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Surface morphologies and electrical properties of antimony-doped tin oxide films deposited by plasma-enhanced chemical vapor deposition
Surface and Coatings Technology 138 Ž2001. 229᎐236
Surface morphologies and electrical properties of antimony-doped tin oxide films deposited by plasma-enhanced chemical vapor deposition Keun-Soo Kima , Seog-Young Yoonb, Won-Jae Lee c , Kwang Ho Kimb,U b
a Hyundai Electronics Industries Co. Ltd., 1, Hyangjeong-dong, Hungduk-gu, Cheongju-si 361-725, South Korea Department of Inorganic Materials Engineering, Pusan National Uni¨ ersity, 30 Changjeon-dong, Keumjeong-gu, Pusan 609-735, South Korea c Korea Electrotechnology Research Institute, P.O. Box 20, Changwon 641-600 South Korea
Received 17 July 2000; received in revised form 12 October 2000; accepted 12 October 2000
Abstract Antimony-doped tin oxide ŽATO. films were deposited on Corning glass 1737 substrates by a plasma-enhanced chemical vapor deposition ŽPE-CVD. technique using a gas mixture of SnCl 4 ᎐SbCl 5 ᎐O 2 ᎐Ar. Electrical properties and surface morphologies of these films were studied by varying the deposition temperature, input gas ratio, Rws ŽPSbCl5rPSnCl4 .x, and RF power. The PE-CVD method effectively enhanced the deposition rate and also improved the surface roughness of the deposit compared with thermal CVD. The antimony doped tin oxide films which had relatively good electrical properties were obtained at a deposition temperature of 450⬚C, an input gas ratio of R s 1.12, and a RF power of 30 W. In addition, the studies on the morphological development of the films by AFM analysis suggested that higher input gas ratio and lower deposition temperature led to a decrease in the surface roughness of the deposited films. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Sb-doped SnO 2 thin films; Plasma-enhanced chemical vapor deposition; Surface roughness
1. Introduction Tin oxide ŽSnO 2 . films are attractive for several optoelectronic devices w1᎐4x due to their unique physical properties such as high electrical conductivity, high transparency in the visible part of the spectrum, and high reflectivity in the IR region. Particularly, they are stable up to high temperatures, have excellent resistance to strong acids and bases at room temperature as well as to mechanical wear and they have very good adhesion to many substrates w5᎐7x.
U
Corresponding author. Tel.: q82-51-510-2391; fax: q82-51-5100528. E-mail address: [email protected] ŽK. Ho Kim..
All these properties have led to intense research for SnO 2 coatings over the past few decades. Tin oxide films have been prepared by various techniques such as chemical vapor deposition ŽCVD. w8x, spray pyrolysis w9x, vacuum evaporation w10x, reactive RF sputtering w11x, and dc glow discharge w12x, each of these techniques has advantages and disadvantages. From a practical point of view, tin oxide films doped with antimony or fluorine have been prepared and their optical and electrical properties have been also intensively investigated w13᎐17x. As a result, it is well described in the literature that the film conductivity increases largely without losing optical transparency with a small amount of dopant. However, large additions diminish the conductivity, optical transparency, and crystallinity of the films. This behavior has been interpreted in various
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 1 4 - 2
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K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
ways, such as increased disorder w15,16x, impurity scattering w13x, and precipitation of a second phase w14x. From our previous study w18x, the electrical and optical properties of antimony-doped tin oxide ŽATO. films prepared by the thermal CVD method are well understood. In the present work, ATO films were deposited by the plasma-enhanced chemical vapor deposition ŽPE-CVD. technique to investigate the effect of RF power on the film growth. Also the electrical properties and surface morphologies of ATO films deposited by PE-CVD were studied with variations of input gas ratio and deposition temperature.
with scanning electron microscopy ŽSEM, Hitachi S4200. and atomic force microscopy ŽAFM, Digital Instrument.. With aid of AFM, root mean square ŽRMS. values of surface roughness were estimated to study the effect of the deposition temperature, input gas ratio, R and film thickness. All measurements were carried out ˚ on films having the same thickness of ; 1400 A. The electrical resistance of the films was measured by a conventional four-point probe method. The concentration and mobility of electrical carriers in the films were obtained from Hall measurements performed at room temperature using the standard van der Pauw technique. Indium paste and platinum wire were used for ohmic contact.
2. Experimental procedure 2.1. Deposition of Sb-doped SnO2 films
3. Results and discussion
ATO films were prepared by using a PE-CVD technique. For preparation of ATO films, a gas phase mixture of SnCl 4 , SbCl 5 and O 2 was used as a precursor. Flow rates of these gases were individually controlled with a mass flow controller. The pre-cleaned substrate ŽCorning glass 1737, 1.5 cm = 1.5 cm. was mounted on an alumina susceptor inside the reaction chamber. The substrate temperature was monitored with a K-type thermocouple. The partial pressure of each source in the reactor was controlled by adjusting the amount of carrier gas ŽAr. passed through bubblers with a mass flow controller. The bubblers were maintained at 0⬚C for SnCl 4 and at 35⬚C for SbCl 5 to get correct vapor pressures. A heating string kept outlet tubes from the bubbler at a temperature of approximately 60⬚C to prevent the reaction gas from condensing during flowing. After ATO films were deposited onto the substrate, the samples were allowed to cool down slowly. During antimony doping into SnO 2 films, the partial pressures of reactive gases, on the basis of our previous study, were kept at PSnCl4 , 3.8= 10y3 torr; PO 2 , 5.6= 10y1 torr; total pressure; 1 torr. The SbCl 5 partial pressures were in the range of zero to 1 = 10y2 torr. The deposition temperatures were varied from 250 to 500⬚C, varying RF power from 15 to 90 W.
3.1. Deposition beha¨ ior With a gas mixture of SnCl 4rO 2rAr, tin oxide films were prepared using two different methods, T-CVD and PE-CVD. As shown in Fig. 1, the deposition rate of tin oxide films increased with increase of deposition temperature. This indicated that deposition process was controlled by thermal activation w19x. However, the deposition rate by PE-CVD was faster than that by T-CVD. It would be suggested that the additional plasma energy clearly enhanced the deposition rate of the deposit w20x. In the PE-CVD method the deposition rate increased by increasing RF power to 15 W as shown in Fig. 2. Further increase of RF power beyond this value caused a saturation of the deposition rate. XRD patterns of the deposited films are shown in Figs. 3 and 4. All films for XRD were prepared to have
2.2. Film characterizations X-Ray diffraction ŽXRD. analysis of the tin oxide films was carried out using a Rigaku DrMax-2400 diffractometer to investigate the structure of the deposited films. CuK ␣ radiation with a wavelength of 0.154 nm was used as the X-ray source. The thickness of film was measured with a stylus Ž ␣-STEP. instrument. For this measurement a step was etched with HClrZn powder slurry. The surface morphology of the films was observed
Fig. 1. Deposition rate of SnO 2 films for different methods as a function of deposition temperature.
K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
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Fig. 2. Deposition rate of SnO 2 films at two different deposition temperatures as a function of RF power during PE-CVD deposition.
the same thickness of ; 0.14 m. Fig. 3 shows XRD patterns of Sb-doped SnO 2 films prepared at different deposition temperatures. Below 300⬚C, no XRD peak was observed, and this result suggested that the deposited films were nearly amorphous. However, the increase of deposition temperature above 350⬚C caused the crystallization of the films, as represented by basically polycrystalline, a tetragonal rutile structure of SnO 2 with highly Ž110. preferred orientation. The effects of input gas ratio and RF power on crystallinity of the deposit have been studied. Fig. 4a shows X-ray peak intensities occurred with a variation of input gas ratio Žat different R values.. Some differences in crystallographic orientation of the films deposited with changing R values as indicated by
Fig. 4. XRD patterns for Sb-doped SnO2 films at different Ža. input gas ratios and Žb. RF powers. The deposition temperature was kept at 450⬚C.
Fig. 3. XRD patterns for Sb-doped SnO2 films at different deposition temperatures. The input gas ratio and RF power were kept at R s 1.12 and 30 W, respectively.
changes in the relative peak intensity. These phenomena with Sb doping were similarly reported by Rohatgi et al. w15x and Carroll et al. w16x. In the present case the relative presence of Ž110. along with other orientations remained invariant for subsequent doping levels. This behavior with Sb doping means that in the present case antimony incorporation in SnO 2 lattice has not af-
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fected the structural properties to a considerable extent. These results can be interpreted using the lattice parameters with the dopant levels. While calculating the lattice parameters it should be remembered that the Sb incorporation in the SnO 2 lattice will depend on the dopant level. According to our previous paper w18x, the calculated lattice parameters Ž a-axis and c-axis. for pure SnO 2 were approxi˚ and 3.185 A, ˚ respectively. In dopedmately 4.735 A SnO 2 , the initial increase of lattice constants is considered to be due to the dissolution of Sb into the SnO 2 lattice; at initial dopant levels one expects the Sb incorporation only at substitutional sites of Sn. For higher dopant levels the incorporation will take place at interstitial sites, and it was suggested that some phase like antimony oxides ŽSb 2 O 3 , Sb 2 O4 and Sb 2 O5 . could appear. Our present work in this context indicated that Sb doping up to approximately Ž R s 2.0. in such films took place at substitutional sites only. Therefore, in the present case antimony incorporation in SnO 2 lattice has not affected the structural properties to a considerable extent. However, the chemical composition of the deposited films at each input gas ratio Ž R values. was analyzed through EPMA. The results showed SbrSn atomic ratio was approximately 0.05 when the R value was 1.12, which was the optimal input gas ratio in the present experiment. The XRD patterns for the films did not vary significantly with changes in RF power as can be seen in Fig. 4b. 3.2. Electrical properties Fig. 5 shows the variation of carrier concentration, mobility, and resistivity for ATO films as a function of deposition temperature. A gradual decrease in the resistivity occurs on raising the temperature of deposition
up to 450⬚C. This behavior was attributed to the increase of both carrier concentration and carrier mobility as shown in Fig. 5. However, further raising the deposition temperature above 500⬚C resulted in a little increase again of the film resistivity. In the present study, the deposition temperature 450⬚C was the critical temperature for deposition at which the lowest resistivity was achieved. Since the deposition temperature of 450⬚C produces the best values of electrical conductivity, all subsequent studies have been carried out on the films prepared at 450⬚C. Electrical properties of ATO films are shown in Fig. 6 as a function of input gas ratio, R, at constant deposition temperature of 450⬚C with RF power of 30 W. As the R value increased, the resistivity first decreased, then reached a minimum and increased again at high R values. It is evident that the earlier decrease of the resistivity resulted mainly from an increase of electron concentration due to Sb doping in SnO 2 w21x, whereas the considerable increase of resistivity at high R value was caused by the reduction of both electron concentration and electron mobility w16x. The increase of carrier concentration with doping is due to the replacement of a Sn atom by the added Sb atom w22x, which gives one extra charge carrier. However, mobility increased up to doping levels of R s 0.7 and decreased continuously for further addition of the dopant. It is known that the grain boundary scattering w23x and ionized impurity scattering w24x play a major role in determining the mobility variation in such extrinsically doped semiconductors. The effect of RF power at constant R value Ž R s 1.12. on the electrical properties was shown in Fig. 7. As RF power increased, mobility increased and carrier concentration decreased. The slight increase of resistivity with RF power was due to the decrease of carrier
Fig. 5. Variation of resistivity, mobility and carrier concentration in Sb-doped SnO 2 films with deposition temperature.
K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
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Fig. 6. Variation of resistivity, mobility and carrier concentration in Sb-doped SnO 2 films with input gas ration Ž R value..
concentration, because mobility exhibited a slight increase with RF power w20x. 3.3. Surface morphology Fig. 8 shows the surface morphologies by SEM and surface roughness by AFM for the films deposited at
various conditions. Fig. 8a,b represents the surface morphologies of the undoped tin oxide films deposited by T-CVD and PE-CVD. The facet ripples were shown in the films deposited by T-CVD, however, the films prepared by PE-CVD consisted of fine ripples. The degree of surface roughness for deposited films prepared by PE-CVD were lower than in the case of
Fig. 7. Variation of resistivity, mobility and carrier concentration in Sb-doped SnO 2 films with RF power. Table 1 RMS-surface roughness values of Sb-doped SnO 2 films at two different deposition methods Method
Input gas ratio Rws ŽPSbCl5rPSnCl4 .x
Deposition temperature Ž⬚C.
RF power ŽW.
T-CVD PE-CVD PE-CVD PE-CVD
0 0 1.12 2.63
450 450 450 450
᎐ 30 30 30
RMS-roughness Žnm. 16 12 6 4
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K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
Fig. 8. SEM and AFM images of undoped and Sb-doped SnO 2 films: Ža. undoped SnO 2 film by T-CVD, Žb. undoped SnO 2 film by PE-CVD, Žc. Sb-doped SnO 2 film by PE-CVD Ž R s 1.12., Žd. Sb-doped SnO 2 film by PE-CVD Ž R s 2.63.. The deposition temperature was kept at 450⬚C.
Fig. 8. Ž Continued..
K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
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Fig. 9. AFM images Ž1 m = 1 m. of the Sb-doped SnO 2 films by PE-CVD at various deposition temperatures Ža. 250⬚C, Žb. 350⬚C, Žc. 450⬚C, and 500⬚C. The input gas ratio and RF power were kept in Rs 1.12 and 30 W, respectively.
T-CVD, as shown in Fig. 8a,b, and in Table 1 where the RMS roughness values are summarized. This behavior can be assumed to be due to the addition of plasma energy w20x. From Fig. 8b᎐d, it was found that as the R value increased, the degree of surface roughness decreased. This phenomenon can be explained by two possible reasons. First, with the increasing of doping concentration the deposited films lost the crystallinity w16x. Secondly, the larger of the doping concentrations induced precipitation such as the antimony oxide phase ŽSb 2 O 3 , Sb 2 O4 and Sb 2 O5 etc.. w18x; and the preferred orientation growth of SnO 2 films may be suppressed by the precipitation. As a result, the surface roughness could be decreased with the increasing of R values. The surface roughness values and surface morphologies at different deposition temperatures are shown in Fig. 9 and Table 2, respectively. From Table 2, the degree of surface roughness increased with increasing deposition temperature. It would generally be considered that the surface morphology of deposited films is related with surface thermal energy and mass diffuTable 2 RMS surface roughness values of Sb-doped SnO 2 films prepared by PE-CVD at various deposition temperatures Input gas ratio Deposition RF power RMS roughness Žnm. Rws ŽPSbCl5 rPSnCl4 .x temperature Ž⬚C. ŽW. 1.12 1.12 1.12 1.12
250 350 450 500
30 30 30 30
4 6 12 16
sion of constituent elements. As shown in Fig. 9 at lower deposition temperature such as 250 and 350⬚C, the fine ripples or grains were formed because of a relatively lower thermal energy. However, at higher temperatures, the surface mobility of constituent elements increases, and the growth of facet grains are preferred w20x. As a result, the degree of surface roughness increased with increasing deposition temperature.
4. Conclusions Sb-doped SnO 2 films have been prepared by the plasma-enhanced chemical vapor deposition technique using a gas mixture of SnCl 4 ᎐SbCl 5 ᎐O 2 ᎐Ar and the following results were obtained. Ž1. PE-CVD method largely increased the deposition rate and smoothed the surface of tin oxide films compared with thermal CVD. Ž2. The highly Ž110. oriented polycrystalline films were grown above 350⬚C, while amorphous-like films were grown below 250⬚C. Ž3. The Sb-doped SnO 2 films which had the relatively good electrical properties were formed at 450⬚C, input gas ratio R s 1.12, and RF power 30 W. Ž4. Higher input gas ratio and lower deposition temperature led to a decrease in the surface roughness of the deposited films. References w1x K.L. Chopra, S. Major, D.K. Pandya, Thin Solid Films 102 Ž1983. 1᎐46. w2x Y.K. Fang, J.J. Lee, Thin Solid Films 169 Ž1989. 51᎐56.
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w3x L.D. Angelis, N. Minnaja, Sensors Actuators B 3 Ž1991. 197᎐204. w4x S.G. Ansari, S.W. Gosavi, S.A. Gangal, R.N. Karekar, R.C. Aiyer, J. Mater. Sci.: Mater. Electron. 8 Ž1997. 23᎐27. w5x Z.M. Jarzebski, J.P. Marton, J. Electrochem. Soc. 123 Ž7. Ž1976. 199᎐205. w6x Z.M. Jarzebski, J.P. Marton, J. Electrochem. Soc. 123 Ž10. Ž1976. 333᎐346. w7x F.J. Yusta, M.L. Hichman, H. Shamlian, J. Mater. Chem. 7 Ž8. Ž1997. 1421᎐1427. w8x C.F. Wan, R.D. McGrath, W.F. Keenan, S.N. Frank, J. Electrochem. Soc. 136 Ž5. Ž1989. 1459᎐1463. w9x S. Reddy, A.K. Malik, S.R. Jawalekar, Thin Solid Films 143 Ž1986. 113. w10x M. Mizuhashi, J. Non-Cryst. Solids 38 Ž1980. 329. w11x D.B. Fraser, H.D. Cook, J. Electrochem. Soc. 119 Ž1972. 1368. w12x D.E. Carlson, J. Electrochem. Soc. 122 Ž1975. 1334. w13x T. Arai, J. Phys. Soc. Jpn. 15 Ž5. Ž1960. 916᎐927.
w14x D. Elliott, D.L. Zellmer, H.A. Laitinen, J. Electrochem. Soc. 117 Ž1970. 1343᎐1348. w15x A. Rohatgi, T.R. Viverito, L.H. Slack, J. Am. Ceram. Soc. 57 Ž6. Ž1974. 278᎐279. w16x A.F. Carroll, L.H. Slack, J. Electrochem. Soc. 123 Ž12. Ž1976. 1889᎐1893. w17x N.S. Murty, S.R. Jawalekar, Thin Solid Films 108 Ž1983. 277᎐283. w18x K.H. Kim, S.W. Lee, J. Am. Ceram. Soc. 77 Ž4. Ž1994. 915᎐921. w19x R.N. Ghoshtagore, J. Electrochem. Soc. 125 Ž1. Ž1978. 110᎐117. w20x S. Shirakata, A. Yokoyama, S. Isomura, J. Appl. Phys. Jpn. 35 Ž1996. L722᎐L724. w21x C.A. Vincent, J. Electrochem. Soc. 119 Ž4. Ž1972. 515᎐518. w22x W.A. Badawy, E.A. El-Taher, Thin Solid Films 158 Ž1988. 277᎐284. w23x E. Shanthi, V. Dutta, A. Banerjee, K.L. Chopra, J. Appl. Phys. 51 Ž12. Ž1980. 6243᎐6251. w24x M.N. Islam, M.O. Hakim, J. Mater. Sci. Lett. 4 Ž1986. 1125.
Surface morphologies and electrical properties of antimony-doped tin oxide films deposited by plasma-enhanced chemical vapor deposition Keun-Soo Kima , Seog-Young Yoonb, Won-Jae Lee c , Kwang Ho Kimb,U b
a Hyundai Electronics Industries Co. Ltd., 1, Hyangjeong-dong, Hungduk-gu, Cheongju-si 361-725, South Korea Department of Inorganic Materials Engineering, Pusan National Uni¨ ersity, 30 Changjeon-dong, Keumjeong-gu, Pusan 609-735, South Korea c Korea Electrotechnology Research Institute, P.O. Box 20, Changwon 641-600 South Korea
Received 17 July 2000; received in revised form 12 October 2000; accepted 12 October 2000
Abstract Antimony-doped tin oxide ŽATO. films were deposited on Corning glass 1737 substrates by a plasma-enhanced chemical vapor deposition ŽPE-CVD. technique using a gas mixture of SnCl 4 ᎐SbCl 5 ᎐O 2 ᎐Ar. Electrical properties and surface morphologies of these films were studied by varying the deposition temperature, input gas ratio, Rws ŽPSbCl5rPSnCl4 .x, and RF power. The PE-CVD method effectively enhanced the deposition rate and also improved the surface roughness of the deposit compared with thermal CVD. The antimony doped tin oxide films which had relatively good electrical properties were obtained at a deposition temperature of 450⬚C, an input gas ratio of R s 1.12, and a RF power of 30 W. In addition, the studies on the morphological development of the films by AFM analysis suggested that higher input gas ratio and lower deposition temperature led to a decrease in the surface roughness of the deposited films. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Sb-doped SnO 2 thin films; Plasma-enhanced chemical vapor deposition; Surface roughness
1. Introduction Tin oxide ŽSnO 2 . films are attractive for several optoelectronic devices w1᎐4x due to their unique physical properties such as high electrical conductivity, high transparency in the visible part of the spectrum, and high reflectivity in the IR region. Particularly, they are stable up to high temperatures, have excellent resistance to strong acids and bases at room temperature as well as to mechanical wear and they have very good adhesion to many substrates w5᎐7x.
U
Corresponding author. Tel.: q82-51-510-2391; fax: q82-51-5100528. E-mail address: [email protected] ŽK. Ho Kim..
All these properties have led to intense research for SnO 2 coatings over the past few decades. Tin oxide films have been prepared by various techniques such as chemical vapor deposition ŽCVD. w8x, spray pyrolysis w9x, vacuum evaporation w10x, reactive RF sputtering w11x, and dc glow discharge w12x, each of these techniques has advantages and disadvantages. From a practical point of view, tin oxide films doped with antimony or fluorine have been prepared and their optical and electrical properties have been also intensively investigated w13᎐17x. As a result, it is well described in the literature that the film conductivity increases largely without losing optical transparency with a small amount of dopant. However, large additions diminish the conductivity, optical transparency, and crystallinity of the films. This behavior has been interpreted in various
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 1 4 - 2
230
K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
ways, such as increased disorder w15,16x, impurity scattering w13x, and precipitation of a second phase w14x. From our previous study w18x, the electrical and optical properties of antimony-doped tin oxide ŽATO. films prepared by the thermal CVD method are well understood. In the present work, ATO films were deposited by the plasma-enhanced chemical vapor deposition ŽPE-CVD. technique to investigate the effect of RF power on the film growth. Also the electrical properties and surface morphologies of ATO films deposited by PE-CVD were studied with variations of input gas ratio and deposition temperature.
with scanning electron microscopy ŽSEM, Hitachi S4200. and atomic force microscopy ŽAFM, Digital Instrument.. With aid of AFM, root mean square ŽRMS. values of surface roughness were estimated to study the effect of the deposition temperature, input gas ratio, R and film thickness. All measurements were carried out ˚ on films having the same thickness of ; 1400 A. The electrical resistance of the films was measured by a conventional four-point probe method. The concentration and mobility of electrical carriers in the films were obtained from Hall measurements performed at room temperature using the standard van der Pauw technique. Indium paste and platinum wire were used for ohmic contact.
2. Experimental procedure 2.1. Deposition of Sb-doped SnO2 films
3. Results and discussion
ATO films were prepared by using a PE-CVD technique. For preparation of ATO films, a gas phase mixture of SnCl 4 , SbCl 5 and O 2 was used as a precursor. Flow rates of these gases were individually controlled with a mass flow controller. The pre-cleaned substrate ŽCorning glass 1737, 1.5 cm = 1.5 cm. was mounted on an alumina susceptor inside the reaction chamber. The substrate temperature was monitored with a K-type thermocouple. The partial pressure of each source in the reactor was controlled by adjusting the amount of carrier gas ŽAr. passed through bubblers with a mass flow controller. The bubblers were maintained at 0⬚C for SnCl 4 and at 35⬚C for SbCl 5 to get correct vapor pressures. A heating string kept outlet tubes from the bubbler at a temperature of approximately 60⬚C to prevent the reaction gas from condensing during flowing. After ATO films were deposited onto the substrate, the samples were allowed to cool down slowly. During antimony doping into SnO 2 films, the partial pressures of reactive gases, on the basis of our previous study, were kept at PSnCl4 , 3.8= 10y3 torr; PO 2 , 5.6= 10y1 torr; total pressure; 1 torr. The SbCl 5 partial pressures were in the range of zero to 1 = 10y2 torr. The deposition temperatures were varied from 250 to 500⬚C, varying RF power from 15 to 90 W.
3.1. Deposition beha¨ ior With a gas mixture of SnCl 4rO 2rAr, tin oxide films were prepared using two different methods, T-CVD and PE-CVD. As shown in Fig. 1, the deposition rate of tin oxide films increased with increase of deposition temperature. This indicated that deposition process was controlled by thermal activation w19x. However, the deposition rate by PE-CVD was faster than that by T-CVD. It would be suggested that the additional plasma energy clearly enhanced the deposition rate of the deposit w20x. In the PE-CVD method the deposition rate increased by increasing RF power to 15 W as shown in Fig. 2. Further increase of RF power beyond this value caused a saturation of the deposition rate. XRD patterns of the deposited films are shown in Figs. 3 and 4. All films for XRD were prepared to have
2.2. Film characterizations X-Ray diffraction ŽXRD. analysis of the tin oxide films was carried out using a Rigaku DrMax-2400 diffractometer to investigate the structure of the deposited films. CuK ␣ radiation with a wavelength of 0.154 nm was used as the X-ray source. The thickness of film was measured with a stylus Ž ␣-STEP. instrument. For this measurement a step was etched with HClrZn powder slurry. The surface morphology of the films was observed
Fig. 1. Deposition rate of SnO 2 films for different methods as a function of deposition temperature.
K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
231
Fig. 2. Deposition rate of SnO 2 films at two different deposition temperatures as a function of RF power during PE-CVD deposition.
the same thickness of ; 0.14 m. Fig. 3 shows XRD patterns of Sb-doped SnO 2 films prepared at different deposition temperatures. Below 300⬚C, no XRD peak was observed, and this result suggested that the deposited films were nearly amorphous. However, the increase of deposition temperature above 350⬚C caused the crystallization of the films, as represented by basically polycrystalline, a tetragonal rutile structure of SnO 2 with highly Ž110. preferred orientation. The effects of input gas ratio and RF power on crystallinity of the deposit have been studied. Fig. 4a shows X-ray peak intensities occurred with a variation of input gas ratio Žat different R values.. Some differences in crystallographic orientation of the films deposited with changing R values as indicated by
Fig. 4. XRD patterns for Sb-doped SnO2 films at different Ža. input gas ratios and Žb. RF powers. The deposition temperature was kept at 450⬚C.
Fig. 3. XRD patterns for Sb-doped SnO2 films at different deposition temperatures. The input gas ratio and RF power were kept at R s 1.12 and 30 W, respectively.
changes in the relative peak intensity. These phenomena with Sb doping were similarly reported by Rohatgi et al. w15x and Carroll et al. w16x. In the present case the relative presence of Ž110. along with other orientations remained invariant for subsequent doping levels. This behavior with Sb doping means that in the present case antimony incorporation in SnO 2 lattice has not af-
K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
232
fected the structural properties to a considerable extent. These results can be interpreted using the lattice parameters with the dopant levels. While calculating the lattice parameters it should be remembered that the Sb incorporation in the SnO 2 lattice will depend on the dopant level. According to our previous paper w18x, the calculated lattice parameters Ž a-axis and c-axis. for pure SnO 2 were approxi˚ and 3.185 A, ˚ respectively. In dopedmately 4.735 A SnO 2 , the initial increase of lattice constants is considered to be due to the dissolution of Sb into the SnO 2 lattice; at initial dopant levels one expects the Sb incorporation only at substitutional sites of Sn. For higher dopant levels the incorporation will take place at interstitial sites, and it was suggested that some phase like antimony oxides ŽSb 2 O 3 , Sb 2 O4 and Sb 2 O5 . could appear. Our present work in this context indicated that Sb doping up to approximately Ž R s 2.0. in such films took place at substitutional sites only. Therefore, in the present case antimony incorporation in SnO 2 lattice has not affected the structural properties to a considerable extent. However, the chemical composition of the deposited films at each input gas ratio Ž R values. was analyzed through EPMA. The results showed SbrSn atomic ratio was approximately 0.05 when the R value was 1.12, which was the optimal input gas ratio in the present experiment. The XRD patterns for the films did not vary significantly with changes in RF power as can be seen in Fig. 4b. 3.2. Electrical properties Fig. 5 shows the variation of carrier concentration, mobility, and resistivity for ATO films as a function of deposition temperature. A gradual decrease in the resistivity occurs on raising the temperature of deposition
up to 450⬚C. This behavior was attributed to the increase of both carrier concentration and carrier mobility as shown in Fig. 5. However, further raising the deposition temperature above 500⬚C resulted in a little increase again of the film resistivity. In the present study, the deposition temperature 450⬚C was the critical temperature for deposition at which the lowest resistivity was achieved. Since the deposition temperature of 450⬚C produces the best values of electrical conductivity, all subsequent studies have been carried out on the films prepared at 450⬚C. Electrical properties of ATO films are shown in Fig. 6 as a function of input gas ratio, R, at constant deposition temperature of 450⬚C with RF power of 30 W. As the R value increased, the resistivity first decreased, then reached a minimum and increased again at high R values. It is evident that the earlier decrease of the resistivity resulted mainly from an increase of electron concentration due to Sb doping in SnO 2 w21x, whereas the considerable increase of resistivity at high R value was caused by the reduction of both electron concentration and electron mobility w16x. The increase of carrier concentration with doping is due to the replacement of a Sn atom by the added Sb atom w22x, which gives one extra charge carrier. However, mobility increased up to doping levels of R s 0.7 and decreased continuously for further addition of the dopant. It is known that the grain boundary scattering w23x and ionized impurity scattering w24x play a major role in determining the mobility variation in such extrinsically doped semiconductors. The effect of RF power at constant R value Ž R s 1.12. on the electrical properties was shown in Fig. 7. As RF power increased, mobility increased and carrier concentration decreased. The slight increase of resistivity with RF power was due to the decrease of carrier
Fig. 5. Variation of resistivity, mobility and carrier concentration in Sb-doped SnO 2 films with deposition temperature.
K. Kim et al. r Surface and Coatings Technology 138 (2001) 229᎐236
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Fig. 6. Variation of resistivity, mobility and carrier concentration in Sb-doped SnO 2 films with input gas ration Ž R value..
concentration, because mobility exhibited a slight increase with RF power w20x. 3.3. Surface morphology Fig. 8 shows the surface morphologies by SEM and surface roughness by AFM for the films deposited at
various conditions. Fig. 8a,b represents the surface morphologies of the undoped tin oxide films deposited by T-CVD and PE-CVD. The facet ripples were shown in the films deposited by T-CVD, however, the films prepared by PE-CVD consisted of fine ripples. The degree of surface roughness for deposited films prepared by PE-CVD were lower than in the case of
Fig. 7. Variation of resistivity, mobility and carrier concentration in Sb-doped SnO 2 films with RF power. Table 1 RMS-surface roughness values of Sb-doped SnO 2 films at two different deposition methods Method
Input gas ratio Rws ŽPSbCl5rPSnCl4 .x
Deposition temperature Ž⬚C.
RF power ŽW.
T-CVD PE-CVD PE-CVD PE-CVD
0 0 1.12 2.63
450 450 450 450
᎐ 30 30 30
RMS-roughness Žnm. 16 12 6 4
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Fig. 8. SEM and AFM images of undoped and Sb-doped SnO 2 films: Ža. undoped SnO 2 film by T-CVD, Žb. undoped SnO 2 film by PE-CVD, Žc. Sb-doped SnO 2 film by PE-CVD Ž R s 1.12., Žd. Sb-doped SnO 2 film by PE-CVD Ž R s 2.63.. The deposition temperature was kept at 450⬚C.
Fig. 8. Ž Continued..
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Fig. 9. AFM images Ž1 m = 1 m. of the Sb-doped SnO 2 films by PE-CVD at various deposition temperatures Ža. 250⬚C, Žb. 350⬚C, Žc. 450⬚C, and 500⬚C. The input gas ratio and RF power were kept in Rs 1.12 and 30 W, respectively.
T-CVD, as shown in Fig. 8a,b, and in Table 1 where the RMS roughness values are summarized. This behavior can be assumed to be due to the addition of plasma energy w20x. From Fig. 8b᎐d, it was found that as the R value increased, the degree of surface roughness decreased. This phenomenon can be explained by two possible reasons. First, with the increasing of doping concentration the deposited films lost the crystallinity w16x. Secondly, the larger of the doping concentrations induced precipitation such as the antimony oxide phase ŽSb 2 O 3 , Sb 2 O4 and Sb 2 O5 etc.. w18x; and the preferred orientation growth of SnO 2 films may be suppressed by the precipitation. As a result, the surface roughness could be decreased with the increasing of R values. The surface roughness values and surface morphologies at different deposition temperatures are shown in Fig. 9 and Table 2, respectively. From Table 2, the degree of surface roughness increased with increasing deposition temperature. It would generally be considered that the surface morphology of deposited films is related with surface thermal energy and mass diffuTable 2 RMS surface roughness values of Sb-doped SnO 2 films prepared by PE-CVD at various deposition temperatures Input gas ratio Deposition RF power RMS roughness Žnm. Rws ŽPSbCl5 rPSnCl4 .x temperature Ž⬚C. ŽW. 1.12 1.12 1.12 1.12
250 350 450 500
30 30 30 30
4 6 12 16
sion of constituent elements. As shown in Fig. 9 at lower deposition temperature such as 250 and 350⬚C, the fine ripples or grains were formed because of a relatively lower thermal energy. However, at higher temperatures, the surface mobility of constituent elements increases, and the growth of facet grains are preferred w20x. As a result, the degree of surface roughness increased with increasing deposition temperature.
4. Conclusions Sb-doped SnO 2 films have been prepared by the plasma-enhanced chemical vapor deposition technique using a gas mixture of SnCl 4 ᎐SbCl 5 ᎐O 2 ᎐Ar and the following results were obtained. Ž1. PE-CVD method largely increased the deposition rate and smoothed the surface of tin oxide films compared with thermal CVD. Ž2. The highly Ž110. oriented polycrystalline films were grown above 350⬚C, while amorphous-like films were grown below 250⬚C. Ž3. The Sb-doped SnO 2 films which had the relatively good electrical properties were formed at 450⬚C, input gas ratio R s 1.12, and RF power 30 W. Ž4. Higher input gas ratio and lower deposition temperature led to a decrease in the surface roughness of the deposited films. References w1x K.L. Chopra, S. Major, D.K. Pandya, Thin Solid Films 102 Ž1983. 1᎐46. w2x Y.K. Fang, J.J. Lee, Thin Solid Films 169 Ž1989. 51᎐56.
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w3x L.D. Angelis, N. Minnaja, Sensors Actuators B 3 Ž1991. 197᎐204. w4x S.G. Ansari, S.W. Gosavi, S.A. Gangal, R.N. Karekar, R.C. Aiyer, J. Mater. Sci.: Mater. Electron. 8 Ž1997. 23᎐27. w5x Z.M. Jarzebski, J.P. Marton, J. Electrochem. Soc. 123 Ž7. Ž1976. 199᎐205. w6x Z.M. Jarzebski, J.P. Marton, J. Electrochem. Soc. 123 Ž10. Ž1976. 333᎐346. w7x F.J. Yusta, M.L. Hichman, H. Shamlian, J. Mater. Chem. 7 Ž8. Ž1997. 1421᎐1427. w8x C.F. Wan, R.D. McGrath, W.F. Keenan, S.N. Frank, J. Electrochem. Soc. 136 Ž5. Ž1989. 1459᎐1463. w9x S. Reddy, A.K. Malik, S.R. Jawalekar, Thin Solid Films 143 Ž1986. 113. w10x M. Mizuhashi, J. Non-Cryst. Solids 38 Ž1980. 329. w11x D.B. Fraser, H.D. Cook, J. Electrochem. Soc. 119 Ž1972. 1368. w12x D.E. Carlson, J. Electrochem. Soc. 122 Ž1975. 1334. w13x T. Arai, J. Phys. Soc. Jpn. 15 Ž5. Ž1960. 916᎐927.
w14x D. Elliott, D.L. Zellmer, H.A. Laitinen, J. Electrochem. Soc. 117 Ž1970. 1343᎐1348. w15x A. Rohatgi, T.R. Viverito, L.H. Slack, J. Am. Ceram. Soc. 57 Ž6. Ž1974. 278᎐279. w16x A.F. Carroll, L.H. Slack, J. Electrochem. Soc. 123 Ž12. Ž1976. 1889᎐1893. w17x N.S. Murty, S.R. Jawalekar, Thin Solid Films 108 Ž1983. 277᎐283. w18x K.H. Kim, S.W. Lee, J. Am. Ceram. Soc. 77 Ž4. Ž1994. 915᎐921. w19x R.N. Ghoshtagore, J. Electrochem. Soc. 125 Ž1. Ž1978. 110᎐117. w20x S. Shirakata, A. Yokoyama, S. Isomura, J. Appl. Phys. Jpn. 35 Ž1996. L722᎐L724. w21x C.A. Vincent, J. Electrochem. Soc. 119 Ž4. Ž1972. 515᎐518. w22x W.A. Badawy, E.A. El-Taher, Thin Solid Films 158 Ž1988. 277᎐284. w23x E. Shanthi, V. Dutta, A. Banerjee, K.L. Chopra, J. Appl. Phys. 51 Ž12. Ž1980. 6243᎐6251. w24x M.N. Islam, M.O. Hakim, J. Mater. Sci. Lett. 4 Ž1986. 1125.