Anomalous Conductivity Tailored by Domain-Boundary Transport in

Feb 21, 2018 - Condensed Matter Physics and Materials Science Department, Brookhaven National Laboratory , Upton , New York 11973 , United States...
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Article Cite This: Chem. Mater. 2018, 30, 1677−1685

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Anomalous Conductivity Tailored by Domain-Boundary Transport in Crystalline Bismuth Vanadate Photoanodes Wenrui Zhang,† Danhua Yan,†,‡ Jun Li,§ Qiyuan Wu,‡ Jiajie Cen,‡ Lihua Zhang,† Alexander Orlov,‡ Huolin Xin,† Jing Tao,§ and Mingzhao Liu*,† †

Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, New York 11973, United States Department of Materials Science and Engineering, Stony Brook University, Stony Brook, New York 11794, United States § Condensed Matter Physics and Materials Science Department, Brookhaven National Laboratory, Upton, New York 11973, United States ‡

S Supporting Information *

ABSTRACT: Carrier transport in semiconductor photoelectrodes strongly correlates with intrinsic material characteristics including carrier mobility and diffusion length, and extrinsic structural imperfections including mobile charged defects at domain boundaries, which collectively determines the photoelectrochemistry (PEC) performance. Here we elucidate the interplay between intrinsic carrier transport, domain-boundaryinduced conductivity, and PEC water oxidation in the model photoanode of bismuth vanadate (BiVO4). In particular, epitaxial single-domain BiVO4 and c-axis-oriented multidomain BiVO4 thin films are fabricated using pulsed laser deposition to decouple the intrinsic and extrinsic carrier transport. In addition to the low intrinsic conductivity that is due to the small-polaron transport within BiVO4 domains, we identify anomalously high electrical conductivity arising from vertical domain boundaries for multidomain BiVO4 films. Local domain-boundary conduction compensates the inherently poor electron transport by shortening the transport distance for electrons diffused into the domainboundary region, therefore suppressing the photocurrent difference between front and back illumination. This work provides insights into engineering carrier transport through coordinating structural domain boundaries and intrinsic material features in designing modulated water-splitting photoelectrodes.

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INTRODUCTION Construction of efficient photoelectrochemical (PEC) devices requires sufficient photon collection, prompt charge carrier separation and transport, fast redox kinetics, and long-term operation stability.1−5 Designing oxygen-evolving photoanodes with relatively low overpotential is particularly challenging because of the slow kinetics associated with four-electron transfer for water oxidation.6,7 Among various semiconductor photoanodes, bismuth vanadate (BiVO4, BVO) exhibits many key advantages as a model system, such as strong visible light absorption and energetically favorable band edge position and relatively long hole diffusion length, therefore attracting intensive research interests.8−12 However, pristine BVO suffers from low electron mobility,13 poor water oxidation kinetics, and excessive electron−hole recombination,14,15 which limit the water-splitting efficiency. Encouraging progress has been recently achieved to improve the water oxidation kinetics by modifying the electrode surface with oxygen evolution reaction (OER) catalysts, such as cobalt-phosphate (Co-Pi),14,16 cobalt oxide,17,18 and NiFe-oxyhydroxides (FeOOH/NiOOH).10 On the other hand, addressing the problem of strong electron © 2018 American Chemical Society

localization in BVO mainly relies on nanostructuring to shorten the carrier transportation length10,19,20 or doping to enhance the photoelectrode conductivity.11,14,16,21 Mechanistic study of electronic structure22 and solid-state carrier transport13 suggest that the conduction band minimum comprising hybridized V 3d/O 2p orbitals favors undesired localization of electrons and creates small polarons with poor electron mobility. Although previous studies provide important insights on understanding the electronic behavior, direct experimental studies that can establish the intimate connection between carrier transport and PEC performance are lacking. Moreover, structural imperfections, such as domain/grain boundaries and point defects, widely exist in most photoelectrodes regardless of being thin films or nanostructures. Such features can also actively contribute to carrier transport and PEC activity, but receive much less attention in previous theoretical or experimental studies. Received: December 6, 2017 Revised: February 20, 2018 Published: February 21, 2018 1677

DOI: 10.1021/acs.chemmater.7b05093 Chem. Mater. 2018, 30, 1677−1685

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Figure 1. (a) Schematic single-crystalline BVO/ITO/YSZ photoanode for PEC water oxidation. (b) Tauc plot of a 160 nm thick BVO film over ITO/YSZ determining an optical band gap of 2.60 eV. The top inset is the corresponding optical absorption spectrum versus photon energy. The bottom inset is the optical image of the BVO/ITO/YSZ heterostructure. (c) High-resolution XPS spectra of Bi 4f, V 2p, and O 1s core levels of the BVO film. (d) X-ray absorption near-edge structure of the BVO film collected at V K-edge. (e) Raman spectra of the BVO film on ITO/YSZ with BVO bulk as a comparison.

been demonstrated with combined solid-state carrier transport and electrochemical characterization.

Here we select BVO as a model system to elucidate the strong correlation between intrinsic carrier transport, domainboundary-induced conductivity, and water oxidation behavior. We fabricate crystalline BVO thin films over epitaxial ITO/YSZ as a unique platform for combined solid-state electronic and photoelectrochemical characterization (Figure 1a). The growth of dense single-crystalline BVO thin films has been highly challenging because of its inherent structural complexity, the limited choice of lattice-matching conductive substrates, and the preferred island formation in monoclinic BVO (a = 5.1956 Å, b = 5.0935 Å, c = 11.7044 Å, γ = 90.383).23−25 Direct growth of BVO over single-crystalline yttria-stabilized zirconia (YSZ, cubic, a = 5.145 Å) substrates produces continuous epitaxial single-domain BVO thin films that exhibit the characteristic highly insulating behavior and small-polaron transport. Incorporation of a lattice-matching indium tin oxide (ITO, cubic, a = 10.1269 Å) interlayer between BVO and YSZ enables direct electrical/electrochemical characterization and facilitates c-axis-orientated domain formation inside BVO films, which generates fast electron transport channels and features anomalous high conductivity. Integrated topography and conductive AFM imaging confirm local domain-boundary conduction in crystalline BVO films, which demonstrates free-carrier behavior and leads to the giant conductivity anisotropy. The strong correlation between convoluted intrinsic/extrinsic carrier transport and the PEC behavior has

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RESULTS AND DISCUSSION Crystalline Growth and Microstructure Characterization. The film deposited over ITO/YSZ exhibits the characteristic yellow color of scheelite BVO (Figure 1b, bottom inset). The direct optical band gap for the BVO film is determined to be 2.60 ± 0.05 eV from its optical absorption spectrum (Figure 1b). An indirect transition at 2.52 ± 0.05 eV is also present from the Tauc plot analysis of both singledomain and multidomain BVO films (Figure S1a,b). The slight optical absorption in the sub-band region is associated with structural defects and scattering loss in multidomain BVO films.26 Such sub-band absorption is more suppressed in singledomain BVO because of higher crystallinity and smoother surface (Figure S1c). The measured optical density suggests that a 160 nm thick BVO film absorbs over 90% of the incident light above its band edge. The underlying ITO/YSZ substrate, on the other hand, shows low optical absorption in the wavelength range of interest (350−500 nm) (Figure 1b, top inset). The 1:1 Bi:V stoichiometry in the BVO film is verified by energy-dispersive X-ray spectroscopy (EDX) and X-ray photoelectron spectroscopy (XPS). High-resolution XPS spectra of Bi 4f and V 2p core levels determine their valence states to be respectively 3+ and 5+ (Figure 1c). The O 1s peak shoulder near 533 eV arises from defective oxygen species or 1678

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Figure 2. (a) θ−2θ XRD scans and (b) SAED pattern of c-axis-oriented crystalline BVO over ITO/YSZ. The bare ITO/YSZ XRD scan is provided as a comparison. AFM images with a scan area of 5 × 5 μm2 of (c) multidomain BVO film over ITO/YSZ and (d) single-domain BVO film over bare YSZ. (e) Cross-sectional low-magnification HAADF-STEM image of the BVO film over ITO/YSZ. The inset shows the high-resolution STEM image of the BVO/ITO interface region demonstrating well-ordered lattice matching. (f) Enlarged STEM image demonstrating continuous vertical domain boundary formed in the BVO film. (g) High-resolution STEM image identifying the domain-boundary region formed by [100]-oriented and [010]-oriented domains.

temperature (∼600 °C), the deposition delivers sufficient film crystallinity while maintaining a reasonably smooth and dense morphology. The BVO film over ITO/YSZ consists of multiple domains with domain width ranging from 200 to 500 nm (Figure 2c). The domain height varies within ±25 nm, and the RMS surface roughness (Ra) is 9 ± 2 nm as measured by atomic force microscopy (AFM). When growing on bare YSZ substrate, the BVO film becomes single-domain with a much smaller Ra of 0.5 ± 0.3 nm (Figure 2d). The multidomain microstructure of the BVO film over ITO/YSZ is revealed by the high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM), using a vertical cross-sectional specimen (Figure 2e). The high-resolution STEM image reveals coherent lattice-matched interfaces of BVO/ITO (inset of Figure 2e) and ITO/YSZ (Figure S3). The EDX elemental mapping of a vertical cross-section of the BVO/ITO/ YSZ heterostructure shows clear phase separation with distinct elemental distribution (Figure S3). A closer examination of domain rotation and domain-boundary formation is presented in Figure 2f, revealing the presence of vertical domain boundaries extending across the entire film thickness. Around the domain-boundary region, 90° in-plane rotation occurs between [100]- and [010]-oriented BVO domains (Figure 2g), therefore forming the vertical domain boundary. The analyses of domain-boundary formation from differently in-planeoriented domains are presented in Figure S4 and demonstrate the same feature as a continuous vertical domain boundary throughout the film. Solid-State Transport and Vertical Conductive Channels along Domain Boundaries. The crystalline BVO film over ITO/YSZ is determined as n-type through electrochemical impedance measurement and corresponding Mott−Schottky (1/C2sc versus E) analysis (Figure S5), with a donor density ND

surface hydroxides, which is frequently observed in metal oxides.27 The X-ray absorption near-edge structure (XANES) collected near V K-edge confirms the oxidation state of 5+ for V in the film. The intense pre-edge peak at ∼5467 eV indicates that the V atoms occupy tetragonal sites (Figure 1d). Raman spectroscopy determines the monoclinic polymorph in the crystalline BVO films, as made evident by the symmetric V−O stretching mode at ∼824 cm−1 (Figure 1e).28 X-ray diffraction (XRD) peaks in the θ−2θ scan of the BVO films over ITO/YSZ substrates can be indexed to monoclinic BVO (space group I2/b, PDF 01-074-4893) (Figure 2a).29 The diffraction pattern is dominated by even numbered (00l) diffraction peaks, reflecting c-axis-oriented growth of BVO film on ITO/YSZ substrates. Rocking curve measurements determine a full width at half-maximum of 0.215° for the (004) diffraction peak of BVO on ITO/YSZ, which is slightly larger than that of 0.156° on bare YSZ (Figure S2a). Similar residual strain is seen between single- and multidomain BVO films (Figure S2b). The high-quality crystalline growth BVO on epitaxial ITO/YSZ is also demonstrated by selected area electron diffraction (SAED) pattern (Figure 2b). In the pattern we observe two BVO (004) diffraction spots separated by a small angle (∼2.5°), which suggests domain rotation of highly textured (001)-oriented BVO domains. The epitaxial orientation relationship of BVO relative to ITO and YSZ is further investigated by XRD ϕ-scan measurements, which show that the BVO growth mainly follows the epitaxial constrains from the underlying ITO/YSZ as (001)BVO∥(002)ITO∥(001)YSZ (outof-plane) with either [100] BVO ∥[200] ITO ∥[100] YSZ or [010]BVO∥[200]ITO∥[100]YSZ (in-plane) (Figure S2c,d). It should be noted that surface morphology and smoothness of crystalline BVO films are extremely sensitive to growth conditions and underlying substrates. At optimized substrate 1679

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Figure 3. (a) Out-of-plane and in-plane current−voltage (I−V) characteristics of crystalline BVO films. Insets show the schematic of in-plane and out-of-plane electrical measurement, respectively. (b) Temperature-dependent in-plane conductivity of the BVO film. Inset shows the resistivity fitting plot based on the small-polaron transport deriving a hopping activation energy of 557 meV. (c) Temperature-dependent out-of-plane conductivity of BVO films with a thickness of 160 and 960 nm. Inset shows room-temperature out-of-plane film conductivity as a function of film thickness. (d) Topography z-height and (e) conductive AFM mapping of the BVO film grown on ITO/YSZ. (f) Local I−V characteristics of selected domain boundary and domain regions demonstrating clear domain-boundary conduction. Top set compares relative conductance between domain boundary and neighboring domain interior of three different areas. Bottom inset shows the schematic setup of the conductive AFM measurement.

= (1.68 ± 0.13) × 1018 cm−3 and a flat band potential EFB = 0.15 ± 0.02 VRHE (volts versus reversible hydrogen electrode). The donor likely arises from native point defects such as oxygen vacancies.30 A similar level of donor density has been previously observed for bulk BVO single crystals grown by a floating zone technique.13 Electronic transport within the BVO films, however, exhibits significant anisotropy between out-ofplane and in-plane directions (Figure 3a). The in-plane conductivity measurements are carried out for single-domain BVO films grown on bare YSZ substrates, which reveal an insulating behavior with an in-plane conductivity (σ∥) of merely 5.81 × 10−7 S cm−1. The single-domain BVO is free of domain boundaries and reflects mainly the intrinsic transport behavior for this material. On the other hand, the out-of-plane conductivity (σ⊥) is determined from a multidomain BVO film grown on ITO/YSZ, which contains vertical domain boundaries and features significantly enhanced conductivity. The measured σ⊥ of 5.69 × 10−3 S cm−1 at 300 K is 4 orders of magnitude larger than σ∥. It should be noted that the observed high σ⊥ is unlikely to arise from the elemental diffusion of bottom ITO into BVO. Within its detection limit, atomicresolution STEM-EDX analysis reveals no observable interdiffusion between BVO and ITO. As the detection limit of EDX is roughly 0.5%, we prepare 0.5% (in atom ratio) In-doped and 0.5% Sn-doped BVO films under the same growth condition for comparison. Through transport measurements it is found that the intentional doping of indium or tin to BVO has negligible impact to BVO in-plane film conductivity (Figure S6), thus

ruling out the cation diffusion as a source for the high out-ofplane conductivity. To further understand the large and unexpected conductivity anisotropy, we study the temperature dependence of electron transport in crystalline BVO films, along both directions. For in-plane transport, σ∥ rises by nearly 4 orders of magnitude from 5.81 × 10−7 to 1.61 × 10−3 S cm−1, as the temperature increases from 300 to 500 K (Figure 3b). Such temperature dependence suggests the presence of a thermal activation process, with an activation energy on the order of a few hundred meV, which is consistent to the small-polaron hopping mechanism that has been proposed for electron transport in BVO single crystals.13 According to this mechanism, electrons interact with the surrounding lattice distortion so that their mobility is suppressed significantly, with a temperature (T)

( ) E

dependent conductivity σ(T ) ∝ T −1 exp − k Th , where Eh is B

the hopping activation energy.31,32 Through a ln[σ(T)T]−1/T linear fit (inset of Figure 3b), the activation energy Eh is determined to be 557 meV for the single-crystalline BVO film, which agrees with previous first principle calculation results.33 On the other hand, the out-of-plane conductivity σ⊥ measured from a 160 nm thick BVO film over ITO/YSZ decreases slightly from 5.69 × 10−3 to 5.11 × 10−3 S cm−1 as the film is heated from 300 to 500 K (Figure 3c), which is atypical for conventional semiconductors but resembles a free-carrier (metallic) behavior.34 1680

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Figure 4. Photocurrent density versus potential curves of (a) BVO and (b) CoOx/BVO photoanodes under front and back illumination in pH 7.0 phosphate buffer without and with 1.0 M Na2SO3. The top inset in part b is the low-magnification STEM image showing the ultrathin CoOx layer on BVO surface. The bottom inset in part b is the high-resolution XPS spectrum of Co 2p core level of the CoOx layer. (c) Charge separation efficiency and charge injection efficiency and (d) IPCEs of crystalline CoOx/BVO photoanode measured under back and front illumination at 1.23 VRHE in pH 7.0 phosphate buffer. (e) Schematics assessing the domain-boundary conduction effect on photocarrier transport under front and back illumination in crystalline BVO thin film photoanodes.

different domain-boundary regions as determined from local AFM measurements (inset of Figure 3f) and solid-state electronic measurements (Figure S7b). Clear conduction signal is obtained when the AFM tip is placed on either peak or valley regions of domain boundaries, which rules out the possible influence of contact area difference in the conduction measurement. Indeed, the high conduction behavior mainly arises from interconnected charged defects at the continuous vertical domain boundary. In fact, as we increase the film thickness to suppress the formation of continuously vertical conducting channels, a continuous reduction of σ⊥ is observed (inset of Figure 3c). For a BVO film as thick as 960 nm, the value of σ⊥ is very similar to σ∥, reflecting the intrinsic, low conductivity of BVO that follows the small-polaron hopping mechanism (Figure 3c). The disappearance of domainboundary conducting channels for thicker BVO films is again captured by the conductive AFM mapping. The mapping of a 960 nm thick film reveals a homogeneous nonconductive behavior that is not correlated to the film morphology (Figure S8). Water-Splitting Efficiencies and OER Catalyst Modification. The high quality and unique electronic transport properties of crystalline BVO thin films have a profound impact on its PEC water-splitting behavior, which is first studied by measuring the photocurrent density−potential (J−V) relations

Since cation diffusion has been ruled out as a source for the high out-of-plane conductivity, we propose that the free-carrierlike transport along this direction likely arises from domainboundary leakage that is supported by accumulated charged defects at structural boundaries.35−37 The continuous, vertical domain boundaries in BVO films provide conducting channels for fast electron transport along the out-of-plane direction and yield the high apparent conductivity. Such conductive domain boundaries are different from common grain boundaries in polycrystalline films, as they need to be continuous and vertical along the out-of-plane direction in order to support fast electron transport between the bottom electrode and the aqueous interface. The domain-boundary-supported vertical conduction channels in the BVO film are revealed through integrated topography AFM and conductive AFM imaging. According to the imaging for a 160 nm thick film that has a σ⊥ of 5.69 × 10−3 S cm−1, the current signal is significantly higher at the domain boundaries as compared to the inner region of a domain (Figure 3d,e). The linear behavior of the out-of-plane I−V curve suggests ohmic contact, and the contact resistance is assumed to be negligible between the contact electrodes and BVO films (Figure S7a). The local I−V characteristics at representative domain and domain boundaries further confirm the drastic conduction difference (Figure 3f). It should be noted that the domain-boundary conduction varies little across 1681

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photocurrent enhancement when compared to the bare BVO photoelectrode. The peak energy conversion efficiencies, calculated on an applied-bias photon-to-current metric,41 are enhanced to 0.24% at 0.89 VRHE and 0.18% at 0.86 VRHE for back and front illumination, respectively, which are smaller than state-of-the-art efficiencies (1−2%)10,20 because of the limited surface area in planar thin film electrodes. The onset potential shifts to the cathodic direction by 0.12 V, and a higher fill factor in the J−V plot is observed. Given that the CoOx/BVO photoanode delivers a Jback of 1.95 mA cm−2 and a Jfront of 1.33 mA cm−2 for sulfite oxidation, its ηaq inj is determined at 69.2% and 71.4%, respectively, for the front and back illumination (Figure 4c). The values are more than 2 times larger than those of bare BVO electrodes, which demonstrates that the surface CoOx catalyst on BVO indeed improves the PEC water oxidation activity. For the even thicker CoOx layer, the photocurrent drops under both illumination directions, which may be due to the increased interfacial resistance from the thick catalyst layer. The water oxidation IPCEs are further measured for the CoOx/BVO photoanode. The illumination direction dependence is also observed, but with larger IPCE values. The IPCEb of the CoOx/BVO photoanode increases up to ∼36% for λ < 440 nm, while a maximum IPCEf of ∼28% is observed near λ = 460 nm (Figure 4d). The combined electronic and PEC characterization for single-crystalline BVO thin film electrodes provides valuable insights into the strong correlation between intrinsic carrier transport, extrinsic domain-boundary transport, and watersplitting activity in this material. Compared to the relatively long minority carrier diffusion length (∼100−200 nm),8,13 the strong electron localization in BVO leads to lower electron conductivity for single-domain BVO films and the illumination direction dependence of photocurrent. Contrary to most photoanode materials, larger photocurrents are produced from BVO under back illumination. In this configuration, the electron−hole pairs are generated near the bottom electrode, which provides shorter electron transport distance to compensate the low electron mobility. On the other hand, vertical domain boundaries provide fast electron transport channels and generate dramatically enhanced conductivity in multidomain BVO films (Figure 4e). Such vertical conductive channels facilitate directional electron transport along the outof-plane direction and reduce carrier recombination. As a result, photoelectrons generated by front illumination shorten the transport path by diffusing across BVO domains into the vertical boundaries and taking this faster route to the back contact. In this sense, vertical domain boundaries partially compensate the slow electron transport inside BVO domains and help reduce the discrimination against Jfront. The observation of higher Jback in multidomain BVO films is likely due to the large domain width, which exceeds the electron diffusion length and limits the number of photoelectrons arriving at the domain boundary. Further improvement in Jfront may be achieved by reducing the diffusion distance of electrons inside domains through directed domain growth. The above results recognize the important and general role of structural defects, including domain boundaries, in modulating photoelectrode performance. The structural boundary effect can be extended beyond the model BVO system and should be coupled with intrinsic carrier transport in designing photoelectrodes for efficient water splitting.

under either front or back AM 1.5G illumination. The dark current density curves measured without and with hole scavenger are largely negligible in the potential range of interest (Figure S9). For BVO films deposited under the optimal condition, the photocurrent densities are maximized at a thickness of 160 nm, with Jback = 0.72 mA cm−2 and Jfront = 0.50 mA cm−2 at 1.23 VRHE (Figure 4a). To the best of our knowledge, this is the first study using c-axis-oriented crystalline BVO thin film photoanodes for efficient PEC water splitting. With the greatly improved quality, the photoanode delivers significantly higher photocurrent when compared to other planar dense BVO films reported previously.24,38 To determine the photoanode charge separation efficiency (ηsep) and charge injection efficiency into the aqueous electrolyte (ηaq inj), we further perform PEC characterization in the presence of 1 M sodium sulfite (Na2SO3), which serves as the hole scavenger to suppress the surface recombination of charge carriers. Integrating the absorbed photo flux Φ based on the photoelectrode optical absorption profile and the AM 1.5G solar irradiance determined a photocurrent of 4.83 mA cm−2 for a 160 nm thick BVO thin film photoelectrode. This yields a ηsep of 44.5% and 30.2% at 1.23 VRHE for the back and front illumination, respectively. Calculation details are provided in the Supporting Information. By comparing the photocurrent for water and sulfite oxidation, ηaq inj is readily determined to be 33.4% and 34.2% at 1.23 VRHE, respectively, for the back and front illuminations (Figure S10a). The incident photon-tocurrent conversion efficiency (IPCE) of the crystalline BVO thin film photoanode is also highly dependent on the illumination direction. At 1.23 VRHE, the IPCE under back illumination (IPCEb) for water oxidation remains steady between 20% and 25% for λ < 440 nm, but quickly drops to zero beyond 500 nm, which closely follows the optical absorption spectrum of the film (Figure S10b). On the other hand, the IPCE values under front illumination (IPCEf) are generally smaller than IPCEb and exhibit more pronounced wavelength dependence. At 1.23 VRHE, the IPCEf is maximized at ∼19% near the optical band gap (460 nm), and decreases significantly for longer or shorter wavelengths. This can be understood from the trade-off between photon collection and carrier delivery depending on the wavelength of the incident light. Under front illumination, carrier recombination is more sensitive to the light wavelength, since photoelectrons need to travel over longer distance to back contact. The light with a wavelength around 460 nm ensures appropriate photon absorption and shorter electron transport distance, therefore resulting in the observed wavelength-dependence IPCE. To further enhance water-splitting efficiencies, we select cobalt oxide (CoOx) as oxygen evolution reaction (OER) catalyst39,40 and deposit an ultrathin surface catalyst layer on BVO in situ by PLD. Nominal thickness of the catalyst layer is optimized to be ∼10 nm with respect to the water-splitting activity (Figure S11). The cross-sectional STEM image shows that the ultrathin catalyst layer consists of uniformly distributed islands (top inset of Figure 4b). High-resolution XPS analysis determines a mixed oxidation state of 2+ and 3+ in the CoOx layer (bottom inset of Figure 4b). The addition of the ultrathin cobalt oxide layer shows very minor effect on the light absorption of the BVO/ITO/YSZ heterostructure, but significantly enhances its water-splitting activity. With this ultrathin CoOx catalyst layer, Jback and Jfront of the modified BVO photoanode, respectively, increase to 1.35 and 0.95 mA cm−2 at 1.23 VRHE (Figure 4b), which represent almost 2-fold 1682

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For out-of-plane measurements, the ITO underlayer is used as the bottom electrode; the square Au/Ti layer bilayer with a side width of 50 μm is deposited as the top electrode. For in-plane measurements, BVO films are deposited over bare YSZ substrates, and only top contacts are used. The electrical measurements are performed on an MMR microprobe station connected with an Agilent 4156C parameter analyzer to measure the current−voltage characteristics and resistivity at different temperatures. The conductive AFM measurements are performed using a Bruker Dimension Icon AFM operating in the PeakForce TUNA mode. Conductive silicon probes coated with doped diamond (model DDESP-10) are used to simultaneously collect both the topography and current mapping images with an applied sample bias.

CONCLUSIONS We have probed the direct correlation of domain-boundary conduction with intrinsic carrier transport and PEC water oxidation in the BVO model system. Crystalline BVO thin films, in two distinct forms of epitaxial single-domain or highly textured c-axis-oriented multidomain, are constructed using pulsed laser deposition. The ultralow conductivity for singledomain BVO films reflects the intrinsic small-polaron transport with an activation energy of 557 meV at 300 K. The multidomain BVO films feature unprecedented high conductivity that resembles free-carrier transport behavior. This anomalous conductivity in BVO is identified to arise from vertical domain boundaries using complementary structural characterization, solid-state electronic measurement, and local conductive AFM. Such vertical conductive channels enable fast transport for electrons diffused into the domain-boundary region and reduce photocurrent difference between front and back illumination. This study presents direct experimental identification of domain-boundary conduction in BVO and provides general design guidelines in engineering structural defects coupled with intrinsic material characteristics for semiconductor photoelectrode design.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b05093. Calculation details, UV−vis absorption spectra and Tauc plot analysis, XRD rocking curves, θ−2θ and Φ scans, low-magnification and high-resolution STEM images, EDX mapping and line scans, Mott-schottky analysis, I− V characteristics, topography and conductive AFM images with corresponding local I−V curves, dark J−V curves, charge separation/injection efficiencies, IPCE curves, and optimization of CoOx as an OER catalyst on BVO photoanodes (PDF)

EXPERIMENTAL SECTION

Thin Film Photoelectrode Preparation. BVO thin films are deposited on YSZ(001) single-crystalline substrates by pulsed laser deposition using a KrF excimer laser (λ = 248 nm) with a laser fluence of 1.5 J cm−2 and a repetition rate of 5 Hz. The BVO laser ablation target is prepared by a conventional ceramic sintering method. For epitaxial single-domain BVO thin film growth, bare YSZ(001) substrate with 2 degree miscut is used. For photoelectrochemical characterization of BVO photoanodes, a 40 nm thick ITO film is first deposited at 600 °C in 5 × 10−7 Torr of oxygen on the YSZ substrate. BVO films are subsequently deposited at 600 °C in 20 mTorr of oxygen. After deposition, the films are cooled down to room temperature under the same atmosphere at a rate of 10 °C min−1. The growth temperature, oxygen pressure, and film thickness are optimized with respect to the PEC activity. Microstructure, Optical, and Chemical Characterization. The film microstructure and crystallinity are characterized by X-ray diffraction (XRD, Rigaku Ultima III) and transmission electron microscopy (TEM, JEOL 2100F, JEOL ARM-200F, FEI Talos F200X). The surface morphology is studied with atomic force microscopy (AFM, Park NX20). Optical absorption spectra are obtained with a UV−vis/NIR spectrophotometer (PerkinElmer, Lambda 950). XPS measurements are carried out in an ultrahigh vacuum (UHV) system using Mg Kα (hν = 1253.6 eV) as the excitation source. Raman spectroscopy was collected on a WiTec Alpha combination microscope with 532 nm laser as an excitation source. Photoelectrochemistry Characterization. PEC measurements are performed with a potentiostat (VersaStat4, PAR) in a threeelectrode configuration with BiVO4 films as the working electrode, a platinum wire as the counter electrode, a Ag|AgCl|3 M KCl electrode as the reference electrode, and a pH = 7.0 phosphate buffer as the electrolyte. An active illumination area of 0.18 cm2 is obtained from a 150 W solar simulator with an AM 1.5G filter (100 mW cm−2). Incident photon-to-current efficiency (IPCE) measurements are conducted with a 300 W xenon arc lamp and a grating monochromator (Newport CS310) equipped with band-pass filters to remove high-order diffractions. The light power for each wavelength is measured by an optical power meter (Newport 1918-C) equipped with a UV-enhanced Si photodiode sensor. Electronic Transport Characterization. For solid-state electronic measurement, Au (100 nm)/Ti (5 nm) bilayer top electrodes are fabricated on top of BVO films by photolithography with mask aligner (Karl Suss MA-6) and e-beam evaporation (Lesker PVD75).

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Wenrui Zhang: 0000-0002-0223-1924 Danhua Yan: 0000-0003-4112-3018 Jiajie Cen: 0000-0002-9446-5161 Mingzhao Liu: 0000-0002-0999-5214 Author Contributions

W.Z. and M.L. conceived the project and designed the experiments. W.Z. fabricated the samples and conducted XRD, AFM, UV−vis, XPS, PEC, and electrical property measurements. D.Y. and J.C. helped in part of the PEC measurements. Q.W. performed Raman measurements. J.L., L.Z., and H.X. performed the TEM characterization. W.Z., M.L., and D.Y. analyzed the results and wrote the manuscript. All authors discussed the results and commented on the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research used resources of the Center for Functional Nanomaterials, which is a U.S. DOE Office of Science User Facility, at Brookhaven National Laboratory under Contract DE-SC0012704. The XAS study used resources at ISS 8-ID beamline of the National Synchrotron Light Source II, a U.S. Department of Energy (DOE) Office of Science User Facility operated for the DOE Office of Science by Brookhaven National Laboratory under Contract DE-SC0012704. A.O. acknowledges support from National Science Foundation (NSF DMR-1254600). J.L. was supported by DOE BES Early Career 1683

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Chemistry of Materials

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DOI: 10.1021/acs.chemmater.7b05093 Chem. Mater. 2018, 30, 1677−1685

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DOI: 10.1021/acs.chemmater.7b05093 Chem. Mater. 2018, 30, 1677−1685