Photochemical Properties, Composition, and Structure in Molecular

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Photochemical Properties, Composition, and Structure in Molecular Beam Epitaxy Grown Fe “Doped” and (Fe,N) Codoped Rutile TiO2(110) Andrew N. Mangham, Niranjan Govind, Mark E. Bowden, V. Shutthanandan, Alan G. Joly, Michael A. Henderson, and Scott A. Chambers* Fundamental and Computational Sciences Directorate, Pacific Northwest National Laboratory, Richland, Washington, 99352 ABSTRACT: We have investigated the surface photochemical properties of Fe “doped” and (Fe,N) codoped homoepitaxial rutile TiO2(110) films grown by plasma-assisted molecular beam epitaxy. Fe does not incorporate as an electronic dopant in the rutile lattice but rather segregates to the film surface. However, codeposition of Fe with N enhances the solubility of Fe, and DFT calculations suggest that codopant complex formation is the driving force behind the enhanced solubility. The codoped films, in which a few atomic percent of Ti (O) are replaced with Fe (N), exhibit significant disorder compared to undoped films grown under the same conditions, presumably due to dopant-induced strain. Codoping redshifts the rutile bandgap into the visible. However, the film surfaces are photochemically inert with respect to hole-mediated decomposition of adsorbed trimethyl acetate. The absence of photochemical activity may result from dopant-induced trap and/or recombination sites within the film. This study indicates that enhanced visible light absorptivity in TiO2 does not necessarily result in visible light initiated surface photochemistry.

1. INTRODUCTION Photocatalysis on oxide semiconductor surfaces has been an area of intensive research for many years. In particular, titanium dioxide (TiO2) has garnered a great deal of attention since it was first shown that rutile surfaces are effective in driving photoelectrolysis of water with UV light.1 Subsequently, the photochemical properties of both rutile and anatase polymorphs of TiO2 have been studied with an eye toward dye-sensitized solar cells2,3 and photodegradation of organic pollutants.4,5 These efforts have resulted in significant insight into the surface structure and photoreactivity of the various phases of TiO2, particularly the rutile(110) surface.6 9 However, the ∼3 eV bandgap that gives rutile much of its stability and photocatalytic power in the UV also hinders its utility in solar applications because of the nature of the solar spectrum, which is much richer in visible light than in UV. Considerable effort has gone into rendering TiO2 active in the visible region, either by sensitizing with the adsorption of dye molecules10 or by narrowing the bandgap through material doping.11 Different ways of anion doping TiO2, most often with nitrogen,12 have been explored, as have similar strategies for cation doping using a broad range of transition metals.13 18 Both approaches have been shown to shift the bandgap into the visible and improve photocatalytic activity in some cases.19 23 However, the majority of these reports have been studies focused on nanoparticles, which are notoriously hard to characterize because of their large surface-to-volume ratio, the multiplicity of exposed surface orientations, and the complex interactions within nanoparticle assemblies. As a result, these studies tend to focus on the effects of doping on the nanoparticles without providing a thorough materials analysis that might help one to understand r 2011 American Chemical Society

any changes in the properties. There are relatively few fundamental studies focused on the nanoscale structure of doped TiO2 and the underlying principles driving bandgap reduction and enhancement of photocatalytic activity. The use of plasma-assisted molecular beam epitaxy (MBE) has made it possible to grow epitaxial single-crystal thin films of N doped TiO2 in both the rutile and anatase polymorphs.24 26 Recent studies have yielded considerable insight into the relationship between the structural environment of the dopant and the associated effect on photocatalytic activity.27,28 It has been shown that the solubility limit of N is ∼1 atom % within the anion sublattice for rutile and anatase.28 In both rutile and anatase polymorphs, this concentration was sufficient to shift the absorption edge into the visible region. However, in the case of rutile, N doping actually reduces the photocatalytic activity, based on the rate of hole-mediated photodecomposition of chemisorbed trimethyl acetate, presumably due to hole trapping at N sites.28 The present study focuses on cation doping of rutile TiO2 with Fe, as well as (Fe,N) codoping. Fe doping of TiO2 for the purposes of bandgap manipulation or the improvement of photocatalytic properties has been of interest for quite some time. Several theoretical studies29 31 have predicted that partially filled Fe 3d states should cause a significant redshift in the bandgap of TiO2, and indeed there have been many recent reports of Fe “doped” TiO2 nanoparticles with bandgaps in the Received: April 1, 2011 Revised: June 13, 2011 Published: July 05, 2011 15416

dx.doi.org/10.1021/jp203061n | J. Phys. Chem. C 2011, 115, 15416–15424

The Journal of Physical Chemistry C visible, as well as improved photocatalytic activity.17,32 36 Dopant concentrations in these systems have been claimed to range from as little as 0.001 mol %16 to as much as 28 mol %.32 In these studies, X-ray diffraction (XRD) was typically used to determine whether or not Fe substituted for Ti in the lattice. However, as acknowledged in some of these papers, XRD may not yield definitive evidence for dopant substitution due to its detection limit and the possibility that the Fe dopants exist in an amorphous phase.13,16,19,36 By growing homoepitaxial single-crystal thin films of Fe “doped” rutile TiO2, we are able to employ a range of materials characterization techniques and explore their structural properties at a deeper level than has been done previously. Additionally, we have used plasma-assisted MBE to grow (Fe,N) codoped films, and have characterized the films using several methods. By comparing the properties of Fe and (Fe,N) codoped TiO2, we can assess the benefit of codoping for enhancing the dopant solubilities, as well as the photochemical properties. There have been several recent investigations of anion cation codoping in TiO2.37 40 Both Naik et al.40 and Yang et. al41 examined the effects of (Fe,N) codoping on anatase TiO2. Both reported enhanced redshifting of the bandgap but had differing results with respect to enhancement of photocatalytic activity. In a combined experimental and theoretical study, Zhu et al.42 presented evidence for mutually enhanced solubility of Cr and N via codoping. The enhanced dopant concentrations were attributed to the formation of dopant pairs stabilized by Coulombic interaction. According to their calculations, these dopant pairs exhibited a lower kinetic barrier to lattice substitution than either dopant alone, leading to enhanced solubilities. Their experimental evidence was based on studies of anatase nanoparticles, and their results suggested not only an increase in concentration relative to the single-dopant approach but also an enhanced ability to retain N during annealing in an oxygen-rich environment.

2. EXPERIMENTAL METHODS All films were grown by plasma-assisted molecular beam epitaxy in a custom chamber.43 As-received TiO2(110) substrates were annealed in air for 8 h at 1000 °C, resulting in atomically flat surfaces with terrace widths of ∼0.3 μm, as judged by tapping-mode atomic force microscopy (AFM). The annealed substrates were sonicated in acetone and isopropanol and then ozone cleaned on the bench. Following insertion into the MBE chamber, the substrates were exposed to activated oxygen from an electron cyclotron resonance (ECR) microwave plasma source for 45 min at an oxygen pressure of 2 10 5 Torr. The resulting surfaces were clean, as judged by XPS, and showed (1 1) RHEED patterns characteristic of a well-ordered, atomically flat surface. High-temperature effusion cells were used for Ti and Fe evaporation, and the ECR plasma source was used for activated N and O. The fluxes of Ti and Fe were calibrated against cell temperature prior to growth with the plasma running using a quartz crystal oscillator (QCO) set to measure TiO2 and Fe2O3 deposition rates. When growing (Fe,N) codoped films or N doped films, a 3:1 mix of N2 and O2 was bled into the ECR plasma cavity. This ratio was needed to maximize N incorporation which was limited to a few atom % of the anion sublattice even under N-rich conditions. The fluxes of activated O and N were difficult to estimate.24 Recent measurements of the flux of

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atomic oxygen from the ECR plasma source under conditions similar to those used here yielded a value of ∼3 1013 atoms O/ cm2 s.44 The substrate temperature during growth was ∼550 °C, as judged by a two-color infrared pyrometer. This temperature was chosen to minimize N loss from the films, as discussed previously.24 26 A buffer layer of 10 15 monolayers of undoped or N doped TiO2 was first deposited, using RHEED oscillations to determine the growth rate. The Fe shutter was then opened, and Fe “doped” or (Fe,N) codoped films were grown to the desired thickness. The Fe concentration ranged from 0.5 to 2 atom % within the cation sublattice. Because these doped TiO2 films were grown on TiO2(110), neither X-ray reflectivity (XRR) nor Rutherford backscattering spectrometry (RBS) could be used to accurately determine sample thickness. Thus, the growth time and the deposition rate obtained from the RHEED oscillations were used to calculate final thickness. Films were typically grown to a thickness of ∼500 Å. Following growth, the films were cooled in the plasma and then transferred under ultrahigh vacuum (UHV) conditions to an appended X-ray photoelectron spectroscopy (XPS) chamber consisting of a GammaData/Scienta SES 200 analyzer and a monochromatic Al KR X-ray source for in situ measurements. Data were taken at normal emission (90°) and at 12° off the surface plane, to vary the surface sensitivity. Static charging occurred due to the insulating nature of these films, resulting in the need to use a low-energy electron flood gun to compensate the positive surface change. As a result, the as-measured spectra exhibited a range of binding energies, depending on the flood gun settings. We used the Ti 2p3/2 peak as an internal standard to correct for charging. This peak typically exhibits a binding energy of 458.9 eV in n-type TiO2(110) rutile. Thus, all spectra for the present films were shifted such that the Ti 2p3/2 peak fell at this energy. Samples were then transferred under UHV conditions to an appended photodesorption chamber. This chamber included a molecular dosing apparatus for trimethyl acetic acid (TMAA, Sigma Aldrich, research grade), a quadrupole mass spectrometer, and a fiber optic feedthrough connected to an external Hg arc lamp. TMAA undergoes acid dissociation on rutile (110) to form densely packed adlayers of trimethyl acetate (TMA) in which each anion bridges two Ti4+ cations along the surface [001] direction.45 49 When illuminated with broadband UV and visible light, electron hole pairs are generated near the surface. The hole destabilizes the TMA anion, resulting in decomposition to form CO2 and an organic fragment which interacts with the surface prior to desorption.50 Detection of the direct product molecule CO2 via photodesorption spectroscopy at mass 44 yields a measurement of surface photoactivity. Following these in situ measurements, Rutherford backscattering spectrometry (RBS), nuclear reaction analysis (NRA), high-resolution X-ray diffraction (XRD), and UV visible spectroscopy in transmission mode were conducted ex situ.

3. RESULTS AND DISCUSSION Figure 1 shows RHEED patterns for undoped (a), Fe “doped” (b), N doped (c), and (Fe,N) codoped (d) films recorded immediately after cooling in the plasma. Both Fe “doped” and (Fe,N) codoped films showed strong RHEED diffraction rods consistent with a reasonably smooth epitaxial film surface, though comparison of the films to the homoepitaxial undoped 15417

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The Journal of Physical Chemistry C

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Figure 1. Typical RHEED patterns for undoped (a), Fe “doped” (b), N doped (c), and (Fe,N) codoped (d) TiO2(110) rutile films. Figure 3. Ti 2p spectra for undoped (a), N doped (b), Fe “doped” (c), and (Fe,N) codoped (d) rutile TiO2(110).

Figure 2. (a) Normal emission Fe 2p spectra for (Fe,N) codoped rutile, Fe “doped” rutile, Fe2O3, and Fe3O4. (b) N 1s spectra for (Fe,N) codoped rutile with detector at takeoff angles of 90° and 12° with respect to the sample surface.

film shows some degradation in surface quality. Compared with the RHEED patterns of Fe “doped” films, (Fe,N) codoped film patterns exhibited higher background intensities and

some intensity modulation along the rods, indicating somewhat rougher and more disordered surfaces. The surface crystallinity and flatness could be improved by increasing the substrate temperature to ∼650 °C. However, the extent of N incorporation was significantly reduced at this temperature. The N doped films grown showed slightly rougher surfaces than either of the films containing Fe. Figure 2a shows Fe 2p XPS spectra for Fe “doped” and (Fe,N) codoped films along with spectra taken from epitaxial R-Fe2O3/ R-Al2O3(0001) and Fe3O4/MgO(001) films. All spectra have been normalized for the sake of clarity. The line shape and binding energies in the doped films match those of Fe2O3, leading to a Fe3+ charge state assignment. There is no evidence of Fe2+ in any of the films. Figure 2b shows the N 1s spectrum of a typical (Fe,N) codoped film. Here again the spectra have been normalized to account for differences in signal intensity. The peak at 396.0 eV is assigned to N3 which substitutes for oxygen within the anion sublattice of TiO2 (NO).24,51 On the basis of its binding energy, the peak at 400.2 eV is most likely due to N2. Angle-resolved measurements show that this species is present only on or near the surface of the film, whereas NO is distributed uniformly throughout the nominal XPS probe depth (50 Å).26 However, N2 is not expected to sorb on this (or any) surface at room temperature due to its inertness. This peak disappears after light annealing, whereas the NO peak remains. We thus tentatively conclude that this peak is due to N2 in subsurface interstitial sties. The oxidation states for the Fe and the substitutional N are favorable for dopant complex formation via Coulomb interaction.42 Moreover, the simultaneous formation of substitutional Fe3+ and N3 insures mutual charge compensation. The N dopant concentration (NO) was estimated using the formula 100(AN1s/σN1s)/(AN1s/σN1s + AO1s/σO1s)24 where {A} are peak areas and {σ} are sensitivity factors.52 N doping affects the Ti 2p peak shape by the addition of a weak low-binding-energy shoulder to the Ti 2p spectrum.24 Figure 3 shows Ti 2p3/2 spectra for undoped (a), N doped (b), Fe “doped” (c), and (Fe,N) codoped (d) films. In all cases, the spectra have undergone a Shirley background removal and were fit using Voigt functions. Fit peaks are included in the graph and offset from the baseline by 10% for easier identification. Doping 15418



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Figure 4. Residual traces from Ti 2p peak fitting for both Fe “doped” and (Fe,N) codoped TiO2. For the codoped film, traces are shown for the cases in which the spectrum is fit with or without including a peak for Ti O5 δN1+δ.

Table 1. Dopant Concentrations from Fluxes (Column 1) and from XPS at Normal (90°) and Shallow-Angle (12°) Emission (Columns 2 5) for Various Doped Films nominal Fe Fe cation Fe cation N anion N anion sample

cation %

% 90°

% 12°

% 90°

% 12°

Fe “doped” rutile TiO2 (Fe,N) codoped

2.0% 2.0%

2.4% 1.8%

9.4% 2.4%

NA 1.8%

NA 0.9%

NA

NA

NA

3.1%

2.0%

rutile TiO2 N doped rutile TiO2

with N results in a transfer of intensity from the Ti 2p3/2 primary lattice peak to a weak second feature shifted ∼1.0 eV to lower binding energy. This feature, which scales in intensity with the N concentration as determined by the method described, has been assigned as Ti bound to ∼5 O ligands and ∼1 NO ligand, as in previous work on N doped rutile and anatase.24 26 This peak is not present in spectra for the undoped and Fe “doped” films. To establish this point, we have included in Figure 4 residual traces that remain from the Ti 2p spectra after fitting. We show three spectra, one for the Fe “doped” film and two for (Fe,N) codoped films. The (Fe,N) codoped traces show the results of fitting both with and without including the low-binding-energy shoulder. When that feature is not included in the fit, we clearly see a peak in the residuals at ∼1 eV lower binding energy relative to the lattice Ti peak. In contrast, we see a relatively flat residual trace when this feature is included in the fit. The residual trace for the Fe “doped” film results from a fit that did not include a lowbinding-energy shoulder, and we do not see a peak arise 1 eV downfield from the lattice peak. These results establish that the peak ∼1 eV to lower binding energy is due to Ti bound to a combination of N and O. Table 1 shows Fe and N concentrations for a set of representative films measured at takeoff angles of 90° and 12°, as well as the target Fe concentrations based on the QCO calibration of the Fe and Ti fluxes (first column). The method for estimating the N concentration is described above; the Fe concentration was estimated in a similar way by the formula 100(AFe2p/σFe2p)/

Figure 5. Fe 2p XPS spectra for Fe “doped” rutile (110) before (a) and after (b) growth of a 10 nm pure TiO2(110) epitaxial cap layer, as well as for (Fe,N) codoped TiO2(110) before (c) and after (d) growth of a 10 nm pure TiO2(110) epitaxial cap layer.

(AFe2p/σFe2p + ATi2p/σTi2p). The XPS probe depths in TiO2 are ∼5 nm and ∼1 nm at 90° and 12°, respectively. As seen in the table, the Fe concentration is greatly increased within the top 1 nm of the Fe “doped” film relative to that within the top 5 nm. This result strongly suggests that Fe segregates to the surface of the Fe “doped” TiO2(110) films to a significant extent. In contrast, the (Fe,N) codoped films show far less Fe surface segregation, and the Fe concentration is closer to the target value based on the metal fluxes throughout the film. Some degree of Fe surface segregation was observed in all ten Fe “doped” films grown for this investigation. A much lower extent of surface segregation was observed in the eight (Fe,N) codoped films grown. This difference is indicative of enhanced Fe stability at cation sites when N is doped into anion sites, perhaps due to dopant-pair formation, and mutually enhanced solid solubility. XPS was used to verify that Fe is significantly less mobile in rutile when Fe and N are simultaneously present as dopants compared to when Fe alone is present. An Fe “doped” film and an (Fe,N) codoped film were grown with the same Fe doping level (∼2%), growth rate, and substrate temperature. The Fe “doped” film showed strong surface segregation of Fe (4.8 cation atom % at 90° vs 10.2 cation atom % at 12°), while the (Fe,N) codoped film showed significantly less Fe segregation and Fe concentrations much closer to the target doping level (2.0 cation atom % at 90° and 3.1 cation atom % at 12°). 10 nm of undoped epitaxial TiO2 was then grown on top of both films to see if Fe was sufficiently loosely bound in either film to diffuse to the surface. Figure 5 shows that Fe diffuses to the top surface of the capping layer for the Fe-only film (Figure 4 b) but not from the codoped film (Figure 5d). The Fe 2p peak intensity is nearly the same after depositing the 10 nm rutile cap (Figure 4b) as it is before (Figure 5a) for the Fe “doped” film. In contrast, the Fe 2p signal is completely attenuated after depositing the cap layer (Figure 5d) on the (Fe,N) codoped film (Figure 5c). High-resolution X-ray diffraction scans for (Fe,N) codoped and Fe “doped” films are shown in Figure 6. Both films show sharp (220) Bragg peaks at 56.62° associated with the underlying rutile substrate. However, the (Fe,N) codoped film gives rise to a distinct second (220) peak at a lower angle (Figure 6b). The same observation was made in previous experiments on N doped rutile films, and the second peak was assigned to the Bragg 15419


The Journal of Physical Chemistry C reflection in the doped epitaxial layer. It was shown that substitution of the larger N3 anion (1.46 Å) for the O2 anion (1.36 Å)53 at the ∼1 anion atom % doping level caused an expansion in the a and c lattice parameters of 0.5% and 0.8%, respectively.26 In the (Fe,N) codoped films, the a lattice parameter is 4.653 Å, which is 1.3% larger than that of pure rutile (4.594 Å). This expansion is consistent with substitution of the larger N3 anion at O sites and substitution of the larger Fe3+ cation (0.785 Å) for Ti4+ (0.745 Å)53 at cation sites. There are no interference fringes surrounding the Bragg peaks for the codoped film. This result could indicate poor crystallographic ordering in the film, a rough film surface, or a strain gradient in the film. Measurements of the in-plane lattice parameter (c, not shown) for this film show a strong peak for the substrate but only very weak peaks for the epitaxial layer, indicating considerable inplane disorder. Examination of the Fe “doped” films shows only the bulk (220) peak with some peak asymmetry consistent with mosaic spread, but with no distinct film peak. Taken together, the XRD measurements suggest that when TiO2 is doped with both Fe and N, the result is successful substitution of both dopants, along with significant film disorder. When the film is doped with only Fe, the XRD is essentially the same as that expected for an undoped film, suggesting that either Fe doping does not alter the

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structure of TiO2 or that Fe does not substitute for Ti at all but rather segregates to the surface, presumably in the form of Fe2O3. Similar conclusions have been drawn previously.54 The absence of Fe2O3 Bragg peaks may result simply because the amount formed is below the detection limit of XRD.13,16,19,36 Film crystallinity and Fe site occupancy were assessed by means of RBS.55 RBS spectra taken in the channeling and random geometries for an Fe “doped” and an (Fe,N) codoped film are presented in Figure 7. The indications of disorder shown in the XRD and RHEED analysis of the (Fe,N) codoped films are confirmed by the higher minimum yield of the Ti signal in the channeling direction for the codoped film (40%, Figure 7b) compared to the Fe “doped” film (12%, Figure 7a). Rocking curves for Fe and Ti backscattering in these two films are shown in Figure 8 a and b. In the case of the Fe “doped” film there is a strong minimum for the Ti signal, consistent with excellent overall crystallinity, but only a weak minimum for the Fe signal indicating little, if any, site correlation between Fe and Ti. Complementary information about N substitution for O is gained from NRA.24 Here we used 0.95 MeV deuterons and the 14N(d,R)12C and the 16O(d,p)17O nuclear reactions to measure rocking curves, shown for the codoped film in Figure 8c. The rocking curve measurements obtained from RBS and NRA for the (Fe,N) codoped film show different behavior from the Fe “doped” film. Fe and N both show minima along the channeling direction which exhibit the same depth. This result indicates that Fe substitutes for Ti and N substitutes for O in this film. The rather shallow minima (∼0.6) reveal significant disorder, consistent with the high minimum yield.

Figure 6. (220) Bragg reflections for Fe “doped” rutile (110) (a) and (Fe,N) codoped (b) rutile (110) epitaxial films.

Figure 7. RBS spectra for: Fe “doped” (a) and (Fe,N) codoped (b) rutile TiO2(110) films.

Figure 9. Valence band XPS spectra for Fe “doped”, (Fe,N) codoped, undoped, and N doped TiO2(110) epitaxial films.

Figure 8. RBS (a and b) and NRA (c) rocking curves about [110] for Fe “doped” (a) and (Fe,N) codoped (b and c) rutile TiO2(110) epitaxial films. 15420



Figure 10. Absorption spectra for Fe “doped”, (Fe,N) codoped, N doped, and undoped rutile TiO2(110).

To determine electronic structure, we have utilized valence band (VB) XPS and optical absorption. The VB XPS results are shown in Figure 9. As can be seen, all four types of films show essentially identical VB spectra. The lack of any features in the VB spectra from N is most likely a result of low sensitivity and low concentration.26 The lack of features from Fe is also due to low concentration as well as the probability that Fe does not contribute much to the VB (see density of states discussion below). Optical absorption spectra for the same film set are shown in Figure 10. These spectra have been normalized by subtracting a horizontal background to bring the low-energy region to zero and then by setting the maximum absorption for each film to unity. The spectra of pure and Fe “doped” rutile are nearly identical and show that the presence of Fe does not affect the band gap. This result is not surprising in light of the fact that Fe does not appear to substitute for Ti in this material, based on the XPS results discussed above. In the case of the N doped and (Fe,N) codoped samples, the absorption edge is red-shifted to ∼2.4 eV. In the case of N doped rutile, this feature has been shown to be a result of excitations from N 2pderived states above the VBM to the CB,26 and the same appears to be true here. The absorbance feature seen in the codoped film is similar to that seen in the N doped film, though lower in intensity. CO2 photodesorption spectra from the photodecomposition of adsorbed TMA with broad band visible/UV irradiation are shown in Figure 11. Each trace was normalized by dividing by the TMA coverage derived from O 1s spectra, as described in an earlier work.28,56 Only the Fe “doped” and undoped films revealed any photochemical response, and only when illuminated with UV light (data for visible light irradiation not shown). The N doped film is inactive in the UV and visible, presumably because of hole trapping on N sites, as observed previously.28 The (Fe,N) codoped film is also photoinactive for the hole-mediated decomposition of TMA. While this material could also be exhibiting hole trapping on N sites, the disorder in the lattice must also be considered as a possible reason for photoinactivity. The disorder resulting from successful codoping undoubtedly leads to a multitude of point defects in the lattice that may increase hole trapping and/or electron hole recombination.

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Figure 11. Photodesorption data for Fe “doped”, (Fe,N) codoped, undoped, and N doped TiO2 under UV + visible irradiation.

4. THEORETICAL MODELING Our investigation reveals that rutile exhibits substantially different properties when codoped with N and Fe than when doped with Fe alone. In concept, Fe (N) should incorporate in the cation (anion) sublattice within rutile. However, substitutional Fe is only stabilized in rutile with substitutional N codoping. The lattice expands in response to the presence of both kinds of dopants, and the strain associated with codoping results in considerable disordering. A critical question not easily answered by experiment is whether the observed stabilization is a result of dopant complex formation. We have addressed this issue using theory. To this end, ground-state periodic density functional theory (DFT) calculations were performed on bulk rutile codoped with Fe and N. All calculations were performed with the Quantum-Espresso package57 using a plane-wave cutoff of 35 Ry and a density cutoff of 300 Ry. Vanderbilt ultrasoft pseudopotentials57 60 were used. Semicore pseudopotentials were chosen for Ti (3s23p63d24s2) and Fe (3s23p63d64s2), and valence only pseudopotentials were chosen for O (2s22p4) and N (2s22p3), respectively. The bulk lattice parameters for the tetragonal rutile TiO2 crystal lattice (space group 4/m 2/m 2/m) were taken from crystallographic data in the literature: a = b = 4.5937 Å, c = 2.9587 Å, and R = β = γ = 90°. A full optimization of lattice parameters and atomic positions for bulk model for rutile TiO2 was performed using the Perdew, Burke, Ernzerhof (PBE) generalized gradient approximation (GGA) exchange-correlation functional61 augmented with the U correction62 of 4.2 eV to account for the correlated nature of the Ti 3d states. This choice of U yielded a band gap of 3.2 eV which is in agreement with the experimental value in bulk TiO2 rutile. The Brillouin zone was sampled with a 5 5 8 Monkhorst Pack k-point grid.63 The optimized lattice parameters at this level of theory are a = 4.639 Å, b = 4.639 Å, c = 2.974 Å, R = 90°, β = 90°, and γ = 90°, respectively. For the rutile TiO2 supercell models, the optimized rutile TiO2 unit cell lattice was expanded into a 2 2 3 simulation cell consisting of 24 Ti and 48 O atoms. Fe and N atoms were placed toward the center of the TiO2 supercell. Five configurations were constructed in which we substituted one Ti with an Fe and one O with a N and varied the distance between dopants. Full optimizations of the lattice parameters and atomic positions were performed. These larger supercell calculations were performed with a 2 2 2 Monkhorst Pack k-point grid. The 15421



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Figure 12. Unrelaxed (a) and relaxed (b) geometries of the codoped TiO2 rutile showing the bond distances in the vicinity of the (Fe,N) dopant pair.

Table 2. Unrelaxed and Relaxed Bond Lengths (in Å) in the Vicinity of the (Fe,N) Dopant Pair for the Lowest-Energy Substitutional Configuration bond N Fe N Ti

unrelaxed, Å

relaxed, Å

1.962 1.962

1.663 2.266

N Ti

2.000

2.302

Fe O (axial)

2.000

2.059 2.059

Fe O (axial)

2.000

Fe N (equi)

1.962

1.663

Fe O (equi)

1.962

1.991

Fe O (equi)

1.962

2.069

Fe O (equi)

1.962

1.956

energy was lowest when the cation and anion were nearest neighbors, consistent with the enhanced Coulomb attraction. The configuration where the next nearest O atom was substituted with N was marginally higher in energy (∼0.15 eV). All the other substituted configurations were considerably higher in energy. The optimized lattice parameters (a and c) were found to be ∼1% larger than those computed for the undoped cell in the lowest-energy configuration, consistent with the XRD data. Figure 12 shows the optimized coordinates along with the bond lengths for the lowest-energy substitutional configuration, and the numbers are listed in Table 2. Compared with the unrelaxed geometry, there is significant bond deformation within the lattice in the vicinity of the dopant pair, as well as in the surrounding lattice. This supports the RBS and NRA results which showed that codoping Fe and N results in substantial structural disorder. Figure 13 shows the total and partial densities of states (DOS) for the pure and codoped systems, as well as the partial densities of states for the Fe 3d and N 2p. The upper and lower part of each plot represents the spin up and spin down DOS, respectively. The ordinate of each DOS plot is scaled for clarity. The character of the upper part of the valence band is dominated by the O 2p states, while the lower part of the conduction band is dominated by Ti 3d states. Previous calculations in N doped TiO2 rutile exhibited some hybridization between N and Ti at the top of the valence band.64 These hybridized states are absent in the codoped system. However, a hybridized N Fe gap state is seen at ∼2.0 eV above the valence band. The appearance of this gap state suggests a band gap lowering of ∼0.9 eV.

Figure 13. Total and partial density of states (DOS) for the pure and (Fe,N) codoped TiO2 rutile system. (a) Total DOS for pure TiO2; (b) total DOS for codoped TiO2; (c) partial Fe 3d DOS; (d) partial N 2p DOS.

A L€owdin charge analysis yielded a charge of approximately +3.5 on the Fe consistent with the experimental result. A total magnetization of 3 μB is observed, which suggests an Fe3+ cation in an intermediate crystal field splitting environment (in between the low (5 μB) and high (1 μB) splitting limits). This result is not surprising considering the deformation seen in the vicinity of the dopant pair as well as the surrounding lattice (Figure 12).

5. CONCLUSIONS Fe “doped” and (Fe,N) codoped rutile TiO2(110) films were grown by plasma-assisted molecular beam epitaxy, and their compositional, structural, and photochemical properties were 15422


The Journal of Physical Chemistry C investigated. Fe alone does not substitute for Ti within the cation sublattice. Rather, Fe shows a strong tendency to segregate to the surface of the film where it most likely forms disordered Fe2O3. These results differ sharply from those in the majority of literature based on nanoparticle studies. On the basis of the significant difference in properties, it may be that solution phase growth of doped nanoparticles results in fundamentally different material than crystalline films grown in vacuum. Alternatively, the apparent differences may result from inadequate characterization of the nanoparticles. We note that materials characterization adequate to determine the compositional, structural, and electronic properties of the dopants was not done in the majority of these nanoparticle studies. One exception that stands out is a report from Asong et al.65 that examined Fe doping in anatase nanoparticles and used EPR to show clear evidence of Fe segregation to the surface of the particles. More commonly, dopant concentrations were determined based on quantities of reagents mixed, and dopant substitution for host ions was tacitly assumed rather than experimentally determined. In contrast, codeposition of Fe and N along with Ti and O results in successful doping of both Fe and N at the appropriate lattice sites, albeit with the introduction of considerable strain. Thus, we conclude that N substitution for O stabilizes Fe substitution for Ti. These experimental observations are supported by DFT calculations which show that Fe N nearest-neighbor dopant complex formation lowers the electronic energy relative to other configurations with the dopants further apart but also introduces significant strain. (Fe,N) codoped films exhibit a red-shifted absorption band most likely derived from excitation from N 2p derived states at the top of the VB into the CB. Fe “doped” films show comparable photochemical activity to undoped films because Fe does not actually substitute for Ti. In contrast, (Fe,N) codoped films are photochemically inactive. Although disappointing from a photochemical point of view, the present study highlights the importance of conducting fundamental investigations to understand the properties of cation anion codoping in metal oxides.

’ AUTHOR INFORMATION Corresponding Author

*Phone: +1 509 371 6517. Fax: +1 509 371 6066. E-mail: [email protected].

’ ACKNOWLEDGMENT 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 Pacific Northwest National Laboratory. All calculations reported were performed on the Chinook supercomputer at EMSL. This work was supported by the US Department of Energy, Office of Science, Division of Chemical Sciences. N.G. also acknowledges support from the EMSL Intramural Program and thanks Dr. Amity Andersen for many useful discussions. ’ REFERENCES (1) Fujishima, A.; Honda, K. Nature 1972, 238, 37. (2) Oregan, B.; Gratzel, M. Nature 1991, 353, 737. (3) Gratzel, M. Nature 2001, 414, 338. (4) Carp, O.; Huisman, C. L.; Reller, A. Prog. Solid State Chem. 2004, 32, 33.

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Photochemical Properties, Composition, and Structure in Molecular

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