Epitaxial Growth and Orientational Dependence of Surface

Zhen Zhang and John T. Yates , Jr. Chemical Reviews 2012 112 (10), ... A. Nakaruk , H. Chen , A. Waibel , P. Koshy , C. C. Sorrell. e-Journal of Surfa...
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J. Phys. Chem. C 2010, 114, 6595–6601

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Epitaxial Growth and Orientational Dependence of Surface Photochemistry in Crystalline TiO2 Rutile Films Doped with Nitrogen Takeo Ohsawa,†,‡ Michael A. Henderson,† and Scott A. Chambers*,† Fundamental and Computational Sciences Directorate, Pacific Northwest National Laboratory, Richland, Washington 99352, and WPI-AIMR, Tohoku UniVersity, Sendai 980-8577, Japan ReceiVed: January 11, 2010; ReVised Manuscript ReceiVed: February 17, 2010

We have prepared and investigated the structural, compositional, morphological, and photochemical properties of N-doped TiO2(110), -(100), and -(001) epitaxial films grown by means of plasma-assisted molecular beam epitaxy. The N solid solubility is limited to ∼1-2 atom % of the total anions in the lattice in films where excellent long-range structural order is maintained throughout growth. The photochemical activity of the resulting surfaces was evaluated by using hole-mediated decomposition of adsorbed trimethyl acetate. Undoped surfaces of the three orientations exhibited comparable photochemical activities. However, the dependence of the photochemical activity on N concentration shows a marked crystallographic dependence. The results are rationalized in terms of the apparent crystallographic anisotropy of hole mobility as well as hole trapping and detrapping probabilities. 1. Introduction Considerable effort has gone into fundamental investigations of heterogeneous photocatalysis on oxide semiconductors. Titanium dioxide, TiO2, has been widely studied since its utility for the photoelectrolysis of water was discovered.1 In addition, TiO2 has also been of interest for the degradation of organic pollutants,2,3 hydrophilic coatings,4 and dye-sensitized solar cells.5,6 The surface structure and reactivity of TiO2, particularly rutile TiO2(110), have been extensively studied by using stateof-the-art microscopic and spectroscopic techniques to provide insight into molecular-level processes on oxide surfaces.7-12 However, one of the major limitations of TiO2 is that with a 3 eV bandgap, titania absorbs relatively little of the solar spectrum. Asahi et al.13 first proposed doping TiO2 with N as a way to shift its band gap into the visible. Following this suggestion, a significant number of studies with various photochemical reactions have been performed with visible-light irradiation with what is claimed to be N-doped TiO2 (TiO2-xNx). However, the origin of visible-light photocatalytic activity and the actual role of nitrogen in this material are still controversial. A number of reports state that bandgap narrowing results from either hybridization of N 2p states with the O 2p band, or localized N 2p state formation above the top of the valence band.14-17 Most previous work on TiO2-xNx has been based on the properties of nanoparticles. It is critically important to determine specifically how N interacts with and incorporates into the TiO2 lattice in order to understand the resulting electronic structure and photocatalytic activity. However, such determinations are exceedingly difficult in light of the heterogeneity and multiple surface orientations found in nanoparticle assemblies. Recent progress in plasma-assisted molecular beam epitaxy (PAMBE) makes it possible to grow high-quality TiO2-xNx epitaxial films for both rutile and anatase polymorphs.18-20 These studies revealed that the N solubility limit for phase-pure material is * To whom correspondence should be addressed. Phone: +1 509 371 6517. Fax: +1-509 371 6498. E-mail: [email protected]. † Pacific Northwest National Laboratory. ‡ Tohoku University.

approximately 2 atom % within the anion sublattice (i.e., x ) 0.04 in TiO2-xNx). We have used trimethyl acetic acid (TMAA) to determine the photochemical activity of the resulting surfaces. TMAA is a useful photochemical probe molecule because it undergoes acid dissociation on rutile(110) and anatase(001) to form densely packed and stable trimethyl acetate (TMA) adlayers at room temperature, in which each TMA anion bridges across two Ti4+ cations.21-25 Exposing these surfaces to broadbandultraviolet(UV)-plus-visiblelightresultsinelectron-hole (e-/h+) pair formation in the substrate, followed by a holemediated reaction in which TMA decomposes to CO2 plus an organic fragment.26 Detection of CO2 via a quadrupole mass spectrometer (QMS) constitutes a direct probe of surface photochemical activity and allows hole availability and mobility to be determined. 2. Experimental Section In the present work, we have examined hole-mediated photochemical decomposition of TMA on surfaces of crystalline TiO2-xNx rutile grown homoepitaxially by plasma-assisted molecular beam epitaxy (PAMBE) in the (110), (100), and (001) orientations. Rutile TiO2(110), -(100), and -(001) single-crystal substrates were etched in buffered HF and rinsed in deionized water to remove impurities. To prepare atomically flat surfaces, the substrates were then annealed in a tube furnace with flowing oxygen at atmospheric pressure. The annealing temperature for (110)-, (100)-, and (001)-oriented crystals was 900, 600, and 900 °C, respectively. Surface morphology was subsequently examined by atomic force microscopy (AFM). Once under vacuum, all TiO2 substrates were cleaned by room temperature exposure to activated oxygen from an electron cyclotron resonance (ECR) plasma source at a pressure of 2 × 10-5 Torr in the system loadlock, resulting in complete removal of adventitious carbon, as evidenced by in situ X-ray photoelectron spectroscopy (XPS). TiO2-xNx films were then grown by using a mix of ECR-activated oxygen and nitrogen, as described elsewhere.18-20 The actual N concentration in the film resulting from a given combination of oxygen and nitrogen partial pressures in the chamber during growth was not always the same

10.1021/jp1002726  2010 American Chemical Society Published on Web 03/10/2010

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Figure 1. AFM images of HF-etched and O2-annealed TiO2(110) (a), TiO2(001) (b), and TiO2(100) (c) single crystal surfaces. Height profiles indicate that step heights correspond to projections of the relevant lattice parameters normal to the surface.

from growth to growth, or orientation to orientation. Thus N 1s and O 1s photoelectron peak areas were used to determine the value of x for each film. This method yields a lower limit for x. Nuclear reaction analysis (NRA) typically yields larger x values than XPS, and we used x values averaged over the two methods in our previous work.19 The (00) specular beam intensity in the reflection high-energy electron diffraction (RHEED) patterns was monitored during deposition, yielding large-amplitude intensity oscillations resulting from layer-bylayer growth. XPS core-level and valence-band measurements were carried out in situ after growth by using a GammaData/ Scienta SES 200 photoelectron spectrometer with monochromatic AlKR X-rays. The film surfaces were dosed with TMAA (Aldrich, research grade) in an appended chamber using a pinhole directional doser, resulting in TMA adlayer formation. These adlayers were then exposed to UV-plus-visible light from a mercury lamp source (ORIEL) in UHV, and CO2 at mass 44 was monitored with a quadrupole mass spectrometer.

Figure 2. RHEED intensity oscillations during the epitaxial growth of undoped and N-doped (N% ) 0.01%) TiO2(110) films.

3. Results and Discussion Figure 1 shows AFM images for HF-etched and annealed TiO2(110), -(001), and -(100) surfaces. All surfaces showed terrace-step structures in which the step heights correlate with the relevant lattice parameter projections normal to the surface, as evidenced by the height distributions shown at the bottom of Figure 1. These surface morphologies are comparable with those reported elsewhere.27,28 The substrate surface preparation described above is essential for successful homoepitaxial growth of high-quality rutile films. Following the transfer of TiO2 substrates into the growth chamber, all surfaces were recleaned at room temperature by exposure to activated oxygen and heated to the desired growth temperature in activated oxygen prior to growth. Figure 2 shows typical RHEED intensity oscillations for TiO2-xNx(110) growth at x ) 0 and 0.02 at a substrate temperature of 550 °C. Large-amplitude intensity oscillations for the (00) beam were observed for all N concentrations up to x ) 0.02, revealing layer-by-layer growth and the formation of rather flat surfaces. On the basis of the RHEED oscillation period, the growth rate was ∼0.05 Å/s in this orientation. The value of x in TiO2-xNx was controlled by varying the nitrogen partial pressure at a fixed oxygen partial pressure. Figure 3a shows the N mole fraction within the anion sublattice (x in TiO2-xNx) as a function of the nitrogen partial

Figure 3. N content (x) as a function of PN2 partial pressure introduced during the growth of rutile (110) (a). The x value was determined using N 1s and O 1s peak areas and atomic photoemission cross sections. Full-width-at-half-maximum (fwhm) for Ti 2p3/2 and O 1s peaks vs PN2 partial pressure (b). The size of the symbol denotes the magnitude of the uncertainty.

pressure fraction, PN2/(PN2 + PO2). Here, x was determined from N 1s and O 1s core-level photoemission peak areas and X-ray photoemission cross sections, as described elsewhere.29 The value of x increased monotonically with PN2/(PN2 + PO2). However, nonideal films resulted when x exceeded 0.02. Previous work shows that excessive nitrogen flux results in secondary phase formation and/or the formation of heavily defected films, rather than structurally excellent TiO2-xNx(110).18,19 The threshold for poor film growth appears to be x > ∼0.02,

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Figure 4. Final RHEED patterns for clean single-crystal (a-c), as well as undoped (d, e) and N-doped (x ) 0.02) TiO2 films (g-i) in the (110), (001), and (100) orientations.

Figure 5. O 1s (a), Ti 2p (b), and valence band (c) XPS spectra for undoped epitaxial TiO2(110), -(100), and -(001) films. The inset shows the leading edge of the valence band.

and our present results corroborate this finding. TiO2-xNx film grown with x > ∼0.02 exhibited a large increase in the fullwidth at half-maximum (fwhm) in the Ti 2p3/2 spectrum, as seen in Figure 3b, and the presence of a Ti3+ feature at lower binding energy. The formation of disordered, N-containing secondary phases, some containing Ti3+, results in a substantial increase in the Ti 2p3/2 peak width.18,19 To avoid the ambiguities associated with these ill-defined films, we limited our investigation to structurally excellent films with x e ∼0.02. Figure 4 shows the final RHEED patterns for (110)-, (001)-, and (100)-oriented single crystals and TiO2-xNx film surfaces. Strong patterns are seen for all surfaces, although there is some intensity modulation in the streaks for the (100) surface at x ) 0.02, indicating some roughness for this surface. These results show that overall our surfaces are relatively flat, which is important because rough surfaces typically have high concentrations of steps and kinks, which may in turn act as electron and/ or hole traps or as recombination centers. Figure 5 shows O 1s, Ti 2p, and valence band spectra for undoped TiO2 epitaxial films in the three orientations. These spectra serve as a baseline to which we will compare the analogous spectra for N-doped films. In pure, homoepitaxial rutile TiO2, interstitial Ti3+ diffuses from the growth front into the bulk during the deposition, resulting in deep blue coloration, n-type conductivity, and interstitial Ti3+, as detected by electron

paramagnetic resonance (EPR).19 The conductivity is correlated with the presence of interstitial Ti3+, which is a shallow donor in rutile. In all samples, there were neither C 1s peaks nor a second O 1s peak 1.5-2.0 eV to higher binding energy (BE) from the lattice oxygen peak, indicating clean surfaces without any measurable organic or hydroxyl contamination. The O 1s and Ti 2p peaks exhibit very similar line shapes for all films. However, there are slight BE differences for the three orientations. These reveal slight differences in band bending. Linear extrapolation to the energy axis for the leading edge of the valence band spectra (Figure 5c, inset) reveals that the (110)oriented film exhibits a valence band maximum (VBM) of 2.98 eV relative to the Fermi level. In contrast, the VBMs for the (100)- and (001)-oriented films are 3.22 and 3.32 eV, respectively. Inasmuch as the bandgap of rutile is nominally 3.0 eV, these results clearly reveal the n-type character for all three orientations resulting from interstitial Ti3+ doping. In no case is the Ti3+ concentration sufficiently high at the surface to result in detectable Ti3+-derived peaks in either the Ti 2p or VB peaks spectra. These fall at ∼2 eV lower binding energy than the Ti4+derived lattice peak and ∼1 eV below the Fermi level, respectively.10,19 Figure 6 shows XPS spectra for our TiO2-xNx(001) film set. These are also typical of the other orientations. Increasing the nitrogen partial pressure at fixed oxygen partial pressure results

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Figure 6. N 1s (a), O 1s (b), and Ti 2p (c) XPS spectra for N-doped TiO2(001) films. A contour plot of Ti 2p is also shown in part c.

Figure 7. Valence band spectra for TiO2-xNx rutile films in the (110) (a), (001) (b), and (100) (c) orientations.

in a monotonic increase in the N 1s peak intensity, as shown in Figure 6a. The N 1s spectrum consists of a single peak at 396.6 eV, which is a characteristic of substitutional N3-.18,19 The O 1s spectra are virtually identical for undoped and N-doped films, resulting in no change in line shape (Figures 5a and 6b). Thus, N incorporation does not perturb the O sublattice in any measurable way other than the substitution event itself. Our expectation is that substitutional N at an O lattice site would exhibit a -2 formal charge. The -3 charge state most likely occur because 2p holes on substitutional N2- are compensated by itinerant electrons from interstitial Ti3+. There is a systematic increase in asymmetry to lower binding energy in the Ti 2p3/2 and 2p1/2 peaks with increasing N concentration, as seen in Figure 6c. This change is due to an unresolved new peak ∼1 eV to lower binding energy previously assigned to Ti ions bound to five oxygen and one nitrogen ligands.18,19 The growth of this feature is clearly seen in the contour plot at the top of Figure 6c. We show in Figure 7 valence band spectra for N-doped rutile films in the three orientations and we compare these to those for the undoped analogues. The differences are minor, bordering on negligible, as noted earlier for the (110) orientation.18 The reason for the very small differences is that the spectra are dominated by contributions from O and Ti due to the limited N solubility in epitaxial TiO2 rutile, Exposure of TiO2-xNx film surfaces to TMAA was carried out immediately after transfer under UHV conditions to the appended photochemistry chamber. Details of this system can be found elsewhere.30,31 It is well-known that TMAA adsorbs on two Ti sites with bidentate manner, resulting in TMA at the surface. Since the cation-cation spacing on the (100) surface is the same as that on (110), we can expect much the same kind of bonding of TMA to both surfaces. This is the case for formic acid, which has been studied on both surfaces. The structure and bonding on the (001) remains unexplored. It has been demonstrated with STM measurements that TMA sorbs

in an ordered fashion on rutile (110) and anatase (001).23,31 In rutile (110), the TMA adlayer exhibits a (2 × 1) reconstruction in keeping with the size of the anion relative to the Ti-Ti spacing along rows of five-coordinate Ti on the surface. Moreover, the TMA O 1s peak area relative to the lattice oxygen peak area is consistent with this structural model.30,31 To make a quantitative comparison of TMA photodecomposition rates for N-doped rutile films as a function of N doping level, CN, we must correct for differences in the initial TMA coverage. A coverage of 1.0 monolayer (ML) is defined in terms of the number of available surface Ti sites to which TMA binds. By this definition, the TMA saturation coverage is 0.5 ML. To determine coverage, we use a combination of C 1s and O 1s peak areas measured at different takeoff angles (θ). We show in Figure 8a C 1s spectra for a saturation TMA dose at room temperature. The spectra consist of peaks characteristic of aliphatic (∼285.0 eV) and carboxyl (∼289.0 eV) carbon, as expected for TMA. The C 1s peak intensity increases with decreasing θ because TMA is a surface-bound species. The associated O 1s spectra are shown in Figure 8b. These spectra consist of the more intense lattice oxygen peak at 530.3 eV, and a shoulder at ∼531.9 eV, containing contributions from the TMA anion and surface OH resulting from the acid proton binding to undercoordinated surface oxygens. Both TMA and OH oxygen exhibit a 1s binding energy that is shifted to higher value from the lattice O 1s position.19,31-33 This weaker peak is enhanced by going to a low takeoff angle θ. We fitted these spectra to a lattice O 1s line shape that is the same as that measured for clean TiO2, and second peak shifted to higher binding energy for the TMA and OH oxygen. We thus define a O 1s peak area ratio Rexp as Aad/(Aad + Alat) where Aad is the TMA-plus-OH peak area and Alat is the lattice oxygen peak area.30 We then correct all photodesorption spectra, such as those shown in Figure 9, for the initial TMA coverage by dividing by Rexp.

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Figure 8. C 1s (a) and O 1s (b) core-level spectra for a range of takeoff angles for a TMA-saturated TiO2(110) film surface.

Figure 9. CO2 photodesorption spectra for clean and N-doped TiO2 epitaxial films in the (110) (a), (001) (b), and (100) (c) orientations. The inset shows the natural logarithm of the CO2 photodesorption rate vs time. The slope equals the rate constant, k.

Panels a-c of Figure 9 show TMA-coverage-corrected CO2 photodesorption spectra for saturation doses on homoepitaxial TiO2-xNx(110), -(001), and -(100) rutile film surfaces with use of the full UV-visible emission spectrum from a 100 W Hg arc lamp for excitation. For all pure films, a sharp desorption spike occurs at t ) 0 when the film is first irradiated, accompanied by a rapid decay as the surface is depleted of TMA. The process is driven by photogenerated holes migrating from the bulk to the surface and reacting with TMA.12 The process is self-limiting when performed in the absence of O2, presumably because of negative charge accumulation on the surface. Such negative charge would be scavenged by O2, if O2 was present, and the photodecomposition process would proceed to completion.12,20,24 However, sufficient activity is observed in the absence of O2 and the photodesorption products are more easily detected if the chamber is free of O2. Any inhibiting effect in the kinetics due to the absence of gas-phase O2 is avoided by comparing initial rates of TMA photodecomposition (i.e., before significant charge imbalance has had a toll on the rate.) The TMA photodecomposition rate is significantly affected by N doping. As CN increases, the initial CO2 photodesorption yield decreases dramatically for rutile, despite the reduction in photoabsorption threshold energy for e-/h+ pair creation to 2.5 eV when N is introduced.18,19 The best explanation for this effect is that holes become trapped at isolated N 2p states, which fall near the top of the O 2p derived valence band. The solubility of N in phase-pure TiO2 precludes N 2p-derived band formation due to the large spatial separation of N dopants (∼1.0-1.5 nm). It is also possible that substitutional N dopants act as e-/h+

pair recombination centers which would also reduce the flux of holes reaching the surface. The initial TMA photodecomposition rates as a function of CN are shown in the insets of Figure 9 for each orientation. These were estimated assuming the functional form r(t) ) r0 exp(-kt), where r(t) and r0 are the photodesorption rates at irradiation time t and at t ) 0 (initial light exposure time), respectively, and k is essentially the overall rate constant for TMA photodecomposition during this initial period. For all three orientations, the data can be fit to lines for the first four seconds of irradiation, with the slopes corresponding to k. The rate constant decreases monotonically as a function of CN for rutile (110). However, the extent of decrease in k with CN is quite different for (001) and (100) than for (110), indicating that e-/h+ pair recombination probabilities are strongly dependent on orientation. The data in Figure 9 also reveal that once a hole is trapped at a N site, the resulting N2- species is not longlived, but rapidly regenerates to N3- on the time scale of our measurements. This trapping and regeneration process appears to be rapid compared to the time scale for hole diffusion to the surface because the rate of CO2 photodesorption does not recover to the level observed in the undoped film as the irradiation period progresses, as seen in Figure 9a. If the trapped hole states were long-lived, we would expect to see the CO2 photodesorption rate increase as more N dopant sites became and remained occupied with holes. In fact, the data in the inset of Figure 9 give no indication of a change in reaction probability that one might expect to see for a long-lived trap state. The regeneration process likely involves occupation of the partially filled 2p orbital of the N2- species by either an excited

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Ohsawa et al. 4. Conclusions The effect of N doping in rutile on the UV and visible-light photoactivity was examined by using epitaxial growth techniques to prepare well-defined crystalline film surfaces of rutile TiO2-xNx (0 e x e 0.02). These surfaces were exposed to saturation doses of trimethyl acetic acid in ultrahigh vacuum, and subsequently irradiated with broadband UV and visible light. Photoactivity was determined by monitoring CO2 photodesorption that results from hole scavenging and photodecomposition. This study reveals that holes either do not trap at substitutional N dopant sites in (100)-oriented films, or that hole trapping in N-doped TiO2(100) is more transient in nature compared to that in N-doped rutile(110) and -(001).

Figure 10. Summary of photodesorption rate constants as a function of x in TiO2-xNx rutile films in the three orientations.

conduction band electron or by an electron residing in the donor level (located at the Fermi level) in epitaxial TiO2-xNx.19 Although N doping strongly affects the photochemical decomposition rates for all three orientations, and in nonequivalent ways, the rate in undoped TiO2 films is essentially independent of orientation. Figure 10 shows the initial rate constants for coverage-normalized TMA decomposition plotted as a function of x. Here k is defined as ln(RCO2)t - ln(RCO2)t)0, where RCO2 is the CO2 desorption rate, and k was determined from the inset plots in Figure 9. The k values for x ) 0 are quite similar. However, k rapidly decreases with increasing x in both (110) and (001). The rate of change is much smaller for the (100) orientation. These differences are presumably due to differences in the probability of hole trapping at N sites for the different orientations. The paths from the point of electron-hole generation in the bulk to the surface vary with the surface orientation. Thus, for a given orientation, we expect that the trend in k with CN would vary according to the trend in hole mobility normal to the surface. While the undoped TiO2 surfaces are clearly more photoactive for TMA decomposition than are the TiO2-xNx surfaces, the most significant insight from Figure 10 is the strong crystallographic dependence of the effect of N doping on photochemical activity. We interpret these results to imply that holes generated in the O 2p-derived VB of rutile TiO2 by UV light are either more readily trapped at localized N 2p states as they diffuse along the 〈110〉 or 〈001〉 directions, or that the detrapping probability out of a N hole trap is greater along 〈100〉 trajectories. This result would suggest that O 2p VB holes diffusing along 〈100〉 directions in rutile are less likely to encounter and be trapped at N-dopant sites. The lack of change in k between x ) 0.01 and 0.02 for the (001) orientation is not understood, but the (001) surface is also not well understood. In our former study, we found that visible-light-induced decomposition of TMA adsorbed on Ti ridges that define the (1 × 4) reconstruction of N-doped anatase (001) readily occurs.30 It is important to note that all three crystal orientations of N-doped rutile do not show visible light photoactivity for TMA (Figure 9), indicating either that e-/h+ pairs generated at N-dopant sites in rutile are rapidly quenched, or that holes are irreversibly trapped at N-dopant sites in rutile, irrespective of orientation. However, as shown in Figure 9, there is no CO2 desorption from TMA-covered rutile TiO2-xNx (x ) 0.02) after visible-light exposure for any orientation.

Acknowledgment. . This work was supported by the U.S. Department of Energy, Office of Basic Energy Sciences, Division of Chemical Sciences, Geosciences, and Biosciences. Pacific Northwest National Laboratory is a multiprogram national laboratory operated for the U.S. Department of Energy by the Battelle Memorial Institute under contract DEAC06-76RLO1830. This work was performed in the Environmental Molecular Sciences Laboratory, a national scientific user facility sponsored by the Department of Energy’s Office of Biological and Environmental Research and located at the Pacific Northwest National Laboratory. References and Notes (1) Fujishima, A.; Honda, K. Nature 1972, 238, 37–38. (2) Hoffmann, M. R.; Martin, S. T.; Choi, W.; Bahnemann, D. W. Chem. ReV. 1995, 95, 69–96. (3) Carp, O.; Huisman, C. L.; Reller, A. Prog. Solid State Chem. 2004, 32, 33–177. (4) Wang, R.; Hashimoto, K.; Fujishima, A. Nature 1997, 388, 431– 432. (5) O’Regan, B.; Gra¨tzel, M. Nature 1991, 353, 737–740. (6) Gra¨tzel, M. Nature 2001, 414, 338–344. (7) Diebold, U. Surf. Sci. Rep. 2003, 48, 53–229. (8) Henrich, V. E.; Cox, P. A. The Surface Science of Metal Oxides; Cambridge University Press: Cambridge, England, 1994. (9) Bikondoa, O.; Pang, C. L.; Ithnin, R.; Muryn, C. A.; Onishi, H.; Thornton, G. Nat. Mater. 2006, 5, 189–192. (10) Wendt, S.; Sprunger, P. T.; Lira, E.; Madsen, G. K. H.; Li., Z.; Hansen, J. Ø.; Matthiesen, J.; Blekinge-Rasmussen, A.; Lægsgaard, E.; Hammer, B.; Besenbacher, F. Science 2008, 320, 1755–1759. (11) Henderson, M. A. Surf. Sci. Rep. 2002, 46, 1–308. (12) Henderson, M. A.; White., J. M.; Uetsuka, H.; Onishi, H. J. Am. Chem. Soc. 2003, 125, 14974–14975. (13) Asahi, R.; Morikawa, T.; Ohwaki, T.; Aoki, K.; Taga, Y. Science 2001, 293, 269–271. (14) Irie, H.; Watanabe, Y.; Hahimoto, K. J. Phys. Chem. B 2003, 107, 5483. (15) Nakamura, R.; Tanaka, T.; Nakato, Y. J. Phys. Chem. B 2004, 108, 10617. (16) Serpone, N. J. Phys. Chem. B 2006, 110, 24287. (17) Livraghi, S.; Paganini, M. C.; Giamello, E.; Selloni, A.; Valentin, C. D.; Pacchioni, G. J. Am. Chem. Soc. 2006, 128, 15666. (18) Cheung, S. H.; Nachimuthu, P.; Joly, A. G.; Engelhard, M. H.; Bowman, M. K.; Chambers, S. A. Surf. Sci. 2007, 601, 1754. (19) Chambers, S. A.; Cheung., S. H.; Shutthanandan, V.; Thevuthasan, S.; Bowman, M. K.; Joly, A. G. Chem. Phys. 2007, 339, 27. (20) Cheung, S. H.; Nachimuthu, P.; Engelhard, M. H.; Wang, C. M.; Chambers, S. A. Surf. Sci. 2008, 602, 133–141. (21) Uetsuka, H.; Onishi, H.; Henderson, M. A.; White, J. M. J. Phys. Chem. B 2004, 108, 10621–10624. (22) Uetsuka, H.; Pang, C.; Sasahara, A.; Onishi, H. Langmuir 2005, 21, 11802–11805. (23) Lyubinetsky, I.; Yu, Z. Q.; Henderson, M. A. J. Phys. Chem. C 2007, 111, 4342–4346. (24) White., J. M.; Henderson, M. A. J. Phys. Chem. B 2005, 109, 12417–12430. (25) Henderson, M. A.; White, J. M.; Uetsuka, H.; Onishi, H. J. Catal. 2006, 238, 153–164. (26) White, J. M.; Szanyi, J.; Henderson, M. A. J. Phys. Chem. B 2004, 108, 3592–3602.

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