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Growth rate and microstructure of copper thin films deposited with metal-organic chemical vapor deposition from hexafluoroacetylacetonate copper(I) allyltrimethylsilane
Thin Solid Films 335 (1998) 229±236
Growth rate and microstructure of copper thin ®lms deposited with metalorganic chemical vapor deposition from hexa¯uoroacetylacetonate copper(I) allyltrimethylsilane Jong-Hoon Son, Man-Young Park, Shi-Woo Rhee* Laboratory for Advanced Materials Processing (LAMP), Department of Chemical Engineering, Pohang University of Science and Technology (POSTECH), Pohang 790-784, South Korea Received 27 January 1998; accepted 6 May 1998
Abstract Metal-organic chemical vapor deposition of copper using the copper(I) compound, (hfac)Cu(ATMS) (hfac hexafluoroacetylacetonate, ATMS allyltrimethylsilane as a precursor, was carried out on TiN surface over a substrate temperature range of 60,2758C. It was found that the deposition temperature could be substantially lower compared with (hfac)Cu(VTMS) (VTMS vinyltrimethylsilane). In the substrate temperatures ranging from 60 to 908C, the Arrhenius plot showed a reaction-rate-limited regime with an activation energy of 15.0 kcal/mol. Above 908C, the deposition rate showed a feed-rate-limited regime with an activation energy of 0.1 kcal/mol. The copper ®lms contained no detectable impurities by Auger electron spectroscopy and gave resistivities below 2.0 mV cm in the temperature range of 125,1708C. As substrate temperature increased, the small-grained, smooth and continuous ®lm structure changed to large-grained and rough ®lm structure that was poorly connected and resulted in high resistivities. The polycrystalline phases with a preferred orientation of (111) and loss of selectivity were observed over a wide range of substrate temperatures. q 1998 Elsevier Science S.A. All rights reserved. Keywords: Copper; Metallization; Metal-organic chemical vapor deposition (MOCVD); (hfac)Cu(ATMS)
1. Introduction Chemical vapor deposition of copper has been extensively studied for interconnect metal of ultra large scale integration (ULSI) circuits. Copper exhibits a lower resistivity (bulk resistivity 1:67 mV cm) than aluminum which has been the choice for the interconnection metal. Low resistivity of interconnect metal can yield signi®cant reduction in (resistance £ capacitance), or RC time delay. Copper is also predicted to display enhanced electromigration and stress resistance, thus leading to higher reliability and improved performance [1]. It can be deposited either by sputtering or by metal-organic chemical vapor deposition (MOCVD). Much attention, however, has been placed on MOCVD because the CVD process has several advantages over physical vapor deposition techniques including the ability to achieve conformal coverage and the possibility of selective deposition [2,3]. Copper CVD precursors can be divided into two groups, i.e. Cu(I) and Cu(II) compounds. Cu(II) compounds such as Cu(b -diketonate)2 were initially used as a source material. * Corresponding author. Tel.:182-562-279-2265; fax:182-562-2793590; e-mail: [email protected].
The derivative Cu(hfac)2 (hfac hexafluoroacetyl ) acetonate deposits pure copper ®lms with quite low deposition rate only over the narrow temperature range of 340,3808C [4] and the addition of a reducing agent such as H2 was needed to lower the substrate temperature [5±8]. Therefore, recent studies have been focused on the series of Cu(I) compounds with the general empirical formula (b -diketonate)CuL, where L is the Lewis base ligand. These include (PMe3trimethylphosphine) [9,10], (hfac)Cu(PMe3) (hfac)Cu(2-butyne) [11±13], (hfac)Cu(1,5-COD) (COD cyclooctadiene) [14±16], (hfac)Cu(VTMS) (VTMS vinyltrimethylsilane) [17±22], and (hfac)Cu(VTMOS) (VTMOS vinyltri ÿ methoxysilane) [23,24]. These compounds have exhibited improved deposition characteristics over Cu(II) compounds including suf®cient volatility, highpurity copper ®lm, low substrate temperature and high deposition rate. It is well known that they deposit highpurity copper ®lms via thermally-induced disproportionation reaction rather than thermal decomposition on a heated substrate according to Eq. (1) [25±27]: ÿ ÿ 2 b 2 diketonate CuI L ! Cu0 1 CuII b 2 diketonate 2 12L (1)
0040-6090/98/$ - see front matter q 1998 Elsevier Science S.A. All rights reserved. PII S0 040-6090(98)008 68-2
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J.H. Son et al. / Thin Solid Films 335 (1998) 229±236
Fig. 1. Schematic diagram of Cu MOCVD apparatus.
Variation of the nature of neutral ligand L is known to affect the precursor properties and the properties of the copper thin ®lms deposited from MOCVD with these precursors. For example, Dubois et al. in their temperature programmed desorption (TPD) studies, showed that CuL bond strength has an signi®cant effect on the deposition rate and selective deposition [28]. In this study, we report a study of Cu MOCVD from hexa¯uoroacetylacetonate copper(I) allyltrimethylsilane in the presence of Ar carrier gas. We investigated the resistivity, purity, crystal structure and surface morphology of the deposited ®lm at various reaction conditions. Kinetic data such as an activation energy for the copper deposition on TiN were also evaluated. 2. Experimental procedures The schematic diagram of the MOCVD reactor is shown in Fig. 1. The precursor was (hfac)Cu(ATMS) manufactured by Mitsubishi Materials (Japan), which was loaded into the precursor vessel in a nitrogen atmosphere dry box. The precursor was carried to the reactor by Ar carrier gas at a ¯ow rate of 30,100 sccm. To supply a suf®cient amount of (hfac)Cu(ATMS) without thermal degradation, the precursor source was maintained at 408C throughout the run and the feed line at 508C to prevent precursor condensation. The susceptor was resistively heated, and then set at the desired value at the range of 60,2758C before each run. The total pressure in the reactor was adjusted to 1,3 Torr by the throttle valve between the pump and the reaction chamber and measured by Gransville±Philips 275 Convectron gauge attached to the reactor. The reactor was evacuated with a rotating-vane pump at its full pumping speed.
The Si wafer coated with sputtered TiN and another coated with thermally oxidized SiO2 were used as a substrate. TiN layer had a face-centered cubic (fcc) structure with a polycrystalline texture. TiN substrate was rinsed with deionized water, and then dried by the blow of N2 gas. No further chemical cleaning method was carried out for TiN substrate. SiO2 substrate was cleaned in H2SO41H2O2 (3:1) solution for 10 min followed by deionized water rinse and N2 gas drying. The substrates were then loaded into the reactor and evacuated to base pressure before the deposition. The deposition rate was calculated from the increase in weight of the substrate, and con®rmed by cross-sectional scanning electron microscopy (XSEM). The sheet resistance was determined by the four-point probe method. To reduce the effect of electron scattering from ®lm surfaces called `size effect' [29], samples with the thickness of over 400 nm were used. Measurements were taken at several positions and then averaged. An X-ray diffractometer (XRD, MAC Science M18XHF) with copper Ka radiation was used to determine the crystal structure and the preferred orientation of Cu thin ®lm. The ®lm purity was evaluated by Auger electron spectroscopy (AES, Perkin±Elmer PHI600). The surface morphology of the ®lm was observed using a scanning electron microscope (SEM, JEOL JSM840A) and atomic force microscope (AFM, Park Scienti®c Instruments Autoprobe-CP). 3. Results and discussion In order to gain insight into the reaction mechanism, the deposition rates were measured over a substrate temperature range of 60,3008C at constant reactor pressure of 3 Torr.
J.H. Son et al. / Thin Solid Films 335 (1998) 229±236
Fig. 2. The Arrhenius plot of the copper deposition rate on TiN using (hfac)Cu(ATMS). Reactor pressure, 3.0 Torr, carrier gas ¯ow rate, Ar 100 sccm.
The substrates used were TiN layer on silicon wafer. Fig. 2 shows the Arrhenius plot at this reaction condition along with the deposition rate with (hfac)Cu(ATMS). The minimum temperature that deposition of copper occurred was as low as 608C and the deposition rate increased with increasing substrate temperature. Jain et al. and many other
231
researchers have proposed that dissociation of (hfac)CuL to L and reaction intermediate (hfac)Cu is the rate-limiting step in the reaction mechanism of copper nucleation [3,15,23]. Therefore, such a low deposition temperature compared with that of other Cu(I) precursors is believed to be related with weaker Cu-L bond of (hfac)Cu(ATMS). Between 60 and 908C, the plot is linear and shows reactionrate-limited regime. The linear plot gives activation energy of 15.0 kcal/mol at this regime. Above 908C, deposition become feed-rate-limited. In the feed-rate-limited regime, the deposition rate become insensitive to the substrate temperature and the activation energy calculated was 0.1 kcal/mol. Fig. 3 shows typical Auger elemental survey (AES) spectra from copper thin ®lms deposited at different substrate temperatures of 75 and 1758C. Some impurities such as carbon, oxygen, and ¯uorine were detected at the surfaces of each sample. However, after argon-ion sputtering to remove the top layers, all evidence of C, O, F disappeared within the detection limits, indicating that the impurities incorporated into the bulk during deposition were minimal. Fig. 4 shows Auger depth pro®les of the copper ®lm deposited at the substrate temperature of 2758C. These pro®les also reveal that carbon and oxygen exist only on the top layers of copper thin ®lm. The atomic concentrations detected in the Cu thin ®lm layer were 97.8±99.6% copper, 0.11±0.57% carbon, and 0.13±0.32% oxygen. The electrical resistivity of the copper ®lm was measured by a four-point probe method. Fig. 5 shows the change of resistivity as a function of the substrate temperature. The ®lms deposited at the temperature between 125 and 1758C showed resistivities as low as the value of bulk copper (1.67 mV cm), which indicates that these copper thin ®lms are
Fig. 3. Typical AES spectra from the deposited copper ®lm (a) T s 758C, (b) T s 1758C.
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Fig. 4. AES depth pro®les from the deposited copper ®lm at the substrate temperature of 2758C.
highly pure and have well-connected grains. However, below 1258C the resistivity slightly increased as the substrate temperature decreased. Such an increase of resistivity below 1258C appears to be due to the grain-boundary (GB) scattering of electrons [5,15,30]. As the crystallite size becomes smaller than the electron mean-free path, a contribution to the ®lm resistivity from GB scattering arises. We believe that above 1258C, grains became large enough to minimize GB scattering and this phenomena disappeared. The increase of resistivity above 1758C may be related to the connection among the copper grains but not with ®lm purity because AES depth pro®les in Fig. 4 proved that the copper ®lms above 1758C are highly pure. To investigate the morphology changes with increasing substrate temperature, SEM, XSEM and AFM studies were evaluated. Fig. 6 shows SEM and XSEM photographs of copper ®lms grown on TiN substrate with various substrate Ê temperatures. The copper ®lm thickness was about 7000 A for all copper thin ®lms. In the low temperature region of 75,1258C, small and ®ne crystallites formed smooth surfaces and well-connected copper ®lms. The grains observed in this temperature region appear to grow by the Volmer±Weber growth mechanism, i.e. the 3D island growth, and these closely packed grains acted as a continuous layer on which further nucleation of copper can occur. However, between 125 and 1758C, the morphology of copper ®lms changed signi®cantly. The 3D island growth mechanism of grains was also obvious in the high temperature region. However, the ®lms consisted of much larger grains which have their bottoms on the Cu/TiN interface and the surface of the ®lm was rougher than those deposited at low temperatures. The root mean square (RMS) value of surface roughness as a function of substrate temperature was calculated by AFM measurement and plotted in Fig. 7.
Fig. 5. Resistivity of the copper ®lm as a function of the substrate temperature.
Below 1258C, RMS surface roughness was almost constant Ê . But when substrate temperature exceeded the at 300 A transition region of 125,1758C, surface roughness increased linearly with increasing substrate temperature. It seems that this morphology change is related to the increased mobility of adatoms or adsorbed species on the substrate [24,31]. At higher temperatures, the copper grains may continuously coarse due to faster surface migration of adatoms, whereas small and ®ne crystallites appear at lower temperatures since the surface diffusion is relatively slow. We also believe that the grains in the high temperature region grew not only by grain growth but also by the coalescence of large grains [22]. Generally, coalescence of grains decreases the island density, resulting in local denuding of the substrate where further nucleation then occurs. Threedimensional AFM images in Fig. 8 shows the surface morphology changes of copper thin ®lms deposited at 75 and 1758C with increasing ®lm thickness. In the case of 758C samples, there was no signi®cant change of surface morphology with various ®lm thickness. On the other hand, at 1758C, surface roughness increased and apparent number density of grains gradually decreased as the ®lm thickness increased, which is a clear evidence for grain coalescence. And we were also able to observe some small grains growing on the substrate between large grains by XSEM analysis in Fig. 6h±j. The ®lm structure observed in the high temperature region can lead to the formation of voids between grains. Thus the high resistivity of the ®lms deposited at these high temperatures is probably due to these poorly-connected ®lm structures, which was indicated also by many other researchers [9,18,23].
Fig. 6. SEM and XSEM photographs of the copper ®lms deposited at various substrate temperatures: (a,f) T s 758C; (b,g) T s 1258C; (c,h) T s 1758C; (d,i) T s 2258C; (e,j) T s 2758C.
J.H. Son et al. / Thin Solid Films 335 (1998) 229±236
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Fig. 9 shows an XRD spectrum of copper ®lms deposited at various substrate temperatures on TiN. All peaks corresponded to those of pure copper. There was no peak associated with the oxidized forms of copper. To quantitatively analyze the preferred orientation, the ratio of I(111)/I(200) was plotted as shown in Fig. 10. The copper thin ®lms had a fcc structure with (111) preferred orientation at all substrate temperatures. Since the preferred orientation is known to keep the metal layer from electromigration, these thin ®lms would show high resistance to failure [32]. 4. Conclusions Fig. 7. RMS value of surface roughness as a function of substrate temperaÊ. ture. Film thickness, 4000 A
Pure copper thin ®lm was deposited from (hfac)Cu(ATMS) in the substrate temperature range of 60,2758C. It was found that the deposition temperature
Fig. 8. Three-dimensional AFM images of the copper ®lms with various ®lm thickness. (a) 758C, 100 nm; (b) 758C, 400 nm; (c) 758C, 700 nm (d) 1758C, 100 nm; (e) 1758C, 400 nm; (f) 1758C, 700 nm.
J.H. Son et al. / Thin Solid Films 335 (1998) 229±236
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logical transition region, small and closely packed grain structure turned into large individual grain structure with dramatic increase of surface roughness. The copper ®lms had a fcc structure with (111) preferred orientation at all substrate temperatures. Acknowledgements The Authors would like to thank LG Semicon Co., for the ®nancial support and Mitsubishi Materials Co. for the precursor. References
Fig. 9. XRD patterns of the copper ®lms deposited at various temperatures. Ê. Film thickness, 4000 A
could be substantially lower compared with (hfac)Cu(VTMS) (VTMS vinyltrimethylsilane). From 60 to 908C, the Arrhenius plot showed a reaction-rate-limited regime with an activation energy of 15.0 kcal/mol. Above 908C, deposition became feed-rate-limited with an activation energy of 0.1 kcal/mol. The copper ®lms obtained were highly pure, and the resistivity reached the value of bulk copper at the substrate temperature of 125,1758C. For the copper thin ®lms deposited below 1258C and above 1758C, the increase of resistivity was observed, and it is thought to be resulted from structural problems such as electron scattering between grains in the ®lm deposited at low temperature and poor electrical connection among the copper grains in the ®lm deposited at high temperature. There was a signi®cant change in the ®lm morphology between 125 and 1758C. In this morpho-
Fig. 10. Variation of I(111)/I(200) ratio with different substrate temperatures.
[1] S.P. Murarka, R.J. Gutmann, A.E. Kaloyelos, W.A. Lanford, Thin Solid Films 236 (1993) 257. [2] J.A.T. Norman, B.A. Muratore, P.N. Dyer, D.A. Roberts, A.K. Hochberg, L.H. Dubois, Mater. Sci. Eng. B 17 (1993) 87. [3] M.J. Hampden-Smith, T.T. Kodas, Polyhedron 14 (6) (1995) 699. [4] D. Temple, A. Reisman, J. Electrochem. Soc. 136 (11) (1989) 3525. [5] A.E. Kaloyelos, A. Feng, J. Garhart, K.C. Brooks, S.K. Ghosh, A.N. Saxena, F. Luehers, J. Electronic. Mater. 19 (3) (1990) 271. [6] W.G. Lai, Y. Xie, G.L. Grif®n, J. Electrochem. Soc. 138 (11) (1991) 3499. [7] D.H. Kim, R.H. Wentorf, W.N. Gill, J. Electrochem. Soc. 140 (11) (1993) 3267. [8] D.H. Kim, R.H. Wentorf, W.N. Gill, J. Appl. Phys. 74 (8) (1993) 5164. [9] H.K. Shin, K.M. Chi, M.J. Hampden-Smith, T.T. Kodas, J.D. Farr, M. Paffett, Angew. Chem. Adv. Mater. 3 (5) (1991) 246. [10] H.K. Shin, K.M. Chi, M.J. Hampden-Smith, T.T. Kodas, J.D. Farr, M. Paffett, Chem. Mater. 4 (4) (1992) 788. [11] A. Jain, K.M. Chi, T.T. Kodas, M.J. Hampden-Smith, J.D. Farr, M.F. Paffett, Chem. Mater. 3 (6) (1991) 995. [12] T.H. Baum, C.E. Larson, Chem. Mater. 4 (2) (1992) 365. [13] T.H. Baum, C.E. Larson, J. Electrochem. Soc. 140 (1) (1993) 154. [14] S.K. Reynolds, C.J. Smart, E.F. Baran, T.H. Baum, C.E. Larson, P.J. Brock, J. Appl. Phys. Lett. 59 (1991) 2332. [15] A. Jain, K.M. Chi, M.J. Hampden-Smith, T.T. Kodas, J.D. Farr, M.F. Paffett, J. Mater. Res. 7 (2) (1992) 261. [16] R. Kumar, F.R. Fronczek, A.W. Maverick, W.G. Lai, G.L. Grif®n, Chem. Mater. 4 (3) (1992) 577. [17] J.A.T. Norman, D.A. Roberts, A.K. Hochberg, Proc. 12th Int. Symp. on CVD, Vol. 93-2 p. (1993) 221. [18] A. Jain, K.M. Chi, T.T. Kodas, M.J. Hampden-Smith, J. Electrochem. Soc. 140 (5) (1993) 1434. [19] E.S. Choi, S.K. Park, H.H. Lee, J. Electrochem. Soc. 143 (2) (1996) 624. [20] W.J. Lee, J.S. Min, S.K. Rha, S.S. Chun, C.O. Park, D.W. Kim, J. Mater. Sci. Mater. Electron. 7 (1996) 111. [21] G. Braeckelmann, D. Manger, A. Burke, J. Vac. Sci. Technol. B 14 (3) (1996) 1828. [22] H.Y. Yoen, Y.B. Park, S.W. Rhee, J. Mater. Sci. Mater. Electron. 8 (1997) 189. [23] E.S. Choi, S.K. Park, H.K. Shin, H.H. Lee, Appl. Phys. Lett. 68 (7) (1996) 1017. [24] C.H. Jun, Y.T. Kim, J.T. Baek, H.J. Yoo, D.R. Kim, J. Vac. Sci. Technol. A 14 (6) (1996) 3214. [25] S.L. Cohen, M. Liehr, S. Kasi, Appl. Phys. Lett. 60 (1) (1992) 50. [26] S.L. Cohen, M. Liehr, S. Kasi, J. Vac. Sci. Technol. A 10 (4) (1992) 863. [27] J. Farkas, M.J. Hampden-Smith, T.T. Kodas, J. Phys. Chem. 98 (27) (1994) 6763.
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[28] L.H. Dubois, B.R. Zegarski, J. Electrochem. Soc. 139 (11) (1992) 3295. [29] E.H. Sondheimer, Adv. Phys. 1 (1951) 1. [30] A.F. Mayadas, M. Shatzkes, Phys. Rev. B 1 (1970) 1382.
[31] R.F. Bunshah, Handbook of Deposition Technologies for Films and Coatings, 2nd ed., Noyes, New Jersey, 1994, p. 693. [32] S.M. Sze, VLSI Technology, 2nd ed., McGraw±Hill, New York, 1988, p. 411.
Growth rate and microstructure of copper thin ®lms deposited with metalorganic chemical vapor deposition from hexa¯uoroacetylacetonate copper(I) allyltrimethylsilane Jong-Hoon Son, Man-Young Park, Shi-Woo Rhee* Laboratory for Advanced Materials Processing (LAMP), Department of Chemical Engineering, Pohang University of Science and Technology (POSTECH), Pohang 790-784, South Korea Received 27 January 1998; accepted 6 May 1998
Abstract Metal-organic chemical vapor deposition of copper using the copper(I) compound, (hfac)Cu(ATMS) (hfac hexafluoroacetylacetonate, ATMS allyltrimethylsilane as a precursor, was carried out on TiN surface over a substrate temperature range of 60,2758C. It was found that the deposition temperature could be substantially lower compared with (hfac)Cu(VTMS) (VTMS vinyltrimethylsilane). In the substrate temperatures ranging from 60 to 908C, the Arrhenius plot showed a reaction-rate-limited regime with an activation energy of 15.0 kcal/mol. Above 908C, the deposition rate showed a feed-rate-limited regime with an activation energy of 0.1 kcal/mol. The copper ®lms contained no detectable impurities by Auger electron spectroscopy and gave resistivities below 2.0 mV cm in the temperature range of 125,1708C. As substrate temperature increased, the small-grained, smooth and continuous ®lm structure changed to large-grained and rough ®lm structure that was poorly connected and resulted in high resistivities. The polycrystalline phases with a preferred orientation of (111) and loss of selectivity were observed over a wide range of substrate temperatures. q 1998 Elsevier Science S.A. All rights reserved. Keywords: Copper; Metallization; Metal-organic chemical vapor deposition (MOCVD); (hfac)Cu(ATMS)
1. Introduction Chemical vapor deposition of copper has been extensively studied for interconnect metal of ultra large scale integration (ULSI) circuits. Copper exhibits a lower resistivity (bulk resistivity 1:67 mV cm) than aluminum which has been the choice for the interconnection metal. Low resistivity of interconnect metal can yield signi®cant reduction in (resistance £ capacitance), or RC time delay. Copper is also predicted to display enhanced electromigration and stress resistance, thus leading to higher reliability and improved performance [1]. It can be deposited either by sputtering or by metal-organic chemical vapor deposition (MOCVD). Much attention, however, has been placed on MOCVD because the CVD process has several advantages over physical vapor deposition techniques including the ability to achieve conformal coverage and the possibility of selective deposition [2,3]. Copper CVD precursors can be divided into two groups, i.e. Cu(I) and Cu(II) compounds. Cu(II) compounds such as Cu(b -diketonate)2 were initially used as a source material. * Corresponding author. Tel.:182-562-279-2265; fax:182-562-2793590; e-mail: [email protected].
The derivative Cu(hfac)2 (hfac hexafluoroacetyl ) acetonate deposits pure copper ®lms with quite low deposition rate only over the narrow temperature range of 340,3808C [4] and the addition of a reducing agent such as H2 was needed to lower the substrate temperature [5±8]. Therefore, recent studies have been focused on the series of Cu(I) compounds with the general empirical formula (b -diketonate)CuL, where L is the Lewis base ligand. These include (PMe3trimethylphosphine) [9,10], (hfac)Cu(PMe3) (hfac)Cu(2-butyne) [11±13], (hfac)Cu(1,5-COD) (COD cyclooctadiene) [14±16], (hfac)Cu(VTMS) (VTMS vinyltrimethylsilane) [17±22], and (hfac)Cu(VTMOS) (VTMOS vinyltri ÿ methoxysilane) [23,24]. These compounds have exhibited improved deposition characteristics over Cu(II) compounds including suf®cient volatility, highpurity copper ®lm, low substrate temperature and high deposition rate. It is well known that they deposit highpurity copper ®lms via thermally-induced disproportionation reaction rather than thermal decomposition on a heated substrate according to Eq. (1) [25±27]: ÿ ÿ 2 b 2 diketonate CuI L ! Cu0 1 CuII b 2 diketonate 2 12L (1)
0040-6090/98/$ - see front matter q 1998 Elsevier Science S.A. All rights reserved. PII S0 040-6090(98)008 68-2
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J.H. Son et al. / Thin Solid Films 335 (1998) 229±236
Fig. 1. Schematic diagram of Cu MOCVD apparatus.
Variation of the nature of neutral ligand L is known to affect the precursor properties and the properties of the copper thin ®lms deposited from MOCVD with these precursors. For example, Dubois et al. in their temperature programmed desorption (TPD) studies, showed that CuL bond strength has an signi®cant effect on the deposition rate and selective deposition [28]. In this study, we report a study of Cu MOCVD from hexa¯uoroacetylacetonate copper(I) allyltrimethylsilane in the presence of Ar carrier gas. We investigated the resistivity, purity, crystal structure and surface morphology of the deposited ®lm at various reaction conditions. Kinetic data such as an activation energy for the copper deposition on TiN were also evaluated. 2. Experimental procedures The schematic diagram of the MOCVD reactor is shown in Fig. 1. The precursor was (hfac)Cu(ATMS) manufactured by Mitsubishi Materials (Japan), which was loaded into the precursor vessel in a nitrogen atmosphere dry box. The precursor was carried to the reactor by Ar carrier gas at a ¯ow rate of 30,100 sccm. To supply a suf®cient amount of (hfac)Cu(ATMS) without thermal degradation, the precursor source was maintained at 408C throughout the run and the feed line at 508C to prevent precursor condensation. The susceptor was resistively heated, and then set at the desired value at the range of 60,2758C before each run. The total pressure in the reactor was adjusted to 1,3 Torr by the throttle valve between the pump and the reaction chamber and measured by Gransville±Philips 275 Convectron gauge attached to the reactor. The reactor was evacuated with a rotating-vane pump at its full pumping speed.
The Si wafer coated with sputtered TiN and another coated with thermally oxidized SiO2 were used as a substrate. TiN layer had a face-centered cubic (fcc) structure with a polycrystalline texture. TiN substrate was rinsed with deionized water, and then dried by the blow of N2 gas. No further chemical cleaning method was carried out for TiN substrate. SiO2 substrate was cleaned in H2SO41H2O2 (3:1) solution for 10 min followed by deionized water rinse and N2 gas drying. The substrates were then loaded into the reactor and evacuated to base pressure before the deposition. The deposition rate was calculated from the increase in weight of the substrate, and con®rmed by cross-sectional scanning electron microscopy (XSEM). The sheet resistance was determined by the four-point probe method. To reduce the effect of electron scattering from ®lm surfaces called `size effect' [29], samples with the thickness of over 400 nm were used. Measurements were taken at several positions and then averaged. An X-ray diffractometer (XRD, MAC Science M18XHF) with copper Ka radiation was used to determine the crystal structure and the preferred orientation of Cu thin ®lm. The ®lm purity was evaluated by Auger electron spectroscopy (AES, Perkin±Elmer PHI600). The surface morphology of the ®lm was observed using a scanning electron microscope (SEM, JEOL JSM840A) and atomic force microscope (AFM, Park Scienti®c Instruments Autoprobe-CP). 3. Results and discussion In order to gain insight into the reaction mechanism, the deposition rates were measured over a substrate temperature range of 60,3008C at constant reactor pressure of 3 Torr.
J.H. Son et al. / Thin Solid Films 335 (1998) 229±236
Fig. 2. The Arrhenius plot of the copper deposition rate on TiN using (hfac)Cu(ATMS). Reactor pressure, 3.0 Torr, carrier gas ¯ow rate, Ar 100 sccm.
The substrates used were TiN layer on silicon wafer. Fig. 2 shows the Arrhenius plot at this reaction condition along with the deposition rate with (hfac)Cu(ATMS). The minimum temperature that deposition of copper occurred was as low as 608C and the deposition rate increased with increasing substrate temperature. Jain et al. and many other
231
researchers have proposed that dissociation of (hfac)CuL to L and reaction intermediate (hfac)Cu is the rate-limiting step in the reaction mechanism of copper nucleation [3,15,23]. Therefore, such a low deposition temperature compared with that of other Cu(I) precursors is believed to be related with weaker Cu-L bond of (hfac)Cu(ATMS). Between 60 and 908C, the plot is linear and shows reactionrate-limited regime. The linear plot gives activation energy of 15.0 kcal/mol at this regime. Above 908C, deposition become feed-rate-limited. In the feed-rate-limited regime, the deposition rate become insensitive to the substrate temperature and the activation energy calculated was 0.1 kcal/mol. Fig. 3 shows typical Auger elemental survey (AES) spectra from copper thin ®lms deposited at different substrate temperatures of 75 and 1758C. Some impurities such as carbon, oxygen, and ¯uorine were detected at the surfaces of each sample. However, after argon-ion sputtering to remove the top layers, all evidence of C, O, F disappeared within the detection limits, indicating that the impurities incorporated into the bulk during deposition were minimal. Fig. 4 shows Auger depth pro®les of the copper ®lm deposited at the substrate temperature of 2758C. These pro®les also reveal that carbon and oxygen exist only on the top layers of copper thin ®lm. The atomic concentrations detected in the Cu thin ®lm layer were 97.8±99.6% copper, 0.11±0.57% carbon, and 0.13±0.32% oxygen. The electrical resistivity of the copper ®lm was measured by a four-point probe method. Fig. 5 shows the change of resistivity as a function of the substrate temperature. The ®lms deposited at the temperature between 125 and 1758C showed resistivities as low as the value of bulk copper (1.67 mV cm), which indicates that these copper thin ®lms are
Fig. 3. Typical AES spectra from the deposited copper ®lm (a) T s 758C, (b) T s 1758C.
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Fig. 4. AES depth pro®les from the deposited copper ®lm at the substrate temperature of 2758C.
highly pure and have well-connected grains. However, below 1258C the resistivity slightly increased as the substrate temperature decreased. Such an increase of resistivity below 1258C appears to be due to the grain-boundary (GB) scattering of electrons [5,15,30]. As the crystallite size becomes smaller than the electron mean-free path, a contribution to the ®lm resistivity from GB scattering arises. We believe that above 1258C, grains became large enough to minimize GB scattering and this phenomena disappeared. The increase of resistivity above 1758C may be related to the connection among the copper grains but not with ®lm purity because AES depth pro®les in Fig. 4 proved that the copper ®lms above 1758C are highly pure. To investigate the morphology changes with increasing substrate temperature, SEM, XSEM and AFM studies were evaluated. Fig. 6 shows SEM and XSEM photographs of copper ®lms grown on TiN substrate with various substrate Ê temperatures. The copper ®lm thickness was about 7000 A for all copper thin ®lms. In the low temperature region of 75,1258C, small and ®ne crystallites formed smooth surfaces and well-connected copper ®lms. The grains observed in this temperature region appear to grow by the Volmer±Weber growth mechanism, i.e. the 3D island growth, and these closely packed grains acted as a continuous layer on which further nucleation of copper can occur. However, between 125 and 1758C, the morphology of copper ®lms changed signi®cantly. The 3D island growth mechanism of grains was also obvious in the high temperature region. However, the ®lms consisted of much larger grains which have their bottoms on the Cu/TiN interface and the surface of the ®lm was rougher than those deposited at low temperatures. The root mean square (RMS) value of surface roughness as a function of substrate temperature was calculated by AFM measurement and plotted in Fig. 7.
Fig. 5. Resistivity of the copper ®lm as a function of the substrate temperature.
Below 1258C, RMS surface roughness was almost constant Ê . But when substrate temperature exceeded the at 300 A transition region of 125,1758C, surface roughness increased linearly with increasing substrate temperature. It seems that this morphology change is related to the increased mobility of adatoms or adsorbed species on the substrate [24,31]. At higher temperatures, the copper grains may continuously coarse due to faster surface migration of adatoms, whereas small and ®ne crystallites appear at lower temperatures since the surface diffusion is relatively slow. We also believe that the grains in the high temperature region grew not only by grain growth but also by the coalescence of large grains [22]. Generally, coalescence of grains decreases the island density, resulting in local denuding of the substrate where further nucleation then occurs. Threedimensional AFM images in Fig. 8 shows the surface morphology changes of copper thin ®lms deposited at 75 and 1758C with increasing ®lm thickness. In the case of 758C samples, there was no signi®cant change of surface morphology with various ®lm thickness. On the other hand, at 1758C, surface roughness increased and apparent number density of grains gradually decreased as the ®lm thickness increased, which is a clear evidence for grain coalescence. And we were also able to observe some small grains growing on the substrate between large grains by XSEM analysis in Fig. 6h±j. The ®lm structure observed in the high temperature region can lead to the formation of voids between grains. Thus the high resistivity of the ®lms deposited at these high temperatures is probably due to these poorly-connected ®lm structures, which was indicated also by many other researchers [9,18,23].
Fig. 6. SEM and XSEM photographs of the copper ®lms deposited at various substrate temperatures: (a,f) T s 758C; (b,g) T s 1258C; (c,h) T s 1758C; (d,i) T s 2258C; (e,j) T s 2758C.
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Fig. 9 shows an XRD spectrum of copper ®lms deposited at various substrate temperatures on TiN. All peaks corresponded to those of pure copper. There was no peak associated with the oxidized forms of copper. To quantitatively analyze the preferred orientation, the ratio of I(111)/I(200) was plotted as shown in Fig. 10. The copper thin ®lms had a fcc structure with (111) preferred orientation at all substrate temperatures. Since the preferred orientation is known to keep the metal layer from electromigration, these thin ®lms would show high resistance to failure [32]. 4. Conclusions Fig. 7. RMS value of surface roughness as a function of substrate temperaÊ. ture. Film thickness, 4000 A
Pure copper thin ®lm was deposited from (hfac)Cu(ATMS) in the substrate temperature range of 60,2758C. It was found that the deposition temperature
Fig. 8. Three-dimensional AFM images of the copper ®lms with various ®lm thickness. (a) 758C, 100 nm; (b) 758C, 400 nm; (c) 758C, 700 nm (d) 1758C, 100 nm; (e) 1758C, 400 nm; (f) 1758C, 700 nm.
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logical transition region, small and closely packed grain structure turned into large individual grain structure with dramatic increase of surface roughness. The copper ®lms had a fcc structure with (111) preferred orientation at all substrate temperatures. Acknowledgements The Authors would like to thank LG Semicon Co., for the ®nancial support and Mitsubishi Materials Co. for the precursor. References
Fig. 9. XRD patterns of the copper ®lms deposited at various temperatures. Ê. Film thickness, 4000 A
could be substantially lower compared with (hfac)Cu(VTMS) (VTMS vinyltrimethylsilane). From 60 to 908C, the Arrhenius plot showed a reaction-rate-limited regime with an activation energy of 15.0 kcal/mol. Above 908C, deposition became feed-rate-limited with an activation energy of 0.1 kcal/mol. The copper ®lms obtained were highly pure, and the resistivity reached the value of bulk copper at the substrate temperature of 125,1758C. For the copper thin ®lms deposited below 1258C and above 1758C, the increase of resistivity was observed, and it is thought to be resulted from structural problems such as electron scattering between grains in the ®lm deposited at low temperature and poor electrical connection among the copper grains in the ®lm deposited at high temperature. There was a signi®cant change in the ®lm morphology between 125 and 1758C. In this morpho-
Fig. 10. Variation of I(111)/I(200) ratio with different substrate temperatures.
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