A Periodically Ordered Nanoporous Perovskite Photoelectrode for

b Department of Materials Science and Engineering, University of Central Florida, 4000 Central Florida Blvd., Orlando,. Florida, 32816, USA. c Departm...
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A Periodically Ordered Nanoporous Perovskite Photoelectrode for Efficient Photoelectrochemical Water Splitting Li Shi, Wei Zhou, Zhao Li, Supriya koul, Akihiro Kushima, and Yang Yang ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.8b03940 • Publication Date (Web): 13 Jun 2018 Downloaded from http://pubs.acs.org on June 14, 2018

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A Periodically Ordered Nanoporous Perovskite Photoelectrode for Efficient Photoelectrochemical Water Splitting Li Shi a, Wei Zhou c, Zhao Li a,b, Supriya Koul b, Akihiro Kushima b, Yang Yang a,b* a

NanoScience Technology Center, University of Central Florida, 4000 Central Florida Blvd., Orlando, Florida, 32816, USA. b Department of Materials Science and Engineering, University of Central Florida, 4000 Central Florida Blvd., Orlando, Florida, 32816, USA. c Department of Applied Physics School of Science, Tianjin University, Tianjin, 300354, P. R. China. *

Corresponding Author

[email protected]

Supporting Information Placeholder ABSTRACT: Non-metallic materials with localized surface plasmon resonance (LSPR) have a great potential for solar energy harvesting applications. Exploring non-metallic plasmonic materials is desirable yet challenging. Herein, an efficient non-metallic plasmonic perovskite photoelectrode, namely SrTiO3, with a periodically ordered nanoporous structure showing an intense LSPR in the visible light region is reported. The crystallinecore@amorphous-shell structure of the SrTiO3 photoelectrode enables a strong LSPR due to the high charge carriers density induced by oxygen vacancies in the amorphous shell. The reversible tunability in LSPR of the SrTiO3 photoelectrode was observed by oxidation/reduction treatment and incident angle adjusting. Such non-metallic plasmonic SrTiO3 photoelectrode displays a dramatic plasmon-enhanced photoelectrochemical (PEC) water splitting performance with a photocurrent density of 170.0 µA cm−2 under visible light illumination and a maximum incident photon-to-current-conversion efficiency (IPCE) of 4.0% in the visible light region, which are comparable to the state-of-the-art plasmonic noble metal sensitized photoelectrodes.

Keywords: perovskite; oxygen vacancy; core-shell structure; plasmonic; photoelectrochemical water splitting

Solar energy harvesting is considered as one of the most promising ways to satisfy the ever-growing demand for human beings as the energy from the sun is abundant and clean.1, 2 Hydrogen production via photoelectrochemical (PEC) splitting of water has been developed as a potential technology for solar energy harvesting.1 The development of efficient photoanodes is desperately needed for practical PEC applications. However, most of the re-

cently developed photoanodes still suffer from unsatisfactory solar energy conversion efficiencies, largely due to their insufficient light absorption, low charge separation, and transfer efficiencies.3 One effective strategy for expanding the light absorption and improving the efficiencies is to combine the photoanodes with plasmonic noble metals due to the localized surface plasmon resonance (LSPR).4-7 For example, dramatically plasmonenhanced PEC water splitting performances have been achieved over some photoanodes, such as TiO2, BiVO4, and ZnO, via decorating them with plasmonic Au nanostructures.4, 8 However, the plasmonic metals, especially noble metals, are rare and expensive, which would limit their wide applications. Recently, LSPR has also been found in semiconductors and transition-metal oxides because of the doping induced high charge carriers density (by the formation of vacancies or doping with exotic atoms).9, 10 The underlying physics of LSPR in the defective semiconductors are very similar to the plasmonic metals, which result from the interaction of the free charge carriers with the light of appropriate frequency.11, 12 Because some semiconductors are inherently more stable and favorable for solution-based catalytic reactions than metals, the non-metallic plasmonic materials should be considered as the promising alternatives to the plasmonic metals in PEC water splitting.13, 14 SrTiO3, a well-known perovskite metal oxide, has received considerable attention because of its potential applications in the photo(electro)catalytic water splitting, batteries, nonvolatile memory, and solar cells.15-18 SrTiO3 even becomes a more favorable candidate for photocatalysis than TiO2 in some cases because its conduction band position is more negative than that of TiO2. However, the wide bandgap (3.2 eV) makes SrTiO3 active under UV light irradiation, which occupies no more than 5% of the solar light energy.19 Thus, numerous works have mainly focused on expanding the light response of SrTiO3 by means of elemental doping and combining with other visible light active photosensi-

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tizers.20, 21 Creating LSPR induced light absorption could be another route to expand the light absorption of SrTiO3, effectively enhancing the solar-to-energy conversion efficiency. However, so far, this has not been experimentally validated yet. Inspired by pioneer works on plasmonic metal oxides, creating oxygen vacancies in SrTiO3 could be an effective method to introduce abundant free charges in SrTiO3, and eventually, it is possible to have obvious LSPR property. Because of many favorable features, SrTiO3 is a good candidate to be applied as a photoelectrode in the PEC water splitting.22-24 Many the state-of-the-art SrTiO3 photoelectrode for PEC water splitting are still made in the form of either bulk or powder materials, which would suffer from the high recombination rates of electron/hole pairs because of the existence of many grain boundaries and poor particle-to-particle contact when they are transferred onto the conducting substrate.25-27 An efficient method to solve the aforementioned problem of bulk or powder materials is to directly fabricate conducting substrate-anchored and selforiented thin-film photoelectrodes.22 Moreover, the higher PEC performance of photoelectrode can be obtained from highly porous nanostructures, which provide the high surface area, the short diffusion length for charge carriers, and favorable charge transport pathways for chemical reactions.28 Although there are some investigations of the fabrication of porous SrTiO3 films,22, 29-32 a highefficiency plasmonic SrTiO3 thin-film catalyst is still not yet been discovered.22, 32 Herein, a non-metallic plasmonic SrTiO3 photoelectrode with a periodically ordered nanoporous structure was hydrothermally converted from TiO2 nanoporous film followed by a reduction treatment. The as-prepared reduced SrTiO3 (r-SrTiO3), having a crystalline-core@amorphous-shell structure with abundant oxygen vacancies in the amorphous shell, can absorb visible light due to the LSPR property. The reversible tunability in the plasmonic resonance of r-SrTiO3 can be observed through simple oxidation/reduction process and adjusting the incident angles. Such rSrTiO3 photoelectrode exhibits a dramatic enhanced PEC water splitting performance than pristine SrTiO3 photoelectrode under visible light irradiation.

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(Figure S2). Transmission electron microscopy (TEM) of rSrTiO3 nanoporous film (Figure 1f) shows a honeycomb-shaped morphology, which is consistent with the SEM. The high-angle annular-dark-field (HAADF) scanning transmission electron microscopy (STEM) image and the corresponding elemental mapping of r-SrTiO3 show a uniform distribution of Ti, Sr, and O, further confirming the formation of SrTiO3 (Figure 1g). Figure 1h shows the high-resolution TEM (HRTEM) of rSrTiO3, indicating that a disordered layer of 4 nm thick is apparently formed on its surface. Moreover, the core of r-SrTiO3 is still highly crystalline with a measured lattice spacing of 0.27 nm, which is consistent with the cubic SrTiO3 (110) plane.26 The HRTEM confirms that the r-SrTiO3 nanoporous film has crystalline-core@amorphous-shell structure. During the preparation of rSrTiO3, NaBH4 was used as an efficient oxygen scavenger by thermally decomposing and generating active hydrogen.35 The active hydrogen has a strong reductive capability, which can react the oxygen atoms on the surface of SrTiO3, thus producing disordered (oxygen vacancies) surface layer. As a comparison, the pristine SrTiO3 without reducing treatment is highly crystallized, as observed from the clear lattice feature displayed in the HRTEM image (Figure S1c).

Results and Discussion The typical synthetic procedure of plasmonic r-SrTiO3 photoelectrode is presented in Figure 1a. The highly-ordered TiO2 nanoporous film was firstly obtained by Ti anodization.33 Scanning electron microscopy (SEM) image reveals that the as-obtained TiO2 nanoporous film shows honeycomb-shaped morphology with a pore size of about 60 nm and the film thickness of 120 nm (Figure 1b, d). Then the as-obtained TiO2 nanoporous film was hydrothermally converted to SrTiO3 nanoporous film followed by annealing treatment. These 120 nm thick TiO2 nanoporous film is rationally chosen in this work because it can be completely converted to SrTiO3 without destroying the porous structure. As shown in Figure S1, the obtained SrTiO3 film shows a highly porous structure, which is very similar to the annealed TiO2 nanoporous film. However, the cavity wall of SrTiO3 (~20 nm) becomes thicker compared with TiO2 nanoporous film (~10 nm), due to the phase transition and lattice expansion.34 Finally, the plasmonic r-SrTiO3 nanoporous film was obtained by reducing the obtained SrTiO3 nanoporous film with NaBH4 in the N2 atmosphere. The r-SrTiO3 nanoporous film shows almost the same morphology as the SrTiO3 nanoporous film (Figure 1c, e). Energy-dispersive spectroscopy (EDS) analysis indicates that only Ti, O, Sr elements are detected from the r-SrTiO3 nanoporous film

Figure 1. (a) Schematic illustration of the preparation procedures of r-SrTiO3. (b, c) SEM images of TiO2 and r-SrTiO3, respectively. (d, e) Cross-sectional SEM images of TiO2 and r-SrTiO3, respectively. (f) TEM image of r-SrTiO3. (g) STEM image and corresponding elemental mapping of r-SrTiO3 (left to right: Ti, Sr, O). The yellow dashed parts in (g) indicate nanocavities. (h) HRTEM image of r-SrTiO3. Scale bars, (b, c) 200 nm, (d, e, g) 100 nm, (f) 50 nm, (h) 5 nm. The crystalline phases of the materials were recorded by X-ray diffraction (XRD). In contrast to the previous reports which usually contained incompletely converted TiO2 via the similar hydrothermal method by using micron-sized TiO2 nanotubes/nanorods as precursors,32, 36-38 the obtained SrTiO3 nanoporous film in this study showed excellent phase purity, indicating a complete phase conversion. The properties of the rationally designed TiO2 film, such as porous structure, ultrathin cavity wall (~10 nm) and thin film thickness (~120 nm), are beneficial for the complete conversion to SrTiO3. The diffraction peaks are ascribed to metallic Ti

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and cubic SrTiO3 phase (JCPDS Card No. 035-0734),26 and no XRD peaks of TiO2 or other composition were found (Figure 2a). The high purity of SrTiO3 nanoporous film was further confirmed by Raman spectra. As shown in Figure S3, the SrTiO3 demonstrates two main regions of Raman peaks in 250–400 cm-1 and 600–750 cm-1, which are in line with the standard Raman spectrum of SrTiO3.39 As for the r-SrTiO3, the XRD peaks and Raman peaks exhibit slightly broadening and intensity decreasing compared with the pristine SrTiO3, which are owing to the existence of oxygen vacancies. In contrast, direct annealing of the TiO2 nanoporous film at 450 oC only leads to the production of crystalline TiO2, which is revealed by the XRD and Raman spectra (Figure S4).

Figure 2. (a) XRD and (b) ESR of SrTiO3 and r-SrTiO3. (c) Density of states (DOS) and (d) charge density distribution of rSrTiO3. Ov refers to the oxygen vacancy. (e) UV-vis spectra and (f) SERS spectra of SrTiO3 and r-SrTiO3. The formation of oxygen vacancy in the r-SrTiO3 nanoporous film can be evidenced by X-ray photoelectron spectroscopy (XPS). The high-resolution XPS of Ti 2p, Sr 3d and O 1s of SrTiO3 and r-SrTiO3 nanoporous films are shown in Figure S5. The O 1s high-resolution XPS of r-SrTiO3 are fitted into two Gaussian peaks centered at 529.6 eV and 531.5 eV, while only one peak at 529.6 eV are found in the pristine SrTiO3. The lower binding energy located at 529.6 eV is ascribed to the lattice oxygen, while the higher binding energy at 531.5 eV is attributed to the oxygen vacancies.40 The oxygen vacancy content of 12.5% was estimated by calculating the integral areas of the corresponding peaks. The Ti 2p3/2 XPS spectra are shown in Figure S5b. Notably, compared with pristine SrTiO3, the r-SrTiO3 exhibits an additional shoulder at lower binding energy, which can be ascribed to Ti3+.41 The formation of Ti3+ indicates the oxygen vacancies in the r-SrTiO3, leading to the LSPR in the r-SrTiO3. The Sr 3d XPS of SrTiO3 and r-SrTiO3 are almost same except that the 3d peaks of r-SrTiO3 becomes broadening, which is due to the

surface disordering created by the oxygen vacancies.41 As an effective technique for probing such oxygen defects, electron spin resonance (ESR) spectroscopy was adopted to provide more direct evidence. As displayed in Figure 2b, the r-SrTiO3 shows a ESR peak at g-value of 2.002, which is ascribed to paramagnetic oxygen vacancies,42 while no ESR signal was detected for the pristine SrTiO3. The presence of oxygen vacancies may affect the electronic structure of SrTiO3, and this can be proved by applying the density functional theory (DFT) calculation. As shown in Figure 2c, the region close to the Fermi level (Ef) of r-SrTiO3 is made up of Ti3d orbitals, presenting a metallic character rather than semiconducting property. Moreover, the free electron distribution probed by calculating the charge density reveals that the free electron density around Ti atoms with oxygen vacancy is much higher than that of Ti atoms without oxygen vacancy (Figure 2d). Due to large amounts of free electrons induced by oxygen vacancies, rSrTiO3 has the possibility to own strong LSPR effect. The optical properties of the materials were investigated by the UV-vis absorption spectra, as shown in Figure 2e. The r-SrTiO3 shows an obviously strong visible light absorption peak centered at 516 nm. This visible light response is obviously different from intrinsic band gap excitation of SrTiO3 whose absorption edge located at around 378 nm (Figure 2e). This distinct optical behavior in the visible region may be assigned to the LSPR effect, which is induced by the formation of free electrons on the surface of the r-SrTiO3. The abundant oxygen vacancies entail localized electrons with an appreciable concentration on the surface of the r-SrTiO3, which would undergo collective oscillations when excited with light of appropriate frequency and support such observed LSPR. In sharp contrast, due to the absence of the free electrons, the pristine SrTiO3 does not exhibit LSPR effect, that is, no absorption peaks can be observed in the visible region (Figure 2e). Surface-enhanced Raman scattering (SERS) was employed to further demonstrate the local electromagnetic field enhancement in the r-SrTiO3. A common probe molecule, Rhodamine 6G (Rh6G), was used to examine the performance of the r-SrTiO3 as SERS substrate. As shown in Figure 2f, the Raman scattering peaks are clearly visible for the r-SrTiO3. The peak at 1180 cm−1 can be attributed to the N−H in-plane vibration, while the other peaks are assigned to the C−C breathing modes. All the peaks are consistent with the SERS of the Rh6G.43 For comparison, when using the pristine SrTiO3 without LSPR effect as the substrate material, no SERS of Rh6G were obtained (Figure 2f). The strong SERS signal on the r-SrTiO3 substrate demonstrates the plasmonic polarization presented on the r-SrTiO3 surface. One intriguing feature of r-SrTiO3 is that its LSPR induced light absorption shows reversible tunability upon oxidation by O2 in air and reduction by NaBH4 in N2, as shown in Figure 3a-b. Figure 3a shows the evolution of the visible light absorption spectra of the as-prepared r-SrTiO3 after exposure to oxidation in air at 400 oC. It is obvious that the r-SrTiO3 exhibits a redshift of absorption peak position and decrease of peak intensity after gradual oxidation by O2 in the air, evolving from a peak centered at 516 nm to a peak centered at 567 nm. Eventually, the LSPR effect of r-SrTiO3 disappeared after oxidation by O2 for 8 h. It should be noted that the oxidation of r-SrTiO3 would decrease the free carrier concentration, hence leading to the redshift of LSPR absorption peak to a longer wavelength, and finally losing LSPR effect. Consistent with the tendency of the LSPR effect influenced by O2 oxidation, the SERS activity decreases with the increase of the oxidation time (Figure S6). Furthermore, the LSPR effect of the oxidized r-SrTiO3 could be recovered through reduction treatment

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by NaBH4. As shown in Figure 3b, after NaBH4 reduction, the reactivated r-SrTiO3 displayed an intense LSPR induced light absorption at 520 nm, approaching the original r-SrTiO3. Another significant finding is that the LSPR absorption properties of the as-prepared r-SrTiO3 nanoporous film can also be tuned by changing the incident angle. Figure 3c displays the incident angle-dependent light absorption spectra of r-SrTiO3 nanoporous film. The angle-resolved absorption measurements were performed by using an integrating sphere with a rotating center mount (RTC-060-SF, Labsphere Inc.). A schematic is shown in Figure 3d, and by continuously rotating the r-SrTiO3 photoelectrode from θ= 0o to 75o in an interval of 15o, absorption spectra from the same dimer structure were respectively recorded. It is obvious that the LSPR absorption properties are dependent of the incident angle, that is, the absorption peak position changes from 516 nm to 592 nm accompanied by the decrease of the intensity gradually with the increase of incident angle from 0o to 75o. The absorption peak position is able to shift back when the incident angle decreases to 0o. The absorbance fluctuation may be due to the nonuniformity of oxygen vacancies at the inner wall of rSrTiO3 nanoporous film. Moreover, the angle-resolved SERS on r-SrTiO3 tested by mapping the SERS signal at a different angle of incident laser shows that the r-SrTiO3 exhibits LSPR properties in a broad range of incident angles (Figure S7). The tunability and reversibility of the LSPR feature obtained through oxidation/reduction process and adjusting the incident angle provide a flexibility for utilizing plasmonic r-SrTiO3 with a targeted LSPR wavelength, and show great potential for PEC applications.

Figure 3. (a) Visible light absorption spectra of r-SrTiO3 after exposure to oxidation in air at 400 oC over time in ambient conditions. (b) Visible light absorption spectra of the original r-SrTiO3 and recovered r-SrTiO3. The recovered r-SrTiO3 was obtained by NaBH4 reduction of the oxidized r-SrTiO3. (c) Incident angledependent light absorption spectra of r-SrTiO3. (d) Schematic illustration of the incident angle-dependent light absorption measurement. The PEC water splitting performances were investigated using the pristine SrTiO3 and r-SrTiO3 photoelectrodes. Figure 4a shows linear-sweep voltammograms (LSV) sweeps for pristine SrTiO3 and r-SrTiO3 photoelectrodes in dark and light under AM 1.5G illumination. Both photoelectrodes display negligible dark currents with respect to their photocurrents. Obviously, the photocurrent of r-SrTiO3 photoelectrode under illumination is much higher

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than that of the pristine SrTiO3 photoelectrode. Furthermore, the onset potential of photocurrent shows a slight blueshift from 0.31 V vs. RHE for pristine SrTiO3 to 0.24 V vs. RHE for r-SrTiO3. The higher photocurrent density and lower onset potential indicate the superiority of r-SrTiO3 for PEC water splitting compared with SrTiO3. To examine the photoresponse of SrTiO3 and r-SrTiO3 photoelectrodes over time, I−t curves were measured at 1.23 V vs RHE (Figure 4b) under AM 1.5G irradiation. Both photoelectrodes have obvious photoresponse in chopped light cycles, which are reproducible in several on/off cycles with an almost constant photocurrent. The r-SrTiO3 photoelectrode exhibits the photocurrent up to 1.71 mA cm-2 at 1.23 V vs RHE under AM 1.5G irradiation, which is about 3.1 times higher than pristine SrTiO3 (0.56 mA cm-2). The r-SrTiO3 photoelectrode also shows good stability after long time reaction (Figure S8). The obtained PEC performances of r-SrTiO3 photoelectrode are much higher than the stateof-the-art SrTiO3 photoelectrodes (Table S1).22-24, 44, 45 To test the PEC performance under visible light irradiation of the non-metallic plasmonic r-SrTiO3 photoelectrode, PEC measurements were further performed. Firstly, I−t curves were measured under illumination of visible light (AM 1.5G combined with a L-42 cutoff filter, λ≥420 nm) at the potential of 1.23 V vs. RHE (Figure 4c). The pristine SrTiO3 photoelectrode shows a negligible amount of photocurrent density, owing to the large bandgap of SrTiO3 that cannot be excited under visible light. In sharp contrast, a greatly enhanced photocurrent density was obtained from the plasmonic r-SrTiO3 photoelectrode, with a reproducible photocurrent density of 170.0 µA cm-2. To evaluate the relationship between LSPR effect and photocurrent density on the r-SrTiO3 photoelectrode, incident photon-to-current-conversion efficiency (IPCE) were measured and calculated in the light region between 430 nm to 730 nm. As shown in Figure 4d, the shape of the IPCE active spectra of the r-SrTiO3 photoelectrode exhibits similarity compared with that of the LSPR absorption spectrum of the rSrTiO3. These results provide a clue that the LSPR effect of the rSrTiO3 makes a contribution to the PEC performance in visible light region. The PEC performance of r-SrTiO3 photoelectrode decreases after air oxidation (Figure S9), which further demonstrates that the LSPR effect promotes the PEC performance of rSrTiO3 photoelectrode. The enhanced PEC performance by plasmonic SrTiO3 under visible light in this work can even be comparable to the reported plasmonic Au nanoparticles enhanced PEC performances,46-49 however, the cost-effective properties of SrTiO3 material should be more promising in practical applications due to its lower cost than plasmonic noble metals. Figure 4e shows a schematic illustration of the mechanism for the LSPR effect of r-SrTiO3 for PEC splitting of water. Under visible light, the plasmonic SrTiO3 disordered shell essentially acts as a photosensitizer to absorb the resonant photons. Then the LSPR-induced energetic hot electron-hole pairs transfer the hot electrons to the conduction band of the adjacent crystalline SrTiO3 core, which flows to counter electrode (Pt photocathode) through external circuit for water reduction reaction. The remained energetic positive charges (hot holes) on the disordered shell participate in water oxidation. It is noteworthy that the enhancement of photocurrent density (0.17 mA cm-2) on the r-SrTiO3 photoelectrode from LSPR effect under visible light irradiation is far less than the overall enhancement of the photocurrent density under irradiation of the AM 1.5G (1.15 mA cm-2; from 0.56 to 1.71 mA cm-2), which is due to the electronic structure modification after formation of disordered layer on the surface of r-SrTiO3 photoelectrode. To validate this

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assumption, electrochemical impedance spectroscopy (EIS) and Mott–Schottky (MS) plots were recorded on the pristine SrTiO3 and r-SrTiO3 photoelectrodes. Figure S10 shows the EIS curves for the pristine SrTiO3 and r-SrTiO3 photoelectrodes. A smaller semi-cycle is observed in the EIS of r-SrTiO3 photoelectrode under illumination than the pristine SrTiO3 photoanode, meaning the interfacial resistance at the disordered SrTiO3/electrolyte interface is much smaller than the that of the crystalline SrTiO3/electrolyte interface. The Mott-Schottky (MS) plots in Figure S11 shows that the r-SrTiO3 photoelectrode shows a substantially smaller slope than SrTiO3 photoelectrode, suggesting increased carrier densities, leading to a lower resistance and higher charge transfer rate.22 The enhanced charge transfer for the r-SrTiO3 photoelectrode is considered as the major reason for the PEC performance enhancement in the UV light region. Under UV-visible light illumination (AM 1.5G), both crystalline SrTiO3 core and amorphous SrTiO3 shell could be exited, the hot holes remained on the amorphous SrTiO3 shell, together with the photogenerated holes transferred from crystalline SrTiO3 core could execute water oxidation, and simultaneously, the photogenerated electrons from the crystalline SrTiO3 core and hot electrons from the amorphous SrTiO3 shell transferred to the Pt electrode through external circuit for water reduction (Figure S12).

Figure 4. (a) Linear sweeps voltammogram (LSV) of SrTiO3 and r-SrTiO3 photoelectrodes. (b, c) Transient photocurrent responses of SrTiO3 and r-SrTiO3 photoelectrodes at 1.23 V vs. RHE under AM 1.5G illumination without and with a L-42 cutoff filter (λ≥420 nm), respectively. (d) IPCE spectra in the region of 430– 730 nm at 1.23 V vs. RHE. (e) Schematic illustration of the proposed mechanism of r-SrTiO3 for PEC water splitting under visible light.

Conclusion In summary, a plasmonic perovskite semiconductor, namely rSrTiO3 photoelectrode, was rationally developed. The r-SrTiO3 photoelectrode has a crystalline-core@amorphous-shell structure

with abundant oxygen vacancies in the amorphous shell, which induces an intense LSPR in the visible-light region. The r-SrTiO3 photoelectrode undergoes reversible tunability in its plasmonic resonance through oxidation/reduction treatment and incident angle adjusting, which provides flexibility for PEC applications in a targeted LSPR wavelength. Under light irradiation, the r-SrTiO3 photoelectrode showed much higher PEC water splitting activity than the pristine SrTiO3 photoelectrode. This work illustrates a type of non-metallic plasmonic perovskite semiconductor with plasmon-enhanced activity toward catalytic reactions. The material developed in this study will open a paradigm for advanced photoelectrode rational design in the field of solar-energy harvesting. Experimental Section Anodic growth of the TiO2 nanoporous film. Titanium foils (0.05 mm thick, 99.7% purity, MTI Corporation) were firstly cleaned by ultrasonication in acetone and ethanol solution for 20 min and then dried in air. The highly-ordered TiO2 nanoporous film was obtained through Ti foil anodization in a solution of 3 M HF/H3PO4 (98%, Alfa Aesar, US) using a constant voltage of 10 V at 80 oC for 4 h. After anodization, the sample was washed with ethanol several times and then dried in air. The as-obtained TiO2 nanoporous film was then directly used as a precursor to fabricate SrTiO3. Synthesis of SrTiO3 nanoporous film. In a typical synthesis, 0.02 M, 60 ml Sr(OH)2 was put into a Teflon-lined stainless steel autoclave (80 ml), then a 3 cm×3 cm pre-obtained TiO2 nanoporous film was added into the above solution with a suitable angle against the inner wall of the autoclave. Then, the autoclave was sealed tightly and put into a pre-heated electric oven (180 oC) and kept for 4 hours. After cooling down to the room temperature, the obtained sample was annealed at 450 oC for 1 hour in air followed by washing with 0.05 M HCl and water. Synthesis of r-SrTiO3 nanoporous film. The prepared SrTiO3 nanoporous film was mixed with 0.5 g NaBH4 and heated in a tube furnace at 320 oC in the N2 atmosphere for 2 hours. After cooling down to the room temperature, the obtained sample was washed with 0.05 M HCl and pure water thoroughly. Characterization. The morphology and EDS of the samples were observed by the field-emission scanning electron microscope (FESEM, ZEISS ultra 55) and the transmission electron microscope (TEM, FEI Tecnai F30). The cross-sectional part of the sample for TEM observation was prepared by cutting off the samples via using a Tescan LYRA-3 Model GMH focused ion beam microscope. The chemical states and composition of samples were characterized by X-ray photoelectron spectroscopy (XPS) measurement conducted on a PHI Quantera Scanning X-ray Microprobe. X-ray diffraction (XRD) patterns were obtained by an XRD machine (PANalytical, equipped with a Panalytical X’celerator detector utilizing Cu Kα radiation, and the λ is 1.54056 Å). SERS and Raman spectra were obtained by measuring the samples on a Renishaw InVia Microscope and the excitation laser was set at 532 nm. The ESR spectra were recorded with JEOL JES-FA-200. The light absorption properties of the prepared samples were recorded on the Cary Win UV−visible spectrometer. The incident angle dependent light absorption spectra of r-SrTiO3 photoelectrode were measured using an integrating sphere (RTC-060-SF, Labsphere Inc.) connected to a spectrometer (Ocean Optics, HR2000+), and a collimated tungsten source is an incident upon samples mounted to a rotational center holder, allowing angle-resolved measurements.

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DFT calculations. The DFT calculations presented here were carried out with the Vienna ab-initio Simulation Package (VASP). The electron-ion interactions were employed by Projectoraugmented wave (PAW) combined with GGA-PBE function. The total energy during the calculation was set within 1×10-5 eV/atom by increasing the cut-off energy together with the number of k points. Thus, a cut-off energy at 450 eV, as well as 4×4×4 k points, were finally set for the 2×2×2 supercell, which included 40 atoms. Moreover, we set the energy convergence tolerance to be below the value of 1×10-5 eV/atom. At the same time, the Hellmann-Feynman force working on each atom should be reduced to no more than 0.01 eV/Å, which was obtained by relaxing the lattice and atomic coordinates. The concentration of oxygen vacancies in the calculation model of r-SrTiO3 was set to be 12.5%, which is same as the experimentally obtained in r-SrTiO3. Photoelectrochemical water splitting experiments. Photoelectrochemical experiments were performed in 1 M KOH solution in a three-electrode configuration under simulated AM 1.5G (light intensity: 100 mW cm-2) irradiation. The SrTiO3 and r-SrTiO3 photo-electrodes serve as the working electrodes, and the Ag/AgCl (3 M KCl) and the platinum foil serve as the reference electrode and the counter electrode, respectively. The experimentally obtained potentials vs. Ag/AgCl (3 M KCl) were transferred to the reversible hydrogen electrode (RHE) via using the formula ERHE = EAg/AgCl + 0.059pH + E0Ag/AgCl, where EAg/AgCl is the experimentally obtained potential and E0Ag/AgCl is the constant of 0.209 V at 25 oC for an Ag/AgCl (3 M KCl) reference electrode. The linear-sweep voltammograms (LSV) sweeps were recorded with a scan rate of 10 mV s-1. The I−t curves were recorded under interrupted light illumination (light on/off) at a constant potential of 1.23 V vs RHE. IPCE measurements were performed by utilizing a Zahner CIMPS-QEIPCE system combined with monochromator in 1 M KOH solution with an applied potential at 1.23 V vs RHE. The IPCE can be calculated according to the equation:1 IPCE =



100%

(1)

where J, P, and λ refer to the photocurrent density (mA cm-2), the incident light intensity (mW cm-2) and the monochromatic light wavelength (nm), respectively. Electrochemical impedance spectroscopy (EIS) experiments were conducted by using 1.23 V vs RHE in the frequency range of 105 Hz to 0.1 Hz with an amplitude of 10 mV under the simulated solar light of AM 1.5G irradiation. Mott-Schottky plots were measured at a frequency of 5000 Hz in the dark. The Mott−Schottky equation can be described as follow:22  

=

    



(V－Vfb－ ) 

(2)

where C refers to the capacitance, A refers to the area, e refers to the electronic charge, ε0 and εr refer to the permittivity of vacuum and dielectric constant of a semiconductor, respectively, ND refers to the charge carrier density, V refers to the applied potential, Vfb refers to the flat band potential, κ refers to the Boltzmann’s constant, and T refers to the temperature. Thus, the ND can be obtained from the slope of the linear fit of Mott−Schottky equation.

ASSOCIATED CONTENT ACKNOWLEDGMENT

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This work is funded by the University of Central Florida through a startup grant (20080741). Dr. Li Shi acknowledges the financial support from the Preeminent Postdoctoral Program (P3) at University of Central Florida.

Supporting Information Supporting Information Available: SEM and TEM images of TiO2 and SrTiO3; EDS of r-SrTiO3; Raman and XPS spectra of SrTiO3 and r-SrTiO3; SERS spectra of r-SrTiO3; longtime photocurrent responses of r-SrTiO3; LSV, EIS and MS curves of r-SrTiO3 photoelectrode. This material is available free of charge via the Internet at http://pubs.acs.org.

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