts to a higher wavenumber as the graphitic component in the film increases w21x. Since the WC1yx phase has centrosymmetric crystal structure of B1-NaCl type which is inactive to Raman excitation w22x, the Raman spectra reflect the atomic bond structure of the carbon
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be varied from 5.2 to 42 at.% by changing the CH4 concentration of the mixture sputtering gas of Ar and CH4. In this concentration range, the films were composed of b-WC1yx particles of a few nanometers in diameter in a continuous amorphous carbon matrix. The present work showed that the physical contact between the carbide particles is an important structural factor that determines the physical properties of the nanocomposite films. The hardness and residual compressive stress of the film increased only when the W concentration was higher than 13 at.%. Structural analysis and electrical conductance measurements evidently showed that the increase in hardness and residual compressive stress occurred as the carbide particles were in contact with each other in the amorphous carbon matrix. Acknowledgments
Fig. 8. Changes in Raman G-peak position with the W concentration in the deposited film.
matrix only. (For a similar reason, we could not obtain a carbon Raman peak when the W concentration was larger than 18 at.%.) The G-peak position shifted from 1532 to 1544 cmy1 with increasing W concentration, which shows the increase in graphitic component. The high momentum of a sputtered W particle seems to cause structural relaxation of the carbon matrix as in aC:H films deposited at very high ion energy w23x. The gradual decrease in the resistivity of Fig. 7 could be due to the higher fraction of the graphitic phase in the amorphous carbon matrix. It is thus expected that the mechanical properties of the amorphous carbon matrix degraded with W incorporation. By comparing data of Figs. 4 and 8, the major structural factor that determined the mechanical property of the film becomes evident. When the W concentration was larger than 13 at.%, the mechanical properties of the film were improved by increasing the physical contact of the carbide particles although the mechanical properties of the amorphous carbon matrix were degraded by the W incorporation. On the other hand, when the W concentration was smaller than 10 at.%, the mechanical properties of the film were dominated by those of the amorphous matrix since the carbide particles were isolated throughout the matrix. The decrease in hardness with W concentration in this concentration range could be well understood by this structural model. 4. Conclusions A hybrid deposition system of r.f.-PACVD and DC magnetron sputtering was used to deposit WC–C nanocomposite films. The W concentration in the film could
The authors gratefully acknowledge the financial support from the Ministry of Science and Technology of Korea and the Korea Science and Engineering Foundation through the Center for Advanced Plasma Surface Engineering at the Sungkyunkwan University. References w1x H. Tsai, Mater. Sci. Forum. 52–53 (1989) 71. w2x K. Miyoshi, R.L.C. Wu, A. Garscadden, Surf. Coat. Technol. 54–55 (1992) 428. w3x C. Donnet, Surf. Coat. Technol. 100–101 (1998) 180. w4x J. Meneve, E. Dekempeneer, J. Smeets, Diamond Films Technol. 4 (1994) 23. w5x M.-G. Kim, K.-R. Lee, K.Y. Eun, Surf. Coat. Technol. 112 (1999) 204. w6x T.W. Scharf, R.D. Ott, D. Yang, J.A. Barnard, J. Appl. Phys. 85 (1999) 6. w7x K.-R. Lee, K.Y. Eun, J.-S. Rhee, Mater. Res. Soc. Symp. 356 (1995) 233. w8x V.Y. Kulikovsky, F. Fendrych, L. Jastrabik, et al., Surf. Coat. Technol. 102 (1998) 81. w9x Y.L. Su, W.H. Kao, Wear 236 (1999) 221. w10x W.-J. Wu, T.-M. Pai, M.-H. Hon, Diamond Rel. Mater. 7 (1998) 1478. w11x T. Zehnder, J. Patscheider, Surf. Coat. Technol. 133–134 (2000) 138. w12x W.J. Meng, T.J. Curtis, L.E. Rehn, P.M. Baldo, J. Appl. Phys. 83 (1998) 6076. w13x S. Veprek, P. Nesladek, A. Niederhofer, F. Glatz, M. Jilek, M. Sima, Surf. Coat. Technol. 108–109 (1998) 138. w14x J. Mussil, Surf. Coat. Technol. 125 (2000) 322. w15x S. Veprek, S. Reiprich, Thin Solid Films 268 (1995) 64. w16x P.Yu. Butyagin, Colloid J. 59 (1997) 112. w17x A. Brenner, S. Senderoff, J. Res. Natl. Std. 42 (1949) 105. w18x J. Robertson, E.P. O’Reilly, Phys. Rev. B 35 (1987) 2946. w19x J. Robertson, Surf. Coat. Technol. 50 (1992) 185. w20x X-Ray Powder Diffraction File no. 20-1316 (JCPDS-ICD1), 1994. w21x M.A. Tamor, W.C. Vassell, J. Appl. Phys. 76 (1994) 3823. w22x H. Poulet, J.P. Mathieu, Vibration Spectra and Symmetry of Crystals, Gordon and Breach, New York, 1976. w23x N.-H. Cho, K.M. Krishnan, D.K. Veirs, et al., J. Mater. Res. 5 (1990) 2543.
Microstructure and mechanical properties of WC–C nanocomposite films Se Jun Parka,b, Kwang-Ryeol Leea,*, Dae-Hong Kob, Kwang Yong Euna a
Future Technology Research Division, Korea Institute of Science and Technology, P.O. Box 131, Cheongryang, Seoul, 130-650, South Korea b Department of Ceramic Engineering, Yonsei University, Seoul, 120-701, South Korea Received 5 February 2002; received in revised form 8 May 2002; accepted 13 May 2002
Abstract WC–C nanocomposite film was prepared by using a hybrid deposition system of r.f.-PACVD and DC magnetron sputtering. W concentration in the film was varied from 5.2 to 42 at.% by changing the CH4 fraction of the mixture sputtering gas of Ar and CH4. Hardness, residual compressive stress and electrical resistivity were characterized as a function of W concentration. Raman spectroscopy, XRD and high resolution TEM were employed to analyze the structural change in the film for various W concentrations. In the present W concentration range, the film was composed of nano-sized WC particles of diameter less than 5 nm and hydrogenated amorphous carbon matrix. Content of the WC particles increased with increasing W concentration. However, the mechanical properties of the film increased only when the W concentration was higher than 13 at.%. Structural analysis and electrical conductance measurements evidently showed that the increase in hardness and residual stress occurred as the WC particles were in contact with each other in the amorphous carbon matrix. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: WC–C nanocomposite films; Hybrid r.f.-PACVD and magnetron sputtering; Microstructure evolution; Mechanical and electrical properties
1. Introduction Although diamond-like carbon (DLC) film has many attractive properties of high hardness, optical transparency, chemical inertness, low friction and high wear resistance, a limited number of applications was implemented due to high residual compressive stress up to 10 GPa and weak film–substrate bonding especially with ferrous substrates w1x. Furthermore, environmental dependence of the tribological properties also limits the applications of pure DLC films w2x. Various third elements including Si, N, B and transition metals were considered to be an alloying element to remedy these disadvantages without losing favorable properties. Silicon, N and Ti were the most common elements used for third element addition w3–9x. It is well known that Si addition to the DLC films can improve their friction *Corresponding author. Tel.: q82-2-958-5494; fax: q82-2-9585509. E-mail address: [email protected] (K.-R. Lee).
behaviors against steel in ambient environment, while losing their wear resistance w5,10x. Among the transition metals, Ti addition effects on the structure and properties of DLC films were investigated in a wide range of deposition conditions w8,9,11,12x. It was observed that in most cases of transition metal incorporation, its carbide particles of various sizes were distributed in the amorphous carbon matrix w11,12x. Recently, Meng et al. reported the formation of nanocomposite films composed of TiC particles in hydrogenated carbon matrix when Ti concentration in the film was higher than 5.5 at.% w12x. Nanostructured films such as nanocomposite or nano-scale multilayer have drawn much attention owing to the possibility of excellent mechanical and thermal properties, which cannot be expected from those of each phase w13,14x. Completely new properties were reported in the nanocrystalline materials with a grain size of approximately 10 nm or less. In this range of grain size, the number of atoms in the grain is comparable to that in the boundary region. Hence, the properties of these materials are determined
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Fig. 1. Schematic of hybrid deposition system used in the present work.
mainly by the behaviors of the grain boundaries. Superhard properties of nanocomposite materials were discussed in terms of prohibited dislocation formation w15x or nanocrystalline subatomic structures w16x. In the case of nano-scale multilayers, the interlayer interface could play a similar role to that of the grain boundary in the nanocomposite materials. Investigations on the relationship between the microstructure and the physical properties would be thus a prerequisite for developing nanostructured thin films. In the present work, hydrogenated WC–C nanocomposite films were deposited by using a hybrid deposition system. We focused on the relationship between their mechanical properties including residual stress and the microstructure of the film in a wide range of W concentration. In the present work, we could identify an important structural factor to determine the mechanical properties of the nanocomposite film. It was observed that the hardness and the residual stress of the film significantly increased when the physical contact occurred between the nano-sized WC particles. On the other hand, when the particles were isolated from each other, the mechanical property of the film was mainly dependent on that of the carbon matrix. 2. Experimental A hybrid deposition system which combined DC magnetron sputtering and a 13.56-MHz r.f.-PACVD was used to deposit the WC–C nanocomposite films. Fig. 1 shows the schematic of the deposition system used in the present work. The substrates were placed on the
water cooled cathode to which an r.f. power was delivered via impedance matching network. During the deposition, the temperature of the sample was estimated to be less than 100 8C. The film was deposited by sputtering a high purity (99.95%) W target using a mixture of Ar and CH4 as the sputtering gas. The fraction of CH4 in the sputtering gas was varied from 33 to 58% to change the W concentration in the deposited film. A constant power density of 4 Wycm2 was used for the sputtering; 13.56 MHz of r.f. power was applied to the substrate during film deposition, which generated a self bias voltage of y150 V. The base pressure of the reaction chamber was less than 10y3 Pa. The pressure during deposition was kept at 1.33 Pa. A Si (100) wafer was used for the substrate. Thin (100"10 mm thick) Si strips of size 5=50 mm2 were also used for residual stress measurement. Before deposition, substrates were cleaned by an r.f. glow discharge of Ar at a bias voltage of y400 V at a pressure of 0.5 Pa for 15 min. Deposition rate increased from 8.5 to 12.6 nmymin as the CH4 fraction in the sputtering gas increased. Films of thickness 270"30 nm were deposited in the present work by adjusting the deposition time. The concentration of W in the film was measured by Rutherford backscattering spectrometry using a collimated 4He2q beam of 2.0 MeV. Because CH4 was used for carbon deposition, hydrogen should be incorporated in the film. In the present work, however, we reported the fraction of W in W and C without considering hydrogen. Film thickness was measured by an Alphastep profilometer using a step formed by masking the substrate during the deposition. All of the film–substrate composites were convex, showing that the residual stress is compressive. The curvature of the strip sample was measured by a laser reflection method. The residual stress of the film was then calculated from the equilibrium equation for a bending plate w17x. The hardness of the film was measured by nanoindentation in a continuous stiffness measurement (CSM) mode. Because of the substrate effect, the measured hardness varied with penetration depth. Hardness values at a penetration depth of 50 nm were reported in the present paper. Although these values could be different from the genuine hardness of the film, they could characterize the hardness if the values were compared between films of identical thickness deposited on the same substrate. The structure of the film was analyzed by employing transmission electron microscopy (TEM) and X-ray diffraction (XRD). XRD was measured by using Cu Ka X-ray radiation. The samples were tilted by 108 to exclude Si diffraction peaks. The TEM plan-view sample was prepared by a typical procedure composed of grinding, dimpling and ion milling. A cold stage was used during the ion milling process to suppress any change in microstructure. The sample was observed under a TEM
S.J. Park et al. / Diamond and Related Materials 11 (2002) 1747–1752
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Fig. 2. W concentration in the film for various CH4 fractions in the sputtering gas.
Fig. 3. Dependence of the residual compressive stress of the film on W concentration.
(Hitachi H-9000 NAR) at the electron acceleration voltage of 300 KV. Macro Raman spectra were obtained using Ar ion laser at a resolution of 1 cmy1. The electrical resistivity of the film was measured by a fourpoint probe system.
a-C:H film (2.4 GPa). This result agrees with the previous observation that the carbon films deposited by sputtering have a graphitic bond structure due to the low energy of the sputtered atoms w18,19x. The residual stress varied in two different ways depending on the W concentration. When the W concentration was smaller than 13 at.%, the residual stress was similar to that of the sputtered pure carbon film. However, beyond 13
3. Results and discussion Fig. 2 summarized the W concentration of the film for various fractions of CH4 in the sputtering gas. W concentration of the film deposited without r.f. bias was also presented for comparison. W concentration was linearly proportional to the Ar fraction in the sputtering gas. As the Ar fraction increased from 0.4 to 0.67, W concentration of the film increased from 5.2 to 42 at.%. This result shows that the W concentration in the film can be systematically controlled by changing the CH4 fraction in the sputtering gas. Composition of the film was independent of the applied bias voltage. Even if the substrate bias can change the energy of the ionized sputtered atoms, its effect on the chemical composition would be negligible in the present experimental condition because of the low ionization ratio of the sputtered atom. Fig. 3 shows the dependence of residual stress on the W concentration. Two values of the residual stress were presented for pure carbon film: one for the hydrogenated amorphous carbon (a-C:H) film deposited by CH4 glow discharge (solid triangle); and the other for the films deposited by sputtering from a highly poisoned W target with carbon (solid circle). The sputtered carbon film had much smaller residual stress (0.4 GPa) than that of
Fig. 4. Dependence of the measured hardness on W concentration.
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Fig. 5. XRD spectra of deposited film for various W concentrations.
at.% of W, the residual stress monotonically increased with increasing W concentration. At a W concentration of 42 at.%, the residual stress was 2.7 GPa. The hardness of the film was summarized in Fig. 4. As in the residual
stress, the hardness of pure carbon film deposited by sputtering was much lower than that of a-C:H film. In the range of W concentration from 0 to 9.5 at.%, the hardness slightly decreased from 10 to 8.4 GPa. However, when W concentration was larger than 10 at.%, the hardness of the film increased from 8.4 to 18.3 GPa as the W concentration increased to 42 at.%. The most interesting behavior observed in the present work would be that the hardness and residual stress of the film significantly increased only when the W concentration was larger than approximately 10 at.%. The XRD spectrum and TEM microstructure of the film showed that the film is composed of nano-sized tungsten carbide particles and an amorphous matrix. Fig. 5 shows the XRD spectra for various W concentrations. The intensity was normalized to the film thickness and shifted upward for ease of comparison. The number on the spectrum is the W concentration of the film in atomic percent. No significant diffraction peak was observed when the W concentration was 5.2 at.%. However, the diffraction peak at a diffraction angle of approximately 37.58 started to appear at a W concentration of 6.9 at.%. The diffraction peak corresponds to the (111) plane of NaCl type b-WC1yx phase w20x. (TEM electron diffraction of Fig. 6 also confirmed the formation of a NaCl type b carbide phase.) The shift of the diffraction peak from 378 might occur due to the residual
Fig. 6. TEM plan view and electron diffraction pattern of the films. (a) 6.9, (b) 14.5 and (c) 26 at.% W concentration.
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stress of the nano-sized carbide particles. The intensity of the diffraction peak increased with W concentration, showing that the W in the film preferred the formation of the carbide phase. The position of the peak also moved to 378. The full width at half maximum (FWHM) of Lorentzian fitting decreased from 9.538 to 6.058 as the W concentration increased from 6.9 to 42 at.%. The value of FWHM was intimately related with the crystallinity of the carbide phase. This spectral change thus showed that the crystallinity of the carbide phase also increased with increasing W concentration. Fig. 6 shows the TEM microstructures and corresponding electron diffraction patterns of the deposited films. As can be seen in Fig. 6a, small crystalline carbide particles of diameter less than 5 nm were uniformly distributed in the amorphous matrix when the W concentration was 6.9 at.%. TEM electron diffraction exhibited only a diffuse ring of WC1yx (111) plane, which reflects the poor crystallinity of the carbide phase. (Because the beam size for the electron diffraction was much larger than the carbide particles, a diffraction ring was observed rather than the diffraction spots.) However, in the sample of W concentration equal to 14.5 at.% (Fig. 6b), the diffraction ring of WC1yx (111) became sharper and the WC1yx (220) diffraction ring could be observed. (A very sharp ring of diameter smaller than that of WC1yx (111) is presumably due to the deep focusing of aperture.) This change shows that the crystal size and crystallinity of the WC1yx particles increased with W concentration. Fig. 6c shows that further increasing W concentration to 26 at.% resulted in highly crystalline WC phase formation as can be judged from two clear diffraction rings. Because of the overlap of particle images in the projected TEM microstructure, the change in particle size could not be clearly observed. However, the sharper diffraction rings show that the increased number of carbide particles of high crystallinity exists in the films of higher W concentration. Figs. 5 and 6 show that the films were nanocomposite of WC1yx particles embedded in an amorphous carbon matrix. The number density and crystallinity of the carbide particles increased with W concentration. The change in residual stress and hardness observed in the present work could be understood in terms of the interlinks between the carbide particles. When the W concentration was low, the carbide particles might be isolated in the amorphous carbon matrix. In this case, the mechanical properties of the film could not be related with those of the carbide particles, but determined by those of the amorphous carbon matrix. This explanation agrees with the experimental observation: when the W concentration was less than approximately 10 at.%, the mechanical properties of the film were essentially the same as those of the sputtered carbon film (see Figs. 3 and 4). However, at higher W
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Fig. 7. Resistivity of the film for various W concentrations.
concentrations where physical contact between the carbide particles could occur, the mechanical properties of the film were affected by those of the carbide particles. As the content of the carbide particles increased, a larger number of carbide particles could contact each other resulting in higher hardness and residual stress. The microstructure of Fig. 6c shows the increased possibility of contact between the carbide particles at higher W concentration. The electrical resistivity of the film definitely showed that physical contact between the carbide particles occurred when the W concentration was larger than 13 at.%. Fig. 7 shows that electrical resistivity varied with W concentration. The resistivity gradually decreased from 7.9 to 2.31 Vcm with increasing W concentration from 5.2 to 9.5 at.%. The gradual change in resistivity could be due to changes in the matrix phase, as will be discussed in the Raman spectrum analysis. However, the resistivity suddenly decreased by approximately two orders of magnitude at a W concentration of approximately 13 at.%. In the range of W concentration from 14.5 to 42 at.%, a gradual decrease in the resistivity was again observed. The sudden decrease in resistivity could be possible only when the efficient conducting paths resulted from the contact of the WC1yx particles. Fig. 8 shows the G-peak position of carbon Raman spectra for various W concentrations. It is empirically known that the G peak position of Raman spectra shifts to a higher wavenumber as the graphitic component in the film increases w21x. Since the WC1yx phase has centrosymmetric crystal structure of B1-NaCl type which is inactive to Raman excitation w22x, the Raman spectra reflect the atomic bond structure of the carbon
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be varied from 5.2 to 42 at.% by changing the CH4 concentration of the mixture sputtering gas of Ar and CH4. In this concentration range, the films were composed of b-WC1yx particles of a few nanometers in diameter in a continuous amorphous carbon matrix. The present work showed that the physical contact between the carbide particles is an important structural factor that determines the physical properties of the nanocomposite films. The hardness and residual compressive stress of the film increased only when the W concentration was higher than 13 at.%. Structural analysis and electrical conductance measurements evidently showed that the increase in hardness and residual compressive stress occurred as the carbide particles were in contact with each other in the amorphous carbon matrix. Acknowledgments
Fig. 8. Changes in Raman G-peak position with the W concentration in the deposited film.
matrix only. (For a similar reason, we could not obtain a carbon Raman peak when the W concentration was larger than 18 at.%.) The G-peak position shifted from 1532 to 1544 cmy1 with increasing W concentration, which shows the increase in graphitic component. The high momentum of a sputtered W particle seems to cause structural relaxation of the carbon matrix as in aC:H films deposited at very high ion energy w23x. The gradual decrease in the resistivity of Fig. 7 could be due to the higher fraction of the graphitic phase in the amorphous carbon matrix. It is thus expected that the mechanical properties of the amorphous carbon matrix degraded with W incorporation. By comparing data of Figs. 4 and 8, the major structural factor that determined the mechanical property of the film becomes evident. When the W concentration was larger than 13 at.%, the mechanical properties of the film were improved by increasing the physical contact of the carbide particles although the mechanical properties of the amorphous carbon matrix were degraded by the W incorporation. On the other hand, when the W concentration was smaller than 10 at.%, the mechanical properties of the film were dominated by those of the amorphous matrix since the carbide particles were isolated throughout the matrix. The decrease in hardness with W concentration in this concentration range could be well understood by this structural model. 4. Conclusions A hybrid deposition system of r.f.-PACVD and DC magnetron sputtering was used to deposit WC–C nanocomposite films. The W concentration in the film could
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