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n‑BiSI Thin Films: Selenium Doping and Solar Cell Behavior Nathan T. Hahn,† Alexander J. E. Rettie,† Susanna K. Beal,† Raymond R. Fullon,‡ and C. Buddie Mullins*,†,‡ †

Department of Chemical Engineering and ‡Department of Chemistry and Biochemistry Center for Electrochemistry, Texas Materials Institute, Center for Nano- and Molecular Science, The University of Texas at Austin, C0400 Austin, Texas 78712, United States S Supporting Information *

ABSTRACT: BiSI (indirect band gap = 1.57 eV) is a recently discovered photoelectrode material possessing promising optical properties for use in alternative thin film solar cells. In this work, we study the effects of selenium doping on BiSI film properties and also demonstrate the incorporation of BiS1−xSexI films into both electrochemical and solid state solar cells. Tuning the band gap of BiS1−xSexI by substituting selenium for sulfur was accomplished by substituting various amounts of SeO2 for thiourea in the BiSI spray pyrolysis precursor solutions. This strategy was employed to reduce the direct band gap of BiS1−xSexI films from 1.63 eV to as low as 1.48 eV, as measured by UV−vis−NIR diffuse reflectance spectroscopy for x = 0.4. Both electrochemical and solid state solar cell devices utilizing n-BiSI as the light absorbing material demonstrated open circuit voltages of nearly 0.4 V. The electrochemical devices showed much higher short circuit currents and power conversion efficiencies than the solid state devices. Power conversion efficiencies of up to 0.25 and 0.012% were measured for electrochemical and solid state devices, respectively, under AM1.5G illumination.

1. INTRODUCTION Growing global energy needs in the face of problems related to fossil fuel consumption have pushed research on renewable energy sources to the forefront. Of these sources, solar energy has the greatest potential to meet a TW-scale demand, but its implementation on such a scale is hampered by high costs related to the manufacture and implementation of photovoltaic devices. Much of this cost is associated with the light-absorbing material, which has traditionally consisted of electronic grade, wafer-based silicon. A modern approach is to replace this thick (∼1 mm) absorber layer of single-crystalline silicon with a thin (∼1 μm) film of a polycrystalline material possessing superior optical absorption coefficients in the visible and NIR range.1 This orders-of-magnitude reduction in material requirement could potentially allow for the manufacture of much cheaper solar cells, even if the thin film material itself is rarer and more expensive than silicon, which is extremely abundant.2 Several thin film absorber materials currently receive extensive attention for this purpose, and the most popular thin film absorbers at present are a-Si, Cu(In,Ga)Se2, and CdTe.3−5 The primary issue with the latter two options is the scarcity of indium and tellurium, which are among the most rare and expensive elements on earth.2 Alternative semiconductor compounds consisting of cheaper and more abundant elements are thus needed for the future growth of solar energy on a large scale, and a few candidates such as CZTS (Cu2ZnSnS4) and pyrite (FeS2) have received widespread research interest.6,7 Recently, several bismuth-containing semiconductor compounds have been discovered to possess promising optical, electrical, and/or photoelectrochemical (PEC) properties.8−12 © 2012 American Chemical Society

Although bismuth is indeed more rare than the elements constituting CZTS and pyrite, bismuth compounds are nevertheless interesting candidates for further research as thin film light absorbers for photovoltaic applications. BiSI in particular has an indirect band gap of 1.57 eV, well suited to solar applications, and polycrystalline films deposited by nonvacuum methods have demonstrated photocurrents of over 5 mA/cm2 in PEC cells under AM1.5G illumination with external quantum efficiencies of up to 40% in the visible range.12 These BiSI films were found to be n-type and to possess hole diffusion lengths of roughly 50 nm, as determined photoelectrochemically. However, their PEC instability makes them unsuitable for electrochemical solar cells, and further study of these films is necessary in order to determine their viability in solid state devices. Additionally, manipulation of these films’ optical and electronic properties by doping has yet to be explored and could provide further improvements in performance. In this work, we investigate both the selenium doping of BiSI thin films produced by spray pyrolysis as well as their utilization in both electrochemical and solid state solar cell devices.

2. EXPERIMENTAL METHODS 2.1. BiSI Film Synthesis. General film synthesis procedures were performed in a like manner to our previous study.12 Briefly, precursor solutions containing bismuth, sulfur, and Received: September 5, 2012 Revised: October 24, 2012 Published: October 31, 2012 24878

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iodine were prepared by dissolving 0.02 M Bi(NO3)3·5H2O (Alfa), 0.04 M NH4I (Acros), and the desired amount of thiourea (Fisher) in ethylene glycol (99+%, Acros). Selenium incorporation was accomplished by adding a solution of 0.25 M SeO2 in ethanol to this solution while stirring. The concentration of thiourea in the glycol solution was adjusted depending on the amount of SeO2 to be added such that the (thiourea + SeO2) concentration was maintained at 0.04 M in each case. The precursor selenium doping levels are given by [SeO2 ] [thiourea] + [SeO2 ]

2.3. Characterization Methods. Scanning electron microscopy (SEM) and energy dispersive X-ray spectroscopy (EDX) were performed with a Quanta FEG 650 electron microscope (FEI). X-ray diffraction (XRD) experiments were performed with a Bruker D8 diffractometer, and a mean grain size was calculated using the Scherrer equation:

D=

0.9λ β cos(θ)

(2)

Here D is the mean crystallite size, λ is the X-ray wavelength employed (0.154 nm), and β is the full width at half-maximum of the diffraction peak of interest. X-ray photoelectron spectroscopy (XPS) and ultraviolet photoelectron spectroscopy (UPS) were performed using a Kratos AXIS X-ray photoelectron spectrometer. The XPS spectra were calibrated to the adventitious carbon peak at 284.5 eV. UPS spectra were recorded using 21.2 eV photons produced by a helium lamp. UV−vis−NIR transflectance spectra were measured with a Cary 500 spectrophotometer attached to a Labsphere DRA-CA-5500 integrating sphere. The true value of transmittance (T) was calculated by correcting for substrate and film scattering, and this value was used to calculate film absorbance (A) according to

(1)

Although initially clear, the mixed solutions developed a red precipitate over time and were thus used immediately after mixing. The precursor solutions were loaded into a syringe and pumped intermittently (15 s on, 25 s off) through an ultrasonic spray nozzle (130 kHz, Sonotech) positioned above a hot plate in a ventilated enclosure under normal atmospheric conditions. Prior to deposition, the FTO-coated glass substrates (TEC15, Pilkington) were ultrasonically cleaned with ethanol, rinsed with water, and dried in air, and for deposition, the substrate surface temperature was set at 275 °C, as calibrated using an infrared pyrometer (Microepsilon). 2.2. Solid State Solar Cell Fabrication. Thin film solid state solar cell devices were fabricated utilizing n-BiSI as the light absorbing layer and p-CuSCN as a p-type window layer. The CuSCN layer was applied to the BiSI films using a solution based drop-casting method similar to that performed in the literature.13−15 Briefly, a saturated solution of CuSCN (Alfa) in n-propyl sulfide (98%+, Alfa) was prepared by adding 0.2 g of CuSCN in 10 mL of n-propyl sulfide and stirring overnight. Following stirring, the suspension was allowed to settle for several hours until needle-like crystals formed on the bottom of the beaker. The yellow supernatant liquid was drawn off using a pipet and diluted with an equal amount of n-propyl sulfide to create solutions of half the saturated concentration. In most cases, the BiSI films were treated with a 0.5 M aqueous NaSCN solution for 30 s prior to deposition,14−16 which resulted in much better performance. For application of the CuSCN solution, BiSI films were placed on a hot plate in a ventilated enclosure, and the substrate surface temperature was set to approximately 80 °C. The CuSCN solution was drop-cast onto the films in 5 μL increments for the desired number of droplets (typically 10−15). These droplets were carefully distributed evenly over the film to maximize uniformity. To make electrical contact with the CuSCN layer and complete the cell, platinum coated FTO-glass (Solaronix) was pressed against the film by steel clips. The cells were masked to yield an active area of 1 cm2. A schematic of the n-BiSI/p-CuSCN solar cells is shown in Figure 1.

A = −log10(T )

(3)

The photoelectrochemical properties of the samples were tested using a three-electrode electrochemical cell with a Ag/ AgCl reference electrode and platinum wire counter electrode. All potentials are given relative to Ag/AgCl. The working electrode (the photoanode consisting of the BiSI film) possessed an illuminated area of 0.21 cm2 which was exposed to the desired electrolyte (typically 0.5 M NaI and 0.01 M I2 in acetonitrile). The relevant solar cell parameters of both liquidjunction and heterojunction devices such as open circuit voltage (Voc), short circuit current density (jsc), fill factor (ff), and power conversion efficiency (η) were extracted from I−V curves. Illumination was accomplished by a 150 W Xe lamp with an AM1.5G filter (Newport), the overall power density of which was calibrated to 100 mW/cm2 at the film surface using a thermopile (Newport). A monochromator (Newport) was employed to study the spectral response of the films and was used in conjunction with a monochromatic power meter and photodiode (Newport) to calculate the external quantum efficiencies (EQE) from the measured photocurrent density (jph) and photon flux (I) at a given wavelength (λ) according to EQE =

jph (λ) I (λ )

× 100%

(4)

For the EQE tests, the monochromatic light power incident on the film was between approximately 100 and 500 μW/cm2 depending on the wavelength. A potentiostat (CH Instruments − CHI660D) was operated by a desktop computer to perform the electrical measurements.

3. RESULTS AND DISCUSSION 3.1. Effects of Selenium Doping on Film Properties. 3.1.1. Crystallography and Composition. Precursor selenium doping levels of 0.2, 0.4, and 0.6 were investigated in addition to undoped solutions. Higher doping levels resulted in poorer uniformity, and the films were determined to be principally BiOI. For this reason only, the preceding compositions were studied in depth. Film crystallography was analyzed using XRD

Figure 1. Schematic diagram of a n-BiSI/p-CuSCN solar cell as fabricated. 24879

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Figure 2. (a) XRD patterns for films deposited with the indicated precursor doping levels (Se/(S + Se)). Miller indices for select BiSI diffraction peaks are indicated. (b) Enlarged view of the (120), (200), and (210) peaks. (c) EDX measured atomic ratios of selenium and sulfur to bismuth for various precursor doping levels. (d) Film atomic ratios measured by EDX and plotted vs precursor doping level.

in order to study the effects of selenium incorporation on film structure and phase composition. All films showed a good degree of crystallinity, and the XRD patterns for films deposited with no selenium matched the expected patterns of polycrystalline BiSI (PDF #01-073-1171) (Figure 2a,b). The films deposited with the highest selenium doping levels showed peaks corresponding to a BiOI phase as well. This was most pronounced for a precursor doping level of 60%, which showed pronounced XRD features at 2-theta values of roughly 29.5 and 31.5°, which are the most intense peaks of BiOI patterns (PDF #01-075-5209).11 EDX was used to confirm the bulk compositions of the BiS1−xSexI films. Although the measured ratio of sulfur to bismuth decreased monotonically as their ratio in the precursor solution decreased, the ratio of selenium to bismuth did not show similarly monotonic behavior, peaking at a precursor doping level of 40% (Figure 2c). This implies that having a significant source of sulfur in the precursor solution somehow facilitates selenium incorporation. A reasonable explanation is that selenium incorporates more easily into BiSI than BiOI due to their structural differences. The size difference between the covalent radii of Se (120 pm) and O (66 pm) is much larger than that between Se and S (105 pm), and this means it would be relatively difficult for Se to substitute for O in BiOI. When relatively high sulfur precursor concentrations are present, the dominant phase formed should be BiSI, but as the sulfur concentration in the precursor is diminished (i.e., such that the ratio of S:Bi is less than 1), BiOI apparently becomes a more prevalent phase and the overall Se:Bi ratio is reduced. As shown in Figure 2d, the actual BiS1−xSexI doping levels in the films measured by EDX followed a qualitatively reasonable trend, increasing from x = 0.25 to 0.4 as the precursor amount was changed from 0.2 to 0.6. The total ratio of sulfur and selenium to bismuth in the films decreased to roughly 0.61 for a

precursor selenium doping level of 0.6, indicating that the secondary BiOI phase may account for nearly 40% of the bismuth in these films, assuming there is minimal selenium incorporation into the BiOI phase. In all other films, this ratio was approximately 0.9, indicating that some BiOI formation is probably unavoidable under these conditions. Film synthesis in an inert atmosphere would likely prevent this, although it is not within our experimental capabilities at present. The XRD patterns for films deposited over a range of precursor doping levels showed notable trends in peak position, texture, and mean crystallite size, illustrating the effect of selenium incorporation on crystal structure. The substitution of selenium for sulfur in BiSI should theoretically expand the lattice and thereby reduce the 2-theta values of the corresponding XRD peaks, and this was observed for the Sedoped films. When the 2-theta values of the (200) and (040) peaks were used in accordance with Bragg’s law to calculate the a and b parameters of the BiSI unit cell as a function of film Se/ (S + Se) composition, a clear trend emerged, indicating that the BiSI lattice was expanded toward that of BiSeI (Figure 3a). Additionally, the crystallographic texture of the films increased with higher Se-doping levels, and their preferred orientation grew toward the (121) plane, which is also consistent with the powder diffraction standards of BiSI and BiSeI (Figure 3b). Another characteristic of the XRD patterns is that the peaks became broader as Se-doping increased, implying a decreasing trend in the average crystallite size as calculated by the Scherrer equation (Figure 3b). 3.1.2. Film Morphology. This trend in crystallite size was also observed in SEM images (Figure 4), which showed that the film features changed from large micrometer-scale rods to smaller, cube-like structures. The sizes of these features are still much larger than the crystallite sizes calculated from the XRD patterns, meaning that these features are not single crystalline, 24880

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although the individual crystallites do appear to be highly oriented within each feature based on the SEM images. This trend toward smaller features suggests that Se-doping disrupts the growth of large microrod features typically observed for BiSI films deposited by spray pyrolysis.12,18 Another striking feature of the SEM images is the formation of large, thin square platelets, reminiscent of BiOI as observed in a previous study.11 Actually, smaller nanoplatelets could also be observed on sections of films deposited with lower Se-doping levels, but their coverages were relatively insignificant. This mild BiOI formation on the surface of BiSI is consistent with the fact that the EDX (Se + S)/Bi ratio does not exceed ∼0.95 for any of the films. 3.1.3. Optical Properties. To determine whether the substitution of selenium for sulfur in the films led to the desired changes in optical absorption, UV−vis−NIR diffuse transmittance and reflectance experiments were performed. Absorbance values were calculated for films of each composition, taking into account the transmittance of the FTO/glass substrate (Figure 5a). Absorbance values of 2−2.5 were measured for wavelengths up to 600 nm, indicating that the visible light absorption was very good. The absorbance onsets of the films showed a pronounced red shift from wavelengths of 800−880 nm, representing a potential 19% increase in solar photon utilization.19 Tauc plots of (αhν)2 vs photon energy were used to estimate the direct band gaps of the films (see the Supporting Information, Figure S-1). The fundamental transitions in V−VI−VII compounds are actually indirect,20,21 but the value of the direct band gap is more relevant for thin film solar cell materials, since high absorption coefficients (as found for energies above the direct gap) are critical. For the range of compositions tested, the direct band gaps showed an almost linear decrease with

Figure 3. Crystallographic characteristics as a function of selenium doping level measured by EDX. (a) Variation in the unit cell parameters a and b calculated from the 2-theta values of the (200) and (040) peak positions (circles) compared to the theoretical values for BiSI and BiSeI (triangles), which are connected by dashed lines. (b) Peak height ratio of the (120):(121) planes and the mean grain size calculated from the (210) peak in the XRD patterns.

Figure 4. SEM images of Se-doped BiSI films deposited with precursor Se/(S + Se) levels of (a) 0.0, (b) 0.2, (c) 0.4, and (d) 0.6. The scale bar applies to all images. 24881

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Se-doping (BiSeI) would result in a direct band gap of about 1.27 eV, which is slightly lower than the reported indirect band gap of 1.29 eV.20,21 Band gap vs composition trends in compound semiconductors typically exhibit “bowing”, meaning the relationship is not linear over the entire composition range,22 resulting in an underestimated band gap for BiSeI using a linear extrapolation over these compositions. 3.1.4. Photoemission Spectroscopy. XPS analyses were performed in order to further study the surface chemistry of the Se-doped BiSI films. Core level spectra for the Se 3d regions were used to examine the chemical nature of the incorporated selenium (Figure 6a). The selenium doped films showed two small but discernible peaks at approximately 53.4 and 54.2 eV, similar to spectra previously observed for Bi2Se3, indicating that the selenium is chemically incorporated as Se2+.16 Significant peaks were not observed at 58−59 eV (slight humps could be observed for the more heavily Se-doped samples), indicating that very little, if any, SeO2 was formed. The peak areas of the elemental XPS spectra were used along with instrument sensitivity factors to quantify the elemental composition of the surfaces. The general trend of Se-doping level observed at the surface of the films was similar to that observed in the bulk by EDX (Figure 6b). It is not clear why the surface Se-doping level decreased as the precursor level increased from 0.2 to 0.4, but this behavior was consistent for two separate samples. Consistent with the previous discussion in section 3.1.1, the XPS spectra of heavily doped films show evidence of surface BiOI formation. Changes in the Bi 4f and O 1s XPS spectra for these films were similar to those observed previously for BiSI films synthesized at higher temperatures (see the Supporting Information, Figure S-2).12

Figure 5. (a) Absorbance spectra calculated from the UV−vis−NIR diffuse reflectance spectra for films with different Se/(S + Se) doping levels as measured by EDX. (b) Direct band gap values determined from plots of (αhν)2 vs hν for these films. The dashed line represents a linear extrapolation of the trend toward BiSeI.

increasing Se-doping level, dropping from 1.63 to 1.48 eV for x = 0.6 (Figure 5b). However, extrapolation of this trend to 100%

Figure 6. Photoemission spectroscopy results for films of various Se-doping levels. (a) XPS Se 3d core level regions for various bulk Se-doping values along with the relevant peak designations. (b) Comparison of the bulk and surface Se-doping values as measured by EDX and XPS, respectively. (c) Normalized UPS spectra aligned such that the Fermi edge is set to 0 eV. (d) Valence band maxima calculated from the UPS spectra. 24882

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unable to diffuse into the depletion region, they will simply recombine with electrons. The transport of deeply generated holes to the film surface might be accomplished by improving material quality (fewer defects, etc.) or by reducing the required transport distance through nanostructuring. Film deposition strategies utilizing inert atmospheres or high vacuum may be capable of producing higher quality films and thereby improve the charge transport properties, although this would increase the cost associated with thin film production. The EQE spectra confirm that changes in absorption brought about by Se-doping lead to the improved conversion of NIR photons, indicating that the band gap of BiSI can be effectively tuned using this method. Unfortunately, this improved conversion of NIR photons came at the cost of the conversion of UV and visible photons. For example, the EQE values at λ = 480 nm exhibited by the Se-doped films were roughly 70% of those exhibited by undoped BiSI. This indicates that the rate of recombination was enhanced by the incorporation of selenium to some degree, possibly due to the inclusion of defects and/or a reduced degree of crystallinity. Linear voltammetric sweeps were performed in the same solution in the dark and under AM1.5G illumination to generate I−V curves (Figure 8). A fill factor (ff) of 0.28 and

UPS measurements were also conducted to determine the influence of Se-doping on the band edge positions of n-BiSI. In general, the films’ valence band spectra did not possess any striking features and were largely similar regardless of selenium content (Figure 6c). The valence band maximum (VBM) of the undoped BiSI film measured by this technique was −6.1 ± 0.2 V vs vacuum level, and selenium incorporation appeared to lower this value by 0.1−0.2 eV at the highest Se-doping levels, which may be convoluted by the presence of BiOI on the surface (Figure 6d). This general behavior is somewhat surprising, since an upward shift in VBM is typically the case for band gap reduction via anion substitution (e.g., CdS1−xSex or CdSexTe1−x).17,18 However, unlike in the aforementioned cadmium chalcogenides, the VBM of BiSI is predicted to consist primarily of I 5p orbitals, rather than those of S 3p, and this may explain the observed behavior.23−25 Apparently, the band gap reduction in BiS1−xSex is caused by lowering of the conduction band minimum (CBM) as selenium is added. An alternative explanation would be that the band gap reduction is actually due to the introduction of defect states just above the VBM of BiSI upon the introduction of selenium, although these defect states would probably not explain the strong shifts in absorbance and EQE spectra. 3.2. Electrochemical Solar Cell Behavior. The PEC performance of the BiS1−xSexI films was tested using a twoelectrode cell containing the 0.5 M NaI/0.01 M I2 couple in acetonitrile and a platinum counter electrode. Although the films were not stable in this electrolyte for long periods of time due to the presence of reactive I3− anions which slowly attacked the films even under open circuit conditions, this electrolyte is a useful benchmark, since I/I2 has been the standard redox couple for layered n-type chalcogenide studies in the past.26 To study the wavelength dependence of the photocurrent response, EQE measurements were performed under shortcircuit conditions (0 V vs Pt) employing illumination through the back sides of the films (Figure 7). Although previous

Figure 8. Photoelectrochemical I−V curves for films with various Se/ (S + Se) levels measured in a two-electrode cell containing I−/I2 in acetonitrile under AM1.5G illumination.

power conversion efficiency of 0.25% were calculated for undoped BiSI along with a Voc between 0.35 and 0.38 V and a jsc of nearly 2.5 mA/cm2 for typical films. The Se-doped films exhibited slightly smaller short circuit currents (jsc) and open circuit voltages (Voc) and, as a result, lower efficiencies. Photocurrents exhibited by the 32% Se-doped films were consistently worse than the other compositions, consistent with the EQE spectra. At present, it is not clear why the photocurrent performance increases again at higher Se-doping levels (i.e., 40%), although the growth of the secondary BiOI phase in these films may somehow benefit charge separation. In general, the I−V curves do not show ideal behavior and give the impression of a large cell resistance. However, this slow increase in photocurrent as the voltage is swept toward reverse bias is consistent with the results previously observed in threeelectrode cell experiments, implying that this phenomenon is related to the BiSI film itself and/or the film-electrolyte interface.12 This previous investigation suggested that poor charge separation and transport within the BiSI films was the main culprit giving rise to this nonideal behavior, and we believe that this is the case in the current study as well.

Figure 7. Short-circuit IPCE values measured for films with various Se/(S + Se) levels measured in a two-electrode cell containing I−/I2 in acetonitrile.

experiments on undoped BiSI revealed that higher EQE values were typically attained using front side illumination, backside illumination was required in this case due to the opacity of the electrolyte solution. The difference in performance is probably related to the large hole effective mass in BiSI in general24 and the relatively short hole diffusion length of spray-pyrolysis deposited n-BiSI of about 50 nm.12 During backside illumination, most of the holes are photogenerated farther from the film-electrolyte interface rather than within the depletion region. If these holes are 24883

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Figure 9. (a) Proposed band edge alignment scheme at open circuit for a n-BiSI/p-CuSCN heterojunction based on the UPS measurements. Estimated semiconductor Fermi levels are denoted by Ef. (b) Cut-away SEM image of a n-BiSI/p-CuSCN heterojunction film deposited on FTO. The scale bar is 1 μm.

A noticeable difference between the I−V curves of the doped and undoped films is in the onset and magnitude of the cathodic dark current, which was much more significant for the doped films, leading to decreased Voc. Cathodic dark currents in electrochemical solar cells tend to be facilitated by midgap surface states, which can act as pathways of electron transfer from the film to the electrolyte (i.e., shunting).27,28 From the dark current behavior, it appears that the influence of these surface states on cell performance follows a consistent trend with increasing Se-doping level. This would suggest that surface state passivation may afford significant improvements in the performance of the heavily Se-doped films. If these surface states can be blocked by some kind of insulating species, as has been accomplished previously with other layered semiconductor photoanodes such as n-WSe2 and n-MoSe2, then improvements in Voc, overall efficiency, and stability could be expected.27−30 This represents an interesting avenue of future work for PEC cells based on n-BiSI, but in the meantime, solid state cells would seem more tenable. 3.3. Solid State Solar Cell Behavior. Due to the electrochemical instability of BiSI, the incorporation of BiS1−xSexI films into solid state solar cells is of particular interest. In this work, we utilized a depleted heterojunction scheme to demonstrate the first such attempt, employing CuSCN as a p-type window layer. CuSCN has received great interest recently as a hole conducting material for extremely thin absorber (ETA) solar cells, and the deposition of CuSCN is fairly straightforward and has given promising results in dyesensitized TiO2 cells.13−15 The valence band edge of p-CuSCN lies at −5.3 V vs vacuum level, whereas the conduction band of BiSI lies at roughly −4.5 V, based on the UPS and UV−vis− NIR analyses. If one assumes that the Fermi level of each semiconductor resides at a typical value of 0.1 V from the VBM and CBM, respectively, then the band alignment of the heterojunction would be capable of providing up to 0.6 V at open circuit, ignoring all losses (Figure 9a). After fabrication, some of the heterojunction films were cut in half and examined using SEM (Figure 9b). Typical heterojunction films consisted of BiSI films of 0.5−1 μm and CuSCN films of 2−4 μm thickness. When the heterojunction solar cells were irradiated from the CuSCN side under AM1.5G illumination, Voc values of up to 0.39 V were measured along with jsc of up to 70 μA/cm2 and η of up to 0.012%. Typical solar cell I−V curves such as those shown in Figure 10 demonstrated fill factors (ff) of roughly 0.4. In comparison to the liquid-junction cells, the heterojunction

Figure 10. I−V behavior of a BiSI/CuSCN solar cell by illuminating from either the (i) BiSI side or (ii) CuSCN side using full spectrum and λ > 420 nm illumination.

cells afforded better ff but similar Voc and much lower jsc. Unfortunately, solar cells fabricated from the Se-doped films performed very poorly. As in the case of the electrochemical solar cells, the Se-doped films exhibited worse defect-mediated shunting behavior than undoped films, which greatly limited Voc and ff. The solid state cells appear to be more severely affected by these selenium-induced defects, indicating that those present at the CuSCN−BiS1−xSexI interface are more deleterious to performance than those present at the electrolyte−BiS1−xSexI interface. This may be due to the fact that surface-trapped holes are more easily transferred to I− (Eo = −5.0 V) than to the valence band of CuSCN (EVBM = −5.3 V). Due to the very poor performance of the solid state solar cells incorporating Sedoped films, we focused primarily on characterizing cells incorporating undoped BiSI films. Although the Voc and ff of the solid state cells could certainly be improved by reducing parasitic losses due to shunting and series resistance through cell optimization, their extremely small photocurrents compared to those of the electrochemical cells (