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6H-SiC interfaces
Surface Science 494 (2001) L805±L810
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Surface Science Letters
The Schottky limit and a charge neutrality level found on metal/6H-SiC interfaces Shiro Hara National Institute of Advanced Industrial Science and Technology (AIST), AIST Tsukuba Central 2, Tsukuba, Ibaraki 305-8568, Japan Received 18 October 2000; accepted for publication 31 August 2001
Abstract We report the Schottky limit and a charge neutrality level (CNL) experimentally demonstrated at metal/6HSiC(0 0 0 1) interfaces. An interface with the Schottky limit was formed by dipping SiC surfaces in boiling pure water before metallization. The total density of interface states, Dit , was a drastically small value of 4:6 1010 states cm 2 /eV, indicating the density of the metal induced gap states was less than this value. In contrast, at incompletely passivated interfaces without the boiling water dipping process, a broad continuum of interface states with Dit of 2:8 1013 states cm 2 /eV was observed with the CNL located at 0.797 eV from the conduction band minimum. The origin of these interface states was found to be in the disordered interface layers. Ó 2001 Elsevier Science B.V. All rights reserved. Keywords: Schottky barrier; Metal±semiconductor interfaces; Interface states; Surface electronic phenomena (work function, surface potential, surface states, etc.); Silicon carbide; Evaporation and sublimation; Oxidation; Etching
Despite ®ve decades of innumerable studies, there is still much controversy about the mechanism of Schottky barrier formation at metal/semiconductor interfaces, especially in terms of the essential factors determining the barrier height [1]. Proposed factors are often categorized into intrinsic and extrinsic factors. A so-called intrinsic factor is the metal induced gap states (MIGS) conceptually proposed by Heine [2]. Computational demonstrations of MIGS have been made [3] and a method for calculating the charge neutrality level (CNL) that dominates the interface Fermi level has been proposed [4]. Extensive calculations have been done [5] and CNL ¯uctuations have been discussed
E-mail address: [email protected] (S. Hara).
[6]. The long-range space-charge contribution to the interface charges has also been considered [7]. Some of the reasons for the controversy arise from the lack of experimental supports for the penetration depth and energy distribution of MIGS in the semiconductor, and for a direct measurement of CNL. To reveal the MIGS eect experimentally, a reduction of the extrinsic factors is a prerequisite. Typical extrinsic factors are defects as pointed out by Spicer et al. [8], disorder that induces gap states (DIGS) [9], and interface geometry as typically shown in the NiSi2 /Si(1 1 1) interface [10]. These factors show the importance of the formation of an abrupt interface. This recognition has led to experiments of low-temperature interface formation [11] and speci®c inert metal selections [12,13]. Imperfect termination of the surface before
0039-6028/01/$ - see front matter Ó 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 3 9 - 6 0 2 8 ( 0 1 ) 0 1 5 9 6 - 5
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the metal deposition is also one of extrinsic factors as shown in the studies of sulfur passivation on GaAs [14] and hydrogen passivation on Si [12]. In the above studies on the extrinsic factors, however, the general trend of the Schottky barrier height (SBH) for various metals is unclear and indirect. Due to the incomplete and partial control of these extrinsic factors, the energy distribution of the interface states becomes complicated, thereby making the SBH unpredictable. Further, the MIGS eect reduces the SBH controllability, which hinders the detection of the extrinsic factors as individual eects on the SBH. According to the controversial history, one can express the total density of interface states Dtotal as it
and exif the interaction between MIGS, DMIGS it trinsic states DEXT is neglected. To observe one it of these two factors, the other factor should be reduced. Here, we use 6H-SiC crystal with a wide band gap that is predicted to have a weaker MIGS eect than Si or GaAs but still predicted to have a remarkable eect with DMIGS of 1014 states cm 2 /eV it 1 [15]. Also, SiC is expected to form an abrupt interface after room-temperature metal deposition [16] because of its strong Si±C bonding. We also apply the monohydride termination technique, developed for the Si(1 1 1) surface using boiling water immersion [17], for the SiC surface. Using these experimental techniques, we form atomby-atom interface connection and we demonstrate the ®rst perfect control of the extrinsic factors and observe the negligible MIGS eect in this wide band gap. Commercially available nitrogen-doped n-type 6H-SiC(0 0 0 1) epitaxial wafers with a carrier concentration of 5 1017 cm 3 were used. All samples were degreased followed by dipping in 5% HF solution. This cleaning process is referred to as
DHF treatment. Some of the degreased samples were RCA cleaned and subsequently thermally oxidized. The oxide layer with a thickness of 10 nm was etched by dipping in 5% HF solution. We will refer to the sequence of the treatment as O/E treatment. Some of the O/E samples were immersed in boiling water at 98 °C for 10 min. The sequence of this treatment will be referred to as BW treatment. All samples were rinsed in the deionized water after each HF process. The treated surfaces were characterized using low energy electron diraction (LEED), Auger electron spectroscopy (AES), and X-ray photoemission spectroscopy (XPS). Metals were deposited onto the SiC surfaces after each surface treatment using an e-beam evaporator with maximum working pressures below 1 10 9 Torr. During deposition, there was no intentional sample heating. In our preliminary report, we used only Ti, Mo, and Ni as contact materials [18]. Here we add Pt and Al to obtain Dit . SBH were measured by I±V and C±V methods. 2 The slope parameter S/ de®ned as o/b =o/m is a parameter required to estimate Dit , where /b is the SBH and /m is the metal work function. In our work, the experimental values of S/ obtained from Al, Ti, Mo, Ni and Pt electrodes are 0.180, 0.549, and 0.754 for DHF, O/E, and BW samples, respectively. Previously reported values of S/ for 6H-SiC(0001) epitaxial ®lms are 0.14 [19] and 0.63 [20], which are similar to the values of our DHF and O/E treatments, respectively. S/ for the BW sample is larger than the reported ones. In a discussion on S/ , one should notice that the leakage current tends to reduce the I±V measured SBH [21], which might cause an underestimation of S/ compared with the real S/ . Fig. 1 is the relation between the ideal factor n and /b0 with Pt electrodes. /b0 is the ¯at-band barrier height canceling the image force lowering. Since all the data plots in each treatment are distributed on each straight line, the lines are extrapolated to ideal unique
1 SX values reported in this paper can be converted to Dit using the formulas of S/ ei = ei q2 dDit and SX @/b = @Xm @/b =@/m @/m =@Xm AS/ , where Xm is the electronegativity of a metal. The conversion factor A @/m =@Xm is 2.27.
2 Detailed experimental conditions and the results of our investigation on the metal/SiC interface will be described in a full paper reported by Teraji and Hara, which will be submitted elsewhere.
MIGS Dtotal E DEXT E; it E Dit it
1
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Fig. 1. Ideal factor n vs. /b0 for Pt/6H-SiC(0001) interfaces. DHF, O/E, BW indicate diluted HF, oxidation and etching, and boiling water dipping process, respectively. Each plot indicates n and /b0 of individual Pt electrodes. The inset shows typical I±V curves.
barrier heights /b0 at n 1, obeying thermionic emission current transport without leakage current. /b0 for DHF, O/E, and BW electrodes are 1:011 0:015 eV, 1:365 0:024 eV, and 1:632 0:007 eV, respectively. The plot concentration in DHF and the plot broadening along the BW line suggest that the DHF interface has a large amount of Dit to form a strong pinning and the pinning is weaker in the BW interface. By obtaining /b0 also from other metal electrodes, relations between /m and /b0 are plotted as shown in Fig. 2. S/ using /b0 for the DHF, O/E and BW electrodes are 0:215 0:050, 0:739 0:079, and 0:994 0:053, respectively. The value of 0.994 is near the Schottky limit of S/ 1 having the maximum SBH controllability. C±V measured barrier heights /CV b0 provide almost the same barrier heights as /b0 and S/ 1:011 for the BW interfaces. (See footnote 2.) This indicates that /CV b0 /b0 6 /b0 at least in our experiments. In other words, the C±V barrier heights coincide with the I±V barrier heights when the interfaces have no leakage current. We estimated Dit from the equation S/ ei = ei q2 dDit [22], where ei is the per-
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Fig. 2. /m vs. /b0 plots. Barrier height controllability is raised going from DHF to BW through O/E treatment. On the BW line, the slope parameter S/ is 0.994, which is approximately in is the observed CNL of these the Schottky limit condition. /exp 0 metal/SiC interfaces. /MIGS is the predicted CNL of MIGS [25]. 0 The inset shows the energy distribution of the interface states. Distributions outside of the hatched areas are unknown in our experiment.
mittivity of the interfacial layer and d the width of the interface layer. We use the bilayer thickness 0.252 nm in the 6H-SiC(0 0 0 1) crystal as d, and use the permittivity in vacuum e0 as ei since no report was found on ei for metal/SiC interfaces. Dit for the DHF, O/E and BW electrodes are 2:8 1013 , 2:7 1012 , and 4:6 1010 states cm 2 / eV, respectively, as indicated in the inset ®gure. Even for the lower error limit of S/ 0:941 0:994 0:053, Dit has a relatively low value of 4:8 1011 states cm 2 /eV. The correlation coecients r for the DHF, O/E, and BW lines are 0.928, 0.983, and 0.996, respectively. These strong correlations suggest there is no remarkable interface reaction generating interface states between the SiC substrate and the various metals. It is important to know the line shape in the relation between /b and /m to understand the distribution Dit E as depicted in Fig. 3. A U-shaped distribution provides a curved line. A straight line is generated by a broad distribution. Local states provide a stepwise line. According to the MIGS model proposed by Terso [4], MIGS should have
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Fig. 3. Types of energy distribution in interface states and corresponding correlations of /m vs. /b0 .
a U-shaped energy distribution. It is believed that the MIGS are not the direct penetration states of the metal wave function but metal-induced states redistributed from the conduction and the valence band states where the total number of the states in the semiconductor is preserved [4]. In our observations, the data plots in Fig. 2 form straight lines at least in the observed individual barrier height ranges that are indicated by hatched areas in the
inset. This broad feature of Dit E suggests that the origin of the interface states is dierent from the MIGS that should have the U-shaped Dit E. In general, energetic broadness in electron states is a typical feature of a disordered structure. Fig. 4 shows cross-sectional lattice images of DHF and BW interfaces by transmission electron microscopy (TEM). The DHF interface has a disordered interface layer of 2 nm thickness. In the BW sample the interface is atomically abrupt and has the atom-by-atom commensurate epitaxial connection with the 1 1 0 26H-SiC == 1 1 1Ti relation without any evidence of a SiO2 formation in our cross-sectional TEM observation and electron energy loss spectroscopy (EELS) of the crosssectional sample. The lattice type of the Ti layer is the face-centred cubic (fcc) lattice. The formation of the fcc-Ti structure has previously been suggested only in an ultra-thin Ti ®lm with a thickness of several angstroms on Al surfaces [23]. We have observed the growth of the fcc-Ti layer with a thickness of 100 nm, which is the ®rst report on this thin layer single-crystal formation [24]. The bulk stable phase of Ti is hexagonal close-packed (hcp) with a lattice mismatch of 4% for the 6HSiC(0001) crystal. The lattice constant a of the observed fcc-Ti structure is 0.438 nm with a lattice mismatch of 0.79% between the 1 1 0 26H-SiC and 1 1 1Ti faces. The phase change at the interface
Fig. 4. High-resolution cross-sectional TEM lattice images with the electron acceleration at 200 eV incident from the [1 1 2 0] direction of (a) DHF and (b) BW interfaces.
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occurs in preference to the incommensurate hcpTi/6H-SiC(0 0 0 1) interface formation. This commensurate connection is strong evidence of direct Ti chemical bonding to the SiC crystal without a hydrogen terminator. In LEED observations BWtreated SiC surfaces exhibit sharp 1 1 patterns whereas DHF-treated SiC surfaces show no diffraction spot indicating that the DHF surface is disordered. The disordered surface and interfaces in DHF samples is consistent with the broad Dit E. The uni®ed disorder-induced gap state (DIGS) model has been proposed as a model to deal with interface disorder [9]. In this model, similar to the branch point mechanism, Dit E is a U-shaped continuum because the origin of the DIGS is de®ned as the penetration tails of the valence and conduction band states. The observed broad Dit E suggests that another type of disorder exists whose origin is dierent from the above mechanism. One of our important observations is the ®rst experimental ®nding of a (CNL, /exp 0 ) indicated in Fig. 2. CNL is the critical parameter to determine the amount of the interface charge, but it has been discussed only for the MIGS theory [4±6]. Here, we extend the CNL phenomenon to a more general concept that is not restricted to a speci®c origin like MIGS. The three lines intersect with each other at the CNL with /b0 of 0.797 eV and /m of 4.65 eV. The observation of one CNL for the three kinds of surface preparations indicates that the origin of Dit is the same for the three interfaces, but the amount of Dit is variable. This /exp is lo0 cated far from the MIGS±CNL /MIGS at 1.43 eV 0 estimated from previous experimental values of /b0 [20] and the value of /m of 5.4 eV assuming that /m is equal to the Miedima's electronegativp ity, XSiC ( XSi XC ) [25]. This is plotted in Fig. 2. In the MIGS theory, the interface Fermi level should be pinned at the branch point EBP where the eects of the conduction and valence band balance [4]. EBP for the 6H-SiC crystal is calculated to be 1.41 eV, which has almost the same value as the above estimation of /MIGS but absolutely dier0 ent from /exp . This dierence supports the assertion 0 that /exp is an extrinsic level, not the intrinsic level 0 based on the branch point mechanism. The disdisorder cussed disordered features give /exp in 0 /0
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this system. In other words, our experimentally , observed CNL is a disorder-related CNL, /disorder 0 not MIGS±CNL /MIGS . Considering the possibil0 ity of a multiple CNL system, another general model with an expression, BP EXT1 Dtotal E DEXT2 E ; it E Dit E Dit it
2 is given, where each term on the right has an individual /0 , and DBP it is Dit with the CNL at EBP , which originates from the MIGS or DIGS theories based on the branch point mechanism. Our metal/ disorder SiC interfaces have Dtotal DBP with it it Dit BP total 10 2 Dit < Dit 4:6 10 states cm /eV for the BW interface in the observed SBH range of 0.45±1.66 eV. This indicates that DMIGS is also lower than it 4:6 1010 states cm 2 /eV. According to the typical MIGS theory of Ref. [4], the penetration length of the metal wave function or the semiconductor layer depth modi®ed by the metal was calculated to be 0.15 nm for ZnS with the band gap of 3.6 eV that is comparable value of 6H-SiC. Nevertheless, Dit calculated for ZnS is 1014 [26]. At least, our ®nding of 6 1010 for the MIGS density strongly suggests that the MIGS penetration depth for SiC is much lower than 0.15 nm. This implies that actual penetration depths for other semiconductors may have much lower than the calculated values [4] and the semiconductors may have Dit with one or two order of magnitude lower than the predictions if one can exclude the extrinsic interface states. From XPS measurements of the core level position of C 1s, the surface potential barrier is 0.19 eV and the surface band bending is nearly ¯at of 0.08 eV for the BW surface. Using the value of 0.08 eV R the ionized total density of the surface states Ds dE is evaluated to be 6:7 1011 states/ cm2 . Let us estimate the density of the surface states Ds by assuming that the surface CNL has the same origin as the interface CNL. The energy width of charged states between the CNL and the Fermi level position is 0.607 eV( 0:797 0:19 eV). Thus, Ds is estimated at 1:1 1012 states cm 2 /eV. This value is one and a half orders of magnitude higher than Dit , implying that the residual surface states are eectively annihilated by
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the electronically ideal interface connection as well as the structurally perfect interface connection. Acknowledgement The author thanks Dr. T. Teraji for his enormous amount of experimental work. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12]
L.J. Brillson, Surf. Sci. 299 (1994) 909. V. Heine, Phys. Rev. 138 (1965) A1689. S.G. Louie, M.L. Cohen, Phys. Rev. B 13 (1976) 2461. J. Terso, Phys. Rev. Lett. 52 (1984) 465. W. M onch, J. Appl. Phys. 80 (1996) 5076. F. Flores, A. Mu~ noz, J.C. Duran, Appl. Surf. Sci. 41 (1989) 144. C. Mailhiot, C.B. Duke, Phys. Rev. B 3 (1986) 1118. W.E. Spicer, I. Lindau, P. Skeath, C.Y. Su, J. Vac. Sci. Technol. 17 (1980) 1019. H. Hasegawa, H. Ohno, J. Vac. Sci. Technol. B 4 (1986) 1130. R.T. Tung, Phys. Rev. Lett. 52 (1984) 461. S. Chang, L.J. Brillson, Y.J. Kime, D.S. Rioux, P.D. Kirchner, G.D. Pettit, J.M. Woodall, Phys. Rev. Lett. 64 (1990) 2551. M. Wittmer, J.L. Freeouf, Phys. Rev. Lett. 69 (1992) 2701.
[13] G. Le Lay, V. Yu, Aristov, K. Hricovini, A. TalebIbrahimi, P. Dumas, R. Gunther, J. Osvald, G. Indlekofer, in: I. Ohdomari, et al. (Eds.), Control of Semiconductor Interfaces, Elsevier, Amsterdam, 1994, p. 39. [14] J.F. Fan, H. Oigawa, Y. Nannichi, Jpn. J. Appl. Phys. 27 (1988) L2125. [15] S. Kurtin, T.C. McGill, C.A. Mead, Phys. Rev. Lett. 22 (1969) 1433. [16] L.M. Porter, R.F. Davis, J.S. Bow, M.J. Kim, R.W. Carpenter, R.C. Glass, J. Mater. Res. 10 (1995) 668. [17] S. Watanabe, M. Shigeno, N. Nakayama, T. Ito, Jpn. J. Appl. Phys. 30 (1991) 3575. [18] S. Hara, T. Teraji, H. Okushi, K. Kajimura, Appl. Surf. Sci. 117 (1997) 394. [19] L.M. Porter, R.F. Davis, Mater. Sci. Engng. B 34 (1995) 83. [20] J.R. Waldrop, R.W. Grant, Y.C. Wang, R.F. Davis, J. Appl. Phys. 72 (1992) 4757. [21] R.F. Schmitsdorf, T.U. Kampen, W. M onch, J. Vac. Sci. Technol. B 15 (1997) 1221. [22] S.M. Sze, Physics of Semiconductor Devices, second ed., Wiley, New York, 1981, Chapter 5. [23] A.A. Saleh, V. Shutthanandan, R.J. Smith, Phys. Rev. B 49 (1994) 4908. [24] Y. Sugawara, N. Shibata, S. Hara, Y. Ikuhara, J. Mater. Res. 15 (2000) 2121. [25] W. M onch et al., in: I. Ohdomari (Ed.), Control of Semiconductor Interfaces, Elsevier, Amsterdam, 1994, p. 169. [26] S.G. Louie, J.R. Chelikowsky, M.L. Cohen, J. Vac. Sci. Technol. 13 (1976) 790.
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Surface Science Letters
The Schottky limit and a charge neutrality level found on metal/6H-SiC interfaces Shiro Hara National Institute of Advanced Industrial Science and Technology (AIST), AIST Tsukuba Central 2, Tsukuba, Ibaraki 305-8568, Japan Received 18 October 2000; accepted for publication 31 August 2001
Abstract We report the Schottky limit and a charge neutrality level (CNL) experimentally demonstrated at metal/6HSiC(0 0 0 1) interfaces. An interface with the Schottky limit was formed by dipping SiC surfaces in boiling pure water before metallization. The total density of interface states, Dit , was a drastically small value of 4:6 1010 states cm 2 /eV, indicating the density of the metal induced gap states was less than this value. In contrast, at incompletely passivated interfaces without the boiling water dipping process, a broad continuum of interface states with Dit of 2:8 1013 states cm 2 /eV was observed with the CNL located at 0.797 eV from the conduction band minimum. The origin of these interface states was found to be in the disordered interface layers. Ó 2001 Elsevier Science B.V. All rights reserved. Keywords: Schottky barrier; Metal±semiconductor interfaces; Interface states; Surface electronic phenomena (work function, surface potential, surface states, etc.); Silicon carbide; Evaporation and sublimation; Oxidation; Etching
Despite ®ve decades of innumerable studies, there is still much controversy about the mechanism of Schottky barrier formation at metal/semiconductor interfaces, especially in terms of the essential factors determining the barrier height [1]. Proposed factors are often categorized into intrinsic and extrinsic factors. A so-called intrinsic factor is the metal induced gap states (MIGS) conceptually proposed by Heine [2]. Computational demonstrations of MIGS have been made [3] and a method for calculating the charge neutrality level (CNL) that dominates the interface Fermi level has been proposed [4]. Extensive calculations have been done [5] and CNL ¯uctuations have been discussed
E-mail address: [email protected] (S. Hara).
[6]. The long-range space-charge contribution to the interface charges has also been considered [7]. Some of the reasons for the controversy arise from the lack of experimental supports for the penetration depth and energy distribution of MIGS in the semiconductor, and for a direct measurement of CNL. To reveal the MIGS eect experimentally, a reduction of the extrinsic factors is a prerequisite. Typical extrinsic factors are defects as pointed out by Spicer et al. [8], disorder that induces gap states (DIGS) [9], and interface geometry as typically shown in the NiSi2 /Si(1 1 1) interface [10]. These factors show the importance of the formation of an abrupt interface. This recognition has led to experiments of low-temperature interface formation [11] and speci®c inert metal selections [12,13]. Imperfect termination of the surface before
0039-6028/01/$ - see front matter Ó 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 3 9 - 6 0 2 8 ( 0 1 ) 0 1 5 9 6 - 5
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the metal deposition is also one of extrinsic factors as shown in the studies of sulfur passivation on GaAs [14] and hydrogen passivation on Si [12]. In the above studies on the extrinsic factors, however, the general trend of the Schottky barrier height (SBH) for various metals is unclear and indirect. Due to the incomplete and partial control of these extrinsic factors, the energy distribution of the interface states becomes complicated, thereby making the SBH unpredictable. Further, the MIGS eect reduces the SBH controllability, which hinders the detection of the extrinsic factors as individual eects on the SBH. According to the controversial history, one can express the total density of interface states Dtotal as it
and exif the interaction between MIGS, DMIGS it trinsic states DEXT is neglected. To observe one it of these two factors, the other factor should be reduced. Here, we use 6H-SiC crystal with a wide band gap that is predicted to have a weaker MIGS eect than Si or GaAs but still predicted to have a remarkable eect with DMIGS of 1014 states cm 2 /eV it 1 [15]. Also, SiC is expected to form an abrupt interface after room-temperature metal deposition [16] because of its strong Si±C bonding. We also apply the monohydride termination technique, developed for the Si(1 1 1) surface using boiling water immersion [17], for the SiC surface. Using these experimental techniques, we form atomby-atom interface connection and we demonstrate the ®rst perfect control of the extrinsic factors and observe the negligible MIGS eect in this wide band gap. Commercially available nitrogen-doped n-type 6H-SiC(0 0 0 1) epitaxial wafers with a carrier concentration of 5 1017 cm 3 were used. All samples were degreased followed by dipping in 5% HF solution. This cleaning process is referred to as
DHF treatment. Some of the degreased samples were RCA cleaned and subsequently thermally oxidized. The oxide layer with a thickness of 10 nm was etched by dipping in 5% HF solution. We will refer to the sequence of the treatment as O/E treatment. Some of the O/E samples were immersed in boiling water at 98 °C for 10 min. The sequence of this treatment will be referred to as BW treatment. All samples were rinsed in the deionized water after each HF process. The treated surfaces were characterized using low energy electron diraction (LEED), Auger electron spectroscopy (AES), and X-ray photoemission spectroscopy (XPS). Metals were deposited onto the SiC surfaces after each surface treatment using an e-beam evaporator with maximum working pressures below 1 10 9 Torr. During deposition, there was no intentional sample heating. In our preliminary report, we used only Ti, Mo, and Ni as contact materials [18]. Here we add Pt and Al to obtain Dit . SBH were measured by I±V and C±V methods. 2 The slope parameter S/ de®ned as o/b =o/m is a parameter required to estimate Dit , where /b is the SBH and /m is the metal work function. In our work, the experimental values of S/ obtained from Al, Ti, Mo, Ni and Pt electrodes are 0.180, 0.549, and 0.754 for DHF, O/E, and BW samples, respectively. Previously reported values of S/ for 6H-SiC(0001) epitaxial ®lms are 0.14 [19] and 0.63 [20], which are similar to the values of our DHF and O/E treatments, respectively. S/ for the BW sample is larger than the reported ones. In a discussion on S/ , one should notice that the leakage current tends to reduce the I±V measured SBH [21], which might cause an underestimation of S/ compared with the real S/ . Fig. 1 is the relation between the ideal factor n and /b0 with Pt electrodes. /b0 is the ¯at-band barrier height canceling the image force lowering. Since all the data plots in each treatment are distributed on each straight line, the lines are extrapolated to ideal unique
1 SX values reported in this paper can be converted to Dit using the formulas of S/ ei = ei q2 dDit and SX @/b = @Xm @/b =@/m @/m =@Xm AS/ , where Xm is the electronegativity of a metal. The conversion factor A @/m =@Xm is 2.27.
2 Detailed experimental conditions and the results of our investigation on the metal/SiC interface will be described in a full paper reported by Teraji and Hara, which will be submitted elsewhere.
MIGS Dtotal E DEXT E; it E Dit it
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Fig. 1. Ideal factor n vs. /b0 for Pt/6H-SiC(0001) interfaces. DHF, O/E, BW indicate diluted HF, oxidation and etching, and boiling water dipping process, respectively. Each plot indicates n and /b0 of individual Pt electrodes. The inset shows typical I±V curves.
barrier heights /b0 at n 1, obeying thermionic emission current transport without leakage current. /b0 for DHF, O/E, and BW electrodes are 1:011 0:015 eV, 1:365 0:024 eV, and 1:632 0:007 eV, respectively. The plot concentration in DHF and the plot broadening along the BW line suggest that the DHF interface has a large amount of Dit to form a strong pinning and the pinning is weaker in the BW interface. By obtaining /b0 also from other metal electrodes, relations between /m and /b0 are plotted as shown in Fig. 2. S/ using /b0 for the DHF, O/E and BW electrodes are 0:215 0:050, 0:739 0:079, and 0:994 0:053, respectively. The value of 0.994 is near the Schottky limit of S/ 1 having the maximum SBH controllability. C±V measured barrier heights /CV b0 provide almost the same barrier heights as /b0 and S/ 1:011 for the BW interfaces. (See footnote 2.) This indicates that /CV b0 /b0 6 /b0 at least in our experiments. In other words, the C±V barrier heights coincide with the I±V barrier heights when the interfaces have no leakage current. We estimated Dit from the equation S/ ei = ei q2 dDit [22], where ei is the per-
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Fig. 2. /m vs. /b0 plots. Barrier height controllability is raised going from DHF to BW through O/E treatment. On the BW line, the slope parameter S/ is 0.994, which is approximately in is the observed CNL of these the Schottky limit condition. /exp 0 metal/SiC interfaces. /MIGS is the predicted CNL of MIGS [25]. 0 The inset shows the energy distribution of the interface states. Distributions outside of the hatched areas are unknown in our experiment.
mittivity of the interfacial layer and d the width of the interface layer. We use the bilayer thickness 0.252 nm in the 6H-SiC(0 0 0 1) crystal as d, and use the permittivity in vacuum e0 as ei since no report was found on ei for metal/SiC interfaces. Dit for the DHF, O/E and BW electrodes are 2:8 1013 , 2:7 1012 , and 4:6 1010 states cm 2 / eV, respectively, as indicated in the inset ®gure. Even for the lower error limit of S/ 0:941 0:994 0:053, Dit has a relatively low value of 4:8 1011 states cm 2 /eV. The correlation coecients r for the DHF, O/E, and BW lines are 0.928, 0.983, and 0.996, respectively. These strong correlations suggest there is no remarkable interface reaction generating interface states between the SiC substrate and the various metals. It is important to know the line shape in the relation between /b and /m to understand the distribution Dit E as depicted in Fig. 3. A U-shaped distribution provides a curved line. A straight line is generated by a broad distribution. Local states provide a stepwise line. According to the MIGS model proposed by Terso [4], MIGS should have
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Fig. 3. Types of energy distribution in interface states and corresponding correlations of /m vs. /b0 .
a U-shaped energy distribution. It is believed that the MIGS are not the direct penetration states of the metal wave function but metal-induced states redistributed from the conduction and the valence band states where the total number of the states in the semiconductor is preserved [4]. In our observations, the data plots in Fig. 2 form straight lines at least in the observed individual barrier height ranges that are indicated by hatched areas in the
inset. This broad feature of Dit E suggests that the origin of the interface states is dierent from the MIGS that should have the U-shaped Dit E. In general, energetic broadness in electron states is a typical feature of a disordered structure. Fig. 4 shows cross-sectional lattice images of DHF and BW interfaces by transmission electron microscopy (TEM). The DHF interface has a disordered interface layer of 2 nm thickness. In the BW sample the interface is atomically abrupt and has the atom-by-atom commensurate epitaxial connection with the 1 1 0 26H-SiC == 1 1 1Ti relation without any evidence of a SiO2 formation in our cross-sectional TEM observation and electron energy loss spectroscopy (EELS) of the crosssectional sample. The lattice type of the Ti layer is the face-centred cubic (fcc) lattice. The formation of the fcc-Ti structure has previously been suggested only in an ultra-thin Ti ®lm with a thickness of several angstroms on Al surfaces [23]. We have observed the growth of the fcc-Ti layer with a thickness of 100 nm, which is the ®rst report on this thin layer single-crystal formation [24]. The bulk stable phase of Ti is hexagonal close-packed (hcp) with a lattice mismatch of 4% for the 6HSiC(0001) crystal. The lattice constant a of the observed fcc-Ti structure is 0.438 nm with a lattice mismatch of 0.79% between the 1 1 0 26H-SiC and 1 1 1Ti faces. The phase change at the interface
Fig. 4. High-resolution cross-sectional TEM lattice images with the electron acceleration at 200 eV incident from the [1 1 2 0] direction of (a) DHF and (b) BW interfaces.
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occurs in preference to the incommensurate hcpTi/6H-SiC(0 0 0 1) interface formation. This commensurate connection is strong evidence of direct Ti chemical bonding to the SiC crystal without a hydrogen terminator. In LEED observations BWtreated SiC surfaces exhibit sharp 1 1 patterns whereas DHF-treated SiC surfaces show no diffraction spot indicating that the DHF surface is disordered. The disordered surface and interfaces in DHF samples is consistent with the broad Dit E. The uni®ed disorder-induced gap state (DIGS) model has been proposed as a model to deal with interface disorder [9]. In this model, similar to the branch point mechanism, Dit E is a U-shaped continuum because the origin of the DIGS is de®ned as the penetration tails of the valence and conduction band states. The observed broad Dit E suggests that another type of disorder exists whose origin is dierent from the above mechanism. One of our important observations is the ®rst experimental ®nding of a (CNL, /exp 0 ) indicated in Fig. 2. CNL is the critical parameter to determine the amount of the interface charge, but it has been discussed only for the MIGS theory [4±6]. Here, we extend the CNL phenomenon to a more general concept that is not restricted to a speci®c origin like MIGS. The three lines intersect with each other at the CNL with /b0 of 0.797 eV and /m of 4.65 eV. The observation of one CNL for the three kinds of surface preparations indicates that the origin of Dit is the same for the three interfaces, but the amount of Dit is variable. This /exp is lo0 cated far from the MIGS±CNL /MIGS at 1.43 eV 0 estimated from previous experimental values of /b0 [20] and the value of /m of 5.4 eV assuming that /m is equal to the Miedima's electronegativp ity, XSiC ( XSi XC ) [25]. This is plotted in Fig. 2. In the MIGS theory, the interface Fermi level should be pinned at the branch point EBP where the eects of the conduction and valence band balance [4]. EBP for the 6H-SiC crystal is calculated to be 1.41 eV, which has almost the same value as the above estimation of /MIGS but absolutely dier0 ent from /exp . This dierence supports the assertion 0 that /exp is an extrinsic level, not the intrinsic level 0 based on the branch point mechanism. The disdisorder cussed disordered features give /exp in 0 /0
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this system. In other words, our experimentally , observed CNL is a disorder-related CNL, /disorder 0 not MIGS±CNL /MIGS . Considering the possibil0 ity of a multiple CNL system, another general model with an expression, BP EXT1 Dtotal E DEXT2 E ; it E Dit E Dit it
2 is given, where each term on the right has an individual /0 , and DBP it is Dit with the CNL at EBP , which originates from the MIGS or DIGS theories based on the branch point mechanism. Our metal/ disorder SiC interfaces have Dtotal DBP with it it Dit BP total 10 2 Dit < Dit 4:6 10 states cm /eV for the BW interface in the observed SBH range of 0.45±1.66 eV. This indicates that DMIGS is also lower than it 4:6 1010 states cm 2 /eV. According to the typical MIGS theory of Ref. [4], the penetration length of the metal wave function or the semiconductor layer depth modi®ed by the metal was calculated to be 0.15 nm for ZnS with the band gap of 3.6 eV that is comparable value of 6H-SiC. Nevertheless, Dit calculated for ZnS is 1014 [26]. At least, our ®nding of 6 1010 for the MIGS density strongly suggests that the MIGS penetration depth for SiC is much lower than 0.15 nm. This implies that actual penetration depths for other semiconductors may have much lower than the calculated values [4] and the semiconductors may have Dit with one or two order of magnitude lower than the predictions if one can exclude the extrinsic interface states. From XPS measurements of the core level position of C 1s, the surface potential barrier is 0.19 eV and the surface band bending is nearly ¯at of 0.08 eV for the BW surface. Using the value of 0.08 eV R the ionized total density of the surface states Ds dE is evaluated to be 6:7 1011 states/ cm2 . Let us estimate the density of the surface states Ds by assuming that the surface CNL has the same origin as the interface CNL. The energy width of charged states between the CNL and the Fermi level position is 0.607 eV( 0:797 0:19 eV). Thus, Ds is estimated at 1:1 1012 states cm 2 /eV. This value is one and a half orders of magnitude higher than Dit , implying that the residual surface states are eectively annihilated by
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the electronically ideal interface connection as well as the structurally perfect interface connection. Acknowledgement The author thanks Dr. T. Teraji for his enormous amount of experimental work. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12]
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