Comparative Study of Titanium Dioxide Atomic Layer Deposition on

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Comparative Study of Titanium Dioxide Atomic Layer Deposition on Silicon Dioxide and Hydrogen-Terminated Silicon Rungthiwa Methaapanon and Stacey F. Bent* Department of Chemical Engineering, Stanford UniVersity, 381 North-South Mall, Stanford, California 94305 ReceiVed: February 11, 2010; ReVised Manuscript ReceiVed: April 27, 2010

Atomic layer deposition (ALD) of amorphous titanium dioxide (TiO2) at 100 °C using the precursors titanium tetrachloride (TiCl4) and water (H2O) was studied on two different surfaces by in situ X-ray photoelectron spectroscopy (XPS). The initial growth rate on hydroxyl-enriched silicon dioxide (SiO2) is found to be higher than on hydrogen-terminated silicon. Moreover, the data show that the growth rate is accelerated during the first several ALD cycles on both surfaces. The interface between the SiO2 substrate and TiO2 is abrupt and composed of Si-O-Ti bonds. On the hydrogen-terminated silicon surface, the XPS results provide evidence of direct Si-Ti bond formation without traces of interfacial oxide. However, a silicon oxide layer forms on this surface after vacuum annealing, concurrent with the reduction of TiO2, suggesting that the TiO2 film is the oxygen source for the silicon oxidation under these conditions. Chlorine incorporates into the TiO2 films on both surfaces and is found to concentrate near the Si/TiO2 interface. Introduction Atomic layer deposition (ALD) is a thin film deposition technique which provides for well-controlled, uniform, and conformal film growth by exploiting a series of alternating selflimiting reactions of gas-phase precursors with the surface.1-4 Because of isolated precursor exposures, gas phase reactions are prevented, and typically, submonolayer coverages of the reaction products form during each ALD cycle. Consequently, the deposition of conformal films with precise thickness at the angstrom scale can be achieved. The process is highly controllable by managing the deposition conditions and number of cycles. The desirable properties of the films produced by ALD make this technique a popular choice for depositing ultrathin materials on both simple and complex structures that required conformal, defect-free coating. ALD can be applied to the deposition of various materials, including metals and metal oxides such as Ru, Pt, HfO2, and Al2O3.5-7 Titanium dioxide (TiO2) is one of the more widely studied ALD systems due to the desirable properties of TiO2, including its chemical stability and outstanding optical and electrical properties, which make it suitable for variety of applications.8 Moreover, TiO2 is photocatalytically active, which makes this material of interest for applications in photocatalysis, such as photodecomposition of water, photodegradation of acids, and photoreduction of carbon dioxide.9-13 It is also important for photoinduced superhydrophilic thin-film coatings for selfcleaning surfaces.14-16 Although there have been previous studies of the TiO2 atomic layer deposition process,8,17-25 the mechanistic details of the nucleation and growth chemistry on different substrates are still not well understood. Nucleation at the substrate can differ from steady state growth of the bulk due to the different chemical characteristics of the substrate and the deposited material. Differences in properties such as growth rate, surface roughness, island formation, or creation of interfacial compounds can be expected during this initial growth period. These initial proper* To whom correspondence should be addressed. E-mail: sbent@ stanford.edu.

ties can in turn dramatically alter the quality of the deposited films, especially when the size of the required devices approaches the nanoscale, i.e., when the thickness of the film is only a few atomic layers from the substrate. It is therefore important to control the initial deposition onto the substrate in order to meet the desired characteristic of the thin films. However, most of the previous studies have focused on the growth of bulk TiO2, not on the nucleation period. The effects of different surface chemistries on the nucleation process and properties of the resulting film are yet to be explored. In this work, we study the characteristics of TiO2 ALD on chemical oxide-covered silicon and hydrogen-terminated silicon. Out of many possible precursor systems for TiO2 ALD, titanium tetrachloride (TiCl4) and water (H2O) are selected as precursors due to their molecular simplicity and broad operating temperature range, resulting in several achievable TiO2 phases.22,26,27 We investigate the ALD growth at 100 °C, a temperature known to result in amorphous TiO2. This low deposition temperature is of interest for film growth on thermally sensitive materials such as polymers.28 This condition also has some potential in the field of photocatalysis since, although the anatase phase is typically the most photoactive, photoactivity of amorphous TiO2 thin films has been reported.29 TiO2 ALD is carried out in an integrated ALD reactor/ ultrahigh vacuum (UHV) chamber that allows for X-ray photoelectron spectroscopy (XPS) analysis after each precursor pulse, without vacuum break. The initial growth of TiO2 is compared for two substrates. The intensities and binding energies of characteristic peaks from the XPS spectra are used to analyze the elemental composition and chemical state of each species as the deposition progresses to lend insight into the nucleation mechanism. Experiment P-type Si (100) wafers (F ) 16.0-30.0 Ω) were used as substrates. All the samples were sequentially rinsed in acetone (OC(CH3)2) and ethanol (C2H5OH) prior to piranha cleaning (30:70, 30% H2O2:H2SO4) to remove organic contaminants. This

10.1021/jp1013303  2010 American Chemical Society Published on Web 05/21/2010

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Figure 1. Simplified schematic diagram of the in situ ALD/XPS system.

process also establishes a hydroxyl-rich chemical oxide surface on the wafer sample. To obtain a hydrogen-terminated silicon surface, samples were subsequently etched in 10% hydrofluoric acid (HF) for 10 min followed by 1 min rinsing in deionized (DI) water to remove silicon dioxide. (Caution: Both processes for surface preparation are hazardous. The piranha cleaning is a rigorous process due to the mixing of the strong acid (H2SO4) and strong oxidizer (H2O2). The reaction is highly exothermic and possibly explosive. Additionally, HF is an extremely corrosive solution, even at low concentration. It can readily penetrate skin, destroy tissue, and decalcify bone. Proper training and extreme care must be taken during these processes.) After surface preparation, the oxide-terminated and hydrogenterminated samples were then transferred to a custom-built, warm-wall, stainless steel, cylindrical ALD reactor for TiO2 deposition. The reactor was connected to an ultrahigh-vacuum chamber (base pressure ∼10-10 Torr) equipped with an X-ray photoelectron spectrometer (XPS) via a load lock, allowing sample transfer without vacuum break. A schematic diagram of the vacuum system is shown in Figure 1. Titanium tetrachloride (TiCl4, 99.90%, Acros Organics) and DI water vapor were alternately introduced into the ALD reactor via stainless steel tubing and VCR fittings. The tubing and fittings were kept at 70 °C to minimize condensation of precursor vapor in the manifold, while the substrate temperature and the chamber wall were kept at 100 °C for controlled deposition. The solenoid valves were controlled through a computer program for precise precursor exposure lengths. During the precursor dosing and reactor purging periods, 50 sccm of nitrogen gas (99.9999% N2, Praxair) was maintained by MKS mass flow controller. After precursor exposure, the samples were transferred under vacuum from the ALD reactor to the UHV chamber at room temperature for quasi-in-situ XPS analysis. The XPS instrument included a hemispherical analyzer (SPECS PHOIBOS100) and a nonmonochromatic X-ray source (SPECS XR-50) generated by exciting Al KR radiation (hν ) 1486.61 eV). Survey scan and fine scan spectra were collected at a normal angle from the surface using 0.5 and 0.05 eV step sizes, respectively. The photoemission peaks were fitted with Voigt profiles and referenced to the bulk Si 2p peak at 99.3 eV after Shirley background subtraction. The Si 2p peak, originating from the underlying substrate, was chosen for calibration. Although a temperature and pressure variation is inevitable during the sample transfer between the ALD reactor and the XPS chamber, the process does not appear to interfere with the

Figure 2. TiO2 growth on oxide-terminated Si (b) and hydrogenterminated Si ([), estimated from XPS spectra. Estimated growth rates are ∼0.16 Å/cycle on oxide-terminated Si and ∼0.09 Å/cycle on hydrogen-terminated Si.

surface chemistry or affect the results. This assertion was tested by comparing samples with and without intermediate transfer steps. The results on samples with many intermediate interventions from sample transfer and XPS scans are consistently the same as on samples subjected to continuous ALD, indicating the stability of adsorbed species on the surfaces. A study by Cheng et al. carried out on the same ALD system at different temperatures in which additional pumping steps were undertaken between precursor exposures supports the reliability of the interventional approach used here.30 To explore thermal effects, the hydrogen-terminated samples were annealed in vacuum after deposition. The surface temperature of the sample was monitored by an infrared pyrometer temperature sensor (Raytek, Marathon MA2SA series) through a glass viewport. Results and Discussion TiO2 growth curves on both oxide-terminated silicon and hydrogen-terminated silicon determined from the XPS measurements are shown in Figure 2. The TiO2 thicknesses shown in Figure 2 are estimated from the attenuation of the Si 2p peak intensities in the XPS spectra using the mean free path of the photoelectron in TiO2 according to the following formula

( )

dTiO2 ) -λTiO2 ln

ISi ISi,0

where dTiO2 is the thickness of TiO2 film, λTiO2 is the photoelectron mean free path in TiO2, and ISi,0 and ISi are the Si 2p peak intensities in the sample before and after deposition, respectively. A mean free path of 20 Å for the Si 2p electron passing through bulk TiO2 is used for all calculations.31 This expression assumes uniform film coverage at the surface; in other words, submonolayer coverage is estimated as an ultrathin TiO2 layer. Consequently, the film thickness at low Ti coverage may be slightly underestimated due to the longer actual electron mean free path through a submonolayer film. The approach therefore may introduce some uncertainty into the thickness values; however, the method is useful for obtaining an approximate thickness

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value in situ. Cross-checking the values of thickness estimated from the Si 2p signal with an estimate of thickness based on the increase of the Ti 2p peak intensities yields similar results. According to Figure 2, TiO2 exhibits nearly linear growth with number of ALD cycles on both surfaces, displaying one of the distinct properties of ALD. Interestingly, TiO2 growth seems to be promoted during the first few cycles compared to later ALD cycles, as can be seen from the abrupt change in estimated thickness from the zeroth cycle to the first half cycle. Moreover, this accelerated growth will be accentuated after correcting for the inelastic mean free path of the submonolayer. The effect is attributed to enhanced TiCl4 adsorption during the first TiCl4 vapor exposure. The result on the SiO2 surface may be explained by the hydrophilic nature of the hydroxyl-rich SiO2 surface obtained from the piranha process compared with the relatively hydrophobic surface of amorphous TiO2, as Gu and Tripp found a direct correlation between the growth rate of TiO2 and water absorption on the surface.32 The accelerated initial growth, although less significant, is also seen on the Hterminated silicon surface; however, the reason for this effect is not known. A comparison of the two data sets in Figure 2 shows that the initial growth rate on the silicon oxide surface is higher than on the hydrogen-terminated surface. The results are similar to HfO2 growth on various silicon surfaces observed by Wang et al. via infrared spectroscopy.33 They found that the IR signal of HfO2 grew linearly on hydrogen-terminated silicon during the first five cycles, with a growth rate lower than that on oxidepassivated silicon. On the other hand, the absence of an inhibited growth period in the data shown in Figure 2, especially at the hydrogen-terminated silicon surface, differs from the incubation period found in previous studies of other of metal oxide ALD systems.34-37 For example, Puurunen et al. studied ZrO2 and Al2O3 growth on hydrogen-terminated silicon and observed that the growth rate during the initial cycles of both materials increased with number of cycles, exhibiting an S-shape growth behavior.37 Although there is no apparent transition from inhibited growth period to the linear steady growth in our study, there is the possibility that all the data obtained are within the incubation period. In this case, the growth rate at higher number of cycles may be larger than what is observed here. In addition to the in situ XPS, we conducted ex situ atomic force microscopy (AFM) to investigate the changing surface roughness after deposition. The roughness of the oxideterminated sample after TiO2 deposition remains similar to the initial SiO2 roughness, indicating a uniform TiO2 film. Only a slight increase in roughness from rms ) 0.219 nm to rms ) 0.245 nm is observed after 50 cycles. On the other hand, the surface roughness of the hydrogen-terminated sample decreases from that of the initially rough HF-etched Si (100) surface with the number ALD cycles. However, the results on this surface are complicated by the changing surface roughness that occurs during the measurement process due to the formation of SiO2 underneath the TiO2 film after air exposure, causing the AFM results on this surface to be less reliable. XPS fine scans of the Si 2p and Ti 2p core level peaks collected after five cycles of deposition on the silicon oxide surface are shown in parts a and c of Figure 3, respectively. The chemical shifts of selected core levels, which provide information about the bonding environment, are plotted in Figure 3b,d as a function of number of ALD cycles. Two peaks appear in the Si 2p region (Figure 3a). The higher intensity peak at lower binding energy is assigned to bulk Si (labeled Si0), while the lower intensity peak at higher binding energy (labeled Si4+)

Methaapanon and Bent

Figure 3. Progress of XPS Si 2p and Ti 2p peak positions during ALD growth on oxide-terminated silicon. (a) and (c) show the XPS spectra from the fine scan for Si 2p and Ti 2p, respectively, while (b) and (d) plot the positions of these elemental peaks versus the number of ALD cycles.

is assigned to surface SiO2. The binding energy of bulk Si, free from the influence of surface reactions, is fixed at 99.3 eV as a reference for all scans, consistent with literature values.38 The position of the Si4+ peak in Figure 3b shifts to lower binding energy as the deposition progresses. The decreasing binding energy of the surface Si peak implies that the interfacial Si in SiO2 formed during TiO2 deposition has a lower binding energy than that of the original bulk SiO2. It is evident from Figure 3c that Ti 2p3/2 and Ti 2p1/2 peaks appear after the first half cycle of TiCl4, confirming that growth of Ti occurs on the surface from the first TiCl4 exposure. Moreover, the high binding energies of the Ti 2p peaks indicate that the Ti atoms are in the Ti4+ state.39 A shift in the Ti 2p peaks to lower binding energies is observed with increasing number of ALD cycles, as shown in Figure 3d, approaching the limiting value of Ti4+ in bulk TiO2 at ∼5 ALD cycles. These results indicate that interfacial Ti species have a higher binding energy than that in bulk TiO2 film. However, the relatively small magnitude of the shift (∼0.5 eV) indicates that Ti remains in the Ti4+ state.39 The behavior of both the Si 2p and Ti 2p peaks at the initial stages of deposition (Figure 3) can be explained by formation of Si-O-Ti bonds at the interface between SiO2 and TiO2. Because Ti is less electronegative than Si, as Ti is added to the Si-O bonds during deposition of TiO2, the Si4+ binding energy should decrease slightly from its bulk SiO2 value, as observed. On the other hand, the Ti4+ binding energy is greater right at the interface due to the proximity of the more electronegative silicon and then decreases for increase number of ALD cycles as the TiO2 becomes more bulklike. As a result, the measured binding energies of both Si and Ti decrease with increasing ALD cycles, according to the change in electron screening of the core levels. The red shift of the Si 2p and Ti 2p peaks observed after deposition thus indicates an interaction between Si and Ti at the interface through a Si-O-Ti bond. In addition to Ti and Si, O is also a key element associated with the TiO2 ALD reactions. Because of the higher electrone-

Titanium Dioxide Atomic Layer Deposition

Figure 4. (a) Peak fitting of O 1s spectra, before and after deposition. (b) Ratio of Ti bound O quantity to Ti quantity as approximated from the O 1s and Ti 2p fine spectra of TiO2 deposition on SiO2-terminated surface.

gativity of Si than Ti, the O 1s binding energy in SiO2 is higher than that in TiO2. The O 1s binding energy for O atoms in SiO2 is separated by more than 2 eV from the binding energy of O atoms in bulklike TiO2, making these two chemical states resolvable from each other in the XPS fine spectra, as can be seen in Figure 4a, collected after deposition. The intensity of the lower binding energy O peak increases with the number of deposition cycles as a result of growth of TiO2, while the intensity of the higher binding energy O peak slightly decreases due to attenuation from the TiO2 film coverage. The ratio of Ti-bound O to Ti in the deposited films can be determined by dividing the normalized integrated area of the O 1s XPS peak at ∼531 eV by the integrated area of the Ti 2p3/2 and 2p1/2 XPS peaks. The result is shown as a function of ALD cycle in Figure 4b. The data reveal a ratio of less than 2 during the initial ALD cycles, rising to the value of 2 expected for stoichiometric TiO2 after ∼5 cycles. This result appears to suggest that substoichiometric TiO2 forms prior to the deposition of stoichiometric oxide. However, as discussed above, the high Ti 2p binding energy indicates that Ti atoms are in the Ti4+ state from the very first precursor exposure, which is inconsistent with the presence of an oxygen-deficient oxide. We therefore propose the following explanation for the low ratios observed at initial ALD cycles: not all of the oxygen in the deposited TiO2 is being accounted for by the peak at ∼531 eV. Namely, the deposited oxygen at the interface is bound to both Si and Ti, as described above. When Si-O-Ti bonds form at the interface, we cannot distinguish Si-O-Ti and SiO2 species as two different peaks in the O 1s spectra because of the similar binding energy of the O 1s core level in these two species. We therefore argue that the interfacial O that is not accounted for in the O-Ti XPS peak is appearing in the O-Si peak, giving rise to the apparent O-deficient TiO2 at the interface. This assertion is corroborated by a slight red shift of the higher binding energy O 1s peak with increasing cycle number, which would result from O sharing between Si and Ti at the SiO2/ TiO2 interface. This trend also agrees well with the XPS results for Si and Ti. Consequently, we conclude that Si-O-Ti bond formation occurs at the interface of SiO2 and TiO2. Examination of Figures 3 and 4 shows that the binding energies of all elements stabilize after ∼5 ALD cycles on oxideterminated silicon, i.e., after a few angstroms of material has been deposited. At this point, the Ti binding energy approaches that in bulk TiO2. The results suggest that the completion of Si-O-Ti bond formation occurs within roughly a monolayer. The interface between SiO2 and TiO2 is therefore abrupt. In contrast to the piranha-cleaned surface, which contains hydroxyl group that are reactive toward TiCl4 adsorption, the

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Figure 5. Progress of Si 2p peak during ALD growth on hydrogenterminated silicon. (a) The XPS spectra from the Si 2p fine scan. (b) Si 2p peak positions as a function of number of deposition cycles.

hydrogen-terminated surface is known to be inert for TiCl4 adsorption.40 Growth mechanisms for other metal oxides on hydrogen-terminated surface have been proposed in several studies. Puurunen et al. suggested growth via precursor chemisorption at oxygen-containing defects that may exist from a small amount of oxygen contamination on the surface.37 Frank et al. observed the growth of Al2O3 and HfO2 via physisorption and the direct replacement of hydrogen by O-containing Hf precursors, respectively.41 Figure 5a shows the Si 2p XPS fine scans following 0 and 5 cycles of TiO2 ALD on hydrogen-terminated silicon. Only a single Si 2p peak can be resolved, which is assigned to bulk silicon (Si0) with a small contribution from H-bonded Si at the surface (Si1+). Interestingly, no oxidized Si is detected even after TiO2 deposition, despite the fact that the samples have been repeatedly exposed to water vapor in the ALD reactor at 100 °C. This result is similar to that reported by Frank et al., in which they observed no interfacial silicon oxide layer at temperatures below 100 °C for HfO2 ALD on hydrogenterminated silicon.41 With the nonmonochromatic X-ray source used in this study, the bulk Si0 peak cannot be distinguished from the interfacial Si1+ of Si-H; instead, one broad peak is observed. Nonetheless, the binding energy of bulk Si 2p is unlikely to be influenced by the changing surface chemistries. Any chemical shift observed in the Si 2p fine scan spectra can therefore be reasonably attributed to a shift of the interfacial Si core level. Figure 5b exhibits the Si 2p peak position plotted versus number of ALD cycles. The data in Figure 5b show that the Si 2p peak shifts to lower binding energy as deposition occurs. Two main conclusions can be drawn from this trend. First, the magnitude of the change in Si 2p binding energy indicates that chemical bonding, rather than only physisorption, occurs upon exposure of the hydrogenterminated silicon surface to the ALD precursors, even at the low temperatures used in the experiments. Second, the direction of the change in Si 2p binding energy indicates the formation of direct Ti-Si bonds. Specifically, the data show a trend opposite to that expected were Si-O bonds to form, in which the highly electronegative O atom would raise the binding energy of Si. Indeed, much higher Si 2p binding energies were observed in Figure 3a for SiO2. Since Ti is the only element involved in the reaction with lower electronegativity than Si, this evidence leads to the inference of direct bond formation between Si and Ti at the interface. Because the signal-to-noise ratio of the Ti 2p spectrum is poor at the very low Ti coverages present during the first few ALD cycles, this bond arrangement could not be confirmed by the Ti 2p chemical shift. However, although to our knowledge there are no previous reports of direct Ti-Si bonding resulting from TiO2 deposition on a hydrogen-

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Figure 6. (a) Atomic percentage of Cl after each precursor exposure on oxide-terminated Si (b) and hydrogen-terminated Si ([) together with (b) ratio of Cl to Ti content after each precursor exposure.

Figure 7. Reduction of Cl content in the deposited TiO2 film: (a) 5 ALD cycles and (b) 25 ALD cycles on hydrogen-terminated silicon after vacuum anneal.

terminated silicon surface, preliminary studies from our group on ALD of TiO2 on halide-terminated Ge by photoemission spectroscopy have also shown evidence for direct Ge-Ti bonding. Moreover, titanium silicides are one of the stable phases in the Ti-Si-O phase diagram,42 further supporting the existence of Ti-Si bond formation at the Si/TiO2 interface. Incorporation of Cl impurities is a known problem for metal oxide films deposited by metal chloride precursors.2 Trace amounts of Cl are commonly observed in TiO2 films when TiCl4 precursor is used.19,43 Similar results are also observed in HfO2 and ZrO2 films deposited by ALD using HfCl4 and ZrCl4, respectively.44 To understand the mechanism of Cl contamination, we have followed the appearance of Cl into the TiO2 film during initial stages of deposition on both the silicon oxide surface and the hydrogen-terminated surface using XPS. Figure 6a shows the atomic percentage of Cl, and Figure 6b shows the corresponding Cl-to-Ti ratio after each precursor exposure on both oxide-terminated silicon and hydrogenterminated silicon. The data in red correspond to the oxideterminated surface. On this surface, the ligand exchange between OH and Cl is apparent, as seen from the fluctuation of Cl composition between half and full ALD cycles. The Cl composition is high after TiCl4 exposure (half ALD cycle) and decreases significantly after water exposure (full ALD cycle). The Cl atomic percentage slightly increases after each full ALD cycle, showing that Cl accumulates in the TiO2 film under normal deposition conditions. The sudden increase of the Cl atomic percentage together with the slight decrease of the Clto-Ti ratio following the first cycle suggest that more Cl is incorporated into the film during the first cycle and that chlorine contamination is more concentrated at the interface. A similar trend is observed on hydrogen-terminated silicon as shown in blue in Figure 6a,b. In Figure 6a, the alternations in chlorine content at each half cycle are less apparent than on the silicon oxide surface because, at each step, less TiCl4 adsorbs on this rather inactive surface and a smaller proportion of Cl appears to be removed during each water pulse. Nevertheless, the trend toward Cl incorporation in the film is evident. The higher Cl-to-Ti ratio on this surface in Figure 6b suggests that the rate of Cl uptake into the film is faster on the hydrogenterminated surface than on the oxide-terminated surface. The reasons for the relatively high Cl composition in the film on the hydrogen-terminated surface may be explained as follows. The direct chemisorption of HCl onto surface Si atoms is kinetically favorable on this surface as reported in the theoretical study by Ghosh and Choi of TiCl4 adsorption on the Si-H (100) 2 × 1 surface.45 In contrast, the direct chlorination of Si atoms by HCl and TiCl4 on oxide surface at the low temperature carried out in our study is unlikely according to the results of Haukka

et al.46 The direct chlorination path may therefore increase the Cl composition on the hydrogen-terminated surface. Additionally, the hydrophobic nature of the hydrogen-terminated silicon surface may cause lower water adsorption and reduce ligand exchange opportunities between Cl and OH. These remaining Cl species, present either from unreacted Cl ligands of TiCl4 precursor or from readsorption of HCl product onto the surface, may in turn contribute to the slower TiO2 growth rate on the hydrogen-terminated surface compared to the oxide-terminated surface.43 Repetitive water exposure diminishes the Cl concentration to some extent but is unable to completely remove Cl from the films characterized in situ, where the chance of water exposure is minimal after the ALD reaction. However, our ex situ XPS results of one sample (7 cycles of TiO2 on SiO2) show no detectable Cl in the fine scan spectra after a prolonged period (2 weeks) in atmosphere. This evidence suggests that complete removal of Cl from a thin TiO2 film is possible if the water exposure is sufficiently long and the deposited film is sufficiently thin to allow water diffusion through the film. We also explored the possibility of reducing the residual Cl content by annealing the TiO2 films. Samples of three different film thicknesses (deposited for 5 cycles, 25 cycles, and 50 cycles on hydrogen-terminated Si) were annealed under ultrahigh vacuum to 350, 500, and 700 °C. The results for two representative thicknesses are shown in Figure 7. At these elevated temperature, regardless of the film thickness, Cl can be completely eliminated from both samples. However, the films studied here are very thin; annealed samples with significantly thicker TiO2 films may leave trace amounts of Cl in the film. The effect of vacuum annealing on the chemical structure of the TiO2 films was also investigated. Figures 8 and 9 show fine scans of the Si 2p and Ti 2p peaks, respectively, of the H-terminated sample before and after vacuum anneal to different temperatures for the different TiO2 film thicknesses. The spectra in Figure 8 exhibit growth of a higher binding energy Si 2p peak at ∼103 eV after annealing to 350 °C, which is indicative of SiOx species that were absent during the ALD reactions. Moreover, the intensity of this SiOx peak increases with both annealing temperature and thickness of the TiO2 films, suggesting further oxidation of Si under these conditions. At the same time, a reduction of TiO2 to TiO2-x is observed on all three TiO2 thicknesses, as seen in Figure 9a-c. The Ti 2p3/2 spectra can be resolved into three chemical states at roughly 459, 458, and 456 eV, which can be assigned to Ti4+, Ti3+, and Ti2+, respectively.47 The corresponding Ti 2p1/2 peaks can be found at 465, 463, and 461 eV. For the unannealed sample, the presence of peaks at 459 and 465 eV indicates the presence of Ti4+ species, as expected for bulk TiO2. After anneal, the appearance of lower binding energy peaks suggests that the

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Figure 8. Si 2p fine spectra of hydrogen-terminated samples (a) after 5 ALD cycles and (b) after 25 ALD cycles before and after vacuum anneal at 350, 500, and 700 °C. The insets show enlarged Si 2p spectra in the SiOx region.

deposited film become substoichiometric TiO2-x. As the annealing temperature is raised higher, Ti loses more of its O and transforms to even lower oxidation states, as seen from the increase in the lower binding energy peak intensities. A larger composition of stoichiometric TiO2 remains on the thicker film, suggesting that the reduction of TiO2 occurs at the interface. The growth of substoichiometric TiO2-x and an interfacial SiOx layer is concomitant for all samples, with the SiOx composition increasing with the extent of TiO2 reduction to TiO2-x. This trend suggests that TiO2 is the likely source of oxygen for Si oxidation during vacuum anneal. The increasing quantity of oxidized Si with increasing TiO2 film thickness and with annealing temperature supports this conclusion, i.e., both conditions enhance the O transport from the TiO2 to the Si. This process is likely thermodynamically limited since extending the annealing time (data not shown) has only a very small influence on the quantity of oxidized/reduced species. A similar annealing approach has been applied to the oxideterminated sample for comparison, as shown in Figure 9d. On this sample, the O transport from TiO2 to Si is apparently inhibited by the SiO2 interfacial layer since we observe stable TiO2 up to at least 500 °C. At higher annealing temperatures, reduction of the TiO2 is detected, and a small amount of additional SiO2 beyond the initial concentration of chemical oxide appears to grow in (data not shown). This result therefore confirms that O from the TiO2 film migrates to the Si/SiO2 interface where it oxidizes Si. Higher temperatures are required to enhance the mobility of O atoms through the SiO2 layer on the oxide-terminated sample compared to the hydrogenterminated sample for which no initial oxide layer is present. However, by comparing the relative change in peak area for the different Si and Ti oxidation states in the Si 2p and Ti 2p spectra, we calculate that the quantity of oxidized Si is not equivalent to the quantity of reduced Ti4+ after anneal. The concurrent decrease of Cl content after annealing suggests that some of the reduction of Ti4+ may occur from desorption of chlorine to the surroundings. In previous studies, an interfacial silicon oxide layer was typically observed between the silicon substrate and TiO2.18,48 Our study shows that a SiO2 interfacial layer between silicon and TiO2 does not form during a vacuum ALD process on a

Figure 9. Ti 2p fine spectra (a) after 5 TiO2 ALD cycles, (b) after 25 TiO2 ALD cycles, (c) after 50 TiO2 ALD cycles on hydrogen-terminated samples before and after vacuum anneal to 350, 500, and 700 °C for 5 min, and (d) after 25 TiO2 ALD cycles on an oxide-terminated sample before and after vacuum anneal to 350, 500, and 650 °C for 5 min.

hydrogen-terminated silicon substrate under moderate temperatures. However, this interface is unstable under elevated temperature or exposure to air. The oxide interlayer forms after vacuum anneal, by using TiO2 as an O source, even in the absence of ambient oxygen. Conclusions TiO2 deposition by ALD on oxide-terminated silicon and hydrogen-terminated silicon is investigated during the initial cycles by XPS. The growth rate on the silicon oxide surface is higher than on the H-terminated surface, likely due to the presence of hydroxyl groups on the surface which serve as reactive sites for TiO2 nucleation. An abrupt layer of Si-O-Ti bonding forms at the SiO2/TiO2 interface. In contrast, at the interface between the hydrogen-terminated surface and TiO2, the presence of direct Si-Ti bond is inferred from the XPS results, and no silicon oxide forms during the ALD process. After vacuum anneal of the hydrogenterminated sample, oxygen in TiO2 oxidizes silicon at the interface to form a silicon oxide interfacial layer. Concomitant with this process, the TiO2 film becomes deficient in oxygen, forming a

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suboxide. Cl contamination is present in the TiO2 film on both surfaces, with the Cl concentration higher closer to the interface between the two materials. Acknowledgment. This work is supported by Stanford School of Engineering and Anandamahidol Foundation. References and Notes (1) Leskela, M.; Ritala, M. Thin Solid Films 2002, 409 (1), 138–146. (2) Leskela, M.; Ritala, M. Angew. Chem., Int. Ed. 2003, 42 (45), 5548– 5554. (3) George, S. M.; Ott, A. W.; Klaus, J. W. J. Phys. Chem. 1996, 100 (31), 13121–13131. (4) George, S. M. Chem. ReV. 2009, 110 (1), 111–131. (5) Aaltonen, T.; Alen, P.; Ritala, M.; Leskela, M. Chem. Vap. Deposition 2003, 9 (1), 45–49. (6) Aaltonen, T.; Ritala, M.; Sajavaara, T.; Keinonen, J.; Leskela, M. Chem. Mater. 2003, 15 (9), 1924–1928. (7) Nyns, L.; Hall, L.; Conard, T.; Delabie, A.; Deweerd, W.; Heyns, M.; Van Elshocht, S.; Van Hoornick, N.; Vinckier, C.; De Gendt, S. J. Electrochem. Soc. 2006, 153 (9), F205–F209. (8) Ritala, M.; Leskela, M.; Niinisto, L.; Haussalo, P. Chem. Mater. 1993, 5 (8), 1174–1181. (9) Kemell, M.; Pore, V.; Tupala, J.; Ritala, M.; Leskela, M. Chem. Mater. 2007, 19 (7), 1816–1820. (10) Linsebigler, A. L.; Lu, G. Q.; Yates, J. T. Chem. ReV. 1995, 95 (3), 735–758. (11) Pheamhom, R.; Sunwoo, C.; Kim, D. H. J. Vac. Sci. Technol. A 2006, 24 (4), 1535–1539. (12) Ni, M.; Leung, M. K. H.; Leung, D. Y. C.; Sumathy, K. Renewable Sustainable Energy ReV. 2007, 11 (3), 401–425. (13) Fujishima, A.; Zhang, X. T.; Tryk, D. A. Surf. Sci. Rep. 2008, 63 (12), 515–582. (14) Paz, Y.; Luo, Z.; Rabenberg, L.; Heller, A. J. Mater. Res. 1995, 10 (11), 2842–2848. (15) Wang, R.; Hashimoto, K.; Fujishima, A.; Chikuni, M.; Kojima, E.; Kitamura, A.; Shimohigoshi, M.; Watanabe, T. Nature 1997, 388 (6641), 431–432. (16) Miyauchi, M.; Kieda, N.; Hishita, S.; Mitsuhashi, T.; Nakajima, A.; Watanabe, T.; Hashimoto, K. Surf. Sci. 2002, 511 (1-3), 401–407. (17) Aarik, J.; Karlis, J.; Mandar, H.; Uustare, T.; Sammelselg, V. Appl. Surf. Sci. 2001, 181 (3-4), 339–348. (18) Mitchell, D. R. G.; Attard, D. J.; Triani, G. Thin Solid Films 2003, 441 (1-2), 85–95. (19) Sammelselg, V.; Rosental, A.; Tarre, A.; Niinisto, L.; Heiskanen, K.; Ilmonen, K.; Johansson, L. S.; Uustare, T. Appl. Surf. Sci. 1998, 134 (1-4), 78–86. (20) Rai, V. R.; Agarwal, S. J. Phys. Chem. C 2008, 112 (26), 9552– 9554. (21) Rahtu, A.; Ritala, M. Chem. Vap. Deposition 2002, 8 (1), 21–28. (22) Aarik, J.; Aidla, A.; Mandar, H.; Sammelselg, V. J. Cryst. Growth 2000, 220 (4), 531–537. (23) Matero, R.; Rahtu, A.; Ritala, M. Chem. Mater. 2001, 13 (12), 4506–4511.

Methaapanon and Bent (24) Rai, V. R.; Agarwal, S. J. Phys. Chem. C 2009, 113 (30), 12962– 12965. (25) Kubala, N. G.; Rowlette, P. C.; Wolden, C. A. J. Phys. Chem. C 2009, 113 (37), 16307–16310. (26) Ritala, M.; Leskela, M.; Nykanen, E.; Soininen, P.; Niinisto, L. Thin Solid Films 1993, 225 (1-2), 288–295. (27) Aarik, J.; Aidla, A.; Uustare, T.; Sammelselg, V. J. Cryst. Growth 1995, 148 (3), 268–275. (28) Triani, G.; Campbell, J. A.; Evans, P. J.; Davis, J.; Latella, B. A.; Burford, R. P. Thin Solid Films 2010, 518 (12), 3182–3189. (29) Sirghi, L.; Hatanaka, Y. Surf. Sci. 2003, 530 (3), L323–L327. (30) Cheng, H. E.; Chen, C. C. J. Electrochem. Soc. 2008, 155 (9), D604–D607. (31) Cho, Y. S.; Heo, J. S.; Kim, J. C.; Moon, S. H. Chem. Vap. Deposition 2006, 12 (11), 659–664. (32) Gu, W.; Tripp, C. P. Langmuir 2005, 21 (1), 211–216. (33) Wang, Y.; Ho, M. T.; Goncharova, L. V.; Wielunski, L. S.; RivillonAmy, S.; Chabal, Y. J.; Gustafsson, T.; Moumen, N.; Boleslawski, M. Chem. Mater. 2007, 19 (13), 3127–3138. (34) Green, M. L.; Ho, M. Y.; Busch, B.; Wilk, G. D.; Sorsch, T.; Conard, T.; Brijs, B.; Vandervorst, W.; Raisanen, P. I.; Muller, D.; Bude, M.; Grazul, J. J. Appl. Phys. 2002, 92 (12), 7168–7174. (35) Gusev, E. P.; Cabral, C.; Copel, M.; D’Emic, C.; Gribelyuk, M. Microelectron. Eng. 2003, 69, 145–151. (36) Hackley, J. C.; Demaree, J. D.; Gougousi, T. J. Vac. Sci. Technol. A 2008, 26 (5), 1235–1240. (37) Puurunen, R. L.; Vandervorst, W.; Besling, W. F. A.; Richard, O.; Bender, H.; Conard, T.; Zhao, C.; Delabie, A.; Caymax, M.; De Gendt, S.; Heyns, M.; Viitanen, M. M.; de Ridder, M.; Brongersma, H. H.; Tamminga, Y.; Dao, T.; de Win, T.; Verheijen, M.; Kaiser, M.; Tuominen, M. J. Appl. Phys. 2004, 96 (9), 4878–4889. (38) Mesarwi, A.; Fan, W. C.; Ignatiev, A. J. Appl. Phys. 1990, 68 (7), 3609–3613. (39) Mayer, J. T.; Diebold, U.; Madey, T. E.; Garfunkel, E. J. Electron Spectrosc. Relat. Phenom. 1995, 73 (1), 1–11. (40) Abeln, G. C.; Hersam, M. C.; Thompson, D. S.; Hwang, S. T.; Choi, T.; Moore, J. S.; Lyding, J. W. J. Vac. Sci. Technol. B 1998, 3874– 3878. (41) Frank, M. M.; Wang, Y.; Ho, M. T.; Brewer, R. T.; Moumen, N.; Chabal, Y. J. J. Electrochem. Soc. 2007, 154 (2), G44–G48. (42) Rahmel, A.; Spencer, P. J. Oxid. Met. 1991, 35 (1), 53–68. (43) Aarik, J.; Aidla, A.; Mandar, H.; Uustare, T. Appl. Surf. Sci. 2001, 172 (1-2), 148–158. (44) Deminsky, M.; Knizhnik, A.; Belov, I.; Umanskii, S.; Rykova, E.; Bagatur’yants, A.; Potapkin, B.; Stoker, M.; Korkin, A. Surf. Sci. 2004, 549 (1), 67–86. (45) Ghosh, M. K.; Choi, C. H. Chem. Phys. Lett. 2008, 461 (4-6), 249–253. (46) Haukka, S.; Lakomaa, E. L.; Root, A. J. Phys. Chem. 1993, 97 (19), 5085–5094. (47) Vanhassel, B. A.; Burggraaf, A. J. Appl. Phys. A: Mater. Sci. Process. 1991, 52 (6), 410–417. (48) Monoy, A.; Brevet, A.; Imhoff, L.; Domenichini, B.; Lesniewska, E.; Peterle, P. M.; de Lucas, M. C. M.; Bourgeois, S. Thin Solid Films 2006, 515, 687-690.

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