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Preferential orientation of titanium carbide films deposited by a filtered cathodic vacuum arc technique
Surface and Coatings Technology 138 Ž2001. 301᎐306
Preferential orientation of titanium carbide films deposited by a filtered cathodic vacuum arc technique Xing-zhao Ding a,U , B.K. Tay a , H.S. Tana , S.P. Lau a , W.Y. Cheung b, S.P. Wong b a b
School of Electrical and Electronic Engineering, Nanyang Technological Uni¨ ersity, Nanyang, 639798, Singapore Department of Electronic Engineering, The Chinese Uni¨ ersity of Hong Kong, Shatin, New Territories, Hong Kong Received 6 November 2000; received in revised form 10 November 2000; accepted 10 November 2000
Abstract Titanium carbide films with a thickness of approximately 100-nm were deposited on SiŽ100. substrates by a filtered cathodic vacuum arc technique. The composition and microstructure of the films were assessed by Rutherford backscattering spectroscopy, X-ray photoelectron spectroscopy, X-ray diffraction, and atomic force microscopy. A negative bias voltage Ž VS ⫽0 ; y1000 V. was applied to the substrate during deposition, and the influence of VS on the crystalline orientation of the as-deposited films was investigated. It was found that the crystallites are randomly oriented in the film deposited at VS s 0 V. In the bias voltage range of VS s y40 ; y500 V, the titanium carbide films exhibited a Ž111. preferential orientation. When VS was increased to y1000 V, however, the film was Ž100. preferentially oriented. The compressive internal stress, determined by the radius of curvature technique, in the titanium carbide films exhibited a minimum value at approximately VS s y80 ; y120 V. The Ž111. preferential orientation can be explained by minimization of elastic energy storage in the films; while the Ž100. preferential orientation in the film deposited at VS s y1000 V is due to the sputter channeling effect, because the Ž100. direction in the TiC lattice shows the most open channeling direction and therefore the lowest sputtering yield. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Preferential orientation; Titanium carbide film; Filtered cathodic vacuum arc deposition
1. Introduction Titanium carbide ŽTiC. is one of the most widely applied hard coating materials. Its chemical bonding is quite special, similar to other interstitial carbides or nitrides, showing a mixed covalent, metallic and ionic character w1x, which results in many advantageous properties of this material, including high hardness, low coefficient of friction, outstanding wear resistance, high melting or decomposition temperature, good electrical conductivity, excellent corrosion resistance and chemical stability. These unique combined properties have made titanium carbide of particular interest in a wide variety of applications. Because of its hardness and wear resistance, TiC has become a major industrial coating material in many tribological applications w1,2x. U
Corresponding author. Tel.: q65-7905454; fax: q65-7933318. E-mail address: [email protected] ŽX. Ding..
Its high electrical conductivity and low work function make it an outstanding candidate for Ohmic contact of some specific SiC-based electronic devices w3,4x and for use as an electron injection electrode in organic light emitting devices ŽOLEDs. w5x. In addition, TiC is presently attracting great attention as a catalyst w6x, a material for the first wall of nuclear reactors w7x, and as diffusion barriers in semiconductor technology w8x. Titanium carbide films have been prepared by a variety of deposition techniques including chemical vapor deposition ŽCVD. w9,10x, sputtering w11,12x, evaporation w13x, and pulsed laser ablation ŽPLD. w5,14x. In the conventional or plasma assisted CVD process, the deposition is carried out in a high substrate temperature above 500⬚C w1x, which often cannot be tolerated by the substrates with a poor thermal resistance in many important technological applications. Physical vapor deposition ŽPVD. processes, Žvarious evaporation and sputtering techniques ., on the other hand, allow
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X. Ding et al. r Surface and Coatings Technology 138 (2001) 301᎐306
coating deposition at lower temperatures than CVD methods. However, lower substrate temperature leads to a lower degree of adhesion between the substrate and the coating, which often results in poorer mechanical properties and delamination of the coatings. In recent years, PLD has been demonstrated to be a suitable technique to deposit TiC films at room temperature. However, it is well known that the generation and inclusion of particulates andror droplets from the target during laser evaporation has been one of the biggest problems of PLD w15x. In the present work, a filtered cathodic vacuum arc ŽFCVA. technique was employed to deposit the titanium carbide films. FCVA deposition is a promising technique for the production of high quality hard thin films w16,17x. The main feature of the FCVA technique is to employ a curved magnetic field to guide the plasma generated from the cathodic vacuum arc to deposit on an out-of-sight substrate. Through this special magnetic filter, most of the unwanted macroparticles and neutral atoms will be removed. Only ions within a defined energy range can reach the substrate, thus producing films with good controllability and reproducibility. Moreover, as the depositing species in the FCVA process are fully ionized, the kinetic energy of these ions can be precisely controlled by adjusting the bias voltage applied to the substrate. It is well known that the low-energy ion bombardment during deposition of thin films may markedly affect the crystal orientation in the films. Often a fibrous texture is generated in which all film crystallites share a crystallographic axis oriented perpendicular to the plane of the substrate. The crystal orientation or texture in thin films has been found to greatly influence the physical properties of materials and consequently their performance in practical applications w18᎐20x. Therefore, controlling the texture of crystalline thin films is important both from the fundamental physics and technological points of view. In our previous paper w21x, some titanium-containing tetrahedral amorphous carbon films with different titanium contents were deposited by a filtered cathodic vacuum arc technique by using compacted titaniumrcarbon mixture cathode targets. It was demonstrated that by using the target with 20 at.% titanium, sub-stoichiometric titanium carbide films could be deposited. In this paper, the evolution of preferential orientation of these titanium carbide films with different negative bias voltage was investigated.
2. Experimental The titanium carbide films were deposited by the FCVA technique. At first, a TirC mixture target with atomic ratio of TirCs 1:4 was prepared. Pure elemental powders of graphite Ž99.5% purity, 325 mesh. and
titanium Ž99.5% purity, 560 mesh., were mixed thoroughly and then compacted into a cylindrical target 60 mm in diameter under an axial pressure of approximately 160 MPa. The FCVA system has been described in detail elsewhere w17,22x. It incorporates an off-plane double-bend ŽOPDB. filter to effectively remove all macro-particles. The arc current was at approximately 110 A. The torridal magnetic field for steering the carbon plasma towards the substrate was maintained at 40 mT. The working chamber was evacuated to a base pressure below 4 = 10y6 torr but rose to 3 = 10y5 torr during deposition due to out-gasing of the cathode. Single crystal Si Ž100. wafers were used as the substrates. Prior to deposition, the silicon substrates were sputtered by an Arq ion beam Ž800 eV, 60 mA. for 5 min in order to remove the native oxide layer. During deposition, a negative DC bias voltage Ž VS ., ranging from 0 to y1000 V, was applied to the substrate. Correspondingly, the energy Ž E ., of the depositing ions can be approximately expressed as Es neŽ VP y VS . q E0 , where VP f y13 V, is the plasma potential w16x, E0 is the original ion energy in the plasma Ž; 28 eV., e is the electron charge, and n is the charge state of the ions in the vacuum arc plasma. According to the experimental results of Anders w23x, carbon arc plasma consists of only single charged state Cq ions, while the titanium arc plasma is composed of 11% Tiqq 75% Ti 2qq 14% Ti 3q with an average charge state of 2.03. In this work the titanium carbide films were deposited at room temperature. The average deposition rate was about 3 nmrmin. The thickness of the films was simply controlled by deposition time and measured to be approximately 100 nm by using a TENCOR P-10 surface profilometer. The composition and microstructure of the asdeposited titanium carbide films were monitored by Rutherford backscattering spectroscopy ŽRBS., X-ray photoelectron spectroscopy ŽXPS., and X-ray diffraction ŽXRD. analyses. In RBS analyses, a beam of 2 MeV He 2q ions were projected into the film at an incident angle of 7⬚ to the normal line of the film plane, while backscattered He 2q ions were detected at a scattering angle of 170⬚. The XPS analyses were carried out using a VG Scientific Microlab 310F with Mg K ␣ radiation as the X-ray source. The XRD experiments were performed on a Siemens D5005 X-ray diffractometer with a glancing incident mode by using CuK ␣ radiation Ž40 kV, 40 mA.. The incident angle was fixed at 1⬚, the 2 scanning step size was set as 0.05⬚, and the counting time at each step was 4 s. The surface morphology of the as-deposited films was directly observed under ambient conditions by a Nanoscope IIIa scanning atomic force microscope ŽAFM.. The system is operated in the non-contact tapping and lift modes. The internal stress Ž s . in the as-deposited titanium
X. Ding et al. r Surface and Coatings Technology 138 (2001) 301᎐306
303
carbide films was determined according to Stoney’s equation w24x: s s
Es t s2 6 Ž1 y ¨ s . t c
ž
1 1 y R R0
/
Where Es , ¨ s and t s are Young’s Modulus, Poisson Ratio and thickness of the silicon substrate, R 0 and R are the curvature radii of the wafer before and after coating with the TiC films, t c is the thickness of the films. R 0 , R and t c are all measured by using the surface profilometer.
3. Results and discussion Fig. 1 shows a typical RBS spectrum of the as-deposited titanium carbide films. The average titanium content in the films was determined to be approximately 46.2" 1.5 at.%, which is more than twice of that in the cathode target. This phenomenon can be interpreted as follows. Due to the lower melting temperature of Ti, in comparison with that of carbon Žthe melting point of titanium and graphite is ; 1660⬚C and ; 3870⬚C, respectively., the titanium particles in the arc spot area during deposition will be completely sublimated and consequently ionized. In contrast, the graphite particles may only be partially evaporated with the formation of some macro-clusters which will be trapped by the OPDB filter and cannot reach the substrate. In addition, the higher titanium concentration in the films can also be partially attributed to the preferential re-sputtering and back-scattering effects. Because the mass of carbon atoms is much lower than that of titanium atoms, carbon atoms on the film surface would be preferentially sputtered by the energetic depositing ions, especially by the heavier titanium ions with higher energies Žin double- and triple-charged states .. On the other hand, the depositing carbon ions could be more easily back-scattered from the heavier titanium atoms on the film surface. In the RBS spectrum, some oxygen andror nitrogen contamination in the films can be evidently detected. This contamination is closely related to the sample preparation process. When compacting the TirC mixed target under ambient conditions, some air bubbles might inevitably be trapped in the target. During deposition, with the erosion of the target, the air contained inside the target will come out and the oxygen and nitrogen components will immediately react with the melting titanium. The oxygen andror nitrogen contamination may also come from the residual gas in the vacuum chamber. Indeed, it is known that TiC is isomorphous with TiN and TiO w1x. Thus, nitrogen and oxygen may incorporate into the
Fig. 1. A typical RBS spectrum of the as-deposited titanium carbide films.
TiC films as impurities during deposition as a result of interactions with residual gases in the vacuum. Typical C 1s and Ti 2p XPS spectra of the as-deposited titanium carbide film surfaces are shown in Fig. 2. The C 1s spectrum is presented as three components. The peak around 282.0 eV can be assigned to the carbidic carbon component ŽC bonded to Ti. w5x, while both the strong and the weak components at 285.0 eV and 289.0 eV could most likely be attributed to adsorbed hydrocarbon groups w25,26x on the film surface. The Ti 2p XPS spectrum can be deconvoluted into two sets of spin-orbit doublets, each separated by 6 eV, corresponding to the carbide ŽTi 2p 3r2 at 455.0 eV. and the oxide components ŽTi 2p 3r2 at 458.9 eV., respectively w5x. The strong TiO 2 spin-doublet was believed to have mainly resulted from the surface oxidization when the samples were kept in air. The oxygen contamination in the films may also be partially responsible for the presence of the oxidic Ti 2p spin-doublet. The internal stress Ž s ., in the as-deposited titanium carbide films as a function of the negative bias voltage is shown in Fig. 3. It was as expected that the internal stress in all of these films would be compressive in nature. It was found that the s exhibited a minimum value at approximately VS s y80 ; y120 V. Fig. 4 shows three typical AFM images of the asdeposited titanium carbide films with different crystal orientations. It was found that the films deposited at lower negative bias voltages, Že.g. VS s 0 ; y160 V. were quite smooth, while the film deposited at a higher bias voltage, e.g. y1000 V, exhibited an evident columnar surface morphology due to the bombardment of high energy depositing ions, especially Ti 2q and Ti 3q ions. The variation of root-mean-square surface roughness of the titanium carbide films with the applied negative bias voltage is shown in Fig. 5. The variation
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Fig. 2. Typical C 1s and Ti 2p XPS spectra of the as-deposited titanium carbide film surfaces: a and c are Ti 2p 3r 2 and Ti 2p1r2 in titanium oxide; b and d are Ti 2p 3r 2 and Ti 2p1r2 in titanium carbide.
trend of the surface roughness was found to be similar to that of the internal stress in the films, i.e. showing a minimum value at approximately VS s y80 ; y120 V. The XRD patterns of the titanium carbide films are shown in Fig. 6. At VS s 0 V, three evident peaks around 2 f 35.9⬚, 41.7⬚, and 60.6⬚ are observed, which can be indexed as the Ž111., Ž200. and Ž220. diffractions of the cubic TiC crystalline phase with a B1 NaCl-type Fig. 4. Three typical AFM images of the as-deposited titanium carbide films with different orientations: Ža. random oriented film deposited at VS s 0 V; Žb. Ž111. preferentially oriented film deposited at VS s y80 V; and Žc. Ž100. preferentially oriented film deposited at VS s y1000 V.
Fig. 3. Compressive internal stress in the titanium carbide films deposited at different negative bias voltage.
structure. This film exhibited a random orientation character. With the increase of negative bias voltage, the Ž200. and Ž220. peaks weakened gradually and disappeared at VS s y160 V. The films showed an evident Ž111. preferred orientation. This preferential orientation could be retained up to VS s y500 V. When the negative bias voltage was elevated further to y1000 V, the Ž111. diffraction peak disappeared; meanwhile the Ž200. peak emerged again, and the film showed a Ž100. preferred orientation. The formation of texture in polycrystalline thin films
X. Ding et al. r Surface and Coatings Technology 138 (2001) 301᎐306
Fig. 5. Root-mean-square surface roughness of the titanium carbide films deposited at different negative bias voltage.
can be affected by many factors, including the surface energy, elastic strain energy, thermal energy, and so on. In the present work, all of the titanium carbide films were deposited at room temperature, therefore the influence of thermal energy can be neglected. To date, the surface energy in the B1 NaCl-type structured crystals, e.g. TiN and TiC, remains ambiguous. Pelleg et al. w27x showed that the Ž100. planes have the smallest surface energy by neglecting metal-metal ŽM᎐M. bond-
Fig. 6. XRD patterns of the titanium carbide films deposited at different bias voltage: Ža. 0 V, Žb. y40 V, Žc. y80 V, Žd. y120 V, Že. y160 V, Žf. y500 V, Žg. y1000 V.
305
ing. However, Quaeyhaegens et al. w28x showed that Ž111. are planes with the smallest surface energy by considering an fcc metal structure. Therefore, the structure evolution of the titanium carbide films cannot be simply explained by the surface energy of the different faces in the TiC lattice. From Fig. 3, it has been noted that the compressive internal stress in our FCVA deposited titanium carbide films is very high, thus the elastic strain energy in the crystal lattice could be a dominating factor in determining the texture of the films. Recently, Mckenzie et al. w29x demonstrated that the preferential orientation in thin films can come about as a result of the synthesis materials under impressed stress conditions. For the materials with the rocksalt crystal structure, including TiC and TiN, it has been revealed w27,29x that the crystallites with the Ž111. orientation have the smallest elastic energy storage in the case of stressed states. Therefore, the experimentally observed Ž111. preferential orientation in the titanium carbide films deposited at a medium bias voltage, i.e. VS s y40 ; y500 V, can be explained by the concept of minimization of the elastic deformation energy. The minimum value of the measured internal stress at approximately VS s y80 ; y120 V in the as-deposited films, as observed in Fig. 3, reasonably supported this viewpoint. It is well known that FCVA deposition is far from a thermodynamic equilibrium process. At VS s 0 V, the energy of the depositing species is relatively small, and the mobility of the surface atoms is very low. Thus the as-deposited film showed a random orientation because of the lower structural relaxation ability. With the increase of negative bias voltage, the depositing ions can be effectively accelerated, and the energy of the ions is approximately proportional to the bias voltage. Under the bombardment of the energetic depositing ion flux, the mobility of the atoms on the growing film surface can be greatly increased. As a consequence of this, the film structure can be relaxed to a lower energy state with the formation of Ž111. preferential orientation. On the other hand, with the increase of ion energy Žor VS ., the lattice damage caused by the ion bombardment could increase gradually, and, as a consequence, the internal stress could also increase gradually, as observed in Fig. 3 at VS s y120; y1000 V, due to the lattice deformation. Moreover, with the increase of ion energy, the sputtering effect could become more and more significant. In the energy range involved in the present work, the sputter yield increases linearly with ion energy w30x. In a study of ion beam assisted deposition processes, Bradley et al. w31x developed a model to explain the development of preferred orientation during film growth. This model is based on the fact that different crystallographic orientations have different sputter yields. In a polycrystalline film, grain orienta-
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tions with higher sputter yields are removed faster by sputtering. The newly deposited layer grows epitaxially on grains with low sputter yield orientation, and these grains will eventually dominate the film. According to Hultman et al. w32x the Ž100. direction in the B1 fcc structure, such as TiN and TiC, shows the most open channeling direction, and consequently the lowest sputtering yield and also the lowest lattice damage. Therefore, the formation of the Ž100. preferential orientation of the titanium carbide film deposited at VS s y1000 V can be attributed to the sputter channeling effect.
4. Conclusions Titanium carbide thin films with a thickness of approximately 100 nm were deposited on SiŽ100. substrates by the FCVA technique. The crystal orientation in the as-deposited films can be significantly influenced by changing the negative bias voltage applied to the substrates. At VS s 0 V, the film showed a random orientation because of the lower surface mobility. In the bias voltage range of VS s y40 ; y500 V, the titanium carbide films exhibited a Ž111. preferential orientation, which can be interpreted as minimization of elastic energy storage in the films. When VS was elevated to y1000 V, however, the film changed to a Ž100. preferential orientation, which can be attributed to the sputter channeling effect because the Ž100. direction in the TiC lattice shows the most open channeling direction and therefore the lowest sputtering yield. References w1x H.O. Pierson, Handbook of Refractory Carbides and Nitrides, Noyes Publications, Westwood, NJ, 1996. w2x G. Radhakrishnan, P.M. Adams, Appl. Phys. A. Mater. Sci. Proc. 69 Ž1999. S33. w3x A.K. Chaddha, J.D. Parsons, G.B. Kruaval, Appl. Phys. Lett. 66 Ž1995. 760. w4x J.D. Parsons, G.B. Kruaval, A.K. Chaddha, Appl. Phys. Lett. 65 Ž1994. 2075. w5x F. Santerre, M.A. El Khakani, M. Chaker, J.P. Dodelet, Appl. Surf. Sci. 148 Ž1999. 24.
w6x I. Kojima, E. Miyazaki, I. Yasumori, J. Catal. 59 Ž1979. 472. w7x A. Kumar, H.L. Chan, J.S. Kapat, Appl. Surf. Sci. 549 Ž1998. 127᎐129. w8x M.L.F. Parames, O. Conde, J. Phys. 3 Ž1993. 217. w9x I.Yu. Konyashin, Thin Solid Films 278 Ž1996. 37. w10x L. Guzman, M. Bonelli, A. Miottello, D.C. Kothari, Surf. Coat. Technol. 100r101 Ž1998. 500. w11x A.A. Voevodin, M.A. Capano, A.J. Safriet, M.S. Donley, J.S. Zabinsky, Appl. Phys. Lett. 69 Ž1996. 188. w12x E. Kusano, A. Satoh, M.K. Tagawa, H. Nanto, A. Kinbara, Thin Solid Films 343r344 Ž1999. 254. w13x O. Knotek, E. Lugscheider, F. Loffler, C. Barimani, Mater. ¨ Res. Soc. Symp. Proc. 403 Ž1996. 247. w14x L.D. Alessio, A.M. Salvi, R. Teghil et al., Appl. Surf. Sci. 134 Ž1998. 53. w15x G. Radhakrishnan, P.M. Adams, D.M. Speckman, Laser applications in microelectronic and optoelectronic manufacturing IV, in: J. Dubowski, H. Helvajian, E.W. Kreeuz and K. Sugioka ŽEds... Proc. SPIE 3618, 1999, pp.487. w16x Shi. Xu, B.K. Tay, H.S. Tan, Z. Li, Y.Q. Tu, J. Appl. Phys. 79 Ž1996. 7234. w17x X. Shi, B.K. Tay, H.S. Tan et al., Thin Solid Films 345 Ž1999. 1. w18x J.F. Nye, Physical Properties of Crystals: Their Representation by Tensors and Matrices, Clarendon Press, Oxford, 1957. w19x J.E. Sanchez, Jr., Mater. Res. Soc. Symp. Proc. 343 Ž1994. 641. w20x H. Ljungcrantz, M. Oden, ´ L. Hultman, J.E. Greene, J.-E. Sundgren, J.Appl.Phys. 80 Ž1996. 6725. w21x Xing-zhao Ding, Y.J. Li, Z. Sun, B.K. Tay, S.P. Lau, G.Y. Chen, W.Y. Cheung and S.P. Wong, J. Appl. Phys. 88 Ž12. Ž2000. Žin press.. w22x X. Shi, D. Flynn, B.K. Tay, H.S. Tan, Filtered Cathodic Arc Source, PCTrGB96r00389, 20 th Feb. 1995 w23x A. Anders, Phys. Rev. E55 Ž1997. 969. w24x G.G. Stoney, Proc. R. Soc. London 82 Ž1909. 172. w25x D. Briggs, M.P. Seah. Practical Surface Analysis. 2nd ed., vol. 1, Wiley, London, Ž1994.. w26x C.D. Wagner ŽEd.., NIST Standard Reference Database version, 2.0, 1997. w27x J. Pelleg, L.Z. Zevin, S. Lungo, N. Croitoru, Thin Solid Films 197 Ž1991. 117. w28x C. Quaeyhaegens, G. Knuyt, J.D Haen, L.M. Stals, Thin Solid Films 258 Ž1995. 170. w29x D.R. McKenzie, M.M.M. Bilek, J. Appl. Phys. 86 Ž1999. 230. w30x H.S. Roosendaal, in: R. Behrisch ŽEd.., Sputtering by Particle Bombardment I-Physical Sputtering of Single-Element Solids, Springer, New York, 1981. w31x R.M. Bradley, J.M. Harper, D.A. Smith, J.Appl.Phys. 60 Ž1986. 4160. w32x L. Hultman, W.-D. Munz, ¨ J. Musil, S. Kadlec, I. Petrov, J.E. Greene, J. Vac. Sci. Technol. A9 Ž1991. 434.
Preferential orientation of titanium carbide films deposited by a filtered cathodic vacuum arc technique Xing-zhao Ding a,U , B.K. Tay a , H.S. Tana , S.P. Lau a , W.Y. Cheung b, S.P. Wong b a b
School of Electrical and Electronic Engineering, Nanyang Technological Uni¨ ersity, Nanyang, 639798, Singapore Department of Electronic Engineering, The Chinese Uni¨ ersity of Hong Kong, Shatin, New Territories, Hong Kong Received 6 November 2000; received in revised form 10 November 2000; accepted 10 November 2000
Abstract Titanium carbide films with a thickness of approximately 100-nm were deposited on SiŽ100. substrates by a filtered cathodic vacuum arc technique. The composition and microstructure of the films were assessed by Rutherford backscattering spectroscopy, X-ray photoelectron spectroscopy, X-ray diffraction, and atomic force microscopy. A negative bias voltage Ž VS ⫽0 ; y1000 V. was applied to the substrate during deposition, and the influence of VS on the crystalline orientation of the as-deposited films was investigated. It was found that the crystallites are randomly oriented in the film deposited at VS s 0 V. In the bias voltage range of VS s y40 ; y500 V, the titanium carbide films exhibited a Ž111. preferential orientation. When VS was increased to y1000 V, however, the film was Ž100. preferentially oriented. The compressive internal stress, determined by the radius of curvature technique, in the titanium carbide films exhibited a minimum value at approximately VS s y80 ; y120 V. The Ž111. preferential orientation can be explained by minimization of elastic energy storage in the films; while the Ž100. preferential orientation in the film deposited at VS s y1000 V is due to the sputter channeling effect, because the Ž100. direction in the TiC lattice shows the most open channeling direction and therefore the lowest sputtering yield. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Preferential orientation; Titanium carbide film; Filtered cathodic vacuum arc deposition
1. Introduction Titanium carbide ŽTiC. is one of the most widely applied hard coating materials. Its chemical bonding is quite special, similar to other interstitial carbides or nitrides, showing a mixed covalent, metallic and ionic character w1x, which results in many advantageous properties of this material, including high hardness, low coefficient of friction, outstanding wear resistance, high melting or decomposition temperature, good electrical conductivity, excellent corrosion resistance and chemical stability. These unique combined properties have made titanium carbide of particular interest in a wide variety of applications. Because of its hardness and wear resistance, TiC has become a major industrial coating material in many tribological applications w1,2x. U
Corresponding author. Tel.: q65-7905454; fax: q65-7933318. E-mail address: [email protected] ŽX. Ding..
Its high electrical conductivity and low work function make it an outstanding candidate for Ohmic contact of some specific SiC-based electronic devices w3,4x and for use as an electron injection electrode in organic light emitting devices ŽOLEDs. w5x. In addition, TiC is presently attracting great attention as a catalyst w6x, a material for the first wall of nuclear reactors w7x, and as diffusion barriers in semiconductor technology w8x. Titanium carbide films have been prepared by a variety of deposition techniques including chemical vapor deposition ŽCVD. w9,10x, sputtering w11,12x, evaporation w13x, and pulsed laser ablation ŽPLD. w5,14x. In the conventional or plasma assisted CVD process, the deposition is carried out in a high substrate temperature above 500⬚C w1x, which often cannot be tolerated by the substrates with a poor thermal resistance in many important technological applications. Physical vapor deposition ŽPVD. processes, Žvarious evaporation and sputtering techniques ., on the other hand, allow
0257-8972r01r$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 0 . 0 1 1 4 0 - 3
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X. Ding et al. r Surface and Coatings Technology 138 (2001) 301᎐306
coating deposition at lower temperatures than CVD methods. However, lower substrate temperature leads to a lower degree of adhesion between the substrate and the coating, which often results in poorer mechanical properties and delamination of the coatings. In recent years, PLD has been demonstrated to be a suitable technique to deposit TiC films at room temperature. However, it is well known that the generation and inclusion of particulates andror droplets from the target during laser evaporation has been one of the biggest problems of PLD w15x. In the present work, a filtered cathodic vacuum arc ŽFCVA. technique was employed to deposit the titanium carbide films. FCVA deposition is a promising technique for the production of high quality hard thin films w16,17x. The main feature of the FCVA technique is to employ a curved magnetic field to guide the plasma generated from the cathodic vacuum arc to deposit on an out-of-sight substrate. Through this special magnetic filter, most of the unwanted macroparticles and neutral atoms will be removed. Only ions within a defined energy range can reach the substrate, thus producing films with good controllability and reproducibility. Moreover, as the depositing species in the FCVA process are fully ionized, the kinetic energy of these ions can be precisely controlled by adjusting the bias voltage applied to the substrate. It is well known that the low-energy ion bombardment during deposition of thin films may markedly affect the crystal orientation in the films. Often a fibrous texture is generated in which all film crystallites share a crystallographic axis oriented perpendicular to the plane of the substrate. The crystal orientation or texture in thin films has been found to greatly influence the physical properties of materials and consequently their performance in practical applications w18᎐20x. Therefore, controlling the texture of crystalline thin films is important both from the fundamental physics and technological points of view. In our previous paper w21x, some titanium-containing tetrahedral amorphous carbon films with different titanium contents were deposited by a filtered cathodic vacuum arc technique by using compacted titaniumrcarbon mixture cathode targets. It was demonstrated that by using the target with 20 at.% titanium, sub-stoichiometric titanium carbide films could be deposited. In this paper, the evolution of preferential orientation of these titanium carbide films with different negative bias voltage was investigated.
2. Experimental The titanium carbide films were deposited by the FCVA technique. At first, a TirC mixture target with atomic ratio of TirCs 1:4 was prepared. Pure elemental powders of graphite Ž99.5% purity, 325 mesh. and
titanium Ž99.5% purity, 560 mesh., were mixed thoroughly and then compacted into a cylindrical target 60 mm in diameter under an axial pressure of approximately 160 MPa. The FCVA system has been described in detail elsewhere w17,22x. It incorporates an off-plane double-bend ŽOPDB. filter to effectively remove all macro-particles. The arc current was at approximately 110 A. The torridal magnetic field for steering the carbon plasma towards the substrate was maintained at 40 mT. The working chamber was evacuated to a base pressure below 4 = 10y6 torr but rose to 3 = 10y5 torr during deposition due to out-gasing of the cathode. Single crystal Si Ž100. wafers were used as the substrates. Prior to deposition, the silicon substrates were sputtered by an Arq ion beam Ž800 eV, 60 mA. for 5 min in order to remove the native oxide layer. During deposition, a negative DC bias voltage Ž VS ., ranging from 0 to y1000 V, was applied to the substrate. Correspondingly, the energy Ž E ., of the depositing ions can be approximately expressed as Es neŽ VP y VS . q E0 , where VP f y13 V, is the plasma potential w16x, E0 is the original ion energy in the plasma Ž; 28 eV., e is the electron charge, and n is the charge state of the ions in the vacuum arc plasma. According to the experimental results of Anders w23x, carbon arc plasma consists of only single charged state Cq ions, while the titanium arc plasma is composed of 11% Tiqq 75% Ti 2qq 14% Ti 3q with an average charge state of 2.03. In this work the titanium carbide films were deposited at room temperature. The average deposition rate was about 3 nmrmin. The thickness of the films was simply controlled by deposition time and measured to be approximately 100 nm by using a TENCOR P-10 surface profilometer. The composition and microstructure of the asdeposited titanium carbide films were monitored by Rutherford backscattering spectroscopy ŽRBS., X-ray photoelectron spectroscopy ŽXPS., and X-ray diffraction ŽXRD. analyses. In RBS analyses, a beam of 2 MeV He 2q ions were projected into the film at an incident angle of 7⬚ to the normal line of the film plane, while backscattered He 2q ions were detected at a scattering angle of 170⬚. The XPS analyses were carried out using a VG Scientific Microlab 310F with Mg K ␣ radiation as the X-ray source. The XRD experiments were performed on a Siemens D5005 X-ray diffractometer with a glancing incident mode by using CuK ␣ radiation Ž40 kV, 40 mA.. The incident angle was fixed at 1⬚, the 2 scanning step size was set as 0.05⬚, and the counting time at each step was 4 s. The surface morphology of the as-deposited films was directly observed under ambient conditions by a Nanoscope IIIa scanning atomic force microscope ŽAFM.. The system is operated in the non-contact tapping and lift modes. The internal stress Ž s . in the as-deposited titanium
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carbide films was determined according to Stoney’s equation w24x: s s
Es t s2 6 Ž1 y ¨ s . t c
ž
1 1 y R R0
/
Where Es , ¨ s and t s are Young’s Modulus, Poisson Ratio and thickness of the silicon substrate, R 0 and R are the curvature radii of the wafer before and after coating with the TiC films, t c is the thickness of the films. R 0 , R and t c are all measured by using the surface profilometer.
3. Results and discussion Fig. 1 shows a typical RBS spectrum of the as-deposited titanium carbide films. The average titanium content in the films was determined to be approximately 46.2" 1.5 at.%, which is more than twice of that in the cathode target. This phenomenon can be interpreted as follows. Due to the lower melting temperature of Ti, in comparison with that of carbon Žthe melting point of titanium and graphite is ; 1660⬚C and ; 3870⬚C, respectively., the titanium particles in the arc spot area during deposition will be completely sublimated and consequently ionized. In contrast, the graphite particles may only be partially evaporated with the formation of some macro-clusters which will be trapped by the OPDB filter and cannot reach the substrate. In addition, the higher titanium concentration in the films can also be partially attributed to the preferential re-sputtering and back-scattering effects. Because the mass of carbon atoms is much lower than that of titanium atoms, carbon atoms on the film surface would be preferentially sputtered by the energetic depositing ions, especially by the heavier titanium ions with higher energies Žin double- and triple-charged states .. On the other hand, the depositing carbon ions could be more easily back-scattered from the heavier titanium atoms on the film surface. In the RBS spectrum, some oxygen andror nitrogen contamination in the films can be evidently detected. This contamination is closely related to the sample preparation process. When compacting the TirC mixed target under ambient conditions, some air bubbles might inevitably be trapped in the target. During deposition, with the erosion of the target, the air contained inside the target will come out and the oxygen and nitrogen components will immediately react with the melting titanium. The oxygen andror nitrogen contamination may also come from the residual gas in the vacuum chamber. Indeed, it is known that TiC is isomorphous with TiN and TiO w1x. Thus, nitrogen and oxygen may incorporate into the
Fig. 1. A typical RBS spectrum of the as-deposited titanium carbide films.
TiC films as impurities during deposition as a result of interactions with residual gases in the vacuum. Typical C 1s and Ti 2p XPS spectra of the as-deposited titanium carbide film surfaces are shown in Fig. 2. The C 1s spectrum is presented as three components. The peak around 282.0 eV can be assigned to the carbidic carbon component ŽC bonded to Ti. w5x, while both the strong and the weak components at 285.0 eV and 289.0 eV could most likely be attributed to adsorbed hydrocarbon groups w25,26x on the film surface. The Ti 2p XPS spectrum can be deconvoluted into two sets of spin-orbit doublets, each separated by 6 eV, corresponding to the carbide ŽTi 2p 3r2 at 455.0 eV. and the oxide components ŽTi 2p 3r2 at 458.9 eV., respectively w5x. The strong TiO 2 spin-doublet was believed to have mainly resulted from the surface oxidization when the samples were kept in air. The oxygen contamination in the films may also be partially responsible for the presence of the oxidic Ti 2p spin-doublet. The internal stress Ž s ., in the as-deposited titanium carbide films as a function of the negative bias voltage is shown in Fig. 3. It was as expected that the internal stress in all of these films would be compressive in nature. It was found that the s exhibited a minimum value at approximately VS s y80 ; y120 V. Fig. 4 shows three typical AFM images of the asdeposited titanium carbide films with different crystal orientations. It was found that the films deposited at lower negative bias voltages, Že.g. VS s 0 ; y160 V. were quite smooth, while the film deposited at a higher bias voltage, e.g. y1000 V, exhibited an evident columnar surface morphology due to the bombardment of high energy depositing ions, especially Ti 2q and Ti 3q ions. The variation of root-mean-square surface roughness of the titanium carbide films with the applied negative bias voltage is shown in Fig. 5. The variation
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Fig. 2. Typical C 1s and Ti 2p XPS spectra of the as-deposited titanium carbide film surfaces: a and c are Ti 2p 3r 2 and Ti 2p1r2 in titanium oxide; b and d are Ti 2p 3r 2 and Ti 2p1r2 in titanium carbide.
trend of the surface roughness was found to be similar to that of the internal stress in the films, i.e. showing a minimum value at approximately VS s y80 ; y120 V. The XRD patterns of the titanium carbide films are shown in Fig. 6. At VS s 0 V, three evident peaks around 2 f 35.9⬚, 41.7⬚, and 60.6⬚ are observed, which can be indexed as the Ž111., Ž200. and Ž220. diffractions of the cubic TiC crystalline phase with a B1 NaCl-type Fig. 4. Three typical AFM images of the as-deposited titanium carbide films with different orientations: Ža. random oriented film deposited at VS s 0 V; Žb. Ž111. preferentially oriented film deposited at VS s y80 V; and Žc. Ž100. preferentially oriented film deposited at VS s y1000 V.
Fig. 3. Compressive internal stress in the titanium carbide films deposited at different negative bias voltage.
structure. This film exhibited a random orientation character. With the increase of negative bias voltage, the Ž200. and Ž220. peaks weakened gradually and disappeared at VS s y160 V. The films showed an evident Ž111. preferred orientation. This preferential orientation could be retained up to VS s y500 V. When the negative bias voltage was elevated further to y1000 V, the Ž111. diffraction peak disappeared; meanwhile the Ž200. peak emerged again, and the film showed a Ž100. preferred orientation. The formation of texture in polycrystalline thin films
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Fig. 5. Root-mean-square surface roughness of the titanium carbide films deposited at different negative bias voltage.
can be affected by many factors, including the surface energy, elastic strain energy, thermal energy, and so on. In the present work, all of the titanium carbide films were deposited at room temperature, therefore the influence of thermal energy can be neglected. To date, the surface energy in the B1 NaCl-type structured crystals, e.g. TiN and TiC, remains ambiguous. Pelleg et al. w27x showed that the Ž100. planes have the smallest surface energy by neglecting metal-metal ŽM᎐M. bond-
Fig. 6. XRD patterns of the titanium carbide films deposited at different bias voltage: Ža. 0 V, Žb. y40 V, Žc. y80 V, Žd. y120 V, Že. y160 V, Žf. y500 V, Žg. y1000 V.
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ing. However, Quaeyhaegens et al. w28x showed that Ž111. are planes with the smallest surface energy by considering an fcc metal structure. Therefore, the structure evolution of the titanium carbide films cannot be simply explained by the surface energy of the different faces in the TiC lattice. From Fig. 3, it has been noted that the compressive internal stress in our FCVA deposited titanium carbide films is very high, thus the elastic strain energy in the crystal lattice could be a dominating factor in determining the texture of the films. Recently, Mckenzie et al. w29x demonstrated that the preferential orientation in thin films can come about as a result of the synthesis materials under impressed stress conditions. For the materials with the rocksalt crystal structure, including TiC and TiN, it has been revealed w27,29x that the crystallites with the Ž111. orientation have the smallest elastic energy storage in the case of stressed states. Therefore, the experimentally observed Ž111. preferential orientation in the titanium carbide films deposited at a medium bias voltage, i.e. VS s y40 ; y500 V, can be explained by the concept of minimization of the elastic deformation energy. The minimum value of the measured internal stress at approximately VS s y80 ; y120 V in the as-deposited films, as observed in Fig. 3, reasonably supported this viewpoint. It is well known that FCVA deposition is far from a thermodynamic equilibrium process. At VS s 0 V, the energy of the depositing species is relatively small, and the mobility of the surface atoms is very low. Thus the as-deposited film showed a random orientation because of the lower structural relaxation ability. With the increase of negative bias voltage, the depositing ions can be effectively accelerated, and the energy of the ions is approximately proportional to the bias voltage. Under the bombardment of the energetic depositing ion flux, the mobility of the atoms on the growing film surface can be greatly increased. As a consequence of this, the film structure can be relaxed to a lower energy state with the formation of Ž111. preferential orientation. On the other hand, with the increase of ion energy Žor VS ., the lattice damage caused by the ion bombardment could increase gradually, and, as a consequence, the internal stress could also increase gradually, as observed in Fig. 3 at VS s y120; y1000 V, due to the lattice deformation. Moreover, with the increase of ion energy, the sputtering effect could become more and more significant. In the energy range involved in the present work, the sputter yield increases linearly with ion energy w30x. In a study of ion beam assisted deposition processes, Bradley et al. w31x developed a model to explain the development of preferred orientation during film growth. This model is based on the fact that different crystallographic orientations have different sputter yields. In a polycrystalline film, grain orienta-
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tions with higher sputter yields are removed faster by sputtering. The newly deposited layer grows epitaxially on grains with low sputter yield orientation, and these grains will eventually dominate the film. According to Hultman et al. w32x the Ž100. direction in the B1 fcc structure, such as TiN and TiC, shows the most open channeling direction, and consequently the lowest sputtering yield and also the lowest lattice damage. Therefore, the formation of the Ž100. preferential orientation of the titanium carbide film deposited at VS s y1000 V can be attributed to the sputter channeling effect.
4. Conclusions Titanium carbide thin films with a thickness of approximately 100 nm were deposited on SiŽ100. substrates by the FCVA technique. The crystal orientation in the as-deposited films can be significantly influenced by changing the negative bias voltage applied to the substrates. At VS s 0 V, the film showed a random orientation because of the lower surface mobility. In the bias voltage range of VS s y40 ; y500 V, the titanium carbide films exhibited a Ž111. preferential orientation, which can be interpreted as minimization of elastic energy storage in the films. When VS was elevated to y1000 V, however, the film changed to a Ž100. preferential orientation, which can be attributed to the sputter channeling effect because the Ž100. direction in the TiC lattice shows the most open channeling direction and therefore the lowest sputtering yield. References w1x H.O. Pierson, Handbook of Refractory Carbides and Nitrides, Noyes Publications, Westwood, NJ, 1996. w2x G. Radhakrishnan, P.M. Adams, Appl. Phys. A. Mater. Sci. Proc. 69 Ž1999. S33. w3x A.K. Chaddha, J.D. Parsons, G.B. Kruaval, Appl. Phys. Lett. 66 Ž1995. 760. w4x J.D. Parsons, G.B. Kruaval, A.K. Chaddha, Appl. Phys. Lett. 65 Ž1994. 2075. w5x F. Santerre, M.A. El Khakani, M. Chaker, J.P. Dodelet, Appl. Surf. Sci. 148 Ž1999. 24.
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