Photoelectrochemical Properties and Behavior of Î±-SnWO4

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Article 4

Photoelectrochemical Properties and Behavior of #-SnWO Photoanodes Synthesized by Hydrothermal Conversion of WO films 3

Zhehao Zhu, Pranab Sarker, Chenqi Zhao, Lite Zhou, Ronald L. Grimm, Muhammad N. Huda, and Pratap Mahesh Rao ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b12640 • Publication Date (Web): 19 Dec 2016 Downloaded from http://pubs.acs.org on December 22, 2016

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ACS Applied Materials & Interfaces

Photoelectrochemical Properties and Behavior of α-SnWO4 Photoanodes Synthesized by Hydrothermal Conversion of WO3 films Zhehao Zhu1,2‡, Pranab Sarker3‡, Chenqi Zhao4, Lite Zhou4, Ronald L. Grimm1, Muhammad N. Huda3*, and Pratap M. Rao4,5*

1

Department of Chemistry and Biochemistry, Worcester Polytechnic Institute, Worcester, MA 01609,

USA 2

Department of Chemical Engineering, Worcester Polytechnic Institute, Worcester, MA 01609, USA

3

Department of Physics, University of Texas at Arlington, Arlington, TX 76019, USA

4

Materials Science and Engineering Graduate Program, Worcester Polytechnic Institute, Worcester, MA

01609, USA 5

Department of Mechanical Engineering, Worcester Polytechnic Institute, Worcester, MA 01609, USA

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ABSTRACT

Metal oxides with moderate band gaps are desired for efficient production of hydrogen from sunlight and water via photoelectrochemical (PEC) water splitting. Here, we report an α-SnWO4 photoanode synthesized by hydrothermal conversion of WO3 films, that achieves photon to current conversion at wavelengths up to 700 nm (1.78 eV). This photoanode is promising for overall PEC water-splitting because the flat-band potential and voltage of photocurrent onset are more negative than the potential of hydrogen evolution. Furthermore, the photoanode utilizes a large portion of the solar spectrum. However, the photocurrent density reaches only a small fraction of that which is theoreticallypossible. Density functional theory-based thermodynamic and electronic structure calculations were performed to elucidate the nature and impact of defects in α-SnWO4 prepared by this synthetic route, from which hole localization at Sn-at-W anti-site defects was determined to be a likely cause for the poor photocurrent. Measurements further showed that the photocurrent decreases over time due to surface oxidation,

which

was

suppressed

by

improving

the

kinetics

of

hole

transfer

at

the

semiconductor/electrolyte interface. Alternative synthetic methods and the addition of protective coatings and/or oxygen evolution catalysts are suggested to improve the PEC performance and stability of this promising α-SnWO4 material.

KEYWORDS: tin tungstate; SnWO4; BiVO4; photoelectrochemical; water splitting; photocatalysis

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INTRODUCTION Energy harvested from sunlight is believed to be a viable option to satisfy the increasing world energy demand.1 Photoelectrochemical (PEC) water splitting enables semiconductors to harness solar energy and simultaneously convert it to a chemical form such as hydrogen for convenient storage, transport, and utilization.1-5 Photoanodes composed of binary metal-oxide semiconductors such as TiO2,6-7 WO3,8 and Fe2O39 are stable in harsh oxidizing and aqueous environments but demonstrate poor efficiency for PEC water oxidation due to unsuitably large band gap energies (e.g. TiO2 and WO3) or poor charge transport properties (e.g. Fe2O3). To date, the complex oxide BiVO4 is the top-performing photoanode material due to a combination of a moderate band gap (Eg ≈ 2.4 eV) and moderate charge transport properties.10 However, the band gap of BiVO4 is still larger than desired. In contrast, some complex oxides including ZnFe2O4 (1.9 eV)11, CuWO4 (2.2 eV)12, and α-SnWO4 (1.64 eV)13 may be more promising photoanode materials due to their smaller band gap and subsequent larger overlap with the solar spectrum.

Among the complex oxides, certain ternary metal oxides (AxByOz) in which A is an ns2 metal cation (Bi3+, Sn2+, Sb3+) and B is an md0 metal cation (Nb5+, Ta5+, Mo6+, W6+), are particularly interesting. Materials in this class often have moderate band gaps due to hybridization of the ns2 cation orbitals with O 2p orbitals, and relatively high charge carrier mobilities due to the formation of large-polaron carriers by the hybridization of the md0 cation orbitals with O 2p orbitals.14-15 One particular ternary metal oxide in this category is α-SnWO4. In α-SnWO4, hybridized Sn 5s and O 2p orbitals contribute to the valence band and the W 5d orbitals principally contribute to the conduction band.13 Earlier electronic structure calculations

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of α-SnWO4 suggested favorable properties for water splitting photoanodes,13 and motivated synthesis and PEC studies.

Several reports detail the synthesis and characterization of α-SnWO4, both as a powder and as thin-film electrodes.13, 16-24 One study utilized a solid-state reaction of SnO and WO3 powders to yield n-type α-SnWO4 powders with an energy band gap of 1.64 eV and flat-band potential of 0.047 V vs VRHE.13 That study analyzed the photocatalytic degradation of rhodamine B dye and did not quantify the water-splitting properties. α-SnWO4 synthesized via co-sputtering of W and SnO2 targets yielded thin films with an indirect band gap of 1.52 eV and a direct transition of 1.95 eV.23 In another study, the hydrothermal conversion of WO3 yielded α-SnWO4 thin-film photoanodes with an indirect band gap of 1.9 eV that demonstrated photocurrent density values of up to 0.050 mA cm–2 in an aqueous potassium phosphate buffer at pH 5.16 Another study reported a general strategy to synthesize complex metal oxides including SnWO4 via chemical and thermal treatments of nanocrystal precursors, but did not explore photoelectrochemical properties of the material.25 Despite these initial results, significant work remains in order to evaluate the viability of α-SnWO4 photoanodes for water splitting applications. Firstly, knowledge gaps and inconsistent reports exist in regards to the values of the band gaps and flat-band potentials of α-SnWO4 thin films. Secondly, the reason for the low photocurrents reported thus far for α-SnWO4 photoanodes remains unknown. Finally, the stability of α-SnWO4 exposed to water oxidation conditions is not understood. Elucidating these knowledge gaps motivates the present study. Overall, the present experimental results compare favorably to the state-of-the-art photoanode material BiVO4 and suggest that α-SnWO4 remains

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promising for photoelectrochemical water-splitting. Moreover, the theoretical results point to directions for further improvement of the performance of α-SnWO4 photoanodes.

EXPERIMENTAL AND THEORETICAL METHODS In this work, we synthesize α-SnWO4 porous film (PF) photoanodes via a hydrothermal conversion of WO3 in order to quantify material properties and photoelectrochemical behavior, in conjunction with density functional theory (DFT) electronic structure calculations. Scanning electron microscopy, X-ray diffraction, UV-visible absorbance, and X-ray photoelectron spectroscopy are performed to quantify material morphology, crystallinity and phase purity, optical properties, and surface chemical composition of the α-SnWO4 photoanodes, respectively. Photoelectrochemical measurements under simulated solar illumination, electrochemical impedance spectroscopy, and incident photon-to-current efficiency (IPCE) experiments are conducted to establish the behavior of the photoanode in contact with the aqueous phase. We further contrast the PEC water oxidation performance of the α-SnWO4 photoanodes in contact with aqueous potassium phosphate (KPi) buffer to that in contact with aqueous hole scavengers and both aqueous and non-aqueous redox couples that undergo facile, one-electron transfer to determine the impact of water oxidation kinetics on the photocurrent and stability of the photoanode. In addition, we report computational studies that extend beyond band gap calculations to elucidate the nature and impact of defects in α-SnWO4.

SnWO4 porous films (PFs) were synthesized by a modification of a previously-reported method, through a combination of sol-gel reaction and hydrothermal conversion. First, WO3 PFs were deposited onto fluorine-doped tin oxide (FTO) substrates (2.0 × 1.5 cm, 2.2 mm thick, TEC -5-



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7, Hartford Glass) by spin-coating an aqueous solution of ammonium metatungstate hydrate (AMT) or tungstic acid (H2WO4) as a tungsten source and polyvinyl alcohol (PVA) polymer as a structure controlling agent.8, 26-27 The AMT/PVA precursor was prepared by dissolving 0.379 g AMT (99.0%, Fluka) in 1 mL DI water, adding a solution containing 1.5 g PVA (87% - 90% hydrolyzed, MWavg 30,000 - 70,000, Sigma-Aldrich) in 9 mL DI water, and mixed by stirring and ultra-sonication. The H2WO4/PVA precursor was prepared by dissolving 0.375 g H2WO4 (99%, Aldrich) into 6 mL aqueous hydrogen peroxide (H2O2, 30 wt%, Sigma Aldrich) and heating at 125 °C until the solution volume was reduced to 1 mL, then adding a solution containing 1.5 g PVA in 9 mL DI water, and mixing by stirring and ultra-sonication. Each layer of the WO3 PF was spin-coated onto FTO using 200 μL of the solution at a spin speed of 2000 rpm for 40s and annealed in air at 550 °C for 10 minutes to remove the polymer and obtain the desired porous nanostructure. After 4, 8 or 12 layers of the precursor were coated, the substrates were finally annealed in air at 550 °C for 2 hours to completely convert the precursors to crystalline WO3.26 Second, WO3 films were converted to α-SnWO4 via a hydrothermal reaction.16 The WO3 films were leaned face down against the wall in a Teflon-lined autoclave (50 mL, Hydrionscientific) filled with 28 mL of 0.5 M SnCl2 (98%, Sigma-Aldrich) aqueous solution. The hydrothermal reaction was performed at 180 °C for 24 hours. Finally, the films were rinsed with 2 M HCl to remove tin hydroxide chloride and tin oxide crystals that formed in the reaction.16 The BiVO4 photoanode used in this study was synthesized by drop-casting and annealing a solution as described in our previous work.28 Briefly, the solution for drop-casting BiVO4 was composed of 0.1225g Bi(NO3)3·5H2O (98%, Sigma Aldrich) and 0.0663g VO(C5H7O2)2 (98%, Sigma Aldrich), dissolved in 5mL acetic acid (≥99.7%, Sigma Aldrich) and 0.25mL acetylacetone (≥99%, Sigma Aldrich) by 30 minutes sonication, resulting in a black-6-


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green solution. 2 drops of 5 µL solution were then dropped onto the FTO substrate, which was then dried on a hotplate set at 500 °C in air for 10 minutes to make one BiVO4 layer. After 3 layers of drop-casting, the sample was annealed in a furnace in air at 550 °C for 2 hours to crystallize the BiVO4.

The morphologies, crystal structures, and chemical compositions of the WO3 PFs and α-SnWO4 PFs were characterized by scanning electron microscopy (SEM, JSM-7000F, JEOL), parallel beam X-ray diffraction (XRD, PANalytical Empyrean, Cu-Kα, 45 kV, 40 mA), and X-ray photoelectron spectroscopy (XPS, PHI 5600, Al-Kα, 13.5 kV, 300 W). WO3 and α-SnWO4 PFs were synthesized on stainless steel substrates for XRD and XPS characterizations to eliminate the presence of SnO2 and Sn4+ in the substrates. The average crystallite size was calculated from the Scherrer equation29 .

= ,

where λ is the X-ray wavelength, β is the measured width of the peak at half-maximum intensity in radians, and θB is the Bragg angle.

The wavelength-dependent optical absorption properties of the samples were obtained with an integrating sphere (Model 4P-GPS-060-SF, Labsphere) using illumination from a Xe lamp (Model 66902, Newport). For both the transmission and reflection measurements, light was incident perpendicular to the back-side (glass) surface of the sample. For the transmission measurements (including both spectral and diffuse components), the samples were placed at a port at the front of the integrating sphere. For the reflection measurements (including both spectral and diffused components), the samples were aligned to a port at the back side of the -7-



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integrating sphere and the reflectance spectra were normalized to the reflection of a white standard with the sample as part of the sphere wall at a side port. A spectrometer (USB 2000+, Ocean Optics) at the top port was used to measure the transmitted (T) and reflected (R) light, respectively. The absorption coefficient (α) was then calculated as 1 100 % − (%) = ln (%) where z is the film thickness measured from cross-section SEM images.30

The PEC measurements were performed in both three-electrode and two-electrode configurations, using a potentiostat (Model SP-200, BioLogic) under either front- or back-side broadband illumination from a Xe lamp. Linear sweep voltammograms (LSVs, i.e. J-V curves) were measured at a scan rate of 10 mV/s. LSVs in aqueous electrolytes were measured in a three-electrode configuration with the photoanode as the working electrode, a Pt wire (0.5 mm diameter) as the counter electrode, and a saturated calomel (SCE) reference electrode. The aqueous electrolyte used was 0.5 M potassium phosphate buffer (KPi) electrolyte at pH 7 with or without 1 M Na2SO3 added as a hole scavenger. LSVs were also measured in aqueous electrolyte with 350 mM potassium hexacyanoferrate(II) (98.5%, Sigma-Aldrich) and 50 mM potassium hexacyanoferrate(III) (99.0%, Sigma-Aldrich) forming a reversible redox couple.31 Potentials in aqueous electrolytes are reported versus RHE using (volt) = $% (volt) + 0.244 (volt) + )0.059 (volt) × pH/. Chronoamperograms (CAs, i.e. J-t curves) were also measured in 0.5 M KPi aqueous solutions in a two-electrode configuration with the photoanode as the working electrode and a Pt wire as the counter electrode, with a voltage of 0 V between the working and counter electrodes. All -8-


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aqueous electrolytes used in this study were thoroughly purged with Ar gas, and measurements were conducted under a blanket of Ar. LSVs in non-aqueous electrolyte were measured in a three-electrode configuration with a Pt wire reference electrode, in anhydrous acetonitrile (99.8%, Sigma-Aldrich, dried further using 3 Å molecular sieves) with 1 M LiClO4 (95.0%, Sigma-Aldrich) electrolyte and 100 mM ferrocene (Fc) (98%, Aldrich) and 0.5 mM ferrocenium (Fc+) tetrafluoroborate (Aldrich) forming a reversible redox couple.32 Potentials in non-aqueous electrolyte are reported versus the Fc/Fc+ couple, which can be approximately compared to potentials on the SCE scale using33

%$)Fc 8 / (volt) = 0 1/0 (volt) + 0.31 (volt) + 40.059 (volt) × log 9. )Fc/

The incident light intensity from the Xe lamp at each wavelength was measured by a spectrometer. The integrated power of the Xe lamp output at wavelengths shorter than 650 nm was 44.5 mW/cm2, as compared to 40.9 mW/cm2 for the standard AM 1.5G spectrum (Figure S1a). The IPCE was measured at 1.23 VRHE using a Xe lamp equipped with a monochromator (Cornerstone 130 1/8 m, Newport). Additionally, for wavelength of 700 nm, monochromatic light was obtained using a combination of long- and short-pass filters (Thorlabs). The spectral irradiance of monochromatic light at each wavelength was measured by a spectrometer and is reported in Figure S1b. The IPCE was calculated using IPCE (%) = G

>?@ (AB⁄ AC ) ×EF

HIJI (AK⁄ A

C ) × (LA)

× 100%,

where Jph is the measured photocurrent density, Pmono is the intensity of the incident monochromatic light, and λ is the wavelength of the monochromatic light.10 Electrochemcial impedance spectroscopy (EIS) was measured from 0.1 Hz to 10 kHz in a three-electrode -9-



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configuration with a SCE reference electrode, in an aqueous electrolyte with 0.5 M potassium phosphate buffer (KPi) of different pH values and 350 mM potassium hexacyanoferrate(II) and 50 mM potassium hexacyanoferrate(III) forming a reversible redox couple to prevent α-SnWO4 from oxidation. The EIS results were then fitted with an equivalent circuit as shown in Figure S2a. Mott-Schottky (MS) plots were generated by calculating the capacitance (C) of the photoanode/water interface from the constant phase element (Q) at different electrode potentials using: M = (N)E/O /, where Q is the constant phase element, R is the resistance in parallel with the constant phase element, and n is the exponent equaling 1 for a capacitor. The flat-band potential (Vfb) was determined by extrapolation to 1/C2 = 0 of the Mott-Schottky equation13 1⁄M F = (2⁄P P QRS )) − TU − (VW ⁄Q )/, where C is the space charge capacitance, εo and εs are the dielectric constants of free space and the film electrode, respectively, q is the electronic charge, Vfb is the flat-band potential, T is the temperature in Kelvin, ND is donor density, V is the applied potential, and kB is Boltzmann’s constant.

DFT calculations were performed within the framework of the standard frozen-core projector augmented-wave (PAW)34-35 method using DFT as implemented in the Vienna ab initio simulation package (VASP)36-37 code. Exchange and correlation potentials were treated in the generalized gradient approximation (GGA) as parameterized by Perdew-Burke-Ernzerhof (PBE).38-39 The basis sets were expanded with plane waves with a kinetic-energy cut-off of 400 eV. Since it is well known that DFT, hybrid-DFT or Hartree-Fock methods fail to account - 10 -


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for the non-bonded dispersion interaction between the layers in layered structures40-42, it is necessary to incorporate empirical vdW interactions in layered structures to compensate for this failure.43-45 Since SnO is a layered structure compound, we employed a vdW correction, namely D2.46 Further, hybrid-DFT (HSE0647, exact exchange = 40% and plane wave energy cut-off = 800 eV) was introduced in conjunction with D2-DFT to reproduce the experimental lattice parameters and band gaps of SnO.48 A 11 x 5 x 11 Monkhorst–Pack49 V-point sampling was used for ion relaxation; however, optical absorption calculations were done with higher 15 x 7 x 15 V-point sampling. For the defect calculations, we constructed a supercell containing 96 atoms. A

V-point mesh of 3×3×5 was employed for all defect-induced calculations. For visualization of the

crystal structures, VESTA (Visualization for Electronic and Structural Analysis)50-51 was used.

RESULTS AND DISCUSSION Figure 1 shows the morphologies of the 8 layers of spin-coated WO3 film and the hydrothermally converted α-SnWO4 film on FTO substrates. The WO3 film has a porous structure that consists of interconnected WO3 nanoparticles with an average diameter of ~150 nm and has a film thickness of 300 nm (Figure 1). The α-SnWO4 film exhibits a partial retention of the porous nanostructure which is composed of elongated nanoparticles with an average minor diameter of ~25 nm and has a film thickness of 300 nm. The porous nanostructure of α-SnWO4 in this work is different from the nanobelt structure in the previous study due to different structures of the WO3 film.16

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Figure 1. SEM images of (a) 8 layers of spin-coated WO3 porous film (PF) and (b) SnWO4 PF hydrothermally converted from 8 layers of WO3 PF on FTO substrates.

The XRD patterns (Figure 2a) indicate that orthorhombic α-SnWO4 (ICDD PDF 04-0110013) can be consistently synthesized by hydrothermal conversion of WO3 PF. The average crystallite size was calculated from the (121) reflection as ~28 nm, which is similar to the average minor diameter of the nanoparticles measured by SEM. The synthesized α-SnWO4 contains a small amount of monoclinic WO3 (ICDD PDF 04-005-4272), tetragonal SnO (ICDD PDF 04-005-4540), and tetragonal SnO2 (ICDD PDF 00-041-1445) phases. The SnO and SnO2 - 12 -


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phases were observed in the previous report as a byproduct from the hydrothermal reaction.16 However, WO3 phase was not observed in that study.

To understand the thermodynamics of formation of α-SnWO4, we further performed DFT calculations. α-SnWO4 crystallizes into the orthorhombic symmetry (Pnna, S. G. : 52, and Z = 4), with DFT-optimized lattice parameters of a = 5.599 Å, b = 11.652 Å, c = 4.991 Å, and α = 90°. These values are in good agreement with experimental results.48, 52 In α-SnWO4, both Sn and W form a layer of SnO4 tetrahedra and WO6 octahedra, respectively, in the ‘b’ direction, as shown in Figure 2b. These layers of cationic polyhedra extend into the ‘c’ direction in each unit cell.

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Figure 2. (a) XRD pattern of α-SnWO4 PF hydrothermally converted from 4 layers of WO3 PF on a stainless steel substrate before and after LSV measurements in an aqueous 0.5 M KPi buffer electrolyte at pH 7. Also shown are standard patterns for α-SnWO4, SnO, SnO2 and WO3. (b) Crystal structure of αSnWO4 represented by ball-and-stick model and polyhedral model. In the figure, green, black, and red balls correspond to Sn, W, and O atoms, respectively.

The chemical potential landscape for α-SnWO4 and related phases was then calculated by DFT (Figure 3a). In this plot, the phase stability of α-SnWO4 is reported as a function of growth - 14 -


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conditions, represented by the chemical potentials of its constituent elements, ∆[\ (i = Sn, W, and O). In general, the phase stability of phase X at constant temperture and pressure can be determined from the free energy of formation of the phase, ∆]^ (X) = ∆_^ (`) − ∆a, which

equals ∑d \ c\ [\ . However, for a crystalline structure of α-SnWO4 with a low level of disorder, a 53 good approximation is to set ∆S ≈ 0, so that now ∆_^ itself equals ∑d \ c\ [\ . Moreover, DFT

calculations are conducted at T = 0 K, so that calculated formation energies are equivalent to formation enthalpies ∆_^ . At thermal equilibrium

∆_^ (acef ) = ∆[gO + ∆[h + 4∆[i .

where ∆[\ are referenced to their standard elemental phases. The above equation is satisfied anywhere in the triangle presented in Figure 3a, so SnWO4 formation is possible anywhere inside the triangle. Now, additional binary oxides of Sn and W can be present within the range of chemical potentials where SnWO4 is formed. To prevent the growth of these secondary binary phases, chemical potentials can be further restricted within a region inside the triangle. For example, to prevent the growth of SnO2, the chemical potentials of Sn and O must satisfy ∆_^ (acfF ) > ∆[gO + 2∆[i . Now these two equations can be solved for the chemical potentials of Sn and W, and those can be used to draw the red line for SnO2 in Figure 3a, which divides the triangle into two parts: SnO2 will not form above the SnO2 line, but will form below it. Similarly, SnO will not (will) form above (below) the blue SnO line, and WO3 will not (will) form below (above) the green WO3 line. The thermodynamic growth conditions that allow the co-existence of SnO and SnO2 phases is due to the relatively Sn-rich and W- and O-poor condition within the grey bounded (DEF) region. However, WO3 cannot be simultaneously present with these two Sn-O phases at thermodynamic equilibrium, suggesting that the presence of WO3 in the synthesized film can be - 15 -



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attributed to the incomplete (kinetically-limited) conversion of the WO3 PF. Note that the coexistence of SnO and SnO2 phases was also observed in the previous hydrothermal synthesis of SnWO4 at pH=11.16

Furthermore, we have calculated the formation energies for various defects in α-SnWO4 as a function of the thermodynamic growth conditions (Figure 3b). The formation energy for a defect δ in a charge-neutral state is53 ∆_k = lk − lmnop + ∑r qr [r ,

where lk and lmnop are the total energies of the SnWO4 with and without defect s, respectively, the index corresponds to different atomic species that constitute the solid, and qr is the

number of atoms removed (added) from (to) the system to form the defect s. qr = 1 if an atom

of species is removed from the system, while qr = −1 if an atom of species is added to the system. Given the relatively Sn-rich and W-poor synthesis condition of the present study, defects such as Sn anti-site at W (SnW) are possible. According to the DFT calculation, SnW is the most probable defect due to its lowest formation energy among all the studied defects. In addition, the anionic VO defect is also probable in the projected experimental growth conditions. Since the VO defect is likely responsible for n-type conduction in α-SnWO4, and this defect is most probable only at point D in Figure 3b, we therefore surmise that the growth conditions at D and nearby it inside DEF (relatively Sn-rich and W/O-poor) mimic the experimental results most closely. At these conditions, relatively high concentrations of the SnW and VO defects are expected.

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Figure 3. The defect formation energies and chemical potential landscape analysis of α-SnWO4 calculated by DFT. (a) Chemical potential landscape showing single-phase growth region (yellow bounded region) of α-SnWO4. ∆µα (α = Sn, W, and O) axes correspond to growth conditions of respective species. (b) Formation energies of different intrinsic defects in α-SnWO4 with respect to its projected experimental growth condition (DEF).

Figure 4a shows the measured optical absorption spectrum of α-SnWO4 prepared from 8 layers of WO3 PF. The sample exhibits strong absorption across most of the visible-light range, with an absorption edge at photon energy of 2.2 eV (565 nm). From this, indirect optical transitions can be determined from the (αhv)1/2 vs hv tauc plot (Figure 4b), while direct optcal transitions can be determined from the (αhv)2 vs hv tauc plot (Figure 4c).23 α-SnWO4 synthesized here was thus determined to possess an indirect transition of ~2.1 eV and a direct transition of ~2.4 eV. These indirect and direct optical transitions are smaller than or comparable to many - 17 -



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other n-type oxide materials such as WO3 (Eg ≈ 2.7 eV)8, BiVO4 (2.4 eV)54, Fe2O3 (2.1 eV)9, and CuWO4 (2.2 eV)12, suggesting that α-SnWO4 has potential to harvest a significant fraction of solar photons compared to other metal oxides, which is promising for its use as a photoanode in a PEC system.

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Figure 4. (a) Optical absorption coefficient of the α-SnWO4 PF prepared from 8 layers of WO3 PF before and after LSV measurements in an aqueous 0.5 M KPi buffer electrolyte at pH 7; (b) Tauc plot to determine indirect optical transitions of the α-SnWO4 sample before LSV measurements; (c) Tauc plot to determine direct optical transitions of the α-SnWO4 sample before LSV measurements. - 19 -



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We next calculated the DFT band structure of pristine α-SnWO4, as well as α-SnWO4 containing either SnW or VO defects. As shown in Figure 5a (left), for pristine SnWO4, the fundamental band gap is found to be indirect and occurs in the ΓZ region of the Brillouin zone, with energy of 1.05 eV, between conduction band states with energy around 1 eV and valence band states with energy at 0 eV. The next largest transition is found to be direct and occurs in the TY region, with energy of 1.1 eV, between conduction band states with energy of around 1 eV and valence band states with energy of around -0.1 eV. The corresponding density of states is shown in Figure 5b. The top of the valence band (at energies of 0 to -0.5 eV) is composed primarily of Sn s states. On the other hand, the bottom of the conduction band, at energy of around 1 eV is composed primarily of W d states. Therefore, these two transitions (direct, 1.05 eV and indirect, 1.1 eV) are most likely to be s to d transitions, which have a difference of orbital angular momentum quantum numbers ∆l = 2. Selection rules dictate that transitions with ∆l = +/1 occur with high probability, while other transitions occur with much smaller probabilities.55 Therefore, these two s to d transitions should have low probability and should produce a small value of absorption coefficient at these energies. The calculated absorption coefficient (due to direct transitions only) is shown in Figure 5c, and indeed shows weak absorption for energies of 1.0 – 1.6 eV. The next significant transition in the band structure is a direct transition at the Γpoint, with energy of 1.6 eV, between valence band states with energy of around -0.6 eV and conduction band states with energy of around 1 eV. As shown in Figure 5b, the valence band states at this energy have nearly equal contribution from Sn s and O p states. Since p to d transitions have ∆l = 1, we expect these transitions to be significantly stronger. Indeed, Figure 5c shows a significant rise in the calculated optical absorption coefficient starting at around 1.6-1.7 - 20 -


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eV, corresponding to this transition. The next major transition occurs at the Γ-point, with energy of around 2.5 eV, between valence band states at around -1.5 eV and conduction band states at around 1 eV. As shown in Figure 5b, these valence band states consist of a high density of O p states. Consequently, we expect the p to d transitions from these states to the conduction band to be especially strong. Indeed, Figure 5c shows an even stronger rise in the absorption coefficient at energy of around 2.5 eV, corresponding to this transition. There is also a direct transition with similar energy at the Y-point, again going from valence band states at -1.5 eV to conduction band states at 1 eV. However, the valence band states here are likely composed of Sn s, leading to an unfavorable s to d transition. Note that our calculations do not include the contribution of indirect transitions to the absorption coefficients, although such contributions are possible.

The band structure and optical absorption spectrum were also calculated for SnWO4 containing VO and SnW defects, as shown in Figure 5a (middle and right) and in Figure 5c. In case of pristine SnWO4 (Figure 5a (left)), the lowermost conduction bands are degenerate. This degeneracy is removed in the presence of VO (Figure 5a (middle)) and SnW (Figure 5a (right)). In the case of VO-α-SnWO4, one of the bands becomes occupied, serving as a donor level, which gives rise to n-type conductivity in α-SnWO4. In contrast, in the case of SnW-α-SnWO4, both bands remain unoccupied and instead a new defect state is introduced beneath those bands. As shown in Figure 5c, the calculated VO - and SnW- α-SnWO4 absorption spectra have absorption onsets at low energies (~0. 2 eV and at ~0.5 eV, respectively) which are due to electron transitions (i) from occupied donor level to unoccupied lowest conduction band and (ii) from highest occupied valence band to unoccupied defect state, respectively. The presence of defects also decreases the energies of the main transitions from 1.6 eV and 2.5 eV to lower values. - 21 -



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We now compare the calculated and measured optical absorption spectra. However, this must be done with caution, because of the well-known tendency of DFT to under-predict band gaps. Firstly, the measured spectrum (Figure 4a) shows that there is weak absorption in the range of 1.5 - 1.9 eV. This is consistent with the prediction of relatively weak absorption for energies below around 2 eV (Figure 5c). Due to the limitations of our measurement apparatus, we are not presently able to accurately measure absorption at energies lower than 1.5 eV. There is then a strong rise in the measured absorption at an energy of around 2.2 eV. This rise is quite similar to the rise in calculated absorption coefficient at energies of 2.0-2.5 eV for pristine and defectcontaining α-SnWO4. To compare the measured and calculated absorption spectra more quantitatively, we plot both the measured and calculated absorption coefficients as (αhv)2 vs hv tauc plots, which permit quantitative extraction of the direct optical transition energy (Figure 4c – experimental, Figure 5d – calculated). The resulting calculated tauc plot shows direct optical transitions of 2.6 eV, 2.5 eV, and 2.3 eV, respectively, for pristine α-SnWO4, VO - α-SnWO4 and SnW - α-SnWO4. The direct transition of 2.4 eV determined from the measured tauc plot is between the 2.3 eV and 2.5 eV direct transitions of Vo - α-SnWO4 and SnW - α-SnWO4, which is consistent with the presence of Vo and SnW defects. However, this should not be taken as evidence of the presence of these defects, for which more direct and sensitive measurements are needed. Moreover, although there seems to be good agreement between the calculated and measured absorption spectra, the agreement and interpretation should only be taken as tentative because of the well-known tendency of DFT to under-predict the value of band gaps. It is possible, for instance, that the DFT calculations are under-predicting the energies of the transitions, and that the observed strong absorption onset at 2.4 eV is actually due to the - 22 -


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transition at the Γ-point that was calculated to cause a rise in absorption at 1.5-1.7 eV rather than the transition at the Γ-point that was calculated to give rise to an absorption onset at energy of 2.0-2.6 eV. A more refined and sensitive measurement of the absorption coefficient over a wider range of energies is required in the future to more thoroughly compare the calculated and measured spectra and assign transitions to particular orbital transitions with greater confidence.

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Figure 5. (a) The DFT electronic band structures of pristine α-SnWO4 (left), VO - α-SnWO4 (middle), and SnW - α-SnWO4 (right). The Fermi level (EF) is set at 0 eV. (b) The DFT projected density of states (pDOS) of α-SnWO4. (c) The DFT-calculated optical absorption spectra of pristine, VO -, and SnW - αSnWO4. (d) The DFT-calculated tauc plot for direct optical transition of pristine, VO -, and SnW - αSnWO4.

The PEC performance of the α-SnWO4 PF photoanode prepared by the conversion of 8 layers of WO3 PF was evaluated by measuring the three-electrode LSV curves and wavelength-dependent IPCE. This film thickness was optimized for maximum photocurrent under white light illumination (Figure S3a). In a chopped-light LSV (Figure 6a), the total current density of the α-SnWO4 PF in 0.5 M KPi buffer electrolyte at pH 7 reaches 0.10 mA/cm2, while the photocurrent density (not including the dark current) reaches 0.08 mA/cm2 at a potential of 1.23 VRHE under simulated AM 1.5 G illumination. This photocurrent density is far below that which is theoretically-possible (several mA/cm2) based on the band gap of the material. The nonzero dark current is likely due to oxidation of Sn2+ to Sn4+ (Eo = 0.15 VNHE) with the participation of water, as have been observed in the previous report.16 Additionally, the anodic dark-current peak between -0.2 and 0.1 VRHE might be attributed to the redox reaction Sn + 2h+ ↔ Sn2+ (Eo = -0.14 VNHE).56 A small of amount of Sn2+ in α-SnWO4 might be initially reduced at the negative potentials and re-oxidized during the anodic scan. This is supported by the LSV of α-SnWO4 with a starting potential of 0 VRHE (Figure S3b), which shows absence of the anodic dark-current peak. Moreover, the photocurrent onset voltage of the photoanode is ~ -0.1 VRHE (Figure 6a), which is approximately equal to the flat-band potential (~ -0.06 VRHE) of the sample determined from the intercept of the potential axis in the MS plot (Figure 6b and equivalent circuit fitting in - 25 -



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Figure S2a). This flat-band potential value was additionally verified by MS plots of α-SnWO4 at different pH values (Figure S2b), which shows that the extrapolated Vfb changes by approximately 0.059 V per pH, as expected. In addition, the flat-band potential of ~ -0.06 VRHE measured here is more negative than the value of 0.047 VRHE reported previously for an electrode prepared by spin-coating a slurry of α-SnWO4 powders.13 The positive value of the MS slope (Figure 6b) indicates that the synthesized α-SnWO4 is an n-type semiconductor, which is likely due to oxygen vacancies in its crystal structure. This is supported by our DFT calculations, which showed that the VO defect has a favorable formation energy (Figure 3b) and contributes an occupied band below the conduction band that behaves as a donor level, as explained earlier (Figure 5a).

The flat-band and photocurrent onset potentials of α-SnWO4 measured here are more negative than those for TiO2 (~0 VRHE)6-7, BiVO4 (~0.1 VRHE)10, WO3 (~0.5 VRHE)10, and Fe2O3 (~0.55 VRHE)9, and more negative than the hydrogen evolution potential (0 VRHE), which indicates that photoexcited electrons from α-SnWO4 should have sufficient energy to reduce water at the same time that photoexcited holes have sufficient energy to oxidize water. This is supported by chopped-light two-electrode CA at 0 V between the working and counter electrodes (Figure 6c), which shows that the α-SnWO4 photoanode exhibits photocurrent density of 0.2 µA/cm2 as compared to ~0 µA/cm2 for a BiVO4 photoanode in the same aqueous electrolyte. Although it is necessary to measure the evolved gases to determine whether this photocurrent strictly corresponds to the overall water splitting reaction, this was not possible in the present study because of the low rates of gas evolution resulting from these modest photocurrents, and

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the lack of stability of the photoanode. Future improvements in the photocurrent density may enable such measurements.

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Figure 6. (a) Chopped light three-electrode linear sweep voltammogram (LSV) of α-SnWO4 PF prepared from 8 layers of WO3 PF measured in 0.5 M KPi buffer electrolyte at pH 7 under back-side illumination. (b) Mott-Schottky (MS) plots of α-SnWO4 PF prepared from 8 layers of WO3 PF measured in 0.5 M KPi buffer with Fe(CN)64-/3- redox couple. (c) Chopped light two-electrode chronoamperograms (CA) of αSnWO4 PF prepared from 8 layers of WO3 PF and BiVO4 control sample measured in 0.5 M KPi buffer electrolyte at pH 7 under back-side illumination.

Although the photocurrent is stable in aqueous electrolyte at short-circuit conditions as shown in Figure 6c, the photocurrent decreases over subsequent three-electrode LSV measurements of the photoanode in aqueous electrolyte, indicating a lack of stability at higher applied voltages (Figure 7a, and raw data in Figure S4a). However, the bulk crystallinity and optical absorption of the photoanode are not degraded by the measurements (Figure 2 - XRD and Figure 4a - optical absorption) and thus are not responsible for the decreased photocurrent of the material. In contrast, the photoanode exhibits stable photocurrent over all voltages during consecutive three-electrode LSV measurements in a non-aqueous electrolyte (Figure 7a, and raw data in Figure S4b). This suggests that the degradation of the photocurrent in aqueous electrolyte is likely due to the oxidation of Sn2+ to Sn4+ at the semiconductor/electrolyte interface by the reaction SnWO4 + H2O + 2h+ → SnO2 + WO3 + 2H+, where h+ is a hole. This is further supported by the XPS spectra of α-SnWO4 photoanodes. It was reported that the binding energies of Sn 3d5/2 for Sn2+ (SnO) and Sn4+ (SnO2) are 485.8 eV and 486.3 eV, respectively.57 As shown in Figure 7b, only the presence of Sn2+ is observed in the Sn 3d5/2 spectra of α-SnWO4 before (485.1 eV binding energy) and after (485.4 eV binding energy) non-aqueous LSV measurements, but both Sn2+ and Sn4+ exist at the surface of α-SnWO4 PF after LSV measurements in aqueous - 28 -


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electrolytes (485.5 eV and 486.6 eV binding energy) (Figure 7b), which suggests that Sn2+ was oxidized to Sn4+ with the participation of water. The resulting SnO2 and WO3 species would form electron trap sites that allow electron-hole recombination, which could be responsible for the photocurrent degradation.58-59 In addition, N 1s spectrum exists in the XPS spectra of the α-SnWO4 photoanodes (398.8 eV binding energy) (Figure S5), which can be attributed to the incomplete conversion of AMT to WO3. However, synthesizing α-SnWO4 photoanodes by using H2WO4, which eliminates the presence of nitrogen in the W-precursor, does not improve the photocurrent density or the stability of the material (Figure S6). The magnitude of photocurrent of the α-SnWO4 photoanode in non-aqueous electrolyte (~30 µA/cm2, Figure S4b) is similar to that in aqueous pH 7 KPi buffer electrolyte, indicating that the surface oxidation of α-SnWO4 is not the reason for the low photocurrents observed.

The stability of the photoanode in aqueous electrolyte can be significantly improved by adding a fast, reversible Fe(CN)64-/3- redox couple to the aqueous electrolyte (Figure 7a and raw data in Figure S7). The stabilizing effect is similar to that achieved by applying oxygen evolution catalysts to the photoanode surface, which enhance the charge dynamics at the semiconductor/electrolyte interface to suppress the competing oxidation of Sn2+ to Sn4+ by photogenerated holes. The photocurrent density in aqueous electrolyte with the Fe(CN)64-/3redox couple reaches ~8 μA/cm2 at 0.6 VSCE, which is far below that in aqueous pH 7 KPi buffer electrolyte and the non-aqueous electrolyte. The photocurrent loss is likely due to fast electron-hole recombination at the exposed FTO surface of the porous film through back reduction of the Fe(CN)63- species.60 The concentration of the Fe(CN)63- species in the aqueous

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electrolyte is higher than that of Fc+ in the non-aqueous electrolyte, which makes the Fe(CN)63species more likely to be reduced and thus lowers the photocurrent.

Figure 7. (a) Difference in photocurrent density (with dark current subtracted) at different potentials between the first and third LSV measurements of α-SnWO4 PF in an aqueous pH 7 KPi buffer electrolyte, an aqueous electrolyte with Fe(CN)64-/3-, and a non-aqueous 1 M LiClO4 electrolyte in anhydrous acetonitrile with Fc/Fc+. A negative value on the y-axis represents degradation of the sample. (b) Sn 3d5/2 XPS spectra of α-SnWO4 PFs on stainless steel as-synthesized (top), after LSV measurements in nonaqueous electrolyte (middle), and after LSV measurement in aqueous electrolyte (bottom).

The IPCE of α-SnWO4 (Figure 8a and raw data in Figure S8) shows photon-to-current conversion up to wavelengths of 700 nm (1.78 eV), while the photocurrent at wavelength of 750 nm (1.65 eV) is within the noise level of the measurement system, and is therefore considered to - 30 -


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be zero. The IPCE for 650 - 700 nm (1.91 - 1.78 eV) is low but non-zero, and increases more sharply for wavelengths shorter than 600 nm (2.07 eV), reaching a value of 2.8% at 350 nm (3.54 eV). These results are consistent with our theoretical predictions of electronic structure and the measured and calculated absorption coefficients, which show a weaker transition at energies of ~1.6-2.0 eV and a stronger transition at energies above ~2.0 eV. The longest wavelength of photon-to-current conversion of α-SnWO4 (700 - 750 nm) compares favorably against that of state-of-the-art BiVO4 which is ~510 nm.10 The photocurrent density obtained by integrating the measured IPCE over both the standard and simulated AM 1.5 G spectrum is approximately 0.14 mA/cm2, which is higher than the ~0.08 mA/cm2 obtained from the LSV curves under broadband illumination at 1.23 VRHE. Possible reasons for this discrepancy might be that the charge transfer at the α-SnWO4/electrolyte interface is slower at the higher current densities in the broadbandilluminated LSV measurements. Such decreases in quantum efficiency with increasing light intensity have also been observed in other oxides such as BiVO4.61

The charge transport within the material was further analyzed by comparing the LSV curves of α-SnWO4 PF under front- and back-side illumination conditions. When the electron-hole pairs are generated in the α-SnWO4 PF under illumination, the holes would only need to travel across the nanoparticles to reach the α-SnWO4/electrolyte interface, the kinetics of which depends primarily on the radius of the nanoparticles. However, the electrons need to travel to the α-SnWO4/FTO interface for current collection. When the illumination is from the back, the excited electrons are generated close to this interface and only need to travel a short distance. However, when illumination is from the front, excited electrons are generated far from the interface, and need to travel across the entire thickness of the film. As shown in Figure 8b, the - 31 -



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photocurrent for α-SnWO4 films at low voltages is similar for front- and back-side illumination, which indicates that the electron diffusion is likely not a limiting factor in the photocurrent. Therefore, hole transport may be the limiting factor.

Figure 8. Photoelectrochemical response of the α-SnWO4 PF prepared from 8 layers of WO3 PF measured in 1 M Na2SO3 solution at pH 7. (a) IPCE measured at 1.23 VRHE under back-side illumination; (b) Chopped light LSVs measured under front- and back- side illumination.

The photocurrent density of the same sample at lower potentials under backside illumination is higher in Na2SO3 solution than in pH 7 buffer due to more efficient surface - 32 -


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charge transfer resulting from the more favorable thermodynamics and faster kinetics of SO32oxidation compared to H2O oxidation (SO32-/SO3-, Eo = 0.73 VNHE; SO32-/S2O62-, Eo = 0.026 VNHE)62-63. Nevertheless, the photocurrent density reaches ~0.08 mA/cm2 at 1.23 VRHE both with and without Na2SO3 (Figure 8b). In both the LSV curves measured with and without Na2SO3, especially at low potentials, the current density sharply increases when the photoanode is illuminated and decays slowly thereafter. Similarly, the current density drops sharply when the illumination is chopped, and then rises slowly. These photocurrent transients could be due to slow surface reaction kinetics, as has been observed for Fe2O3 and BiVO4,9, 61, 64 but their persistence in the presence of the hole scavenger Na2SO3 (in the aqueous electrolyte) and in the presence of a fast, reversible redox couple (in both the aqueous and non-aqueous electrolytes) suggest that the transients might be instead due to charge localization or trapping within the bulk of the material. This possibility of charge localization could be due to strong exciton binding between photo-exited electrons and holes, as was mentioned in a recent study.18 On the other hand, Ziani et al. measured the dielectric constant, and based on this they claimed that exciton dissociation in α-SnWO4 will be very efficient.23 Combined with our observations from front/back illumination studies (Figure 8b), which suggest that the photocurrent is limited by slow hole transport, we may conclude that hole localization or trapping is a likely explanation for the low photocurrents.

The possibility of hole localization within the SnWO4 films needs some attention in this case. In pristine α-SnWO4, hole localization would involve Sn-s or O-p bands, which requires significant strain in the structure. Alternatively, given the growth condition of the present study discussed in conjunction with Figure 3 and Figure 5, at relatively Sn-rich and W-poor synthesis - 33 -



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condition, defects such as Sn anti-site at W (SnW) are likely to be present with large concentration. Such defects can contribute localized empty bands in the band gap, which can contribute to the charge localization. A charge density plot is shown in Figure 9, which indicates that localization of a charge deficient state does indeed occur at the SnW site. This is interpreted to mean that the SnW defect forms a deep, localized hole state. Comparing band structures of VO - and SnW - α-SnWO4 (Figure 5a), the defect band for SnW has less dispersion than that of VO. This may suggest that the SnW defect may correspond to small polaron conduction. We plan to study this issue further in the future. Alternatively, although this also requires further investigation, the presence of SnO, SnO2 and WO3 impurities in the α-SnWO4 material may play a role as charge trapping sites.

Figure 9. Charge density difference plot for SnW defect in α-SnWO4. The localization of a charge deficient state is clearly seen around the defect site. This is interpreted to mean that the SnW defect forms a deep, localized hole state.

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CONCLUSIONS We have synthesized α-SnWO4 PF photoanodes by hydrothermal conversion of a WO3 porous film. SEM and XRD show that the porous nanostructure was retained after hydrothermal reaction and the product contained primarily orthorhombic α-SnWO4 phase. This photoanode exhibits photon-current conversion at wavelengths up to 700 nm, an indirect optical transition at 2.1 eV, a direct optical transition at 2.4 eV, and a low photocurrent onset voltage of ~-0.1 VRHE. These results on α-SnWO4 compare favorably against the state-of-the-art photoanode material BiVO4 which has a more positive photocurrent onset voltage (~0.1 VRHE) and photon-current conversion up to shorter wavelengths of around 510 nm.10, 62, 65 The photocurrent onset voltage of α-SnWO4 also compares favorably to that of TiO2 (~0 VRHE).6-7 Therefore, α-SnWO4 may have future promise for overall photoelectrochemical water-splitting via water oxidation at its surface and water reduction at a counter electrode. However, the photocurrent generated by the α-SnWO4 PF is very small, which does not match the expected formation of large-polaron carriers by the hybridization of the md0 cation orbitals with O 2p orbitals and the resulting high charge carrier mobilities expected in this material. Our measurements indicate that hole (as opposed to electron) transport may be the limiting factor for the photocurrent. The low photocurrents observed may be due to hole trapping at localized states within the band gap of SnWO4, which are likely caused by Sn-at-W antisite defects. Furthermore, electron trapping at SnO2 and WO3 impurity phases may further decrease the photocurrent. While the bulk impurities and defects seem to originate from the hydrothermal synthesis due to Sn-rich conditions and incomplete conversion of WO3, additional impurity phases are formed during the PEC measurements at higher voltages due to oxidation of α-SnWO4 by water. However, this oxidation

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can be completely suppressed by improving the kinetics of hole transfer at the α-SnWO4/water interface.

The most important strategy to increase the photocurrent of α-SnWO4 photoanodes may be to develop modified synthesis conditions or alternative synthesis methods that avoid the formation of deleterious defects such as SnW and obtain higher phase purity. Other possible strategies, which have previously been applied to BiVO4, include nanostructuring to increase charge separation and transfer,10, 28, 66 inserting interfacial layers to passivate the semiconductor/FTO interface,54 doping or alloying α-SnWO4 with other metals or performing reduction treatments to improve electron concentration,15, 67-68 and adding oxygen evolution catalysts to the surface of α-SnWO4 to enhance the kinetics of water oxidation.62, 69 Moreover, the stability of the α-SnWO4 film could also be improved by adding oxygen evolution catalysts to suppress the competing oxidation of Sn2+ to Sn4+ by photogenerated holes, or by coating a thin layer of protective oxide on the surface.16

ASSOCIATED CONTENT Supporting Information. Spectral output of illumination sources, EIS results, additional results on PEC performance and XPS, and raw data used to generate Figure 7a and Figure 8a. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION - 36 -


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Corresponding Authors *Email: [email protected], Tel.: 817-272-1097 *Email: [email protected], Tel.: 508-831-4828 Author Contributions

Z. Z. carried out the photoanode synthesis, optical, EIS and PEC measurements, and XPS characterization. P. S. and M. N. H. performed all the DFT calculations. C. Z. assisted with the photoanode synthesis. L. Z. performed the SEM and XRD characterizations. R. L. G. advised on XPS characterization. Z. Z., M. N. H. and P. M. R. wrote the manuscript. All authors discussed the results and revised the manuscript. ‡These authors contributed equally. Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT Z. Z. acknowledges support from the WPI Summer Undergraduate Research Fellowship (SURF). This material is based upon work supported by the National Science Foundation under Grants No. DMR1609538 and DMR-1609811.

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Photoelectrochemical Properties and Behavior of Î±-SnWO4

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