Effects of Thermochemical Treatment on CuSbS2 Photovoltaic

Aug 1, 2016 - These improvements also lead to more reproducible CuSbS2 PV devices, with performance currently limited by a large cliff-type interface ...
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Effects of Thermochemical Treatment on CuSbS2 Photovoltaic Absorber Quality and Solar Cell Reproducibility Francisco Willian de Souza Lucas,†,‡ Adam W. Welch,†,§ Lauryn L. Baranowski,†,§ Patricia C. Dippo,† Hannes Hempel,⊥ Thomas Unold,⊥ Rainer Eichberger,⊥ Beatrix Blank,∥ Uwe Rau,∥ Lucia H. Mascaro,‡ and Andriy Zakutayev*,† †

National Renewable Energy Laboratory, 15013 Denver West Parkway, Golden, Colorado 80401, United States, Federal University of Sao Carlos, Road Washington Luiz, km 235, São Carlos, SP 13565-905, Brazil § Colorado School of Mines, 1500 Illinois Street, Golden, Colorado 80401, United States, ⊥ Helmholtz Zentrum Berlin für Materialien und Energie GmbH, Hahn Meitner Platz 1, 14109 Berlin, Germany ∥ IEK5-Photovoltaik, Forschungszentrum Juelich GmbH, Wilhelm-Johnen-Straße, 52428 Juelich, Germany ‡

S Supporting Information *

ABSTRACT: CuSbS2 is a promising nontoxic and earth-abundant photovoltaic absorber that is chemically simpler than the widely studied Cu2ZnSnS4. However, CuSbS2 photovoltaic (PV) devices currently have relatively low efficiency and poor reproducibility, often due to suboptimal material quality and insufficient optoelectronic properties. To address these issues, here we develop a thermochemical treatment (TT) for CuSbS2 thin films, which consists of annealing in Sb2S3 vapor followed by a selective KOH surface chemical etch. The annealed CuSbS2 films show improved structural quality and optoelectronic properties, such as stronger band-edge photoluminescence and longer photoexcited carrier lifetime. These improvements also lead to more reproducible CuSbS2 PV devices, with performance currently limited by a large cliff-type interface band offset with CdS contact. Overall, these results point to the potential avenues to further increase the performance of CuSbS2 thin film solar cell, and the findings can be transferred to other thin film photovoltaic technologies. electricity conversion efficiency (certified 12.6%),8,20 but its chemical complexity may hinder its further improvements. These considerations have led to a resurgent interest in lesscomplex nontoxic earth-abundant ternary copper sulfides, such as Cu3P(S,Se)4,21 Cu2SnS3,22,23 and CuSbS2.24 In particular, chalcostibite CuSbS2 is quite interesting since it is crystallographically very different than sphalerite CTS, chalcopyrite CIGS and kesterite CZTSSe. This layered CuSbS2 material features a slightly indirect band gap of 1.4−1.5 eV (falling within the optimum range for a single junction terrestrial solar cell),23,25−29 and strong optical absorption with an extinction coefficient of more than 105 cm−1 at 1.8 eV (slightly better than the aforementioned sphalerite, chalcopyrite and kesterite materials).23,25,28 The CuSbS2 also has tunable hole concentration in the 1016 − 1018 cm−3 range, controlled by the Sb2S3 overflux in a self-regulated growth process.25,28 In addition, CuSbS2 has a melting point of 535 °C, amenable to grain growth at relatively low temperatures.23,25−29

1. INTRODUCTION In the present world development scenario, one of the great challenges is the search for abundant and sustainable sources of energy and fuel.1 Solar energy has been highlighted as one of the most important renewable sources to meet the future global demands. Different kinds of solar cells have been developed, including organic bulk heterojunction,2 dye-sensitized,3 hybrid organic inorganic perovskite,4,5 for both photovoltaic (PV) energy generation and photoelectrochemical (PEC) fuel production.6 Among these, thin film PV are promising technologies to further decrease the present cost and address the economic scalability of Si-based solar cells that dominate the market.7 Some of the most studied thin film solar cell technologies are Cu(In,Ga)Se2 (22.3% energy conversion efficiency) and CdTe (22.1% efficiency).8 However, the scarcity of raw materials (In, Ga) and the elemental toxicity (Cd) may limit environmental benefits and potential terawatt scaling of these solar cell technologies.9 Thus, emerging light absorber materials with less toxic and more common elements are of increasing interest, including Cu2O,10 Sb2Se3,11 Cu3N,12,13 SnS, 14,15 ZnSnN 2 ,16 BiFeO 3 , 17 Cu 2 SnS 3 (CTS), 18 and Cu2ZnSn(S,Se)4 (CZTSSe).19 Among these, the quaternary CZTSSe is the most promising in terms of reported PV solar to © 2016 American Chemical Society

Received: April 26, 2016 Revised: June 28, 2016 Published: August 1, 2016 18377

DOI: 10.1021/acs.jpcc.6b04206 J. Phys. Chem. C 2016, 120, 18377−18385

Article

The Journal of Physical Chemistry C

(Alfa, 99.999%), which was placed close to the CuSbS2 film and on a glass plate. Since the powder inside the alumina boat takes a longer time to heat up and cool down compared to the plate and the sample, small Sb2S3 crystallites (1 g) were placed outside of the boat on the glass plate for maintaining constant Sb2S3 partial pressure during the heating. The sample and the Sb2S3 sacrificial sources were covered with a larger boat (50 mL), as shown in Figure 1A, so the Sb2S3 vapor would quickly

Despite these promising material properties, the highest reported efficiency for a single CuSbS2-based PV device is 3.1%, with an open-circuit voltage (VOC) of 490 mV, short circuit current (JSC) of 14.7 mA/cm2, and fill factor (FF) of 0.44.26 Similar performance (>3%) of a larger set of related CuSbSe2 devices has also been reported.30 Several groups27,28 have shown that the CuSbS2 PV performance is improved by incorporating an intrinsic Sb2S3 interlayer between the p-type CuSbS2 (absorber) and n-type CdS (buffer) layers (from 0.5% to 1%), but with little information about the underlying optoelectronic properties of the absorber. Altogether, there seems to be two factors that currently limit the progress of the CuSbS2 PV technology. First, literature reports on CuSbS2 materials for PV application often provide limited characterization of the properties relevant for photovoltaic device performance, in particular as a function of material structure and morphology. Second, the control of homogeneity of the CuSbS2 thin films, and hence reproducibility of the CuSbS2 devices is challenging.25 In this paper, we demonstrate how thermochemical treatment of CuSbS2 improves the quality of this absorber material and its PV device reproducibility. The thermal treatment in Sb2S3 vapor is followed by a selective KOH chemical etch for removing possible Sb2S3 (and Sb2O3) surface layers due to the Sb2S3-rich experimental conditions. We describe the improvements in solar cell device performance, and most importantly homogeneity, as a result of these thermochemical treatments. The observed effects are correlated with the improvements of the CuSbS2 structural and optoelectronic properties relevant to its PV device performance. Finally, we discuss the performance limiting factors for the thermochemically treated CuSbS2 photovoltaic devices and suggest future steps for further efficiency improvements.

Figure 1. (A) Optimized experimental apparatus and (B) temperature/flow profile graph of the thermal treatment experiments.

achieve equilibrium conditions. Overall, this Sb2S3 annealing approach has an advantage over the S annealing widely used for Cu2ZnSn(S,Se)4 processing,32,33 since it eliminates the chances of decomposition of CuSbS2 into binary competing phases and hence enables long annealing times necessary for the grain growth. For improving the temperature homogeneity along the sample surface, the sample was placed on a thermally conductive Si plate (as shown in Figure 1A), and the part of the tube furnace in contact with the sample plate was wrapped with aluminum foil (not shown in Figure 1A). The film temperature was calibrated by a thermocouple attached to the Si-plate (Omega Engineering Inc., Model 115 KC). The temperature profile used for the TT (Figure 1B) consisted of a 30 min of purge with high flux of N2(g) at room temperature (O2(g) removing step), followed by 30 min at 100 °C (H2O(g) removing step), and by a heating ramp from 100 °C to optimized temperature (450−500 °C). The furnace remained at this optimized temperature during the determined annealing time (5−15 h) before a free cooling step back to room temperature (∼5 h). In order to achieve good vapor equilibrium, but still maintain the atmosphere inside the larger boat free of O2(g), the flux of N2(g) was adjusted as low as possible after the purging step; for all steps, the gas exited the tube furnace through a bubbler filled with mineral oil. 2.3. Film Characterization. All CuSbS2 films were characterized before and after thermal treatment. Phase identification was performed via X-ray diffraction (XRD) (Bruker D8 Discover). The composition and thickness of the films were obtained by X-ray fluorescence (XRF) (Fischerscope X-ray XDV-SDD, Helmut Fischer GmbH). The XRF measurements were taken in air, so it was difficult to resolve a sulfur signal, thus full stoichiometry was extrapolated from determined Cu/Sb ratio combined with XRD results. Scanning electron microscope (SEM) images were taken using a FEI Quanta 600 instrument. Optical data was taken on a Cary 5000 UV/vis/NIR spectrometer in diffuse transmission and reflectivity modes. Room temperature photoluminescence (RT-PL) was performed with a 632.8 nm laser line at 16 mW, 280 μm slit, 695 nm long pass filter and 5 s exposure time. The mobility and concentration of electric charger carriers were obtained by Hall effect measurements (Accent HL5500PC) using van der Pauw configuration at room temperature. The

2. EXPERIMENTAL SECTION 2.1. Absorber Film Deposition. CuSbS2 films were deposited by a self-regulated growth approach25,31 using radio frequency (RF) magnetron cosputtering of one Cu2S and two Sb2S3 targets (99.999%). Phase-pure stoichiometric films were deposited on 50 × 50 mm Eagle XG substrates (for optical and electrical characterization), or on Mo/SiO2 soda lime glass (SLG) substrates. Depositions were performed in 3 mTorr of Argon (99.99%) in a chamber with a base pressure of 10−7 Torr. The resulting absorber film thickness was in the range of 1.4−1.6 μm, and was controlled by the deposition time (∼4 h). Sputter gun power was 40 W for all targets. More details of these CuSbS2 thin film deposition have been previously reported.25 Individual CuSbS2 samples were obtained by cleaving the as-deposited 50x50 mm films into 12.5 × 50 mm pieces before thermal treatment. The resulting pieces were further subdivided by mechanical scribing into 11 mm × 4 mm sections (44 mm2) for property device measurements after the thermal treatment. Note that the as-deposited films subjected in this manuscript to Sb2S3 thermal treatment are already crystalline, phase pure and stoichiometric CuSbS2; this is in contrast to other reports in literature where the as-deposited Cu/Sb metal precursor stacks are annealed in S atmosphere to form the CuSbS2 phase. 2.2. Thermal Treatment. The CuSbS2 thin film thermal treatment (TT) experiments were performed in a one-zone tube furnace in N2 (99.99%) at ambient atmosphere. The volume around the film was maintained rich with Sb2S3 vapor by the use of an alumina boat (5 mL) with 5 g of Sb2S3 powder 18378

DOI: 10.1021/acs.jpcc.6b04206 J. Phys. Chem. C 2016, 120, 18377−18385

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The Journal of Physical Chemistry C conductivity (in dark) and photoconductivity (PC) of the films were characterized by a custom four-point probe coupled to a red laser pointer (∼5 mW output power, 650 nm). The charge carrier mobility and dynamics were characterized by optical pump terahertz probe (OPTP) spectroscopy. In the OPTP measurements an 805 nm pump pulse induced a peak carrier concentration of 1.8 × 1018 cm−3 in the sample. The decay of this carrier concentration was sampled by the reflection of a 0.3−3 THz probe pulse, which is sensitive to the conductivity of the sample. Further a transfer matrix method was used to model the THz reflection spectrum, which yields as optimized parameter the complex mobility spectrum of the excited charge carriers. The obtained mobility spectrum comprises the sum of electron and hole mobilities.34,35 As such, this methods is similar to the time-resolved absorption spectroscopy that is often used to study photophysics of organometallic complexes,36 except it uses THz rather than visible/infrared femtosecond pulse as a probe. 2.4. PV Device Experiments. A PV devices were made from as-deposited and annealed CuSbS2 films on the SLG/Mo substrates, finished by chemical bath deposition (CBD) of CdS, an RF sputtered intrinsic/conductive transparent i-ZnO/ ZnO:Al (iZO/AZO), and an electron-beam evaporated aluminum grid, as previously described.37 Prior to the PV device fabrication, a selective chemical etch for Sb2S3/Sb2O3 was developed and applied to all CuSbS2 absorber films (both as-deposited and thermally treated), in order to eliminate the potential influence of a possible Sb2S3 or Sb2O3 surface layer on the PV device performance.27,28 The absorber films were immersed in a 0.1 mol L−1 KOH bath for 30 min, then washed with deionized water, dried with N2(g) and quickly passed to the next step of device fabrication (CdS deposition). This procedure enables us to study the thermochemical treatment effect on the films’ properties without the possible effect of secondary uncontrolled parameter (a possible thin Sb2S3 surface layer formed on film after the annealing). As shown in the Supporting Information, this chemical etch is selective to Sb2S3 in agreement with our previous publication;38 in addition, the surface etch changes the shape of the photoluminescence spectrum, likely due to decrease in the surface recombination (Figure S1). The resulting devices, with an active area of 0.44 cm2 were characterized by cross-sectional SEM, external quantum efficiency (EQE), and current−voltage (J−V) measurements in dark and under simulated AM 1.5G illumination (100 mW cm2) at 25 °C. The reference solar spectral irradiance AM 1.5G39 was used as base for integrating the EQE curves. The dark−J−V curves were fitted using a one-diode model including series and parallel resistance (four parameters model). The equation used in the modeling is as follows:40−43 ⎡ q J = J0 × ⎢exp (V − JRS) ⎣ nkT

{

⎤

} − 1⎥⎦ + V −R JR P

and the temperature range from 120 to 320 K was covered by 21 equidistant steps with an accuracy of 0.5 K.

3. RESULTS AND DISCUSSION 3.1. Structure and Morphology. Thermal treatment of sputtered CuSbS2 thin films in vapor of Sb2S3 sacrificial powder was performed in N2 flow furnace, described in more details in the Experimental Section. First, we performed thermal treatments of CuSbS2 thin films for 1.5 h at varying temperatures of 455, 490, 515, and 580 °C. A decrease in thickness of the film and formation of pinholes were observed for the samples treated at 580 and 515 °C, indicating sublimation of the CuSbS2 film. At these high temperatures, we did not observe the CuSbS2 film decomposition to Cu12Sb4S13, which was expected according to the published theoretical phase,25 and based on the tabulated Sb2S3 vapor pressure at 515−580 °C (10−3−10−4 Torr44). This difference between the experimental result and the theoretical prediction might be due to the 1 atm of N2 present during the thermal treatment, which was not taken into account by the previous calculations.25 It appears that this N2 is effective in suppressing sublimation of Sb2S3 from the CuSbS2 film, preventing it from decomposition into Cu12Sb4S13. This suggests that further theoretical studies of the CuSbS2 phase diagram and Sb2S3 chemical potential may be needed, similar to what was recently performed for the S vapor.45 After determining the highest temperature at which the CuSbS2 films could be thermally treated without affecting their thicknesses (455 °C, 85% of the CuSbS2 melting point), we performed thermal treatments (TT) for different durations (5, 8, and 11 h). A 14 h condition was also tested, but produced high pinhole density and morphological nonuniformities that would later limit photovoltaic performance. On the basis of XRF characterization, we concluded that the thermal treatment in Sb2S3 vapor did not change the stoichiometry or thickness of the films. The film maintained slightly Cu-poor cation ratio, Cu/(Cu + Sb) = 0.48, for both as-deposited and annealed films, within the error of the XRF measurement from the ideal Cu/ (Cu + Sb) = 0.50 stoichiometry. The sulfur composition was difficult to determine from the XRF measurements in a quantitative way due to limitation of this technique, so the overall CuSbS2 film stoichiometry was deduced from the XRD pattern match with the CuSbS2 reference pattern. XRD patterns of the CuSbS2 films on Mo/SLG thermally treated at 455 °C for different lengths of times can be seen in Figure S2 (Supporting Information). The films were composed of phase-pure CuSbS2 (no secondary phases appear), even after 11 h of thermal treatment, indicating that single-phase CuSbS2 is the thermodynamically stable state at the thermal treatment conditions used in these experiments. This, along with the constant chemical composition, demonstrates the significant advantage of Sb2S3 vapor annealing over the annealing in H2S gas and S-containing vapor, for which CuSbS2 (with Cu1+ and Sb3+) can turn into other phases with higher valence states of the metallic elements (Cu2+ and/or Sb5+). From the CuSbS2 XRD patterns (Figure S2), the unit cell volume could be calculated, as shown in Figure 2. It is clear that the unit cell volume changes toward equilibrium, which indicates a decrease in the lattice stress with increased thermal treatment time. Evaluating the SEM images, shown in the inset of Figure 2, it can also be seen that an increase in grain size occurs with annealing for 8 h, but the grain size does not increase further with increasing annealing time (Figure S3,

S

(1)

Here J0 is the diode saturation current, q is the electron charge, kT is the thermal energy (0.02569 eV at 25 °C), n is the diode ideality factor, V is the applied bias voltage (in forward direction), and RS and RP are the series and parallel resistances, respectively. Finally, the temperature-dependent J−V measurements were performed under laser illumination with a centerwavelength of 462 nm. The intensity of the laser light was chosen to match the JSC of the respective cell under AM 1.5G, 18379

DOI: 10.1021/acs.jpcc.6b04206 J. Phys. Chem. C 2016, 120, 18377−18385

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In the inset in Figure 3B, the Hall measurement data for the pre-TT and TT-11 h films is provided. As seen in Figure 3A, the TT-8 h and TT-11 h films have more symmetric peak shapes and smaller full width at half maxima (fwhm) than pre-TT and TT-5 h films. This indicates that the TT eliminates tail states at the band edge in CuSbS2, or leads to more homogeneous properties of the grains in this polycrystalline material. The decrease in tail states may also explain why the PL peak moves closer to the band edge after annealing. Additionally, the near-band edge PL intensity for the TT-11 h film was three times higher, perhaps due to less nonradiative recombination, which can be caused by midgap or surface states. From Figure 3B, we can see the TT-11 h film had the highest photoconductivity. The conductivity more than doubled under red laser light, and was almost 2 times higher than the PC observed for the TT-8 h film, indicative of better optoelectronic properties. This behavior is in agreement with the observed results from the PL spectra (Figure 3A), where the TT-11 h films also showed higher optoelectronic quality. It was also observed that the TT-11 h film had a higher conductivity in dark (by 1 order of magnitude) than the other films, even compared to the TT-8 h film that had almost the same grain size (see Figure S3 in the Supporting Information). To understand the origin of this increase in conductivity, Hall effect measurements at room temperature (inset in Figure 3B) were performed on pre-TT and TT-11 h films on glass substrate. From these measurements, it was concluded that the TT-11 h film had a higher hole concentration and mobility (ρ = 1017 cm−3 and μ = 4.1 cm2 V−1 s−1, respectively) compared to the pre-TT film (ρ = mid-1016 cm−3 and μ = 2.5 cm2 V−1 s−1). The small increase in carrier concentration can be related to an increase in Cu vacancy (VCu), since the thermal treatment was performed under low Cu chemical potential (Sb2S3-rich). This kind of VCu defect in CuSbS2 is a shallow hole-producer (acceptor defect) with lowest formation energy in all chemical potential conditions.23,28,47,48 At the Sb2S3-rich condition, another important expected defect is S vacancy (VS),23 which has levels closer to midgap. The VS donor-like ionization level is calculated to be 0.2 eV23 above valence band, so it may behave as a very deep recombination-center in CuSbS2. The electrons that are supposed to go to the CdS contact conduction band can instead recombine at this defect in CuSbS2. This observation suggests the annealing conditions should be wellcontrolled: the initial small increase in VCu density may improve the absorber quality, but further increases of this defect can

Figure 2. Unit cell volume for the CuSbS2 films on SLG/Mo as a function of annealing time. The inset shows the SEM surface images of the pre-TT and TT-11 h CuSbS2 films. As shown by the dashed red line, the unit cell volume of nonstressed orthorhombic CuSbS2 powder (ICSD-85133) is 331.39 Å3.46

Supporting Information). However, for the longest annealing times (11 h) some pinholes and delamination were also present. It was not possible to determine the grain size of the CuSbS2 films from the XRD peak widths and to compare them with the SEM results, because the XRD instrumental broadening was larger than crystallite size broadening. 3.2. Optoelectronic Properties. The absorption coefficient (α) and the band gap energy (Eg) of the CuSbS2 films on EXG glass were obtained from diffuse transmission and reflection spectra, coupled with film thickness measurements from XRF. The TT-11 h annealed films showed an optical Eg of 1.5 eV (Figure 3A), consistent with other research.25,28 The graphs of α and (αhν)2 as a function of the photon energy for all other annealing times can be seen in Figure S4 (Supporting Information). No significant variation of the films’ band gap was observed as a function of the thermal treatment time, which agrees with the fact that this treatment does not change the films’ phase and elemental composition. To complement the evaluation of the optical properties of the CuSbS2 films, the photoluminescence (PL) and photoconductivity (PC) experiments were performed on these films at room temperature. The higher intensities of the band-edge PL and larger values of PC are usually indicative of higher optoelectronic quality of the photovoltaic absorber material. Parts A and B of Figure 3 show the PL and the PC for CuSbS2, respectively, where PC is defined as the percentage increase of conductivity under light compared to the conductivity in dark.

Figure 3. (A) Room temperature absorption- and photoluminescence spectra, and (B) Photoconductivity (red solid squares, left scale) and conductivity in dark (blue open circles, right scale) for the CuSbS2 films. Inserted: Room temperature hole concentration (pRT) and mobility (μRT) for the pre-TT and TT-11 h films. 18380

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the layered CuSbS2 contrast with the significantly different hole and electron mobilities in CuInS2 and related 3D-bonded materials (CIGS, CZTS, CTS). This difference can be attributed to the Sb p-states at the conduction band minimum (CBM) of CuSbS2, compared to the In s-states at the CBM on CuInS2. Also, the absolute mobility determined here for our CuSbS2 samples (4 cm2/(V s)) is much lower than that for high-quality CIGS samples (up to 1200 cm2/(V s)).50 Another important factor that sets the layered CuSbS2 aside from the multinary tetrahedrally bonded semiconductors is the large structural difference between the tetrahedral Cu sites and the distorted octahedral-like Sb sites. This difference makes CuSbS2 less prone to cation disorder on the tetrahedral sites, and the band tailing or potential fluctuations that can result from such disorder. Cation site disorder is a major problem for the lifetime and transport of the photoexcited charge carriers in CZTS51−53 and CTS.53,54 The resulting carrier localization has recently been demonstrated by a negative imaginary part of the complex mobility obtained from frequency domain OPTP measurements in CTS, even after annealing.54 Similar measurements performed on the annealed CuSbS2 samples (Figure 4B) show a positive imaginary component, indicating charge carrier localization is not as significant issue in CuSbS2 as in CZTS and CTS absorbers. This supports our hypothesis that CuSbS2 is not as prone to cation disorder as Cu2SnS3 and related multinary tetrahedrally bonded materials. 3.4. Device Fabrication and Characterization. On the basis of all previous material characterization, we could conclude that thermal treatment in Sb2S3 vapor, under the conditions described in this work, improves the structural and optoelectronic properties and homogeneity of sputtered CuSbS2 films. This annealing followed by previously described selective chemical etch using KOH38 ensured that the CuSbS2 surface is free of the residual Sb2S3, which is also a known photovoltaic material,55 that can mask the CuSbS2 device performance. Thus, the next question is how these CuSbS2 material improvements correlate with the device performance. Perhaps even more important question is how the thermochemical treatment influences the PV device reproducibility, since having reproducible devices is critical for improving the device performance at early stages of PV technology development. During the initial CuSbS2 PV device fabrication experiments, we noticed that the increase in grain size for the TT-11 h film promoted a microscale film delamination in some parts of the samples, causing increased series resistance and a decreased parallel (shunt) resistance. To solve this problem without negating the favorable optoelectronic properties of the CuSbS2 absorber films annealed for 11 h, the TT temperature- time profile was changed by adding a step at the intermediate temperature of 435 °C for 30 min, and taking 30 more minutes to heat up to the optimal temperature (455 °C). This stepanneal temperature profile can be seen in Figure 1B (red curve). Note that the total annealing time was not changed. The step-anneal did not significantly change the optoelectronic properties of the film (Step TT-11 h), compared to the films annealed without a step (TT-11 h). The gain in the film adhesion and continuity achieved by step-anneal and the comparative cross-section SEM of the devices is shown in the Figure S6 (Supporting Information). Further, these results show no MoS2 layer formation at the CuSbS2/Mo interface, within the resolution of the SEM technique.

stimulate compensating, n-type VS defects, which may lead to interface recombination.47 The increase in the mobility (Figure 3B) can be associated with a decrease in charge-trap density, or structural defects and, to a lesser extent, with an increase in grain size. Similar increase in the mobility (and in carrier concentration) has also been observed by another group,48 when studying the effect of different temperatures and times of sintering on the electrical properties of CuSbS2. It has also been reported48 that large grains do not guarantee high mobility, suggesting that the electronic defects are the major limiting factor controlling the electrical properties of this material, in agreement with our observations. 3.3. Photoexcited Charge Carrier Dynamics. Photoexcited charge carrier properties is one of the most important factors that determines the optoelectronic quality of the photovoltaic absorber materials, but so far they have not been reported for CuSbS2. In order to examine the photoexcited charge carrier dynamics in CuSbS2 films, and compare the results to the results obtained from PL spectroscopy and Hall measurements, the pre-TT and TT-11 h films were characterized by optical pump terahertz probe (OPTP) spectroscopy. The time-domain OPTP measurements (Figure 4A) indicate that annealing reduces the amplitude of the short-

Figure 4. (A) Time-domain optical pump-terahertz probe (OPTP) differential reflectance for pre-TT and TT-11 h CuSbS2 films on glass substrate. (B) Frequency-domain OPTP mobility spectra for the TT11 h films on glass 10 ps after excitation.

decay component magnitude by 50%, and enhances the longdecay component time from 0.5 to 0.7 ns. This suggests the annealed films have less trapping and longer lifetime of the photoexcited charge carriers. Similar measurements performed on annealed CuSbS2 samples coated with CdS buffer layer show shorter long-decay component times of 0.4 ns (Supporting Information, Figure S5), which is somewhat unexpected. The decrease of this decay component may be attributed to either charge transfer to the wider-gap CdS layer, where the photoexcited charge carriers cannot be detected by OPTP, or increased recombination at the CuSbS2/CdS interface compared to the CuSbS2/air interface. More measurements would be necessary to distinguish between these two possibilities. The frequency-domain OPTP measurement results for the annealed CuSbS2 films (Figure 4B, real part) lead to the conclusion that the mobility of the photoexcited charge carriers is ∼4 cm2/(V s). Since this value is similar to the CuSbS2 hole mobility determined from the Hall effect measurements (Figure 3B inset), it appears that the hole- and electron mobility in CuSbS2 are similar to each other. This is consistent with the prior theoretical calculations of the effective masses in this material.49 Note that the matched electron/hole mobilities in 18381

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Figure 5. (A) Representative CuSbS2 PV device J−V curves under simulated AM1.5G illumination (100 mW cm2) at 25 °C. Inset: Histogram analysis of the TT-11 h device efficiencies. (B) Representative EQE (left scale, full lines) and integrated current (right scale, dashed lines) of the CuSbS2 devices.

Table 1. Short Circuit Current (JSC), Open-Circuit Voltage (VOC), Fill Factor (FF), Efficiency (Eff), Series (RS), and Parallel (RP) Resistances Obtained from Illuminated J−V Curves (Figure 5A) as Well as Dark Saturation Current (J0), Ideality Factor (n), and Fit-Adjusted R2 (adj. R2) Obtained from Dark−J−V Fittings (Figure S7)a

a

property

pre-TT (7)

TT-8 h (7)

step TT-11 h (16)

JSC (mA cm‑2) VOC (mV) FF (%) eff (%) RP (Ω cm2) RS (Ω cm2) J0 (mA cm‑2) n adj. r2

3.31 ± 10.8% 121 ± 8.40% 28.1 ± 0.28% 0.113 ± 0.028% 218 ± 34.6% 71.30 ± 9.33% 4.51 ± 2.67% 10.8 ± 7.67% 0.998 ± 0.100%

4.65 ± 7.53% 281 ± 16.4% 38.2 ± 2.47% 0.503 ± 0.075% 245 ± 24.0% 76.80 ± 4.23% 0.489 ± 5.11% 5.33 ± 6.94% 0.999 ± 0.050%

5.20 ± 5.77% 350 ± 6.86% 55.0 ± 1.57% 1.02 ± 0.0175% 1490 ± 5.97% 72.3 ± 1.87% 0.035 ± 8.57% 2.90 ± 2.41% 0.999 ± 0.030%

The numbers in parentheses are the quantity of CuSbS2 devices used for calculating the standard deviations within 95% confidence interval.

Representative J−V curves under simulated AM1.5G illumination (100 mW cm2) and external quantum efficiency (EQE) at 25 °C can be seen in Figure 5A and Figure 5B, respectively. Figure 5A includes J−V curves for 16 devices with annealed CuSbS2 absorbers, and shows a histogram of their efficiencies in the inset, both of which highlight the homogeneity and reproducibility of the device performance. The more detailed view of the J−V curves in the power generating quadrant is provided in Figure S7A in Supporting Information. In Figure 5A, it is possible to see that the devices made with TT-11 h annealed films have 2−3x higher fill factor (FF), short circuit current (JSC) and open-circuit voltage (VOC), compared to devices made with as-deposited absorbers. These results are summarized numerically in Table 1, along with the statistical analysis of the deviations. Overall, the devices made with annealed CuSbS2 films showed approximately 10 times higher efficiency than pre-TT devices (1.0% vs 0.1%). For the pre-TT and TT-8 h absorber films, seven devices were produced, and they were less homogeneous than 16 of the step TT-11 h devices, as can be seen by the larger variation in device performance in the Table 1. From Figure 5B, we note that the integrated current from EQE measurements agrees with the JSC values obtained from the J−V curves (Table 1), with less than 10% error (comparing the same device for both methods), supporting the J−V measurement results. Comparing the EQE profiles, it can be seen that the devices with the thermally treated CuSbS2 absorbers have higher quantum efficiency at lower photon energies. This indicates higher minority carrier diffusion

lengths, likely due to the larger grains and fewer defect states, consistent with the OPTP measurement results (Figure 4). Finally, the optical band gap estimated at 1.5 eV from EQE spectra agrees with the previously discussed optical absorption onset and PL values (Figure 3A). 3.5. Performance Fitting and Analysis of the PV Devices. To understand better the effect of the thermal treatment on device performance, we performed the diode parameter fitting of the dark−J−V curves (one-diode model with four adjustable parameters, eq 1). The representative fitting of the dark−J−V curves can be seen in Figure S7 (Supporting Information). The average results and uncertainties of this dark J−V fitting are summarized in Table 1, along with the device photoresponse parameters obtained from J−Vcurves under simulated AM1.5G illumination (100 mW cm2). On the basis of the adjusted least-squares fitting coefficient, (R2), it can be concluded that the dark−J−V curve behavior was well fitted by a one-diode model of the eq 1. From the fitting, it was observed that the TT promotes a significant (>100×) decrease in the reverse dark current (J0) losses. Lower J0 is often related to the increase in the VOC, and can also be associated with the increase in the JSC due to less recombination via defects.40−42 Well-developed photovoltaic materials, like Si, have a J0 on the order of 10−7 mA cm−2,40 whereas emerging polycrystalline materials, such as SnS, have J0 close to 10−2 mA cm−2,28 similar to what was observed here for the CuSbS2 devices with annealed absorbers. A substantial (∼5×) increase in the parallel (shunt) resistance after 11 h annealing can also be seen from Table 1, reflecting improved uniformity of the thermally treated absorber. The series 18382

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In contrast, the temperature-dependent performance of the PV devices with the thermochemically treated CuSbS 2 absorbers shows a behavior expected from the classical pn junction models. The JSC remains relatively stable as a function of temperature, as expected for a diode under illumination for temperatures higher than the freeze-out temperature. At 130 K, JSC drops significantly which comes along with the typical sshaped J−V curve, see Figure S8 B (Supporting Information). We therefore attribute this drop in JSC to the freeze-out of dopants. Consequently, the conventional VOC(T) analysis is applicable for this device. The extrapolation of VOC to 0 K is considered to give insight into the dominant recombination mechanism.59 A schematic band diagram with a cliff-type interface band offset of the CuSbS2 absorber with the CdS buffer layer is depicted in the inset of Figure 6, where Φ ≈ 0.7 eV. This diagram shows a larger cliff-type conduction band offset between CuSbS2 and CdS, which is proposed to be the dominant factor that limits the efficiency of the PV devices reported here. An alternative explanation that the VOC(0 K) ≈ 0.7 eV result is dominated by defects can be ruled out by fitting the JV−T curves (Figure S8 B) for the ideality factor “n” using eq 1. The fitting results show that n does not exhibits significant temperature dependence with changing temperature: only a slight variation of 0.2 in the temperature range of 180 to 300 K is observed. This suggests that Φ is not dominated by interface recombination via defects.41 This interpretation of the JV-T data is supported by highsensitivity external quantum efficiency (EQE) measurements using Fourier transform photocurrent spectroscopy (FTPS). If the low VOC(0 K) were caused by bulk recombination via defect states, we would expect to see these defect states in the EQE measured using FTPS methods. Whereas the results of the FTPS measurements (Figure S9 Supporting Information) show some amount of band tailing in the 1.0−1.4 eV energy range, they do not indicate a significant number of defect states close to 0.7 eV. This supports our interpretation that VOC(0 K) ≈ 0.7 eV is indicative of a 0.7−0.8 cliff-type CuSbS2/CdS interface band offset. This interpretation is also consistent with recently published theoretical calculations supported by indirect experimental evidence.49,60 Such a large cliff can lead to significant photoexcited electron recombination at the interface, limiting the current of the CuSbS2/CdS device, as well as the Voc. So it appears that the interface cliff might be the efficiencylimiting factor for the thin film PV devices with annealed CuSbS2 absorbers. Thus, alternative buffer layers should be developed to further improve the chalcostibite device performance.

resistance remained relatively unchanged after annealing. The high diode ideality factor (n > 2), observed for all the CuSbS2 devices (Table 1), is usually attributed to extended defects in the depletion region. Multiple models can be used to explain such large ideality factors, e.g., multilevel coupled defect levels,42,56 tunneling enhanced recombination,57 or strong spatial fluctuations of the contact line resistance.58 A comprehensive review on defect models causing ideality factors greater than 2 has recently been published.40 The decrease in the ideality factor with annealing (from n = 11 to n = 3) indicates that the TT decreases the extended defects that were causing its high value in first place, which is consistent with the aforementioned grain growth (Figure 2). Moreover, this improvement of the dark JV-characteristic of the device is correlated with the improvement of the FF under illumination (Table 1). 3.6. Temperature-Dependent Characterization of the PV Devices. Despite of the significant improvements in material quality of the CuSbS2 absorber layers discussed above, the efficiencies of the CuSbS2 devices reported here do not exceed 1%. In order to elucidate the efficiency limiting factors, we performed temperature-dependent J−V measurements (JV−T) on the PV devices with as-deposited and step-TT-11 h annealed CuSbS2 absorbers. Figure 6 shows JSC and VOC for

Figure 6. Short circuit current (JSC) and open-circuit voltage (VOC) versus measurement temperature for as-deposited (pre-TT) and annealed (TT-11 h) CuSbS2 device. The inset sketches the band diagram of the CdS/CuSbS2 interface with Fermi energy EF, conduction band energy EC, the valence band energy EV, and interfacial barrier Φ.

both of these solar cells as a function of measurement temperature. We observed that in case of PV devices with asdeposited absorbers (pre-TT) the JSC decreases significantly with temperature even around 300 K. A classic pn-junction shows this strong JSC dependence on temperature if the freeze-out of dopants occurs. This leads eventually to a breakdown of the pn-junction for very low temperatures and to s-shaped J−V curves. However, this typical s-shape could not be observed in these preannealed devices (Figure S8 A in Supporting Information). The strong decrease in JSC, which is already evident at 300 K, and the missing sshape, which comes along with a classic freeze-out of dopants, indicates a significant problem with the carrier collection for these devices. This can be related to a number of materials properties such as a high localized defect concentration, poor carrier mobility, insufficient doping levels or an interface barrier that has to be overcome thermally. Thus, the conventional analysis of the VOC(T) extrapolation to 0 K cannot be used in this case to determine the interface band offset.

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SUMMARY AND CONCLUSIONS In summary, thermochemical treatment of CuSbS2 improved its structural and optoelectronic quality, and these improvements translated into enhanced reproducibility of the photovoltaic devices. The thermal treatment in Sb2S3 vapor atmosphere increased CuSbS2 grain size, hole concentration, hole mobility, near-band edge photoluminescence, and photoconductivity. The annealing also improved the photoexcited charge carrier dynamics in CuSbS2 absorbers, leading to longer lifetimes and higher mobilities, with no evidence of charge carrier localization. As a result of the PV device experiments, improvements in dark current (from 4.51 to 0.035 mA cm−2), ideality factor (from 10.8 to 2.90), JSC (from 3.31 to 5.20 mA cm−2), VOC (from 121 to 350 mV), FF (from 28.1% to 55.0%) and 18383

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efficiency (from 0.113 to 1.02%) were observed, with particularly good homogeneity of the annealed devices. We conclude that in general the thermochemical treatment of thin film absorbers in the vapor of their volatile components (Sb2S3 in the case of CuSbS2) is a promising approach to improve the quality of thin film absorber materials and enhance the performance of the solar cell devices. In the specific case of the materials used to illustrate this approach in this paper, the current performance of the resulting thin film PV devices with annealed CuSbS2 absorbers is limited by a large 0.8 eV cliff-type band offset with the CdS buffer layer. Thus, future work on chalcostibite PV devices, as well as other novel thin film absorbers, should explicitly address this issue by considering alternative buffer layers specifically designed for each individual promising absorber material.

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ASSOCIATED CONTENT

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.6b04206. Selective chemical etch (Figure S1), XRD patterns (Figure S2), SEM surface (Figure S3) and cross-section (Figure S6) images, Graphs of absorption coefficient (α) and (αhν)2 vs photon energy (Figure S4), OPTP results (Figure S5), J−V curves under light, dark−J−V fitting (Figure S7), T-dependent J−V curves (Figure S8), and EQE (Figure S9) (PDF)

AUTHOR INFORMATION

Corresponding Author

*(A.Z.) E-mail: [email protected]. Telephone: − Tel.: +1-303-384-6467. Author Contributions

F.W.d.S.L. performed the thermochemical treatment experiments, analyzed and discussed most of the data, and wrote the paper with contribution from all coauthors and under the supervision of L.H.M. and A.Z., who also planned and coordinated the entire study. Precursor thin films were deposited by A.W.W. and L.L.B. Photoluminescence characterization was performed by P.C.D. Optical pump THz probe spectroscopy was performed by H.H., T.U., and R.E. Temperature-dependent J−V measurements were performed by B.B. and U.R. Notes

The authors declare no competing financial interest.

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REFERENCES

(1) Lewis, N. S.; Nocera, D. G. Powering the planet: Chemical challenges in solar energy utilization. Proc. Natl. Acad. Sci. U. S. A. 2006, 103, 15729−15735. (2) Walker, B.; Kim, C.; Nguyen, T. Q. Small molecule solutionprocessed bulk heterojunction solar cells. Chem. Mater. 2011, 23, 470− 482. (3) O’Regan, B.; Lenzmann, F.; Muis, R.; Wienke, J. A solid-state dye-sensitized solar cell fabricated with pressure-treated P25-TiO2 and CuSCN: Analysis of pore filling and IV characteristics. Chem. Mater. 2002, 14, 5023−5029. (4) Lee, Y. L.; Huang, B. M.; Chien, H. T. Highly efficient CdSesensitized TiO2 photoelectrode for quantum-dot-sensitized solar cell applications. Chem. Mater. 2008, 20, 6903−6905. (5) Colella, S.; Mosconi, E.; Fedeli, P.; Listorti, A.; Gazza, F.; Orlandi, F.; Ferro, P.; Besagni, T.; Rizzo, A.; Calestani, G.; et al. MAPbI3‑xClx mixed halide Perovskite for hybrid solar cells: The role of chloride as dopant on the transport and structural properties. Chem. Mater. 2013, 25, 4613−4618. (6) Reece, S. Y.; Hamel, J. A.; Sung, K.; Jarvi, T. D.; Esswein, A. J.; Pijpers, J. J. H.; Nocera, D. G. Wireless solar water splitting using silicon-based semiconductors and earth-abundant catalysts. Science 2011, 334, 645−648. (7) Dross, F.; Baert, K.; Bearda, T.; Deckers, J.; Depauw, V.; El Daif, O.; Gordon, I.; Gougam, A.; Govaerts, J.; Granata, S.; et al. Crystalline thin-foil silicon solar cells: Where crystalline quality meets thin-film processing. Prog. Photovoltaics 2012, 20, 770−784. (8) Research Cell Efficiency Records. http://www.nrel.gov/ncpv/ images/efficiency_chart.jpg (acessed January 7, 2015). (9) Wadia, C.; Alivisatos, A. P.; Kammen, D. M. Materials availability expands the opportunity for large-scale photovoltaics deployment. Environ. Sci. Technol. 2009, 43, 2072−2077. (10) Minami, T.; Nishi, Y.; Miyata, T. Heterojunction solar cell with 6% efficiency based on an n-type Aluminum−Gallium−oxide thin film and p-type Sodium-doped Cu2O sheet. Appl. Phys. Express 2015, 8, 022301−022304. (11) Zhou, Y.; Wang, L.; Chen, S.; Qin, S.; Liu, X.; Chen, J.; Xue, D.; Luo, M.; Cao, Y.; Cheng, Y.; et al. Thin-film Sb2Se3 photovoltaics with oriented one-dimensional ribbons and benign grain boundaries. Nat. Photonics 2015, 9, 409−415. (12) Caskey, C. M.; Richards, R. M.; Ginley, D. S.; Zakutayev, A. Thin film synthesis and properties of copper nitride, a metastable semiconductor. Mater. Horiz. 2014, 1, 424. (13) Zakutayev, A.; Caskey, C. M.; Fioretti, A. N.; Ginley, D. S.; Vidal, J.; Stevanovic, V.; Tea, E.; Lany, S. Defect tolerant semiconductors for solar energy conversion. J. Phys. Chem. Lett. 2014, 5, 1117−1125. (14) Steinmann, V.; Jaramillo, R.; Hartman, K.; Chakraborty, R.; Brandt, R. E.; Poindexter, J. R.; Lee, Y. S.; Sun, L.; Polizzotti, A.; Park, H. H.; et al. 3.88% efficient Tin sulfide solar cells using congruent thermal evaporation. Adv. Mater. 2014, 26, 7488−7492. (15) Sánchez-Juárez, A.; Tiburcio-Silver, A.; Ortiz, A. Fabrication of SnS2/SnS heterojunction thin film diodes by plasma-enhanced chemical vapor deposition. Thin Solid Films 2005, 480-481, 452−456. (16) Lahourcade, L.; Coronel, N. C.; Delaney, K. T.; Shukla, S. K.; Spaldin, N. A.; Atwater, H. A. Structural and optoelectronic characterization of RF sputtered ZnSnN2. Adv. Mater. 2013, 25, 2562−2566. (17) Tiwari, D.; Fermin, D. J.; Chaudhuri, T. K.; Ray, A. Solution processed Bismuth ferrite thin films for all-oxide solar photovoltaics. J. Phys. Chem. C 2015, 119, 5872−5877. (18) Yang, M.; Zakutayev, A.; Vidal, J.; Zhang, X.; Ginley, D. S.; DiSalvo, F. J. Strong optical absorption in CuTaN2 nitride delafossite. Energy Environ. Sci. 2013, 6, 2994−2999. (19) Gershon, T.; Shin, B.; Bojarczuk, N.; Hopstaken, M.; Mitzi, D. B.; Guha, S. The role of Sodium as a surfactant and suppressor of nonradiative recombination at internal surfaces in Cu2ZnSnS4. Adv. Energy Mater. 2015, 5, 1400849.

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Article

ACKNOWLEDGMENTS

The “Rapid Development of Earth-abundant Thin Film Solar Cells” project is supported by the U.S. Department of Energy, Office of Energy Efficiency and Renewable Energy, as a part of the SunShot initiative, under Contract No. DE-AC3608GO28308 to the NREL. For execution of this work, F.W.d.S.L was funded by the Sao Paulo Research Foundation (FAPESP), Grant 2014/12166-3. The optical pump terahertz probe absorber measurements and temperature-dependent device performance characterization were supported by the Helmholtz Association Initiative and Network Fund (HNSEI project SO-075). B.B. acknowledges support from the Hans L. Merkle Foundation via a Ph.D. scholarship. 18384

DOI: 10.1021/acs.jpcc.6b04206 J. Phys. Chem. C 2016, 120, 18377−18385

Article

The Journal of Physical Chemistry C (20) Ramasamy, K.; Malik, M. A.; O’Brien, P. The chemical vapor deposition of Cu2ZnSnS4 thin films. Chem. Sci. 2011, 2, 1170−1772. (21) Itthibenchapong, V.; Kokenyesi, R. S.; Ritenour, A. J.; Zakharov, L. N.; Boettcher, S. W.; Wager, J. F.; Keszler, D. A. Earth-abundant Cu-based chalcogenide semiconductors as photovoltaic absorbers. J. Mater. Chem. C 2013, 1, 657−662. (22) Kanai, A.; Araki, H.; Takeuchi, A.; Katagiri, H. Annealing temperature dependence of photovoltaic properties of solar cells containing Cu2SnS3 thin films produced by co-evaporation. Phys. Status Solidi B 2015, 252, 1239−1243. (23) Yang, B.; Wang, L.; Han, J.; Zhou, Y.; Song, H.; Chen, S.; Zhong, J.; Lv, L.; Niu, D.; Tang, J. CuSbS2 as a promising earthabundant photovoltaic absorber material: A combined theoretical and experimental study. Chem. Mater. 2014, 26, 3135−3143. (24) Zakutayev, A.; Baranowski, L. L.; Welch, A. W.; Wolden, C. A.; Toberer, E. S. Comparison of Cu2SnS3 and CuSbS2 as potential solar cell absorbers. In 2014 IEEE 40th Photovoltaic Specialist Conference (PVSC); IEEE: 2014; pp 2436−2438. (25) Welch, A. W.; Zawadzki, P. P.; Lany, S.; Wolden, C. A.; Zakutayev, A. Self-regulated growth and tunable properties of CuSbS2 solar absorbers. Sol. Energy Mater. Sol. Cells 2015, 132, 499−506. (26) Septina, W.; Ikeda, S.; Iga, Y.; Harada, T.; Matsumura, M. Thin film solar cell based on CuSbS2 absorber fabricated from an electrochemically deposited metal stack. Thin Solid Films 2014, 550, 700−704. (27) Rodríguez-Lazcano, Y.; Nair, M. T. S.; Nair, P. K. Photovoltaic p-i-n structure of Sb2S3 and CuSbS2 absorber films obtained via chemical bath deposition. J. Electrochem. Soc. 2005, 152, G635−G638. (28) Krishnan, B.; Shaji, S.; Ernesto Ornelas, R. Progress in development of Copper Antimony sulfide thin films as an alternative material for solar energy harvesting. J. Mater. Sci.: Mater. Electron. 2015, 26, 4770−4781. (29) Ramasamy, K.; Sims, H.; Butler, W. H.; Gupta, A. Selective nanocrystal synthesis and calculated electronic structure of all four phases of Copper−Antimony−sulfide. Chem. Mater. 2014, 26, 2891− 2899. (30) Welch, A. W.; Baranowski, L. L.; Zawadzki, P.; Lany, S.; Wolden, C. A.; Zakutayev, A. CuSbSe2 Photovoltaic Devices with 3% Efficiency. Appl. Phys. Express 2015, 8, 082301. (31) Welch, A. W.; Baranowski, L. L.; de Souza Lucas, F. W.; Toberer, E. S.; Wolden, C. A.; Zakutayev, A. Copper Antimony chalcogenide thin film pv device development. In 2015 IEEE 42nd Photovoltaic Specialist Conference (PVSC); IEEE: 2015; pp 1−4. (32) Tai, K. F.; Fu, D.; Chiam, S. Y.; Huan, C. H. A.; Batabyal, S. K.; Wong, L. H. Antimony doping in solution-processed Cu2ZnSn(S,Se)4 solar cells. ChemSusChem 2015, 8, 3504−3511. (33) Jeon, J.-O.; Lee, K. D.; Seul Oh, L.; Seo, S.-W.; Lee, D.-K.; Kim, H.; Jeong, J.; Ko, M. J.; Kim, B.; Son, H. J.; et al. Highly efficient Copper-Zinc-Tin-selenide (CZTSe) solar cells by electrodeposition. ChemSusChem 2014, 7, 1073−1077. (34) Strothkämper, C.; Schwarzburg, K.; Schütz, R.; Eichberger, R.; Bartelt, A. Multiple-trapping governed electron transport and charge separation in ZnO/In2S3 core/shell nanorod heterojunctions. J. Phys. Chem. C 2012, 116, 1165−1173. (35) Katsidis, C. C.; Siapkas, D. I. General transfer-matrix method for optical multilayer systems with coherent, partially coherent, and incoherent interference. Appl. Opt. 2002, 41, 3978. (36) McCusker, J. K. Femtosecond absorption spectroscopy of transition metal charge-transfer complexes. Acc. Chem. Res. 2003, 36, 876−887. (37) Contreras, M. A.; Romero, M. J.; To, B.; Hasoon, F.; Noufi, R.; Ward, S.; Ramanathan, K. Optimization of CBD CdS process in highefficiency Cu(In,Ga)Se2-based solar cells. Thin Solid Films 2002, 403404, 204−211. (38) De Souza Lucas, F. W.; Welch, A. W.; Baranowski, L. L.; Dippo, P. C.; Mascaro, L. H.; Zakutayev, A. Thermal treatment improvement of CuSbS2 absorbers. In 2015 IEEE 42nd Photovoltaic Specialist Conference (PVSC); IEEE: 2015; pp 1−5.

(39) Reference Solar Spectral Irradiance: Air Mass 1.5. http://www. rredc.nrel.gov/solar/spectra/am1.5/ (accessed May 15, 2015). (40) Breitenstein, O. Understanding the current-voltage characteristics of industrial crystalline Silicon solar cells by considering inhomogeneous current distributions. Opto-Electronics Rev. 2013, 21, 121−136. (41) Kandil, K. M.; Altouq, M. S.; Al-Asaad, A. M.; Alshamari, L. M.; Kadad, I. M.; Ghoneim, A. A. Investigation of the performance of CIS photovoltaic modules under different environmental conditions. Smart Grid Renewable Energy 2011, 02, 375−387. (42) Breitenstein, O.; Bauer, J.; Altermatt, P. P.; Ramspeck, K. Influence of defects on solar cell characteristics. Solid State Phenom. 2009, 156−158, 1−10. (43) Sze, S. M.; Ng, K. K. Physics of Semiconductor Devices, 3rd ed.; Wiley: New Jersey, 2007. (44) Piacente, V.; Scardala, P.; Ferro, D. Study of the vaporization behaviour of Sb2S3 and Sb2Te3 from their vapour pressure measurements. J. Alloys Compd. 1992, 178, 101−115. (45) Jackson, A. J.; Tiana, D.; Walsh, A. A universal chemical potential for sulfur vapours. Chem. Sci. 2016, 7, 1082−1092. (46) FIZ Karlsruhe Inorganic Crystal Structure Database. http:// www.fiz-Karlsruhe.de/icsd_web.html (accessed December 12, 2015). (47) Perniu, D.; Duta, A.; Schoonman, J. Defect Chemistry of CuSbS2. In 2006 International Semiconductor Conference; IEEE: 2006; pp 245−248. (48) Wubet, W.; Kuo, D.-H. Process limitation for p-type CuSbS2 semiconductor with high electrical mobility of 20 cm2V−1s−1. Mater. Res. Bull. 2014, 53, 290−294. (49) High Performance Computing Center − Materials Database at the National Renewable Energy Laboratory. http://materials.nrel.gov/ (accessed January 12, 2015). (50) Unold, T.; Hempel, H.; Strothkamper, C.; Kaufmann, C. A.; Eichberger, R.; Bartelt, A. Charge carrier mobilities and dynamics in thin film compound semiconductor materials from transient THz absorption. In 2014 IEEE 40th Photovoltaic Specialist Conference (PVSC); IEEE: 2014; pp 2066−2069. (51) Gokmen, T.; Gunawan, O.; Todorov, T. K.; Mitzi, D. B. Band tailing and efficiency limitation in kesterite solar cells. Appl. Phys. Lett. 2013, 103, 103506. (52) Scragg, J. J. S.; Choubrac, L.; Lafond, A.; Ericson, T.; PlatzerBjö rkman, C. A low-temperature order-disorder transition in Cu2ZnSnS4 thin films. Appl. Phys. Lett. 2014, 104, 041911. (53) Zawadzki, P.; Zakutayev, A.; Lany, S. Entropy-driven clustering in tetrahedrally bonded multinary materials. Phys. Rev. Appl. 2015, 3, 034007. (54) Baranowski, L. L.; McLaughlin, K.; Zawadzki, P.; Lany, S.; Norman, A.; Hempel, H.; Eichberger, R.; Unold, T.; Toberer, E. S.; Zakutayev, A. Effects of disorder on carrier transport in Cu2SnS3. Phys. Rev. Appl. 2015, 4, 044017. (55) Deringer, V. L.; Stoffel, R. P.; Wuttig, M.; Dronskowski, R. Vibrational properties and bonding nature of Sb2Se3 and their implications for chalcogenide materials. Chem. Sci. 2015, 6, 5255− 5262. (56) Schenk, A.; Krumbein, U. Coupled defect-level recombination: Theory and application to anomalous diode characteristics. J. Appl. Phys. 1995, 78, 3185. (57) Rau, U. Tunneling-enhanced recombination in Cu(In,Ga)Se2 heterojunction solar cells. Appl. Phys. Lett. 1999, 74, 111−113. (58) Van der Heide, A. S. H.; Schonecker, A.; Bultman, J. H.; Sinke, W. C. Explanation of high solar cell diode factors by nonuniform contact resistance. Prog. Photovoltaics 2005, 13, 3−16. (59) Rau, U.; Schock, H. W. Electronic properties of Cu(In,Ga)Se2 heterojunction solar cells-recent achievements, current understanding, and future challenges. Appl. Phys. A: Mater. Sci. Process. 1999, 69, 131− 147. (60) Welch, A. W.; Baranowski, L. L.; Zawadzki, P.; DeHart, C.; Johnston, S.; Lany, S.; Wolden, C. A.; Zakutayev, A. Accelerated development of CuSbS2 thin film photovoltaic device prototypes. Prog. Photovoltaics 2016, 24, 929. 18385

DOI: 10.1021/acs.jpcc.6b04206 J. Phys. Chem. C 2016, 120, 18377−18385