Properties of structurally excellent N-doped TiO2 rutile

Chemical Physics 339 (2007) 27–35 www.elsevier.com/locate/chemphys

Properties of structurally excellent N-doped TiO2 rutile a,*

S.A. Chambers

b

, S.H. Cheung a, V. Shutthanandan a, S. Thevuthasan a, M.K. Bowman b, A.G. Joly a

a Pacific Northwest National Laboratory, Richland, WA, USA Department of Chemistry, University of Alabama, Tuscaloosa, AL, USA

Received 5 March 2007; accepted 28 April 2007 Available online 22 May 2007

Abstract We have used plasma-assisted molecular beam epitaxy to synthesize structurally near-perfect crystalline films of TiO2xNx rutile for the first time. These materials allow the properties of TiO2xNx to be elucidated without the interfering effects of bulk substoichiometric defects that have characterized previous investigations. In the absence of such defects, the extent of N incorporation in the lattice is limited to 2 ± 1 at.% of the anions. Substitutional N (NO) exhibits a 3 formal charge due to charge transfer from shallow-donor interstitial Ti(III), which forms during epitaxial growth. Hybridization between NO and adjacent lattice Ti ions occurs, resulting in new states at the top of the rutile valence band and an apparent band gap reduction of 0.6 eV.  2007 Elsevier B.V. All rights reserved. Keywords: Molecular beam epitaxy

1. Introduction Bandgap reduction in TiO2 with the goal of enhancing visible light absorption is of significant current interest for important processes such as photocatalytic water splitting. Anion doping, particularly with N, has been a popular approach [1–13]. A number of experimental efforts have sought to determine the properties of intrinsic N-doped TiO2 using materials prepared by either film growth or ion implantation. A common feature in virtually all previous studies is the use of materials with high defect densities. As a result, what is not clear is the extent to which defects modify the interaction of the N dopant with the host lattice. In order to determine the properties of intrinsic Ndoped TiO2, it is necessary to prepare the material in such a way that defect densities are minimized or eliminated. Ideally, one would grow the material one atomic layer at a time such that the atomic fluxes are precisely monitored

*

Corresponding author. Tel.: +1 509 376 1766; fax: +1 509 376 1044. E-mail address: [email protected] (S.A. Chambers).

0301-0104/$ - see front matter  2007 Elsevier B.V. All rights reserved. doi:10.1016/j.chemphys.2007.04.024

so as to generate structurally perfect, crystalline layers with controlled stoichiometries. Using this approach, we have grown N-doped TiO2 homoepitaxially on TiO2(1 1 0) rutile using plasma-assisted molecular beam epitaxy (PAMBE) [14]. The use of rutile substrates results in rutile film growth by interface stabilization [15]. In contrast, the use of LaAlO3(0 0 1) substrates with the same fluxes and substrate temperatures results in the growth of epitaxial N-doped TiO2(0 0 1) anatase because of an accidental lattice match in the (0 0 1) orientation [16]. In the present work, we find that considerably less N can be incorporated into the rutile lattice when defects are absent than when defects are present. In the absence of oxygen vacancy (Ov) defects, the extent of N incorporation is limited to 2 ± 1 at.% within the anion sublattice because of the much stronger thermodynamic drive to form Ti–O bonds compared to Ti–N bonds [17]. We have used structurally near-perfect films to determine the properties of N-doped TiO2 rutile. N substitutes for O in the lattice (NO), and is compensated by itinerant electrons originating from interstitial Ti(III) (Tii) that forms during rutile homoepitaxy, resulting in a formal charge of 3 for NO. Even at 2 at.%, N incorporation

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results in substantial oscillator strength in the bandgap region 2.5–3.0 eV, suggesting bandgap reduction. 2. Experimental The PAMBE system used to carry out this work utilizes high-temperature effusion cells and electron beam evaporators for metal evaporation, an electron cyclotron resonance (ECR) plasma source for activated N and/or O, and an ozone source for activated O. N-doped TiO2 film growth was carried out in two different ways. One method involved mixing O2 and N2 gases in varying proportions in the ECR plasma unit and then delivering a mix of activated N and O to the substrate, along with a Ti atom beam. The second method consisted of delivering activated N to the substrate from the ECR plasma and activated O using the ozone source, along with Ti flux. The second approach allowed us to investigate the possibility of N/O recombination within the ECR plasma unit. A Ti high-temperature effusion cell was used for lower deposition rates whereas a Ti electron beam evaporator was used for higher growth rates. The Ti deposition rate from the effusion cell was calibrated in ultrahigh vacuum (UHV) prior to growth using a quartz crystal oscillator (QCO) positioned at the sample position. The TiO2 growth rate was then monitored during growth using reflection high-energy electron diffraction (RHEED) intensity oscillations. The electron beam evaporator Ti flux was monitored by atomic absorption (AA) spectroscopy during growth. The AA was calibrated with the QCO in the same way described above. The thickness uncertainty incurred with the use of the AA has been found to be 10% [14]. The total film thickness was measured following growth ex situ using X-ray diffraction (XRD), and this value agreed with that determined from the RHEED oscillations to within a few percent. TiO2(1 1 0) substrates were prepared by etching overnight in buffered HF and then annealing in an oxygen tube furnace at 900 C for 1 h. Once in the PAMBE chamber, carbon contamination was removed by exposing the substrates to a oxygen plasma at 2 · 105 Torr for 1 h. The resulting surfaces were clean, as judged by XPS, and atomically flat with terrace widths of 0.25 lm, as judged by AFM. Prior to growth using method one, the partial pressures of O2 and N2 were set by admitting the gases into the ECR plasma source tube via leak valves and using the MBE chamber ionization gauge (IG) and residual gas analyzer (RGA) for monitoring. Growths were typically performed with approximately equal partial pressures of O2 and N2. Following the ECR oxygen plasma cleaning step described above, the O2 pressure was reduced to half the desired total pressure for growth. An equal partial pressure of N2 was then admitted into the plasma tube, the plasma remaining lit throughout the process. Neither the IG nor the RGA were positioned directly in the beam emanating from the ECR plasma source. Therefore, the pressure readings were

indirect, and the absolute fluxes of O and N are not known. The gas fluxes were roughly estimated using the rule of thumb that the arrival rate at the substrate is 1 monolayer (ML) per second at P = 1 · 106 Torr. Although the pressure is higher directly in the beam than at the IG, resulting in an underestimation of the activated species flux, the ECR plasma source does not completely dissociate the O2 and N2. Based on the estimated dissociation efficiencies of 30% for O2 and 10–20% for N2 supplied by the ECR plasma source vendor, order-of-magnitude estimates of the activated O and N fluxes in ML/s are fO ¼ ð3  105 Þ  P O2 and fN ¼ ð1:5  105 Þ  P N2 , where P O2 and P N2 are measured at the IG position. The magnetron current in the ECR power supply was typically 70 mA. Films with the fewest defects and best crystal properties typically resulted from the use of highly anion-rich conditions. In our experiments, these conditions consisted of a ˚ /s, which would result in a Ti atom flux (fTi) of 0.037 A growth rate of 0.02 ML/s TiO2xNx(1 1 0) under anionrich conditions. The minimum total anion flux required to achieve this growth rate and produce fully stoichiometric material is 0.04 ML/s. We found that the best material resulted from using P O2 ¼ P N2 ¼ 5  106 Torr, which is equivalent to fO = 1.2 ML/s and fN = 0.6 ML/s. The maximum x-value achievable under these conditions was found to be 0.04, as described below. Doubling fTi resulted in poorer crystallinity, detectable Ti(III) in the near-surface region, and an enhanced x-value of 0.12. For all growths, the substrate temperature was held at 550 C in order to minimize N loss from the films. When in method two the ozone source was used to deliver activated O, the pressures were matched by first setting the chamber pressure to 5 · 106 Torr in O3, and then increasing the chamber pressure to 1 · 105 Torr with nitrogen from the ECR plasma source. Following growth, specimens were transferred under UHV conditions to an appended X-ray photoelectron spectrometer consisting of a GammaData/Scienta SES 200 analyzer, a monochromatic Al Ka X-ray source and a He resonance lamp for in situ core-level and valence band measurements. Ex situ measurements consisted of high-resolution XRD, nuclear reaction analysis (NRA), electron paramagnetic resonance, and optical absorption in transmission mode. 3. Results We show in Fig. 1 RHEED intensity oscillations during a typical growth under optimized anion-rich conditions. Strong oscillations reveal layer-by-layer growth via nucleation on terraces for the first 30 monolayers (ML), followed by a transition to step-flow growth thereafter. Also shown in Fig. 1 is a typical RHEED pattern following ˚ after cool down. The Bragg rods are growth of 500 A essentially streaky with only weak intensity modulations along each streak, indicating a reasonably smooth single crystal surface. Although flatter surfaces were achieved

S.A. Chambers et al. / Chemical Physics 339 (2007) 27–35

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˚ of Fig. 1. RHEED beam intensities vs time during growth of 500 A TiO1.96N0.04 on TiO2(1 1 0) rutile, along with a RHEED pattern after growth.

by growing at higher substrate temperatures, N incorporation was significantly reduced. In the interest of maximizing N incorporation, we limited the substrate temperature to 550 C. Fig. 2 shows high-resolution N 1 s and Ti 2p core-level spectra for a typical film grown under optimized anion-rich conditions. The N 1s spectrum exhibits peaks at binding energies of 407.2 eV and 396.0 eV. These are assigned to surface bound NO3 [18] and N that has substituted for O in the lattice (NO) [19], respectively. Angle-resolved measurements (not shown) revealed that the NO3 species is present only on the surface whereas NO is distributed approximately uniformly throughout the XPS probe depth ˚ ) [15]. Comparison with spectra for N-containing (50 A reference compounds reveals that the formal charge state of NO is 3 [1,10,11,19]. The N concentration was obtained in two ways from the XPS data. First, we used the areas of the N 1s and O 1s peaks, measured with the same instrumental settings and assuming no change in spectrometer transmission over the kinetic energy region containing the two peaks (950–1080 eV). The N percentage in the anion sublattice is then given by 100(AN1s/rN1s)/ (AN1s/rN1s + AO1s/rO1s), where AN1s and AO1s are the peak areas and rN1s and rO1s are the XPS cross sections, as computed for free atoms by Yeh and Lindau [20]. This method typically gives 1% N concentration as a percentage of the total anions (TiO1.98N0.02) for growth under optimized anion-rich conditions. The second method involves doing a statistical analysis on the high-resolution, spin–orbit split Ti 2p line shape. As seen in Fig. 2b, doping with N results in transfer of intensity from the Ti 2p primary lattice peak into a weak second feature that is shifted 1.0 eV to lower binding energy from the primary peak. This observation is most apparent in the Ti 2p3/2 peak due to its inherently narrower peak width. However, the intensity transfer is also visible in

Fig. 2. Normal emission N 1s (a) and Ti 2p (b) core-level spectra for ˚ of pure TiO2 on ˚ of TiO1.96N0.04 on TiO2(1 1 0) rutile (A) and 500 A 500 A TiO2(1 1 0) rutile (B), along with the difference spectrum (A  B).

the Ti 2p1/2 peak. The intensity of the second Ti 2p3/2 peak scales with the N concentration, as judged by the N 1s peak area. The observed chemical shift is smaller by 1 eV than that associated with Ti(III) defect states in TiO2. We tentatively assign this peak to Ti bound to 5 O ligands and 1 NO ligand. This assignment is reasonable in light of the lower electronegativity of N compared to O. Lower electronegativity in N ligands is expected to result in slightly more valence charge surrounding the Ti core relative to the situation when Ti is bound to six O ligands. Higher valence charge results in more screening of electrons in the L2 and L3 shells and, thus, lower 2p binding energies, as observed. If this assignment is correct, the two peaks areas (A1 at 458.8 eV and A2 at 457.8 eV) should be useful for independently determining the N concentration, as described below. Photoelectron diffraction effects will not skew the concentration determinations by either method because they affect the relevant core-level intensities in the same way. For instance, since N substitutes for O (as substantiated below), 1s photoelectron intensities associ-

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ated with both elements will undergo the same structurally induced diffraction modulation. Likewise, Ti 2p photoelectrons from Ti–O6 and Ti–O5N1 environments will diffract the same way due to the high degree of similarity of their structural parameters. If the NO dopants are randomly distributed in the lattice, the probabilities of finding particular bonding configurations are governed by Poisson statistics. Specifically, if y is the fraction of the anion sublattice sites occupied by N, the probability of finding n N dopants in the first coordination sphere containing m sites surrounding a given Ti can be expressed as P(y, n) = {m!/[n!(m  n)!]}yn(1  y)mn. A plot of P(y, n) for 0 6 y 6 0.1, n = 0, 1, and 2, and m = 6 is shown in Fig. 3. P(y, 0) represents the probability of a given Ti not having any N ligands in its first coordination sphere. Thus, 1  P(y, 0) represents the probability of a given Ti being bound to at least one N. Moreover, we can equate 1  P(y, 0) with A2/(A1 + A2) from the Ti 2p3/ 2 spectrum to find y, and this value of y should agree with that found using the N 1s and O 1s peak areas if this assignment is correct. Indeed, doing so leads to excellent agreement on the N concentration, as seen in Table 1. In contrast, the N concentrations found using NRA are typically slightly higher. Table 1 shows the values of x (note: x = 2y) determined for three representative films using N 1s and O 1s peak areas, Poisson analysis of the Ti 2p3/2 line shape, and comparison with a standard in ˚ thick Si3N4 film grown on Si was NRA. A 1175 ± 50 A used to calibrate the 14N(d, a)12C nuclear reaction to quantify the N content. The first two films were grown using ˚ /s and optimized anion-rich conditions (fTi = 0.037 A P O2 ¼ P N2 ¼ 4  106 TorrÞ. fTi was doubled in the third film and P O2 and P N2 were unchanged relative to the values used for films 1 and 2. Based on these three independent methods, we conclude that the N at.% within the anion

Fig. 3. Poisson distribution for the probability of a given Ti having zero, one or two nearest neighbors N ligands in TiO2xNx with randomly distributed N as a function of y = 2x.

Table 1 Determination of x in TiO2xNx by different methods x from N 1s O1s

A2/(A1 + A2) from Ti 2p3/2

x from Poisson analysis of Ti 2p3/2a

x from NRA

0.020 0.024 0.10

0.058 0.090 0.26

0.020 0.034 0.10

0.060 0.080 0.15

a

x = 2y, and y is determined by equating 1  P(y, 0) with A2/(A1 + A2).

sublattice is typically 2 ± 1% (x = 0.04 ± 0.02) for films grown under optimized anion-rich conditions, but increases to 6 ± 1% (x = 0.12 ± 0.02) for films grown with twice the Ti flux. It has thus far been tacitly assumed that the N 1s peak at 396.0 eV originates from N substituting for O in the lattice. We have used NRA to verify this structural assignment. We used 0.95 MeV deuterons and the 14N(d, a)12C and 16O(d, p)17O nuclear reactions in this experiment. NRA spectra were measured in random and channeling directions, and rocking curves were obtained about [1 1 0]. Typical NRA spectra and rocking curves are shown in ˚ TiO1.96N0.04/ Figs. 4 and 5, respectively, for 500 A TiO2(1 1 0). Minimum yields taken from Fig. 4 are similar for O and N, 12.7% and 13.8%, respectively, establishing that 98% of the N is in O lattice sites. Likewise, the rocking curves show common deep minima along [1 1 0], with similar half widths of 0.38 and 0.33 for O and N, respectively. Thus, N substitutes for O in the lattice with minimal lattice disorder. High-resolution XRD (not shown) reveals that the a and c lattice parameters expand by +0.5% and +0.8%, respectively, in epitaxial rutile films grown under optimized anion-rich conditions relative to pure rutile [15]. These results are consistent with the slightly larger ˚ ) compared to that for O2 ionic radius for N3 (1.46 A ˚ ) [21]. (1.36 A

˚ of TiO1.96N0.04 on TiO2(1 1 0) rutile. The Fig. 4. NRA spectra for 500 A deuteron energy and the angle between incident and outgoing particles were 0.95 MeV and 150, respectively.

S.A. Chambers et al. / Chemical Physics 339 (2007) 27–35

Fig. 5. NRA rocking curves for O and N with the incident deuteron ˚ of TiO1.96N0.04 on TiO2(1 1 0) rutile. The aligned along [1 1 0] in 500 A deuteron energy and the angle between incident and outgoing particles were 0.95 MeV and 150, respectively.

Growth method 2 produces the same extent of N incorporation and film properties as growth method 1. Therefore, recombination and/or reaction of activated N and O within the ECR plasma tube do not have any appreciable effect on film properties or quality. Electronically, TiO1.96N0.04 films grown under optimized anion-rich conditions are consistently highly resistive, but never p-type. This result is somewhat surprising because in principle, NO should be an acceptor in an oxide semiconductor, provided NO maintains the same formal charge as the host anion. However, the N 1s binding energy is consistent with NO being in a 3 formal charge state, rather than 2 like the host anion [22]. N3 O has a filled valence shell and is therefore not able to act as an acceptor. The question at hand is why does N assume a formal charge of 3 rather than 2? A plausible answer comes from an intriguing aspect of pure TiO2(1 1 0) homoepitaxy under optimized ˚ /s and P O ¼ anion-rich conditions (e.g. fTi = 0.037 A 2 2  105 Torr). This experiment yields a rather surprising result. The entire specimen (film and substrate) takes on blue coloration and become semiconducting (n-type). Moreover, EPR measurements indicate the presence of interstitial and substitutional Ti(III), although no Ti(III) is detectable by either core-level or valence band photoemission in the near-surface region [15]. Fig. 6 shows the EPR spectrum from such a pure homoepitaxial TiO2(1 1 0) film (curve A). Several resonances, corresponding to different local structural environments for Ti(III), are visible in the g-factor range 1.9–2.0 (3300–3600 G). Interestingly, if a TiO2(1 1 0) substrate is brought to growth temperature (550 C) in the PAMBE chamber and exposed to activated O for a length of time equal to that required for a typical growth, it becomes neither blue nor semiconducting. Therefore, the onset of blue coloration and n-type behavior is not due to reduction by Ov creation during annealing. Taken together, these data suggest that some

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˚ of homoepitaxial TiO2(1 1 0) rutile Fig. 6. EPR spectra at 10 K for 500 A ˚ of TiO1.96N0.04 (1 0 0) rutile on a-Al2O3(0 0 1) (B). (A) and 500 A

fraction of the incident Ti, after at least partial oxidation at the growth front, diffuses into the bulk and occupies interstitial or substitutional sites. In order to test this hypothesis, we have performed the following experiment. Optimized homoepitaxial growth was commenced under the conditions noted above. The O flux was then reduced in discrete steps, keeping fTi constant. For each O flux value, a few MLs were grown and the instantaneous growth rate, as measured by RHEED intensity oscillations, was determined. This sequence was continued until fO had been reduced by a factor of 40. Even after a factor-of-forty reduction in fO, fO was still in excess of fTi, resulting in fully stoichiometric TiO2 formation at the growth front. The results are summarized in Fig. 7 and Table 2. The numbers in the ‘‘experimental condition’’ column of Table 2 correspond to those superimposed on the RHEED oscillations in Fig. 7a. Below a flux ratio (fO/fTi) of 75, the RHEED intensity oscillation period (s) begins to increase, and the growth rate (1/s) begins to fall (Fig. 7b). The corresponding increase in (0 0) beam intensity seen in Fig. 7a is an artifact of the diminishing O2 pressure in the chamber, which results in an increase in the primary beam electron mean free path. The process is fully reversible in that once fO is increased to near its original value (condition 14), the growth rate also recovers to its original value. This phenomenon can only occur if an increasing fraction of the incident Ti does not remain incorporated in new TiO2 layers when fO drops. In this scenario, at least partially oxidized Ti ions diffuse away from the growth front and into the bulk in increasing proportion, resulting in a diminished growth rate. The relationship between Ti(III) in the interior of the film and N3 at lattice sites is as follows. Ti(III) is a shallow donor in TiO2. Some fraction of the Ti(III) species are thermally ionized at room temperature, resulting in itinerant conduction band (CB) electrons. NO results in new states just off the top of the TiO2 valence band (VB), as seen in VB photoemission (not shown) [11,15]. As Ti 3d-derived

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Fig. 7. RHEED (0 0) beam intensity oscillations (a) and associated frequencies (b) for the experiment described in Table 2.

Table 2 Experimental conditions for homoepitaxial growth of TiO2(1 1 0) in which ˚ /s (0.02 ML TiO2(1 1 0) per second with sufficient oxygen) fTi = 0.037 A Experimental condition

P O2 (Torr)

fO (ML/s)

1 2 3 4 5 6 7 8 9 10 11 12 13 14

1.9 · 105 1.8 · 105 1.5 · 105 1.4 · 105 1.2 · 105 1.0 · 105 8.4 · 106 6.9 · 106 5.6 · 106 4.2 · 106 2.8 · 106 1.4 · 106 4.8 · 107 1.5 · 105

5.7 5.4 4.5 4.2 3.6 3.0 2.5 2.1 1.7 1.3 0.84 0.42 0.14 4.5

CB electrons traverse the material and encounter deep-level traps associated with N2 O acceptor states, they decay into these holes, thereby compensating the associated N acceptors. Thus, a Ti(III) + N2 ! Ti(IV) + N3 charge transfer

(CT) process is suggested to occur by this mechanism. One test of this mechanism is to look for Ti or N EPR signals in TiO2xNx grown on a substrate other than TiO2(1 1 0). Using a different substrate is important because bulk rutile is a sink for Ti(III) and growth of TiO2xNx on TiO2(1 1 0) may result in many more Ti(III) than NO throughout the entire specimen volume. By growing films on a-Al2O3(0 0 1) and performing EPR measurements, we can investigate the interaction between Ti(III) and N2 in a volume confined to the film. If the above CT process occurs, and if Ti(III) and NO are present in approximately equal concentrations, there should be no Ti or N EPR signals, because both Ti(IV) and N3 are EPR silent. Indeed, such is the case, as seen in Fig. 6, spectrum B. N2 should show prominent hyperfine splittings from it is I = 1 nuclear spin. For example, substitutional N 2 in ZnO exhibits a multiplet centered at g = 2.0 resulting from the interaction of the electron with two equivalent I = 1 N nuclei [23]. Similarly, we would expect a well-defined triplet with a 1:1:1 intensity ratio at g = 2.0 for isolated, substitutional N2, if such a species exists [13]. However, no such multiplet is seen. The weak, featureless line seen at g = 1.94 (3500 G) has a width of 100 G and lacks the hyperfine splitting expected for NO. The g-factor falls in the range expected for Ti(III) but is well away from the g  2.0 expected for NO. The signal also lacks the sharp, well-resolved lines observed for Ti(III) in Fig. 6a. Based on its intensity, this EPR signal corresponds to a spin concentration of less than 1015 cm3, orders of magnitude less than the NO concentration in this film (6.4 · 1020 cm3). This feature may be an impurity or Ti(III) at the TiO2/sapphire interface. If it is the latter, the broadening may be due to lattice-mismatch-induced strain. Despite its low concentration, NO has a marked effect on the density of states in TiO2(1 1 0). The effect that NO has on the Ti 2p3/2 line shape (Fig. 2b) reveals charge transfer and, is therefore consistent with hybridization between substitutional N and lattice Ti. N 2p-derived states form at the top of the VB, as seen in Fig. 8 by comparing optical absorption spectra for N-doped (curve A) and pure (curve B) films. These states are also seen over the same energy range in VB photoemission for which spectra are referenced to the Fermi level, which is in turn quite close to the CB minimum in the n-type samples investigated [11,15]. Theoretical partial densities of states show that the top of the VB is dominated by Ti 4p and O 2p character in pure rutile [24], so it is conceivable that Ti 4p–NO 2p hybridized states might lay nearby, at the top of the VB. These states extend up to 2.4 eV, as they do in VB photoemission, suggesting a bandgap reduction. It remains to be seen if these new states result in a bona vide bandgap reduction, with the attendant photoconductivity at hm P 2.4 eV, or whether these states are localized and do not enhance itinerant electron–hole pair creation at photon energies in the visible. We are currently building apparatus to carry out photoconductivity measurements and explore these possibilities.

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˚ of TiO1.96N0.04 on TiO2(1 1 0) Fig. 8. Optical absorbance spectra for 500 A ˚ of pure TiO2 on TiO2(1 1 0) rutile (B), along with the rutile (A) and 500 A difference spectrum (A  B).

4. Discussion The principal conclusions from the present study are: (i) N substitutes for O in the lattice and is limited to 2% of the total anions, (ii) N is formally N3 rather than N2, as in TiN, (iii) N incorporation in the lattice results in a substantial red shift in the bandgap and induces new states at the top of the VB, and, (iv) Tii formation occurs during Ndoped TiO2 homoepitaxy and results in N2 O compensation via electron transfer from Tii(III), leading to the formation of EPR silent Tii(IV) and N3 O . The first three of these conclusions are consistent with those drawn in previous studies, although higher N concentrations can be achieved when high defect levels are present [11,15]. However, the fourth differs from previous conclusions. Other authors claim that O vacancies (Ov), rather than Tii are the prevailing species that compensate N2 O and indeed stabilize sub-

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stitutional N [5,11,13]. We can rule out Ov in favor of Tii as the dominant shallow donor in our material for several reasons. However, before we list these reasons, it is noteworthy that previous EPR [25–27], TEM [28], ion channeling [29], and transport measurements [30] all show that the dominant defect in slightly reduced bulk TiO2 rutile is Tii and not Ov. These experimental results are corroborated by very recent ab initio calculations which show that the migration barrier for Tii is 1 eV lower than that for Ov in rutile [31]. Significantly, these calculations predict that Ti indiffusion should strongly dominate over O outdiffusion during pure and possibly also N-doped TiO2 epitaxial growth. Therefore, the dark coloration present in all film specimens, a few of which are shown in Fig. 9, is most likely due to Tii. The presence of discoloration on the corners of the rutile substrate (Fig. 9b), which are masked from the atomic beams by a Ta retainer ring during growth, along with the darker appearance of the specimen grown on rutile (Fig. 9b) relative to sapphire (Fig. 9c), suggests that rutile is a natural sink for Tii. We now summarize the reasons why we conclude that Tii formation occurs preferentially over Ov creation in oxygen-rich TiO2 homoepitaxy, with and without N doping. First, the experimental conditions used in our growth chamber utterly preclude reduction and, therefore, the likelihood of Ov creation. The substrate temperature (550 C) is too low to thermally reduce bulk TiO2(1 1 0) over the time scale of a growth (2 h), particularly when the surface is being bombarded by activated O throughout the growth process. Indeed, bringing a TiO2(1 1 0) substrate to 550 C for 2 h while exposing to atomic O resulted in the crystal remaining transparent and insulating (Fig. 9a). Second, Ti indiffusion is the most reasonable way to explain the pronounced drop in growth rate as the O flux was reduced at constant Ti flux (e.g. Fig. 7 and Table 2). Ti re-evaporation at low O flux is negligible because the vapor pressure of Ti does not reach 1011 Torr (equivalent to desorption of 1 ML in 24 h), much less a value of sufficient magni-

˚ TiO1.98N0.02/ Fig. 9. (a) Typical TiO2(1 1 0) substrate after heating for 2 h at 550 C in activated O, but with no incident Ti beam exposure, (b) 500 A ˚ TiO1.98N0.02/a-Al2O3(0 0 1). The circle and cross outline the portion of the substrate exposed to the atomic beams during TiO2(1 1 0), (c) and 700A growth; the corners are covered by a Ta retainer ring assembly during growth and are, therefore, bare substrate. The absence of blue discoloration in (a) establishes that the substrate temperature used for growth is not sufficiently high to reduce rutile. Notice that in (b) the discoloration characteristic of Tii donors is clearly visible even where growth did not occur, indicating that Ti diffuses throughout the rutile substrate during growth. In contrast, this phenomenon does not occur when N:TiO2 is grown on sapphire (c); Tii remains in the N:TiO2 film.

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tude to result in measurable Ti loss, until the surface temperature reaches 900 C. Moreover, one cannot explain the data in Fig. 7 by invoking a decrease in the Ti deposition rate from the effusion cell as the O2 pressure is reduced because if anything, the deposition rate increases as the chamber pressure drops. For instance, at an effusion cell temperature of 1560 C, the Ti deposition rate in high vacuum, as measured by a quartz crystal oscillator at the sub˚ /s. This value would give rise to a strate position, is 0.058 A TiO2(1 1 0) growth time of 32 s per monolayer (s/ML) under O-rich conditions assuming no Ti loss to either indiffusion or re-evaporation. The actual growth time for undoped TiO2(1 1 0) at a Ti cell temperature of 1560 C, based on RHEED oscillation periods, ranges from 53 s/ML at P O2 ¼ 1:9  105 Torr to 162 s/ML at P O2 ¼ 4:8  107 Torr. This trend is clearly going in the wrong direction to be due to any variation in Ti sublimation rate from the effusion cell with chamber pressure. Third, the EPR clearly shows several Ti(III) resonances corresponding to different interstitial and substitutional environments for pure TiO2 films (Fig. 6) [25–27]. In contrast, Ti(III) is not visible in core-level or valence band photoemission because it is distributed throughout the film and the substrate and is well below the detection limit for Ti(III) in TiO2 by photoemission (estimated to be 1% of a monolayer using the Ti 2p core-level). There is no a priori reason to think that N doping would inhibit Ti indiffusion and Tii formation. Since N substitutes for O and is coordinated with lattice Ti ions, it is not expected to trap Tii. Therefore, a growing TiO2xNx layer is not expected to be a diffusion barrier for Tii. As in the case of pure TiO2 homoepitaxy, the growth conditions are utterly oxidative; the fO/fTi ratio is 75 under optimized anion-rich growth conditions. There is ample O to satisfy every incoming Ti, and this fact is what limits the extent of N incorporation to 2% of the anions. In fact, in growths for which P O2 ¼ P N2 ¼ 1  105 Torr and ˚ /s (fO/fTi = 285), the thermodynamic preferfTi = 0.020 A ence of Ti for O relative to N is so strong that no N incorporates in anion sites at the XPS detection limit, despite the comparable fluxes of N and O. Some shallow donor is clearly compensating N accep2 tors. This compensation results in N3 O rather than NO , a marked reduction in, but not complete elimination of n-type conductivity compared to similar prepared films with no N [15], and the absence of p-type behavior. The presence of Ov is concluded to be highly unlikely in light of the oxidative conditions. Even though there is no direct evidence for the presence of Tii in N-doped films, its conversion to Tii(IV) by electron transfer to N2 O constitutes a more satisfactory explanation for the totality of observed behavior in this material than compensation of N2 O by Ov. 5. Summary N-doped TiO2 rutile with very high crystalline quality and low defect density has been grown homoepitaxially

by plasma-assisted molecular beam epitaxy and investigated with a variety of techniques to generate a rather thorough description of the material. N substitutes for lattice O to a sparing extent, but does not result in p-type electronic behavior. Rather, N acceptors are compensated by electrons form interstitial Ti(III) that diffuses into the bulk of the film during growth. N doping results in new occupied states near the top of the valence band and an apparent reduction in band gap of 0.6 eV. If electron–hole pairs generated by photoabsorption at hm P 2.4 eV are itinerant, this material may be useful for enhanced photochemistry in the visible. Acknowledgements 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. This work was supported by the US Department of Energy, Office of Science, Division of Chemical Sciences. The authors thank M.A. Henderson for useful discussions and a critical reading of this manuscript. References [1] R. Asahi, T. Morikawa, T. Ohwaki, A. Aoki, Y. Taga, Science 293 (2001) 269. [2] O. Diwald, T.L. Thompson, E.G. Goralski, S.D. Walck, J.T. Yates Jr., J. Phys. Chem. B 108 (2004) 52. [3] O. Diwald, T.L. Thompson, T. Zubkov, E.G. Goralski, S.D. Walck, J.T. Yates Jr., J. Phys. Chem. B 108 (2004) 6004. [4] C. Di Valentin, G. Pacchioni, A. Selloni, Phys. Rev. B 70 (2004) 085116. [5] C. Di Valentin, G. Pacchioni, A. Selloni, S. Livraghi, E. Giamello, J. Phys. Chem. B 109 (2005) 11414. [6] J. Lee, J. Park, J. Cho, App. Phys. Lett. 87 (2005). [7] A.R. Gandhe, J.B. Fernandes, J. Solid State Chem. 178 (2005) 2952. [8] T. Okato, T. Sakano, M. Obara, Phys. Rev. B 72 (2005) 115124. [9] Z. Lin, A. Orlov, R.M. Lambert, M.C. Payne, J. Phys. Chem. B 109 (2005) 20948. [10] M.-S. Wong, H.P. Chou, T.-S. Yang, Thin Solid Films 494 (2006) 244. [11] M. Batzill, E.H. Morales, U. Diebold, Phys. Rev. Lett. 96 (2006) 026103. [12] U. Koslowski, K. Ellmer, P. Bogdanoff, T. Dittrich, T. Guminskaya, H. Tributsch, J. Vac. Sci. Technol. A 24 (2006) 2199. [13] S. Livraghi, M.C. Paganini, E. Giamello, A. Selloni, C. DiValentin, G. Pacchioni, J. Am. Chem. Soc. 128 (2006) 15666. [14] S.A. Chambers, Surf. Sci. Rep. 39 (2000) 105. [15] S.H. Cheung, P. Nachimuthu, A.G. Joly, M.H. Engelhard, M.K. Bowman, S.A. Chambers, Surf. Sci. 601 (2007) 1754. [16] S.H. Cheung, P. Nachimuthu, A.G. Joly, M.H. Engelhard, M.K. Bowman, S.A. Chambers, unpublished. [17] J.A. Dean (Ed.), Lange’s Handbook of Chemistry, 15th ed., McGraw-Hill, Inc., 1998. [18] J.A. Rodriguez, T. Jirsak, G. Liu, J. Hrbek, J. Dvorak, A. Maiti, J. Am. Chem. Soc. 123 (2001) 9597. [19] F.C. Voogt, P.J.M. Smulders, G.H. Wijnja, T. Fujii, M.A. James, T. Hibma, L. Niesen, Phys. Rev. B 63 (2001) 125409. [20] J.J. Yeh, I. Lindau, Atomic Data and Nuclear Data Tables 32 (1985) 1. [21] R.D. Shannon, Acta Crystallogr. A 32 (1976) 751.

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