Photocarrier Transport Mechanisms in Amorphous and Epitaxial TiO2

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C: Physical Processes in Nanomaterials and Nanostructures

Photocarrier Transport Mechanisms in Amorphous and Epitaxial TiO/SrRuO Heterojunction Photocatalysts 2

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David Eitan Barlaz, and Edmund G. Seebauer J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b12659 • Publication Date (Web): 16 Jun 2018 Downloaded from http://pubs.acs.org on June 19, 2018

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The Journal of Physical Chemistry

Photocarrier Transport Mechanisms in Amorphous and Epitaxial TiO2/SrRuO3 Heterojunction Photocatalysts

D. Eitan Barlaz and Edmund G. Seebauer*

Department of Chemical and Biomolecular Engineering, University of Illinois, Urbana, Illinois 61801, USA

Abstract Heterojunction photocatalysts in thin film form offer the possibility for improved optical absorption of solar radiation but have found limited use due to various material challenges.

Epitaxial structures based upon strontium ruthenate (SRO) and TiO2 have

demonstrated unexpectedly high activity under visible-only illumination because of strong absorption by the SRO, high electrical conductivity, and the ability to inject hot electrons into the active TiO2 photocatalyst. The role of photoholes, the mechanisms of carrier transport to the TiO2 surface, and the necessity of an epitaxial structure remain unclear. The present work helps to fill these gaps through rate measurements of methylene blue (MB) photooxidation on SROTiO2 under visible light, together with photoemission measurements of interfacial and free surface band edges of the TiO2. Diffusive transport of thermalized holes appears sufficient to explain the results, and surprisingly, a heterojunction based on amorphous SRO and TiO2 retains a great deal of the metallic properties of crystalline SRO and provides MB degradation rates comparable to an equivalent heteroepitaxial structure reported previously. These findings relax considerably the constraints on translating heterojunctions based on correlated metal oxides into photocatalytic applications.

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*

To whom all correspondence should be addressed. Electronic address: [email protected]. Phone number: (217) 244-9214

Introduction Thin-film semiconducting photocatalysts currently find applications in self-cleaning windows and other weather-exposed surfaces1,2, and in anti-microbial3 and anti-fouling coatings4. Reaction rates on these surfaces would often benefit greatly if minority photocarriers could be driven more efficiently to the surface5 through improved optical absorption5,6 and the manipulation of electric fields within the semiconductor7. Heterostructures offer the possibility to address both challenges5 as well as the need to spatially separate oxidation from reduction processes, but have found very limited use in photocatalysis8,9,10. In thin films, the limitations have several origins, including the complexities of polycrystalline wide-bandgap semiconductors with electrically active grain boundaries11, the difficulty in controlling background carrier concentration reliably11,12,13, and the need to accomplish oxidation and reduction simultaneously at the surface to preserve charge balance14. This last requirement presents acute problems for endothermic reactions wherein the surface catalyzes both forward and reverse reactions, and for semiconductors whose band edges are not positioned optimally with respect to the redox potential of a key reactant species15. One candidate structure that has demonstrated16,17 significant photocatalytic activity under full-spectrum solar illumination for the test reaction of methylene blue (MB) degradation is a heterojunction consisting of single-crystal strontium ruthenate (SrRuO3, “SRO”) with anatase TiO2 grown epitaxially on top. Although the quantum efficiency of this structure fell well below that of common photocatalysts, the activity surpassed the thermodynamic limit 2


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normally imposed by single-component layers. Anatase offers many favorable photocatalytic properties except for weak absorption of the solar spectrum.

SRO behaves as a strongly

correlated metal oxide – a class of materials with unusual optical and electronic properties that shows promise for all-oxide magnetic tunnel junctions18 and other spin-electronic devices19, and electrode materials20. SRO absorbs strongly in the visible spectrum yet has low reflectivity and electrical conductivity close that of a metalloid. Yet the photocatalytic efficiency of this heterostructure is counterintuitive because of the large electronic barriers (≥ 1.3 eV) for electron and hole transfer from SRO into TiO2. Injection of hot electrons has been proposed16 as the primary mechanism, although the literature has been silent about the holes other than to suggest that hot electron injection may reduce reliance on minority carrier transport. No information addresses whether the reaction-driving carriers are still hot upon reaching the TiO2 surface, or the roles of drift or diffusion of thermalized carriers. The necessity of a crystalline epitaxial structure remains unclear. The present work helps to fill these gaps through rate measurements of methylene blue photooxidation under visible-only illumination, together with photoemission measurements of interfacial and free surface band edges of the TiO2. Diffusive transport of thermalized holes appears sufficient to explain the results, and a heterojunction based on amorphous SRO and TiO2 provides rates comparable to the epitaxial structure. Additional ultraviolet illumination from a full solar spectrum actually decreases the photocatalytic rate somewhat. Overall, these findings relax considerably the constraints on translating heterojunctions based on correlated metal oxides into photocatalytic applications.

Experiment Synthesis

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Heterojunction devices were grown on undoped SrTiO3 (STO) single crystals (MTI Corporation) of lateral dimensions 5×5 mm. Before deposition of SRO thin films, the crystals were subjected to successive ultrasonic baths in reagent grade acetone and isopropanol followed by drying with flowing N2. SRO was deposited in two ways depending upon the desired phase; direct-current sputtering yielded the amorphous phase, while pulse laser deposition yielded epitaxy. Sputtering employed a single strontium-ruthenium oxide ceramic target in a chamber with a base pressure below 1×10-5 Torr. The process took place at room temperature in 4 mTorr of a 60:40 Ar:O2 mixture at a power of 15 W.

Pulsed laser deposition took place at Brookhaven National

Laboratory in a PVD Products PLD/MBE 2300 chamber that employed a 248 nm KrF excimer laser. Deposition employed a single strontium-ruthenium oxide ceramic target at 750 °C in 200 mTorr of oxygen. Epitaxial recrystallization of the amorphous SRO was attempted by annealing in air up to 900 °C. Small, highly stressed, single crystal islands formed leaving the overall film highly discontinuous and bearing a surface roughness on the same order as film thickness. These specimens were not used for photocatalytic studies as activity per unit area would be difficult to quantify. TiO2 was grown by a layer-by-layer pseudo-ALD process using titanium tetraisopropoxide (Strem Chemicals Inc., 98%) and deionized water (18 Mohm-cm, Barnstead E-pure D4641) at a rate of 1 nm/cycle. At a deposition temperature of 200 °C all specimens grown were amorphous. Details of the chamber used may be found in previous publications21. Characterization X-ray diffraction measurements of crystal structure employed a PANalytical X-ray diffractometer. The copper source operated at 45 kV and 40 mA. For X-ray reflectivity (XRR) 4


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scans, a 1/32 degree slit is used with the divergence slit, X-ray diffraction scans use a 1/2 degree slit. XRD data were analyzed using MDI JADE software*. XRR data were analyzed for bulk density22,23 using the critical angle position, as well as thickness and roughness23 by modeling of fringe spacing on the PANalytical XRR software package. An Asylum Research MFP-3DTM AFM atomic force microscope (AFM) was used for surface roughness measurements. Tips were Budget Sensors BS-Tap300Al tapping tips with a lateral resolution of approximately 8 nm. All measurements were performed in tapping mode. Data analysis employed the Igor Pro software package. X-Ray Photoelectron Spectroscopy (XPS) measurements were performed with a Kratos Axis ULTRA instrument that employed a pass energy of 40 eV and excitation by monochromatic Al Kα radiation. Specimens were mounted with copper clips on the surface as the semiconducting nature of the TiO2 can lead to charging. All specimens were grounded, and data were collected with charge compensation off. The electrical conductivity of the SRO was sufficient to prevent significant charging when grounded despite the insulating nature of the STO substrates. All data fitting and analysis was performed using the Casa software package. In addition to chemical composition analysis, XPS was employed to estimate the difference in electrostatic potential between the TiO2 free surface and its interface with the underlying SRO. Examples of this uncommon method appear in the literature24,25, and rely upon the variations in binding energy that result from corresponding variations of electrostatic potential within the space charge region of a semiconductor.

If that potential changes

appreciably within the escape depth of photoelectrons, the XPS peaks will exhibit broadening. Although a priori determination from the broadening of potential variation requires extensive

* JCPDS PDF cards: Anatase TiO2 21-1272, SrRuO3 43-0472, SrTiO3 35-0734 5



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modeling, rough estimates may be made by comparison to a reference specimen without the buried interface that anchors the potential close to the surface.

In the present work, the Ti 2p

peak was employed for the TiO2 overlayer. The Ru 3p peak provided a useful aid for deconvoluting the Ti peaks, as the highly conducting SRO does not support internal electric fields and therefore does not experience potential-induced broadening. Ti 2p and Ru 3p peaks were respectively fit using SGL(20) and A(0.2,0.6,0)SGL(50) line shapes, with Shirley background in both cases. Other constraints to aid fitting included constancy of the Ti 2p and Ru 3p peaks areas and the Ti 2p spin-orbit peak spacing. Exact overlayer thicknesses varied slightly among the TiO2 overlayers, but not enough to affect the electrostatic potential differences at the level of accuracy sought here. Optical transmittance measurements of SRO films were taken over the visible region were made using a Varian Cary 5G Spectrophotometer. Transmission measurements were taken in diffuse scattering mode, rather than direct, due to the effect of the unpolished substrate backside. Specimens were mounted against a 2.5 mm through hole to ensure that only the film area was measured. Sheet resistance measurements of STO were made by four point probe (Jandel) under dark conditions. The reported values average measurements at 100 µA, 500 µA and 1 mA at 2-3 distinct locations on each specimen, with typical variations on the order of 1%.

Photocatalysis Rate Determination Photocatalysis experiments were performed using a Newport solar simulator as the light source. The simulator was equipped with a 300 W Xe arc lamp producing 3,760 mW/cm2 full AM1.5G spectrum light with a condenser lens focusing to a 3×4 mm oval shaped focus spot size. Additionally, a λOD2 = 416 nm longpass glass filter (which removes more than 99.97% of 6


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light with energy above the band gap of TiO2) was used to prevent ultraviolet photostimulation of the TiO2 unless otherwise noted. The intensity of the light incident on the surface has been estimated at approximately 3280 mW/cm2 by other users16. Absorbance was monitored at the peak absorbance of 665.03 nm. Each specimen’s kinetic measurement used 15 mL (by weight) of a 1 ppm (3.17×106

M) MB solution diluted with the same deionized water described for TiO2 synthesis.

Absorbance measurements were processed from the raw intensity signal from an Ocean Optics USB4000 Fiber Optic Spectrometer. The dye solution was recirculated continuously. With the reservoir open to the atmosphere, the dye solution was allowed to equilibrate for 2 hours to ensure that the solution was fully saturated with oxygen. Primary sources of error originated from a combination of slight variations in intensity. The detection limit for the experimental setup was a rate of approximately 0.004 µm of MB decolored/hr*cm2 based on a reference sample of TiO2 under filtered light with thickness much greater than the mean free path of a photocarrier.

Results Material Properties Structural characterization of SRO and TiO2 consisted of XRD and AFM measurements (supplemented by XRR in the case of TiO2). XRD of epitaxial SRO exhibited peaks having clean separation from nearby STO substrate peaks. Amorphous SRO exhibited no XRD peaks at all. No signal change is observed upon deposition of amorphous TiO2, shown in Figure 1, as expected.

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Amorphous Annealed Epitaxial

Intensity (AU)

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2θ θ ( °) Figure 1: XRD 2θ-Ω plots of TiO2/SrRuO3 thin films. Visible SrTiO3 substrate peaks are at 22.8° (100) and 46.5 (200); SrRuO3 peaks are at 22.6° (020) and 46.2° (040). The annealed SRO with crystalline islands creates a shoulder on the left side of the STO peak. Crystalline islands on a discontinuous film can be seen in the SEM image insert. Insert image is 2.5 µm across. Small peaks on the amorphous sample have been attributed to carrier wafer and bonding material, see Ref [26]. AFM micrographs of representative specimens showed a roughness on order of 2 nm or less for both epitaxial and amorphous material (Figure 2). TiO2 films of varying thickness were examined by XRR, and roughness estimates from subsequent modeling were consistently on the order of 1-2 nm. Most of the TiO2 roughness probably propagated from that of the underlying SRO.

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Figure 2: AFM micrographs of a) amorphous and b) epitaxial SrRuO3 thin films. Optical transmittance measurements (Figure 3) of 60 nm amorphous and epitaxial SRO films showed that in both cases, transmittance is very low below about 400 nm and then increases sharply to a rough plateau. However, the plateau is nearly a factor of three higher for the epitaxial material than amorphous. Since the interfacial and free surface roughness are 9



comparable in the two cases and are not major factors in diffuse transmittance measurements, the difference probably arises from larger absorbance by the amorphous phase. The resistivity of the amorphous material (9.2×10-4 Ω•cm) lies above that of epitaxial material (2.8×10-4 Ω•cm) by a factor of about three. The expected insulating character of the underlying STO was verified by control measurements and did not pass a stable current. 70

SrTiO3 SrRuO3 - Amorphous

60 % Transmittance

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SrRuO3 - Epitaxial

50 40 30 20 10 0 400

500 600 700 Wavelength (nm)

800

Figure 3: Optical transmittance of epitaxial and amorphous SrRuO3. Potential drop across TiO2 The amorphous SRO was produced from a single target; as a result little control over stoichiometry was possible. XPS peak areas showed a stoichiometric ratio of Sr:Ru:O that approximately equaled 0.2:1:3. A reference specimen of amorphous TiO2 for surface and interface potential measurements was produced by growing 90 nm directly on an Sb-doped Si(100) substrate (Silicon Quest International) with n-type resistivity of 0.013 Ω•cm. The expected potential change within TiO2 on SiO2 is likely to be negligible, and certainly far less than that in the 10


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heterojunction structure.

Reference specimens of amorphous and epitaxial SrRuO3 were

acquired by using the usual STO substrates but without growing the TiO2 overlayer. Figure 4 shows spectral data including the Ti 2p and Ru 3p regions. The epitaxial Ru 3p data has been published previously as a database reference entry27.

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Figure 4: Reference spectra of a) amorphous TiO2 2p region, b) amorphous SrRuO3 3p region and, c) epitaxial SrRuO3 3p region. Figure 5 shows the resulting fittings. The broadening in the full width at half maximum of the later peak equals 0.118 and 0.105 eV for the amorphous and epitaxial SrRuO3 substrates, respectively. We therefore take the potential drop across both these specimens to be on the order of 0.1 eV. Despite the slight residual differences on the high energy side of the convoluted Ru 3p 3/2 and Ti 2p 1/2 peaks, the fits for the Ti 2p 3/2 peaks showed consistent residual values. Several alternate sets of fitting constraints produced very similar determinations of broadening, lending confidence to the veracity of the fitting. Because the Ti and Ru peaks are heavily convoluted, exact peak positions were not determined as was done in Reference [24]. More accurate estimates of the potential drop may be possible if the SRO layer was replaced by another metallic oxide lacking an XPS signature overlapping with the Ti 2p region. SrVO3 would be a useful candidate material for such a comparison as it offers comparable optical absorption in the visible, and TiO2 overlayers may be grown epitaxially or amorphously as with SRO17.

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Figure 5: XPS spectra of a) amorphous TiO2/amorphous SrRuO3, and b) amorphous TiO2/epitaxial SrRuO3.

Photocatalytic Rate Determination Figure 6 compares concentration vs time for various TiO2 thicknesses for both epitaxial and amorphous structures. The rates are roughly comparable for the two kinds of heterojunction.

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Figure 6: Decoloration of MB on various thicknesses of amorphous TiO2 on a) amorphous SrRuO3 and b) epitaxial SrRuO3. Decoloration rates for the epitaxial structures were extracted assuming first order kinetics and converted to a µmol of MB decolored/hr*cm2 unit basis for comparison to systems reported in the literature, Figure 7 shows these values. The rate decreases by about a factor of four with increasing thickness, with a dependence that is adequately fit with a simple exponential functional form. The exponential decay length λ derived from the slope is λ = 38 ± 2 nm.

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µmol/hr⋅⋅cm2) MB Consumption (µ


0.022 0.02 0.018 0.016 0.014

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Visible Light Full Spectrum

0.012 0.01 0.008 0.006

0.004 0

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Thickness (nm)

Figure 7: Experimental Methylene blue consumption rate values for epitaxial SrRuO3 specimens normalized by illuminated surface area. Typical error in rate is likely to be roughly ±5% based on the signal/noise in the absorbance data. The amorphous and epitaxial structures induce identical temporal evolution, seen in Figure 6, in MB concentration to within a few percent.

Discussion The substantial photocatalytic activity of the SRO-TiO2 heterostructure runs counter to the large energy barriers for transfer of both electrons and holes from SRO into TiO2. One proposed mechanism involves the injection of hot electrons16, but with no commentary about the commensurate number of holes that are also required to preserve electrical neutrality of both the SRO and TiO2 layers. In addition, it remains unclear whether the reaction-driving carriers remain hot by the time they arrive at the surface. The roles of electric field drift effects and the epitaxial crystalline structure remain ill-defined. The sections below suggest that diffusive transport of thermalized holes adequately explains the results, and a heterojunction based on amorphous SRO and TiO2 suffices to yield significant photocatalytic activity. 16


Virtually all degradation

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mechanisms for methylene blue photooxidation on TiO214,28 and other n-type oxide semiconductors29,30,31 consider holes to be a rate limiting carrier, as they exist in the minority compared to electrons. MB photocatalysis by TiO2 does not require the presence of hot carriers, and simple calculations suggest that both holes and electrons have thermalized well before reaching the surface. Hot carriers typically last in semiconductors between 10-13 and 10-11 seconds32,33,34 before thermalization (compared to a recombination lifetime of 10-9 to 10-6 seconds35,36). Distances of travel are known only imperfectly, although recent first principles calculations have put this number near 15 nm for good-quality crystalline silicon32. We expect smaller distances in defected polycrystalline or amorphous TiO2 − less than the thicknesses employed here. Nevertheless, the possibility of hot holes or electrons reaching the surface cannot be ruled out and may merit investigation in the future. A suitable series of reactants with redox potentials lying just outside the band edges of TiO2 would be needed in order to create a system where any reactant loss could only be the result37,38 of hot carriers reaching the TiO2 free surface. The rate vs. TiO2 thickness data in Figure 7 yield a phenomenological decay length of λ = 38 nm. Such an exponential decay would be expected from a constant flux of ratelimiting carriers injected at the SRO-TiO2 interface, and subsequently traversing the TiO2 while suffering a spatially invariant loss to volume recombination. Most models of photocatalysis presuppose that light absorption occurs throughout the bulk of the semiconductor. If a depletion region exists near the surface, standard models assume the surface photocurrent comprises two components. One component includes the drift of photocarriers generated within the region and swept to the surface by the local field. Typically all carriers such carriers are assumed to arrive at the surface39,40. The second component includes carriers generated outside the depletion region that manage to diffuse into it. For the present 17



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heterojunction with a non-absorbing overlayer, however, all photocarriers originate at one boundary of the film. For such a film, volume recombination would occur while carriers move across the film regardless of the transport mechanism – drift or diffusion. Hence the functional form of the decay cannot discriminate among transport mechanisms, and guidance must be sought from estimates of the decay length. For hole migration to be dominated by diffusion, λ would equal the minority carrier diffusion length Lp. Estimates for Lp typically presuppose a field-free semiconductor bulk. From the suggestion of Butler for TiO2

39

, Lp for a diffusion-controlled bulk recombination process

obeys: ‫ܮ‬௣ = ൬

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ଵൗ ଶ

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ቀ

ఌఌ೚ ௞்

ସగ௤ మ ேವ

ቁ

ଵൗ ଶ

(1)

where µe and µp respectively denote the mobility of electrons and holes. These mobilities in amorphous TiO2 does not appear to have been reported, so we employ corresponding behavior in amorphous Si and anatase TiO2 as guides. In amorphous Si, µe is about two thirds of µp 41. In anatase, electron and hole mobilities differ by only a few percent42. For rough estimation of Lp in amorphous TiO2, we therefore set µe = µp. The dielectric constant, ε, for amorphous TiO2 has been reported to have values of 10 43, 13.7 44, 18 45 and 33 46. With ε chosen to have a mid-range value of 18, and Nd estimated to be ~ ½ an order of magnitude lower than typical anatase TiO2 (based upon the measured resistivity), Eq. (1) yields an Lp value of about 5 nm. Although this number is nearly an order of magnitude smaller than the observed value of λ, Eq. (1) presupposes a simple Langevin model for recombination within field free bulk material. The Langevin framework fundamentally requires an electron, rather than a donor, concentration to estimate recombination rate47. Within a depletion region, the electron concentration falls below the donor concentration, which thereby increases the corresponding value of Lp − perhaps by an order of magnitude. Thus, it seems quite possible that photohole transport is primarily diffusive. 18


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For hole migration dominated by drift, λ would equal the drift distance of holes in the presence of an electric field E. In anatase, this distance is typically on the order of several micrometers40, and is orders of magnitude larger than Lp for diffusion.

We cannot easily

estimate the drift distance in amorphous TiO2, as values for the hole mobility and lifetime would be required. However, it seems plausible that the drift distance would considerably exceed Lp as in anatase. If so, the estimate given above for Lp would place the drift distance well above the measured value of λ. It is noteworthy that the amorphous TiO2 in this work exhibits considerable photoactivity.

Indeed, the photoactivity is comparable to previously reported work with

heteroepitaxial anatase TiO2 on SRO16. Published literature indicates that volume-absorbing amorphous TiO2 has much lower photocatalytic activity than anatase48. A principal cause of this phenomenon is that anatase TiO2 has an invariant43 bandgap of 3.2 eV, while the bandgap of amorphous TiO2 tends to be greater by at least 0.2 eV. The gap is indirect in both49,50,51. The difference in bandgap causes the amorphous TiO2 to absorb significantly less light, producing fewer photogenerated carriers. In the heterojunction case, the SRO underlayer absorbs the light, and injects photocarriers at about the same rate for both amorphous and anatase SRO. Loss of those photocarriers by recombination apparently does not greatly differ between the amorphous and anatase phases. Figure 7 shows that exposure of the heterostructure to ultraviolet and visible light simultaneously has the counterintuitive effect of greatly depressing the photocatalysis rate. Related phenomena have been reported for TiO2/InP and TiO2/CaFe2O4 heterojunction photoelectrodes52,53.

In the latter case, ultraviolet illumination decreases the rate of H2

production by water splitting by increasing the amount of charge residing in the TiO2 surface trap states. This extra charge produces an electric field that retards the arrival of holes at the surface. The absorption of light by the TiO2 produces additional minority carriers closer to the 19



surface than the SrRuO3. The same absorption also produces additional majority carriers which are then free to recombine with minority carriers from the SrRuO3 or affect the material in other ways. Given the large energy barriers for electron and hole transfer from SRO to TiO2, questions remain about how these carriers are transported across the TiO2/SrRuO3 interface. Hot electrons created within the SRO have been proposed for the electron transfer16,17 based upon study of other SRO-TiO2 electronic devices and remains the most plausible explanation. Concomitant generation of hot holes within the SRO is also a possibility, although no direct evidence exists in the literature for this effect. Figure 8 shows a qualitative schematic band diagram of the heterojunction sketching the carrier generation and transfer. The presence of a defect band at a suitable energy level within the TiO2 band gap would aid such a phenomenon. This defect state causes the TiO2 to act as a leaky dielectric54 for minority carriers, in which holes transport through a defect band located near mid-gap. Such a model rationalizes hole transport across Si/SiO2/Si55 as well as a number of Si/TiO2 heterostructures56,57. A future effort to directly observe such defects using Ballistic Electron Emission Microscopy (BEEM) may offer useful confirmation of this model as well as offer additional evidence for or against the presence of hot carriers reaching the TiO2 free surface. Direct tunneling of thermal holes though the interfacial barrier adequately explains the photocatalytic activity observed for TiO2 capping layers on ZnO58. However, this mechanism does not seem likely in the SRO-TiO2 case, as tunneling path lengths for holes in TiO2 have an upper limit of 2 nm59 and the expected barrier height for holes is over 2.5 eV. The tunneling limit falls far below the TiO2 thicknesses employed here leaving the leaky dielectric defect as the more plausible explanation.

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Figure 8: Proposed qualitative band diagram for TiO2/SrRuO3 heterojunction under illumination. The red line represents hot electron transport, the blue line represents defect mediated minority carrier transport, and the dashed black line represents recombination. The improved optical absorption of SRO in the amorphous phase compared to crystalline probably traces to changes in the density of states. The band structure of SRO is complicated and not easily represented on simple band diagrams. However, by analogy to conventional amorphous semiconductors, the relevant band edges likely smear somewhat in the amorphous phase. Fortunately, the amorphous SRO retains most of the crystalline material’s electrical conductivity – lower by only a factor of three. For comparison, amorphous silicon exhibits a conductivity up to six orders of magnitude60 below that of the polycrystalline form. The improved optical absorption partly compensates the loss in electrical conductivity, making the amorphous SRO quite suitable for heterojunction use. From a technological perspective, the amorphous form offers significant advantages in processing cost. Both the amorphous and crystalline heterojunctions exhibit rates at or above previously published results16. This performance is particularly encouraging considering that the TiO2 in the present work is amorphous instead of crystalline. Creation of amorphous TiO2, especially in thin film form, should offer advantages in processing cost.

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Conclusion The SRO-TiO2 thin film heterostructure overcomes many of the limitations that seem to inhere in such configurations. The present work suggests that hot carriers play an important role, but only in transferring electrons from the correlated metal oxide into the TiO2. Holes are believed to be transported across interfacial barriers primarily by a leaky dielectric mechanism.

Once the carriers enter the TiO2, the photocatalytic rate exhibits a strong

dependence upon thickness that may attributed to diffusive transport even in the presence of modest built-in electric fields.

Furthermore, amorphous materials retain many of the key

advantages of the crystalline forms. Whether such findings generalize to other correlated metal oxide substrates remains unclear for now. Improved photocatalytic rates may be obtained by strategies to improve the minority carrier lifetime in the TiO2. Manipulation of the internal electric fields within the TiO2 will be challenging as the magnitude of the internal field varies only slightly between heterostructured and bulk TiO2.

Acknowledgements This work was partially supported by the National Science Foundation (DMR 13-06822 and DMR 17-09327) as well as the Strategic Research Initiatives Program in the College of Engineering at the University of Illinois. X-Ray, electrical, optical, and photocatalytic characterization was performed at the Center for Microanalysis of Materials at the Frederick Seitz Materials Research Laboratory University of the Illinois at Urbana-Champaign. D.E.B. gratefully acknowledges fellowship support from the Dow Chemical Company. We thank Tiffany Kaspar for assistance interpreting the XPS results.

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TOC Graphic

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Photocarrier Transport Mechanisms in Amorphous and Epitaxial TiO2

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