Article pubs.acs.org/JPCC
Controlling the Atomic Layer Deposition of Titanium Dioxide on Silicon: Dependence on Surface Termination Stephen McDonnell,† Roberto C. Longo,† Oliver Seitz,† Josh B. Ballard,‡ Greg Mordi,† Don Dick,† James H. G. Owen,‡ John N. Randall,‡ Jiyoung Kim,† Yves J. Chabal,† Kyeongjae Cho,† and Robert M. Wallace*,† †
Materials Science and Engineering Department, University of Texas at Dallas, 800 W. Campbell Road, Richardson, Texas 75080, United States ‡ Zyvex Laboratories, LLC, 1321 North Plano Road, Richardson, Texas 75081, United States S Supporting Information *
ABSTRACT: Patterned fabrication depends on selective deposition that can be best achieved with atomic layer deposition (ALD). For the growth of TiO2 by ALD using TiCl4 and H2O, X-ray photoelectron spectroscopy reveals a marked difference in growth on oxidized and hydrogen-terminated silicon surfaces, characterized by typical and predictable deposition rates observed on SiO2 surfaces that can be 185 times greater than the deposition rates on hydrogen-terminated Si(100) and Si(111) surfaces. Large-scale patterning is demonstrated using wet chemistry, and nanometer-scale patterned TiO2 growth is achieved through scanning tunneling microscopy (STM) tip-based lithography and ALD. The initial adsorption mechanisms of TiCl4 on clean, hydrogen-terminated, and OH-terminated Si(100)-(2 × 1) surfaces are investigated in detail through density functional theory calculations. Varying the reactive groups on the substrate is found to strongly affect the probability of precursor nucleation on the surface during the ALD process. Theoretical studies provide quantitative understanding of the experimental differences obtained for the SiO2, hydrogen-terminated, and clean Si(100) and Si(111) surfaces. metal−oxide−semiconductor field-effect-transistor (MOSFET) structures, passivation layers in optoelectronic devices, and even diffusion barriers for food packaging.6−8 ALD can be carried out under a wide range of processing conditions, which can significantly impact both the growth characteristics and the physical properties of the resulting films.8,9 Most importantly, ALD is typically carried out at temperatures low enough for the hydrogen termination on silicon to remain stable (i.e., ≤ 300 °C). Also, the very nature of ALD suggests that surface chemistry will play a critical role in the deposition, or lack of deposition, on a given substrate. Previous work10 showed that hydrogen-terminated Si(100) surfaces are unfavorable for the deposition of TiCl4 at room temperature under ultra-high-vacuum (UHV) conditions. These surfaces were prepared by in situ high-temperature flashing to 1200 °C to remove native SiO2 from the silicon substrates and exposure to atomic hydrogen at 380 °C to passivate the surface. This finding was the basis for selective deposition and patterning using STM tip-based lithography techniques,1 with good titanium growth on bare silicon and almost no detectable deposition on hydrogen-terminated
I. INTRODUCTION Advances in scanning tunneling microscopy (STM) tip-based hydrogen lithography have provided a means for changing the surface reactivity of a substrate with atomic precision. Starting with hydrogen-passivated silicon surfaces, atomically precise areas of bare silicon can be prepared.1 Dagata et al.2 reported the surface modification of Si(111) by the incorporation of ambient oxygen into the substrate stimulated by biasing an STM tip relative to a hydrogen-passivated silicon surface in air or oxygen ambient. The oxygen incorporation was found to be dependent on tip bias, current, and oxygen partial pressure. It was proposed that such surface modification could generate regions that could be selectively etched using STM-based nanolithography.2 Since then, atomic force microscopy has been used to stimulate the patterned oxidation of a number of substrates, including silicon, GaAs, and silicon carbide.3−5 However, the atomic precision demonstrated using STM hydrogen depassivation lithography promises to control atomic-scale patterning. Utilizing these techniques of tip-based surface modification as true nanolithography tools requires the use of deposition methods that can be selective over this narrow range of surface terminations. Atomic layer deposition (ALD) is a useful tool for the highly controllable deposition of a range of materials. Metal oxides deposited by this technique are used as gate dielectrics in © 2013 American Chemical Society
Received: June 17, 2013 Revised: September 4, 2013 Published: September 9, 2013 20250
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(98% H2SO4/30% H2O2, 3:1, v/v) at 90 °C for at least 30 min and rinsed with DIW to provide a clean SiO2 surface free of hydrocarbons or metal contamination prior the etching step. The Si−H surface was then prepared by 1 min of etching in diluted hydrofluoric acid (1:5, v/v with DIW) followed, in the case of Si(111), by 2.5 min in ammonium fluoride.25 The sample was then briefly rinsed with DIW, dried with nitrogen, and loaded as quickly as possible (≤5 min) into the ALD reactor. [Caution: Both of these chemical processes are hazardous. The reaction between H2SO4 and H2O2 is exothermic and possibly explosive. HF is a highly corrosive material that requires the use of Teflon, rather than glassware, and can easily penetrate the skin, bond with Ca2+, and cause nerve damage. As such, even a small exposure (e.g., 2−10% of the body) can be fatal. Proper training is required before handling or working with these chemicals, and appropriate personal protection equipment should be worn at all times when carrying out these sample preparations.] The commercial ALD reactor used in this work was a hotwalled Savannah 100 apparatus from Cambridge NanoTech Inc. TiO2 was deposited at substrate temperatures of 30, 100, and 150 °C using titanium tetrachloride (TiCl4) and water (H2O) as precursors. The chamber was purged in flowing Ar at a pressure of ≤2 × 10−1 mbar between precursor pulses. The TiCl4 precursor pulse was 0.1 s long and resulted in a chamber pressure of 3 × 10−1 mbar, whereas the H2O pulse was 0.05 s long and resulted in a chamber pressure of 8 × 10−1 mbar. The purge time between precursor pulses was 60 s to maximize the removal of mobile precursors (noncovalently bonded) on the surface, which could otherwise still be present during the next precursor pulse. The thickness was controlled by varying the number of ALD cycles, and the resulting thickness on each surface was then estimated by X-ray photoelectron spectroscopy (XPS). Following ALD, the samples were loaded into a multideposition/characterization system (base pressure < 2 × 10−11 mbar) described elsewhere.26,27 XPS analysis was carried out using a monochromatic Al Kα X-ray source described previously.28 For spectral analysis, the curve-fitting software AAnalyzer was used.29 Spectra were fitted with a Voigt line shape, and a Shirley background subtraction was used. For estimations of TiO2 thickness and coverage from the XPS spectra, the ratio of the Ti 2p to the Si 2p core levels was used. Details of this calculation can be found in the Supporting Information, as well as previous publications.30,31 All of the DFT calculations were performed using plane-wave basis sets and projector-augmented-wave (PAW) pseudopotentials, as implemented in the VASP (Vienna Ab Initio Simulation Package) code.32,33 The electronic wave functions were represented by a plane-wave basis with a cutoff energy of 500 eV. The exchange correlation interactions were included by using the semilocal Perdew−Burke−Ernzerhof (PBE) functional of the generalized gradient approximation (GGA).34 The PAW pseudopotentials represent the effects of the inner electrons. The 3p6 electrons of Ti were included explicitly as valence electrons in the calculations (together with the 3d2 and 4s2 electrons). For Cl, Si, and O atoms, the common 3s2p5, 3s2p2, and 2s2p4 valence states, respectively, were used. The unit cell of the ideal Si(100)-(2 × 1) surface is made of two Si atoms per layer. The Si(100)-(2 × 1) surface was modeled with a periodic slab containing six Si atomic layers, with the bottom Si layer passivated by two H atoms per Si atom to account for the effects of the dangling bonds. Depending on
silicon.10 A similar result was reported by Mitsui et al.,11 who highlighted the self-limiting nature of the TiCl4 adsorption on clean silicon. In contrast, a similar study12 using HF to prepared hydrogen-terminated silicon resulted in a much lower selectivity between H- and SiO2- terminated Si(001) surfaces (only a factor of 2 in growth rate) upon sequential exposures of TiCl4 and H2O. There are two possible reasons for the difference between this result and the earlier in situ studies: (1) the ALD environment (i.e., 100 °C, ∼ 1 mbar, alternating H2O pulse) lowers the stability of a hydrogen-terminated surface and (2) in situ atomic hydrogen passivation is required. Although HF-prepared surfaces have been shown to be stable in ALD environments (H2O pulses at 100 °C),13,14 these results indicate that a careful examination of the TiCl4 reactions with hydrogen-terminated and oxidized silicon surface, and the subsequent TiO2 growth by ALD is needed. In this study, the selectivity of TiO2 deposition using TiCl4 and H2O on hydrogen-terminated, oxidized, and clean silicon surfaces was studied. Initially, oxidized and hydrogenterminated Si(100) and Si(111) surfaces were prepared, with H termination being obtained by atomic H dosing of clean surfaces, HF etching, or use of NH4F. Depositions were carried out at 30, 100, and 150 °C. Depositions were also carried out on Si(100) surfaces with a range of different patterns at 100 and 150 °C. To obtain deeper insight into the fundamental mechanisms leading to the different rates of deposition of TiO2 on Si surfaces, density functional theory (DFT) calculations were used to model the initial adsorption mechanisms of TiCl4 on clean, H-passivated, and OH-terminated Si(100)-(2 × 1) surfaces. Such modeling is important because there have been few theoretical studies predicting reaction mechanisms and associated energetics of ALD reactions and film growth15−24 and all of them used cluster models to represent the substrate. For instance, Hu and Turner21 studied the initial reactions involved on the ALD of TiO2 on SiO2 surfaces from TiCl4, but also from TiI4.22 They showed that the effects of tunneling were negligible for all of the adsorption mechanisms proposed. More recently, Ghosh and Choi23 studied the energetics of the initial reaction mechanisms of TiCl4 onto H-passivated and OHterminated cluster models,24 representing the Si(100) surface. Their main conclusions were that the initially generated HCl molecule is easily trapped on the OH-terminated Si cluster,24 thus facilitating HCl readsorption and the presence of Cl impurities. On the contrary, the H-passivated Si cluster model provides a less re-activating chemical environment, making TiCl4 adsorption more difficult.23 A detailed study of the energetics and kinetics of the initial adsorption mechanisms of TiCl4 on a surface model of the Si(100):(2 × 1) system, terminated with different reactive groups, is thus necessary to understand the diverse deposition rates obtained in the experiments presented herein.
II. EXPERIMENTAL AND THEORETICAL METHODS For the experiments carried out in this work, Si(111) and (100) wafers, purchased from WRS Materials (Cz, n-type, double-side polished with a nominal resistivity of 1−5 Ω·cm), were used. Wafers were cut into rectangular samples of approximately 1.0 cm × 1.0 cm and were cleaned by sequential rinsing in a flowing stream of deionized water (DIW), ethyl acetate, acetone, and DIW. This step was mainly used to eliminate organic contamination and particles due to the sample cutting. Subsequently, the samples were immersed in piranha solution 20251
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the specific type of calculation, surface dimers were kept clean, passivated with hydrogen, or terminated with hydroxyl OH groups to model a previous adsorption of H2O molecules. Because VASP implements a periodic-boundary-conditions scheme, we introduced a vacuum region in the direction perpendicular to the slab with a 15-Å thickness to avoid spurious interactions between replica images. The TiCl4 molecules were initially adsorbed on the dimer-reconstructed side of the slab. To make their interaction sufficiently weak, a (4 × 4) unit cell with four dimers along the dimer row was employed. Each Si layer comprised 16 atoms, for a total amount of 96 Si atoms in the model system. All atoms, except for the two bottom Si and H-passivating layers, were allowed to relax. The kinetic barriers, transition states, and reaction paths were obtained using the climbing image−nudged elastic band (CINEB) method.35,36 This method was previously shown to be reliable in obtaining the minimum-energy path (MEP) between a set of two different states37 where the reaction path is divided into a set of images “connected with a spring” as described by Lelis-Sousa and Caldas.38 During the relaxation, the initial and final states were kept frozen as the images moved according to the constraint of the “elastic band.” The MEP was then found when the components of the forces perpendicular to the elastic band vanished, with the relative positions of the images and the barrier being determined by the parallel components of the forces. The CI-NEB method is a slight modification of the original NEB in which, after the initial relaxation of the images, the highest-energy image is driven up to the saddle point by inverting the force acting on it along the tangent to the spring connecting the different images. In this way, the highest-energy image maximizes its energy along the elastic band and minimizes it in all other directions. When the energy converges, it is at the exact saddle point. In all cases, a (3 × 3 × 1) k-point mesh (that gives 10 irreducible k points) within the Monkhorst−Pack scheme was used to ensure a convergence of 10 meV/unit cell. Structural relaxations were performed without including any type of symmetry, to a tolerance of 10−4 eV in the total energy and 0.01 eV/Å in the forces on every atom, for both standard and CINEB structural minimizations. The adsorption energies at T = 0 K were obtained according to the equation
Figure 1. XPS Ti 2p core-level spectra after 45 cycles of TiCl4/H2O on (A) OH-terminated Si(111), (B) OH-terminated Si(100), (C) hydrogen-terminated Si(111), and (D) hydrogen-terminated Si(100). The three spectra in each panel corespond to depositions at 30, 100, and 150 °C (top to bottom or black, red, and blue, respectively, as indicated in panel A).
determined to be in an oxidized chemical state. It is not surprising that the TiO2 grows equally well on both the Si(100) and Si(111) surfaces, because precursor reactions occur with the overlying amorphous SiO2 layer and should have little to no dependence on the underlying substrate orientation. In panels C and D of Figure 1 are shown the Ti 2p spectra for the ALD of TiO2 at 30, 100, and 150 °C on hydrogen-terminated Si(111) and Si(100) surfaces, respectively. It is clear that the deposition on the hydrogen-terminated Si(111) surface at 30 °C is lower than that on SiO2 and that this concentration reduces to trace amounts when the deposition is carried out at 100 and 150 °C. This contrasts with the oxidized Si(111) surface and suggests that highly selective growth using TiCl4 and H2O on oxidized and hydrogen-terminated Si(111) is possible. More surprising are the almost identical observations on hydrogen-terminated Si(100) relative to the hydrogenterminated Si(111), which appears to conflict with earlier reports.12 Possible reasons for these apparent contradictions are discussed below; however, from these results, it is clear that equally selective depositions should be possible on both the Si(100) and Si(111) surfaces. Comparisons of the calculated coverages (for films of 1 ML) for the TiO2 grown on the four different surfaces at all three temperatures are summarized in Table 1. It is clear that hydrogen-terminated Si(111) and hydrogen-terminated Si(100) are equally resistant to TiCl4/H2O ALD at temperatures of 100 and 150 °C. The increase in TiO2 thickness for TiCl4/H2O deposited at 30 °C (when compared to those for the 100 and 150 °C depositions) on all four surfaces is consistent with other studies that have shown an increase in ALD growth per cycle for a range of metal
Eads = [E(slab + molecule) − E(slab) − NmolE(mol)] /Nmol
where E(slab + molecule) is the energy of the supercell with Nmol molecules adsorbed (TiCl4 in this case), E(slab) is the energy of the Si slab with its corresponding terminations (clean surface, H-passivated, or OH-terminated), and E(mol) is the energy of the gas phase of TiCl4.
III. RESULTS A. TiO2 Deposition Selectivity Dependence on Temperature, Surface Orientation, and Surface Termination. Because previous experiments 12 showed poor selectivity for hydrogen-terminated Si(100) compared to oxidized Si(100), this work examined Si(111) substrates upon which a more complete hydrogen passivation was obtained, and Si(100) samples were used for comparison with early studies. Panels A and B of Figure 1 show the Ti 2p spectra for the ALD of TiO2 at 30, 100, and 150 °C on oxidized Si(111) and Si(100) surfaces, respectively. From the binding energy position of the Ti 2p3/2 feature (456.4 eV), the titanium was 20252
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patterning by simple chemical preparation, a rectangular sample was “striped” with four regions of alternating hydrogen and oxide termination (see Figure 2). To demonstrate nanometer-
Table 1. Thicknesses (Å) and Coverages (%) of TiO2 Films Grown on Silicon at Various Temperaturesa
a
sample
30 °C
Si(111)-Ox Si(100)-Ox Si(111)-H Si(100)-H
19.5 Å 19.5 Å 1.6 Å (59%) 5.1 Å
100 °C 12.7 13.1 0.05 0.08
Å Å Å (2%) Å (3%)
150 °C 10.2 Å 9.7 Å 0.14 Å (6%) 0.1 Å (4%)
After 45 cycles of ALD.
precursors at deposition temperatures close to room temperature.39−42 This increased growth rate is attributed to insufficient purging of the H2O precursor, which was shown to take as long as 600 s at room temperature.39 The deposition rate on oxidized silicon is ∼0.28 Å/cycle at 100 °C and 0.22 Å/ cycle at 150 °C; however, the deposition rate on hydrogenterminated silicon is 0.0015 Å/cycle at 100 °C and 0.0025 Å/ cycle at 150 °C. These deposition rates indicate a selectivity for deposition on oxidized over hydrogen-terminated silicon of >185 and >85 for the 100 and 150 °C substrate temperatures, respectively. The demonstrated high selectivity by ALD of TiO2 using TiCl4 and H2O on SiO2 and clean silicon when compared to a hydrogen-terminated surface differs significantly from the results reported previously by Methaapanon and Bent,12 which showed much less selectivity. First, the differences in that work compared to this study should be highlighted. To chemically prepare the hydrogen-terminated surface, the work described here utilized a 1-min etch follow by a short (∼5−10s) DIW rinse, whereas Methaapanon and Bent used a 10-min etch followed by a 1-min DIW water rinse, which can produce very different surfaces, as shown in a previous study.43 Also, in this study, the ALD was carried out ex situ, and the samples were transferred in air to the UHV chamber for XPS analysis. Methaapanon and Bent utilized an in situ ALD/XPS system, which allowed for analysis with no vacuum break but required the sample to be initially pumped from air (∼1000 mbar) to UHV pressures prior to ALD. Hersam et al.44 previously showed how a hydrogen-terminated surface can be deteriorated when transitioning between UHV and atmospheric pressures if the sample is not properly shielded to prevent lineof-sight bombardment with kinetically energetic free radicals. Here, 45 cycles were carried out in one single, 90-min process, whereas the experiments described by Methaapanon and Bent most likely required considerable hours (and possibly days) to allow for analysis between each cycle. The total duration coupled with the repeated exposure to the ALD environment might have led to an increased degradation in the hydrogen termination. Finally, to calculate the thickness of TiO2, Methaapanon and Bent used the attenuation of the Si 2p core level, whereas, in this work, the ratio of the Ti 2p to Si 2p core levels was used. When only the Si 2p attenuation is used, any spurious surface contaminant that adsorbs during the course of the experiment (e.g., hydrocarbons or water) will attenuate the Si 2p signal further and cause an overestimation of the TiO2 thickness. Attempts to replicate the conditions reported by Methaapanon and Bent are discussed in the Supporting Information. B. TiO2 Depositions on Patterned Si(100) Substrates. To demonstrate the potential application of the substrate termination dependence of TiCl4-based TiO2 ALD, hydrogenterminated samples were prepared with patterned areas of either clean or oxidized silicon. To demonstrate a large-area
Figure 2. XPS Ti 2p core-level spectra after 45 cycles of TiCl4/H2O on a patterned sample. The predeposition surface condition of the sample is illustrated on the right, highlighting both SiO2- and hydrogen-terminated regions. The arrows indicate the region from which each of the three XPS spectra were taken. TiO2 was detected on the SiO2 region, but only trace amounts were seen on the hydrogenterminated regions. This is a practical demonstration of patterned TiO2 deposition.
scale patterned deposition, the previously described STM patterning was carried out on a hydrogen-terminated Si(100) sample to expose regions of clean silicon. B.1. TiO2 Depositions on Large-Area Chemically Patterned Si(100) Substrates. A striped rectangular sample with regions of both hydrogen and SiO2 termination was fabricated by initially growing a ∼10-nm thermal oxide. Two regions were covered with Kapton tape to protect those SiO2 regions from a HF etch, which resulted in two oxide-free strips. After removal of the Kapton tape, a piranha (H2SO4/H2O2) etch served the dual purpose of removing tape residues from the SiO2 regions and growing a thinner SiO2 layer on the HF-etched silicon. Finally, a second HF etch was carried out using the minimum time required to remove the thin SiO2 layer and replace it with a hydrogen-terminated surface. The result was two regions with
