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MOCVD growth of (100)-oriented CeO2 thin films on hydrogen-terminated Si(100) substrates
Materials Science and Engineering B54 (1998) 84 – 91
MOCVD growth of (100)-oriented CeO2 thin films on hydrogen-terminated Si(100) substrates Takaaki Ami *, Masayuki Suzuki SONY Corporation Research Center, 174, Fujitsuka-cho, Hodogaya-ku, Yokohama 240, Japan
Abstract Ceria (CeO2) with a fluorite structure is supposed to be an ideal buffer layer for fabricating epitaxially grown perovskite materials on silicon substrates because of its thermal stability and excellent lattice match with Si. On the other hand, there are still technical difficulties in forming CeO2(100), which is more attractive for applications compared to that of (111) orientation. We report the preparation of CeO2 thin films on hydrogen-terminated Si(100) substrates by metal-organic chemical vapor deposition (MOCVD) using Ce(DPM)4 as an MO-source. The surface microroughness of the substrates was investigated using Fourier-transform infrared-attenuated total reflection (FT-IR-ATR) and atomic force microscopy (AFM). Si-wafers cleaned by RCA method were immersed in H2O2-added 0.5% HF solutions for hydrogen-termination, and preferential Si – H2 vibrational peaks accompanied with weak Si–H3 peaks were observed. The root mean square roughness of the substrates estimated by AFM was around 0.2 nm. The crystallinity and orientation were characterized by X-ray diffraction (XRD). The films were polycrystalline, and preferentially orientated to the 100 direction. The preferential orientation factor defined by Lotgering was 0.83. The microstructures were characterized by a field-emission scanning electron microscopy (FE-SEM), an AFM and a transmission electron microscopy (TEM). The grains were ‘non-equiaxed’ columnar, growing perpendicular to the surface at the expense of other columns. The surface texture of a 50 nm-thick film characterized by AFM was rectangular in shape, with a typical size of 100×200 nm. They were aligned in the same direction, and the edges of the rectangular are parallel to the Si(011) facets, suggesting a possible in-plane orientation. Cross-sectional HR-TEM analyses of a 150 nm thick film verified the thickness of the amorphous layer formed at the CeO2/Si(100) interface to be around 2 nm. © 1998 Elsevier Science S.A. All rights reserved. Keywords: CeO2; Thin films; Si substrates; MOCVD
1. Introduction Introducing oxide materials except for SiO2 into the semiconductor industry is very attractive, because they exhibit various high performance properties. In particular, some oxides with a perovskite (-related) type structure show very important properties for semiconductor devices: superconducting, ferroelectric, high-o, and colossal magnetoresistance [1,2]. Ceria (CeO2) with a fluorite structure is supposed to be an ideal buffer layer for fabricating epitaxiallygrown perovskite materials on silicon substrates because of its thermal stability and excellent lattice match with Si, especially around 600°C [3]. The excellent epitaxy of ceria thin films was achieved on a Si(111) substrate without any amorphous layer in the vicinity of the boundary by pulsed laser deposition in an ultra* Corresponding author. 0921-5107/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5107(98)00133-0
high vacuum at room temperature [4]. However, the preferential epitaxial growth-orientation of ceria on Si(100) substrates reported was not CeO2(100) but CeO2(110) [5–7]. In order to fabricate a buffer layer epitaxially on Si(100) substrates, a solid solution of zirconia and ceria ((Zr1 − x Cex )O2) [8] or CeO2/SrTiO3/ Si structure [9] have also been proposed. Metal-organic chemical vapor deposition (MOCVD) is a widely used technique in the field of semiconductor industry because of its high uniformity in large wafers, growth rate of film, and conformal step coverage. For this reason, it is also utilized for growing ceria films on oxide single crystal substrates: SrTiO3 (STO), yttria stabilized zirconia (YSZ), magnesia (MgO) [10] and r-plane sapphire [11,12]. (100)-oriented ceria films were obtained on STO, YSZ and sapphire. Nevertheless, there still are technical difficulties in forming (100)-oriented ceria directly on silicon substrates, especially using MOCVD [13–16]. Although this technique was
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Fig. 1. The experimental setup of the MOCVD system used.
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described in Refs. [13] and [14], experimental data were not given. Meanwhile, atomic scale flatness of the surface and removal of the native oxides of silicon substrates are supposed to be important for epitaxial growth of films on these substrates. Hydrogen-termination by hydrofluoric acid (HF) or HF-based aqueous solution is one promising technique to remove the native oxide and passivate the silicon surface at the same time. A suggested solution proposed for smoothing the Si(100) surface is H2O2-added diluted HF, in which selective oxidation of the defects and etching of the oxide occurs [17,18]. In the present study, the concentrations of the base hydrofluoric acid and H2O2 in the solution were optimized.
2. Experimental
2.1. Hydrogen termination The silicon wafers used in this study were Borondoped Si(100) grown by CZ method with an acceptor concentration of 0.1 – 1.5× 105 cm − 3 (resistivity: 8–12 V cm). Wafers chemically cleaned by conventional SC-1 process were immersed into H2O2-added HF with various concentrations for 2 min. At the final stage of the optimization, the concentration of the HF base solution was fixed to 0.5 wt.% (pH 2.5), as the FT-IR-ATR spectra showed a reasonably good signal-to-noise ratio and absorbance, compared to higher (50% HF [18]) and lower (0.1% HF [17]) concentrations. The concentrated H2O2 (30%) was added to this concentration, and the volume of added H2O2 was optimized. The hydrogen-terminated surfaces were characterized by Fourier-transform infrared attenuated total reflection (FT-IR-ATR; JEOL JIR-6500) with Ge inner reflective crystal in a few minutes after the treatment. The surface roughness was characterized by atomic force microscopy (AFM; Park Scientific Instruments, Autoprobe).
source bottles, and mixed with oxygen gas in the reactor. The substrates were set on the SiC-coated graphite susceptor, and heated by a focused IR lamp.
2.3. Characterization The film thickness and the refractive index were measured by ellipsometry (Gaertner ellipsometer L115B) with He–Ne laser (632.8 nm). The film thickness was also characterized by cross-sectional images taken by a field emission-scanning electron microscope (FE-SEM; Hitachi S4500) with the microstructure of the films. The film composition on the surface and its depth-profile were characterized by Auger electron spectroscopy (AES; JEOL JAMP-30) and Rutherford back-scattering spectroscopy (RBS), respectively. The X-ray diffraction (XRD) was performed on a Rigaku RADIII diffractometer using Cu Ka radiation. For u−2u scans, a fine collimation with divergence slit (DS)-1/2°, scattering slit (SS)-1/2° and receiving slit (RS)-0.05 mm was selected in order to obtain high resolution, when needed. A glancing angle-diffractometer with an x-ray incident angle of 5° was also used in order to eliminate the diffraction peaks from the Si substrate. Surface morphologies were analyzed by AFM. The interface between CeO2 and Si substrates was observed by transmission electron microscope (TEM; JEOL 2000FX-II; 200kV).
3. Results and discussion
3.1. Optimization of the amount of the added H2O2 in 0.5% HF Fig. 2 shows the volume of the added H2O2 dependence of the absorbance of the SiH2, SiH3 peaks and full width of half maximum (FWHM) of the SiH2 peak. The absorbance of SiH2 is increased by adding a slight
2.2. Thin film deposition The experimental setup of the MOCVD system is schematically shown in Fig. 1. Metal-organic (MO)sources are set in Pyrex bottles, which are mounted in the box-type furnace and can be heated up to 250°C homogeneously. A reactor tube, its head fringe, and the manifold are also set in a box-type furnace with valves and heated up homogeneously in order to avoid the condensation of vaporized MO-sources. The MOsource used was Ce(DPM)4 with a vapor-pressure of 0.1 Torr at around 200°C. It was introduced by argon gas into the quartz hot-wall reactor through the
Fig. 2. The volume of the added H2O2 dependence of the absorbance of the SiH2, SiH3 peaks and FWHM of the SiH2 peak.
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Fig. 3. The surface AES spectra of films deposited with and without oxygen.
amount of H2O2. The FWHM of the dihydride peak was decreased with the increase of the absorbance, implying a homogeneous surface. When an excess amount of H2O2 was added to 0.5% HF, the absorbance decreased. As a result, the optimized amount of H2O2 was determined to be 0.2 vol.%. This result was supported by AFM, showing a minimum root mean square roughness of 0.2 nm at the optimized amount.
3.2. MO-source transport and film growth The transport of the MO-source and film deposition mechanism were investigated with the CVD parameters. Experiments for the investigation of the growth rate were performed with parameters of evaporation temperature of the MO-source and Ar-carrier flow rate at a substrate temperature (Tsub) of 700°C and a total reaction pressure (Ptotal) of 2 Torr. The partial pressure of oxygen (pO2) was fixed to 87.5% of the total pressure. The growth rate increased exponentially with increasing evaporation temperature in a range between 170 and 210°C. On the other hand, with the increase of Ar-carrier in a range between 10 and 100 sccm, the growth rate increased linearly showing a slight saturation. These results clearly show that the film growth is masstransport limited under these conditions. On the other hand, the Tsub-dependence of the growth rate is slight between 600 and 800°C. In contrast, the pO2 affects the film growth. Fig. 3 shows the surface AES spectra of films deposited with and without oxygen. The peak intensities of Ce and O are weaker in the film deposited without oxygen, while, the carbon content in this film is much more than that in the film deposited with oxygen. The results of XRD for the film deposited with 87.5% of pO2 showed the formation of single phase-polycrystalline ceria, as will be discussed later. From the result of RBS, the Ce/O ratio of the ceria was calibrated to be
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2.090.1. The refractive index of the films was 2.1–2.4, which is almost the same range as the reported value for bulk ceria [6]. On the other hand, no crystallized phases were detected in films deposited without oxygen by XRD. These results show that oxygen is required to decompose the Ce(DPM)4 and to form the CeO2. The film deposited without oxygen is supposed to be the Ce(DPM)4 itself or this material partially decomposed. However, the growth rate in the pO2 range between 12.5 and 87.5% was almost constant at a Tsub of 700°C and a Ptotal of 2 Torr. In order to suppress the formation of a-SiO2 on the Si surfaces, the initial stage of deposition was performed in argon gas, typically for a minute, then the argon was switched to oxygen to form CeO2 films. Typical growth conditions are shown in Table 1.
3.3. Crystallinity and orientation 3.3.1. XRD analysis by u− 2u scan Because of the excellent lattice matching of ceria to silicon, an unfortunate diffraction peak from the Si(100) substrates shows up in the XRD data collected by a u− 2u scan which is very close to the CeO2(200) peak. However, with the fine collimation described previously, the Si(200) peak was observed to be split from the CeO2(200) peak. The overlap of the Si(200) Cu Ka2 line and CeO2(200) Cu Ka1 line is still critical, but the contribution of the Si(200) Cu Ka2 line can be subtracted fairly accurately. After the reduction of the Cu Ka2 line, the peak separation can be performed with reasonable accuracy. Blown up XRD spectra (32°B 2uB 34°) for a ceria film with a Si substrate and it for a Si substrate during these procedures are presented in Figs. 4 and 5. These techniques reveal the diffraction peaks only from ceria films. Fig. 6 shows the XRD pattern of a ceria film on a Si(100) substrate by u− 2u scan without contributions from the substrate (bottom), together with the XRD of the CeO2 ceramics wafer in the same size as the ceria film (top). The film were grown at 700°C, 2 Torr, with a pO2 of 87.5%. The contribution Table 1 Typical growth conditions for MOCVD CeO2 thin films on Si Substrate Substrate temperature (°C) Total pressure (Torr) Partial pressure of oxygen (%) Source evaporation temperature (°C) Ar carrier gas flow rates (sccm) Ar dilution gas flow rates (sccm) O2 gas flow rates (sccm)
Si(100) 500 – 1000 1 – 10 0 – 87.5 170 – 210 10 – 100 0 – 600 0 – 700
The initial stage of deposition was performed in argon dilution gas just for a short time, then the oxygen was switched to the argon to form CeO2 films.
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Fig. 4. Blown up XRD spectra (32° B2u B 34°) for a ceria film with a Si substrate and it for a Si substrate before Cu Ka2 reduction. The vertical tick marks indicate the expected peak locations for CeO2(200) and Si(200).
Fig. 6. The XRD pattern of a ceria film on a Si(100) substrate by u −2u scan without contributions from the substrate (bottom), together with the XRD of the CeO2 ceramics wafer in the same size as the ceria film (top).
from the silicon substrates were subtracted in the XRD. Only a very broad (111) peak and a weak (311) peak were observed in 10° B2u B 65°, except for the intense (200) peak. The preferred orientation factor F is defined by Lotgering as follows [19]:
3.3.2. XRD analysis by glancing angle diffractometer The intensity measured by the glancing angle diffractometer is affected by the instrumentation, especially the solar slit collimator (SL), which restricts the diffracted X-ray [20]. For a preliminary analysis of the orientation, the orientation factor G for the glancingangle diffractometer was simply defined as:
F = (P− P0)/(1 − P0), where P= I(h00)/ I(hkl) and P0 is the P-value for a non-oriented specimen. P0 was calculated from the XRD of a CeO2 ceramics wafer. Diffraction peaks in 10°B 2u B65° were used for the calculation. The Fvalue for the sample shown in Fig. 6 was 0.83, showing a fairly strong (100) preferred orientation. Nevertheless, the peak intensity and FWHM of the Si(200) peak strongly depend on how the sample is set in the diffractometer. For systematic characterization of the orientation of the ceria films, glancing angle XRD measurements were performed in which no contribution from the silicon substrates appeared.
Fig. 5. Blown up XRD spectra (32° B2u B 34°) for a ceria film with a Si substrate and it for a Si substrate after Cu Ka2 reduction.
G= Q/Q0 where Q=I(200)/I(111), and Q0 is the Q-value for a non-oriented specimen. In order to minimize the error caused by the SL and absorption of X-ray, the two peaks: (200), (111) with the difference of small 2u angle (4.534°) were used. Fig. 7 shows the Tsub-dependence of the I(200), I(111) and orientation factor G. The I(200) and the G are maximized at around 700°C, where the I(111) is lowest. This result clearly shows that Tsub affects the crystallinity and orientation. In the high temperature region, the I(111) increased with Tsub, showing that the orientation changes from (100) to random or possible
Fig. 7. The substrate temperature dependence of the I(200), I(111) and orientation factor G.
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sure. Hereinafter, characterizations on the films grown at 700°C, 2 Torr, with a pO2 of 87.5% will be described.
3.4. Surface morphology of the ceria thin film (FE-SEM, AFM)
Fig. 8. A cross-sectional FE-SEM image of a 150 nm thick ceria film.
(111) orientation. On the other hand, the drastic decrease of I(200) and weak I(111) imply that the crystallinity goes down at low temperatures. The crystallinity and the orientation were also affected by pO2. At a fixed Tsub (700°C) and a Ptotal (2 Torr), the G increased with increasing pO2 in a range between 12.5 and 87.5%, demonstrating the decomposition of the MO-source and crystallization of the (100)oriented CeO2. On the other hand, at a fixed Tsub (700°C) and a pO2 (87.5%), the G decreased with increasing Ptotal in a range between 2 and 10 Torr, suggesting the decreasing crystallinity with excess pres-
Fig. 8 shows a cross-sectional FE-SEM image of a 150 nm thick ceria film. As can be seen, in the first stage of the growth, the film grows in columnar shape, perpendicular to the surface of the substrates, and the column width increases with the film growth. This columnar structure is very similar to that observed by Becht for a ceria film on STO substrate [10], and is called ‘non-equiaxed’ [21], as also discussed in Ref. [10]. An AFM image (scanning area: 1× 1 mm) of a 50 nm thick ceria film is shown in Fig. 9. The surface texture is rectangular in shape, with a typical size of 100×200 nm. All rectangular grains in the observed area are aligned in the same direction, and the edges of the rectangular are parallel to the Si(011) facets, suggesting a possible in-plane orientation. This is notably different from AFM images reported for (111) and (110)-oriented CeO2 on Si [22]. An AFM image of a 150 nm thick film showed similar in-plane alignment, with a larger and roundish rectangular shape with a typical size of 150×300 nm.
3.5. The interface of ceria and Si The high-resolution cross-sectional image of the in-
Fig. 9. An AFM image (scanning area: 1 ×1 mm) of a 50 nm thick ceria film.
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Fig. 10. The high-resolution cross-sectional image of the interface of a 150 nm thick ceria film and a silicon substrate with a beam axis of 011 direction.
terface of a 150 nm thick ceria film and a silicon substrate with a beam axis of 011 direction is shown in Fig. 10. Disoriented crystals are observable, showing that the ceria film was not epitaxially grown. However, the thickness of the amorphous layer was estimated to be 2 nm. This is comparable with the interface prepared using the high-vacuum process [23].
4. Conclusion (100)-oriented CeO2 films were grown on hydrogenterminated Si(100) substrates with an orientation factor of 0.83 defined by Lotgering. The crystallinity and orientation were affected by the partial pressure of oxygen as well as the substrate temperature. The films exhibited columnar growth at the initial stage, and the column width increased with the film growth. The grains of the film observed by AFM were rectangular in shape, with a typical size of 100×200 nm. The grains were aligned parallel to the (011) facets of the Si substrates, suggesting a possible in-plane orientation. The thickness of the amorphous layer at the interface of CeO2 and the Si substrate was about 2 nm. It seems that a better understanding of the initial stage of film growth is needed to achieve epitaxial growth of ceria on Si(100) substrates.
Acknowledgements The authors wish to thank Y. Kawate, Y. Haga, Y. .
Ikeda, H. Masuya, N. Kasahara, Y. Hirano and H. Kobayashi of the Center for Materials Analysis of the SONY Research Center for characterization of the ceria films and the surface of silicon wafers. They also thank A. Machida, N. Nagasawa and K. Mikami of the SONY Research Center for helpful discussions and experimental assistance. They are indebted to H. Ohki for promoting research on oxide electronics.
References [1] M. Suzuki, T. Ami, Mater. Sci. Eng. B41 (1996) 166. [2] M. Suzuki, J. Ceram. Soc. Jpn. 103 (1995) 1088. [3] J.R. Sims Jr., R.N. Blumenthal, High Temp. Sci. 8 (1976) 99 – 110. [4] M. Yoshimoto, K. Shimozono, T. Maeda, et al., Jpn. J. Appl. Phys. 35 (1995) 688. [5] T. Inoue, T. Ohsuna, L. Luo, et al., Appl. Phys. Lett. 59 (1991) 3604. [6] T. Inoue, Y. Yamamoto, S. Koyama, S. Suzuki, Y. Ueda, Appl. Phys. Lett. 56 (1990) 1332. [7] T.S. Chin, J.Y. Huang, L.H. Perng, T.W. Huang, S.J. Yang, S.E. Hsu, Phys. C 192 (1992) 154. [8] T. Hirai, K. Nagashima, H. Koike, S. Matsuno, Y. Tarui, Jpn. J. Appl. Phys. 35 (1996) 5150. [9] H. Nagata, T. Tsukahara, S. Gonda, M. Yoshimoto, H. Koinuma, Jpn. J. Appl. Phys. 30 (1991) 1136. [10] M. Becht, T. Morishita, Chem. Vap. Deposition 2 (1996) 191. [11] Z. Lu, R. Hiskes, S.A. DiCarolis, A. Nel, R.K. Route, R.S. Feigelson, J. Crystal Growth 156 (1995) 227. [12] S. Liang, C.S. Chern, Z.Q. Shi, P. Lu, Y. Lu, B.H. Kear, J. Crystal Growth 151 (1995) 359. [13] M. Becht, T. Gerfin, K.-H. Dahmen, Chem. Mater. 5 (1993) 137.
T. Ami, M. Suzuki / Materials Science and Engineering B54 (1998) 84–91 [14] K.-H. Dahmen, M. Becht, T. Gerfin, High-Tc Supercond. Thin Films, Proc. Symp. A1 Int. Conf. Adv. Mater, North-Holland, Amsterdam, 1992, p. 715. [15] J. McAleese, J.C. Plakatouras, B.C.H. Steele, Thin Solid Films 280 (1996) 152. [16] H.L. Taylor, J.D. Trotter, Proc. Int. Conf. Chem. Vap. Deposition 3 (1972) 475. [17] M. Hirose, M. Hiroshima, T. Yasaka, M. Takakura, S. Miyazaki, Microelectron. Eng. 22 (1993) 3.
.
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[18] C.H. Bjorkman, M. Fukuda, T. Yamazaki, S. Miyazaki, M. Hirose, Jpn. J. Appl. Phys. 34 (1995) 722. [19] F.K. Lotgering, J. Inorg. Nucl. Chem. 9 (1959) 113. [20] I. Yasui, T. Nanba, Diffusion Defect Data 53 (1987) 439. [21] C.V. Thompson, R. Carel, Mater. Sci. Eng. B32 (1995) 211. [22] T. Inoue, Y. Yamamoto, M. Satoh, A. Ide, S. Katsumata, Thin Solid Films 281-282 (1996) 24. [23] T. Inoue, T. Ohsuna, Y. Obara, Y. Yamamoto, M. Satoh, Y. Sakurai, Jpn. J. Appl. Phys. 32 (1993) 1765.
MOCVD growth of (100)-oriented CeO2 thin films on hydrogen-terminated Si(100) substrates Takaaki Ami *, Masayuki Suzuki SONY Corporation Research Center, 174, Fujitsuka-cho, Hodogaya-ku, Yokohama 240, Japan
Abstract Ceria (CeO2) with a fluorite structure is supposed to be an ideal buffer layer for fabricating epitaxially grown perovskite materials on silicon substrates because of its thermal stability and excellent lattice match with Si. On the other hand, there are still technical difficulties in forming CeO2(100), which is more attractive for applications compared to that of (111) orientation. We report the preparation of CeO2 thin films on hydrogen-terminated Si(100) substrates by metal-organic chemical vapor deposition (MOCVD) using Ce(DPM)4 as an MO-source. The surface microroughness of the substrates was investigated using Fourier-transform infrared-attenuated total reflection (FT-IR-ATR) and atomic force microscopy (AFM). Si-wafers cleaned by RCA method were immersed in H2O2-added 0.5% HF solutions for hydrogen-termination, and preferential Si – H2 vibrational peaks accompanied with weak Si–H3 peaks were observed. The root mean square roughness of the substrates estimated by AFM was around 0.2 nm. The crystallinity and orientation were characterized by X-ray diffraction (XRD). The films were polycrystalline, and preferentially orientated to the 100 direction. The preferential orientation factor defined by Lotgering was 0.83. The microstructures were characterized by a field-emission scanning electron microscopy (FE-SEM), an AFM and a transmission electron microscopy (TEM). The grains were ‘non-equiaxed’ columnar, growing perpendicular to the surface at the expense of other columns. The surface texture of a 50 nm-thick film characterized by AFM was rectangular in shape, with a typical size of 100×200 nm. They were aligned in the same direction, and the edges of the rectangular are parallel to the Si(011) facets, suggesting a possible in-plane orientation. Cross-sectional HR-TEM analyses of a 150 nm thick film verified the thickness of the amorphous layer formed at the CeO2/Si(100) interface to be around 2 nm. © 1998 Elsevier Science S.A. All rights reserved. Keywords: CeO2; Thin films; Si substrates; MOCVD
1. Introduction Introducing oxide materials except for SiO2 into the semiconductor industry is very attractive, because they exhibit various high performance properties. In particular, some oxides with a perovskite (-related) type structure show very important properties for semiconductor devices: superconducting, ferroelectric, high-o, and colossal magnetoresistance [1,2]. Ceria (CeO2) with a fluorite structure is supposed to be an ideal buffer layer for fabricating epitaxiallygrown perovskite materials on silicon substrates because of its thermal stability and excellent lattice match with Si, especially around 600°C [3]. The excellent epitaxy of ceria thin films was achieved on a Si(111) substrate without any amorphous layer in the vicinity of the boundary by pulsed laser deposition in an ultra* Corresponding author. 0921-5107/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5107(98)00133-0
high vacuum at room temperature [4]. However, the preferential epitaxial growth-orientation of ceria on Si(100) substrates reported was not CeO2(100) but CeO2(110) [5–7]. In order to fabricate a buffer layer epitaxially on Si(100) substrates, a solid solution of zirconia and ceria ((Zr1 − x Cex )O2) [8] or CeO2/SrTiO3/ Si structure [9] have also been proposed. Metal-organic chemical vapor deposition (MOCVD) is a widely used technique in the field of semiconductor industry because of its high uniformity in large wafers, growth rate of film, and conformal step coverage. For this reason, it is also utilized for growing ceria films on oxide single crystal substrates: SrTiO3 (STO), yttria stabilized zirconia (YSZ), magnesia (MgO) [10] and r-plane sapphire [11,12]. (100)-oriented ceria films were obtained on STO, YSZ and sapphire. Nevertheless, there still are technical difficulties in forming (100)-oriented ceria directly on silicon substrates, especially using MOCVD [13–16]. Although this technique was
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Fig. 1. The experimental setup of the MOCVD system used.
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described in Refs. [13] and [14], experimental data were not given. Meanwhile, atomic scale flatness of the surface and removal of the native oxides of silicon substrates are supposed to be important for epitaxial growth of films on these substrates. Hydrogen-termination by hydrofluoric acid (HF) or HF-based aqueous solution is one promising technique to remove the native oxide and passivate the silicon surface at the same time. A suggested solution proposed for smoothing the Si(100) surface is H2O2-added diluted HF, in which selective oxidation of the defects and etching of the oxide occurs [17,18]. In the present study, the concentrations of the base hydrofluoric acid and H2O2 in the solution were optimized.
2. Experimental
2.1. Hydrogen termination The silicon wafers used in this study were Borondoped Si(100) grown by CZ method with an acceptor concentration of 0.1 – 1.5× 105 cm − 3 (resistivity: 8–12 V cm). Wafers chemically cleaned by conventional SC-1 process were immersed into H2O2-added HF with various concentrations for 2 min. At the final stage of the optimization, the concentration of the HF base solution was fixed to 0.5 wt.% (pH 2.5), as the FT-IR-ATR spectra showed a reasonably good signal-to-noise ratio and absorbance, compared to higher (50% HF [18]) and lower (0.1% HF [17]) concentrations. The concentrated H2O2 (30%) was added to this concentration, and the volume of added H2O2 was optimized. The hydrogen-terminated surfaces were characterized by Fourier-transform infrared attenuated total reflection (FT-IR-ATR; JEOL JIR-6500) with Ge inner reflective crystal in a few minutes after the treatment. The surface roughness was characterized by atomic force microscopy (AFM; Park Scientific Instruments, Autoprobe).
source bottles, and mixed with oxygen gas in the reactor. The substrates were set on the SiC-coated graphite susceptor, and heated by a focused IR lamp.
2.3. Characterization The film thickness and the refractive index were measured by ellipsometry (Gaertner ellipsometer L115B) with He–Ne laser (632.8 nm). The film thickness was also characterized by cross-sectional images taken by a field emission-scanning electron microscope (FE-SEM; Hitachi S4500) with the microstructure of the films. The film composition on the surface and its depth-profile were characterized by Auger electron spectroscopy (AES; JEOL JAMP-30) and Rutherford back-scattering spectroscopy (RBS), respectively. The X-ray diffraction (XRD) was performed on a Rigaku RADIII diffractometer using Cu Ka radiation. For u−2u scans, a fine collimation with divergence slit (DS)-1/2°, scattering slit (SS)-1/2° and receiving slit (RS)-0.05 mm was selected in order to obtain high resolution, when needed. A glancing angle-diffractometer with an x-ray incident angle of 5° was also used in order to eliminate the diffraction peaks from the Si substrate. Surface morphologies were analyzed by AFM. The interface between CeO2 and Si substrates was observed by transmission electron microscope (TEM; JEOL 2000FX-II; 200kV).
3. Results and discussion
3.1. Optimization of the amount of the added H2O2 in 0.5% HF Fig. 2 shows the volume of the added H2O2 dependence of the absorbance of the SiH2, SiH3 peaks and full width of half maximum (FWHM) of the SiH2 peak. The absorbance of SiH2 is increased by adding a slight
2.2. Thin film deposition The experimental setup of the MOCVD system is schematically shown in Fig. 1. Metal-organic (MO)sources are set in Pyrex bottles, which are mounted in the box-type furnace and can be heated up to 250°C homogeneously. A reactor tube, its head fringe, and the manifold are also set in a box-type furnace with valves and heated up homogeneously in order to avoid the condensation of vaporized MO-sources. The MOsource used was Ce(DPM)4 with a vapor-pressure of 0.1 Torr at around 200°C. It was introduced by argon gas into the quartz hot-wall reactor through the
Fig. 2. The volume of the added H2O2 dependence of the absorbance of the SiH2, SiH3 peaks and FWHM of the SiH2 peak.
T. Ami, M. Suzuki / Materials Science and Engineering B54 (1998) 84–91
Fig. 3. The surface AES spectra of films deposited with and without oxygen.
amount of H2O2. The FWHM of the dihydride peak was decreased with the increase of the absorbance, implying a homogeneous surface. When an excess amount of H2O2 was added to 0.5% HF, the absorbance decreased. As a result, the optimized amount of H2O2 was determined to be 0.2 vol.%. This result was supported by AFM, showing a minimum root mean square roughness of 0.2 nm at the optimized amount.
3.2. MO-source transport and film growth The transport of the MO-source and film deposition mechanism were investigated with the CVD parameters. Experiments for the investigation of the growth rate were performed with parameters of evaporation temperature of the MO-source and Ar-carrier flow rate at a substrate temperature (Tsub) of 700°C and a total reaction pressure (Ptotal) of 2 Torr. The partial pressure of oxygen (pO2) was fixed to 87.5% of the total pressure. The growth rate increased exponentially with increasing evaporation temperature in a range between 170 and 210°C. On the other hand, with the increase of Ar-carrier in a range between 10 and 100 sccm, the growth rate increased linearly showing a slight saturation. These results clearly show that the film growth is masstransport limited under these conditions. On the other hand, the Tsub-dependence of the growth rate is slight between 600 and 800°C. In contrast, the pO2 affects the film growth. Fig. 3 shows the surface AES spectra of films deposited with and without oxygen. The peak intensities of Ce and O are weaker in the film deposited without oxygen, while, the carbon content in this film is much more than that in the film deposited with oxygen. The results of XRD for the film deposited with 87.5% of pO2 showed the formation of single phase-polycrystalline ceria, as will be discussed later. From the result of RBS, the Ce/O ratio of the ceria was calibrated to be
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2.090.1. The refractive index of the films was 2.1–2.4, which is almost the same range as the reported value for bulk ceria [6]. On the other hand, no crystallized phases were detected in films deposited without oxygen by XRD. These results show that oxygen is required to decompose the Ce(DPM)4 and to form the CeO2. The film deposited without oxygen is supposed to be the Ce(DPM)4 itself or this material partially decomposed. However, the growth rate in the pO2 range between 12.5 and 87.5% was almost constant at a Tsub of 700°C and a Ptotal of 2 Torr. In order to suppress the formation of a-SiO2 on the Si surfaces, the initial stage of deposition was performed in argon gas, typically for a minute, then the argon was switched to oxygen to form CeO2 films. Typical growth conditions are shown in Table 1.
3.3. Crystallinity and orientation 3.3.1. XRD analysis by u− 2u scan Because of the excellent lattice matching of ceria to silicon, an unfortunate diffraction peak from the Si(100) substrates shows up in the XRD data collected by a u− 2u scan which is very close to the CeO2(200) peak. However, with the fine collimation described previously, the Si(200) peak was observed to be split from the CeO2(200) peak. The overlap of the Si(200) Cu Ka2 line and CeO2(200) Cu Ka1 line is still critical, but the contribution of the Si(200) Cu Ka2 line can be subtracted fairly accurately. After the reduction of the Cu Ka2 line, the peak separation can be performed with reasonable accuracy. Blown up XRD spectra (32°B 2uB 34°) for a ceria film with a Si substrate and it for a Si substrate during these procedures are presented in Figs. 4 and 5. These techniques reveal the diffraction peaks only from ceria films. Fig. 6 shows the XRD pattern of a ceria film on a Si(100) substrate by u− 2u scan without contributions from the substrate (bottom), together with the XRD of the CeO2 ceramics wafer in the same size as the ceria film (top). The film were grown at 700°C, 2 Torr, with a pO2 of 87.5%. The contribution Table 1 Typical growth conditions for MOCVD CeO2 thin films on Si Substrate Substrate temperature (°C) Total pressure (Torr) Partial pressure of oxygen (%) Source evaporation temperature (°C) Ar carrier gas flow rates (sccm) Ar dilution gas flow rates (sccm) O2 gas flow rates (sccm)
Si(100) 500 – 1000 1 – 10 0 – 87.5 170 – 210 10 – 100 0 – 600 0 – 700
The initial stage of deposition was performed in argon dilution gas just for a short time, then the oxygen was switched to the argon to form CeO2 films.
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Fig. 4. Blown up XRD spectra (32° B2u B 34°) for a ceria film with a Si substrate and it for a Si substrate before Cu Ka2 reduction. The vertical tick marks indicate the expected peak locations for CeO2(200) and Si(200).
Fig. 6. The XRD pattern of a ceria film on a Si(100) substrate by u −2u scan without contributions from the substrate (bottom), together with the XRD of the CeO2 ceramics wafer in the same size as the ceria film (top).
from the silicon substrates were subtracted in the XRD. Only a very broad (111) peak and a weak (311) peak were observed in 10° B2u B 65°, except for the intense (200) peak. The preferred orientation factor F is defined by Lotgering as follows [19]:
3.3.2. XRD analysis by glancing angle diffractometer The intensity measured by the glancing angle diffractometer is affected by the instrumentation, especially the solar slit collimator (SL), which restricts the diffracted X-ray [20]. For a preliminary analysis of the orientation, the orientation factor G for the glancingangle diffractometer was simply defined as:
F = (P− P0)/(1 − P0), where P= I(h00)/ I(hkl) and P0 is the P-value for a non-oriented specimen. P0 was calculated from the XRD of a CeO2 ceramics wafer. Diffraction peaks in 10°B 2u B65° were used for the calculation. The Fvalue for the sample shown in Fig. 6 was 0.83, showing a fairly strong (100) preferred orientation. Nevertheless, the peak intensity and FWHM of the Si(200) peak strongly depend on how the sample is set in the diffractometer. For systematic characterization of the orientation of the ceria films, glancing angle XRD measurements were performed in which no contribution from the silicon substrates appeared.
Fig. 5. Blown up XRD spectra (32° B2u B 34°) for a ceria film with a Si substrate and it for a Si substrate after Cu Ka2 reduction.
G= Q/Q0 where Q=I(200)/I(111), and Q0 is the Q-value for a non-oriented specimen. In order to minimize the error caused by the SL and absorption of X-ray, the two peaks: (200), (111) with the difference of small 2u angle (4.534°) were used. Fig. 7 shows the Tsub-dependence of the I(200), I(111) and orientation factor G. The I(200) and the G are maximized at around 700°C, where the I(111) is lowest. This result clearly shows that Tsub affects the crystallinity and orientation. In the high temperature region, the I(111) increased with Tsub, showing that the orientation changes from (100) to random or possible
Fig. 7. The substrate temperature dependence of the I(200), I(111) and orientation factor G.
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sure. Hereinafter, characterizations on the films grown at 700°C, 2 Torr, with a pO2 of 87.5% will be described.
3.4. Surface morphology of the ceria thin film (FE-SEM, AFM)
Fig. 8. A cross-sectional FE-SEM image of a 150 nm thick ceria film.
(111) orientation. On the other hand, the drastic decrease of I(200) and weak I(111) imply that the crystallinity goes down at low temperatures. The crystallinity and the orientation were also affected by pO2. At a fixed Tsub (700°C) and a Ptotal (2 Torr), the G increased with increasing pO2 in a range between 12.5 and 87.5%, demonstrating the decomposition of the MO-source and crystallization of the (100)oriented CeO2. On the other hand, at a fixed Tsub (700°C) and a pO2 (87.5%), the G decreased with increasing Ptotal in a range between 2 and 10 Torr, suggesting the decreasing crystallinity with excess pres-
Fig. 8 shows a cross-sectional FE-SEM image of a 150 nm thick ceria film. As can be seen, in the first stage of the growth, the film grows in columnar shape, perpendicular to the surface of the substrates, and the column width increases with the film growth. This columnar structure is very similar to that observed by Becht for a ceria film on STO substrate [10], and is called ‘non-equiaxed’ [21], as also discussed in Ref. [10]. An AFM image (scanning area: 1× 1 mm) of a 50 nm thick ceria film is shown in Fig. 9. The surface texture is rectangular in shape, with a typical size of 100×200 nm. All rectangular grains in the observed area are aligned in the same direction, and the edges of the rectangular are parallel to the Si(011) facets, suggesting a possible in-plane orientation. This is notably different from AFM images reported for (111) and (110)-oriented CeO2 on Si [22]. An AFM image of a 150 nm thick film showed similar in-plane alignment, with a larger and roundish rectangular shape with a typical size of 150×300 nm.
3.5. The interface of ceria and Si The high-resolution cross-sectional image of the in-
Fig. 9. An AFM image (scanning area: 1 ×1 mm) of a 50 nm thick ceria film.
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Fig. 10. The high-resolution cross-sectional image of the interface of a 150 nm thick ceria film and a silicon substrate with a beam axis of 011 direction.
terface of a 150 nm thick ceria film and a silicon substrate with a beam axis of 011 direction is shown in Fig. 10. Disoriented crystals are observable, showing that the ceria film was not epitaxially grown. However, the thickness of the amorphous layer was estimated to be 2 nm. This is comparable with the interface prepared using the high-vacuum process [23].
4. Conclusion (100)-oriented CeO2 films were grown on hydrogenterminated Si(100) substrates with an orientation factor of 0.83 defined by Lotgering. The crystallinity and orientation were affected by the partial pressure of oxygen as well as the substrate temperature. The films exhibited columnar growth at the initial stage, and the column width increased with the film growth. The grains of the film observed by AFM were rectangular in shape, with a typical size of 100×200 nm. The grains were aligned parallel to the (011) facets of the Si substrates, suggesting a possible in-plane orientation. The thickness of the amorphous layer at the interface of CeO2 and the Si substrate was about 2 nm. It seems that a better understanding of the initial stage of film growth is needed to achieve epitaxial growth of ceria on Si(100) substrates.
Acknowledgements The authors wish to thank Y. Kawate, Y. Haga, Y. .
Ikeda, H. Masuya, N. Kasahara, Y. Hirano and H. Kobayashi of the Center for Materials Analysis of the SONY Research Center for characterization of the ceria films and the surface of silicon wafers. They also thank A. Machida, N. Nagasawa and K. Mikami of the SONY Research Center for helpful discussions and experimental assistance. They are indebted to H. Ohki for promoting research on oxide electronics.
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