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Low-Temperature Atomic Layer Deposition of CuSbS2 for Thin-Film Photovoltaics Shannon C. Riha,*,† Alexandra A. Koegel,† Jonathan D. Emery,‡,∥ Michael J. Pellin,§,∥ and Alex B. F. Martinson*,§,∥ †

Department of Chemistry, University of Wisconsin-Stevens Point, Stevens Point, Wisconsin 54481, United States Department of Materials Science and Engineering, Northwestern University, Evanston, Illinois 60208, United States § Materials Science Division, Argonne National Laboratory, Argonne, Illinois 60439, United States ∥ Argonne-Northwestern Solar Energy Research (ANSER) Center, Argonne National Laboratory, Argonne, Illinois 60439, United States

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‡

S Supporting Information *

ABSTRACT: Copper antimony sulfide (CuSbS2) has been gaining traction as an earth-abundant absorber for thin-film photovoltaics given its near ideal band gap for solar energy conversion (∼1.5 eV), large absorption coefficient (>104 cm−1), and elemental abundance. Through careful in situ analysis of the deposition conditions, a low-temperature route to CuSbS2 thin films via atomic layer deposition has been developed. After a short (15 min) postprocess anneal at 225 °C, the ALD-grown CuSbS2 films were crystalline with micron-sized grains, exhibited a band gap of 1.6 eV and an absorption coefficient >104 cm−1, as well as a hole concentration of 1015 cm−3. Finally, the ALD-grown CuSbS2 films were paired with ALD-grown TiO2 to form a photovoltaic device. This photovoltaic device architecture represents one of a very limited number of Cd-free CuSbS2 PV device stacks reported to date, and it is the first to demonstrate an open-circuit voltage on par with CuSbS2/CdS heterojunction PV devices. While far from optimized, this work demonstrates the potential for ALDgrown CuSbS2 thin films in environmentally benign photovoltaics. KEYWORDS: copper antimony sulfide, CuSbS2, thin film, atomic layer deposition, photovoltaics, thin-film solar cell, ternary metal sulfide

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INTRODUCTION Solar photovoltaic research has made significant strides in the past decade, yet solar energy still only constitutes a small fraction of our present energy utilization. While the cost of traditional solar panels has been reduced considerably, further market penetration will require additional reductions in cost, improved flexibility in design, and more effective energy storage. Thin films based on earth-abundant materials such as copper-based sulfides may address some of these challenges because of their unique optoelectronic properties.1,2 For example, thinner films absorb the same amount of sunlight as a silicon-based solar cell, and lower crystalline quality can still be used to produce high-efficiency devices. This results in lower fabrication cost and the ability to use more flexible substrates. Copper-based sulfide thin films have a rich history in photovoltaics, with Cu2S/CdS heterojunction devices reaching 10% efficiency in the early 1980scomparable to the best Sibased technology of the day.3−5 However, instabilities due to Cu diffusion into the CdS layer led to rapid device degradation and attention quickly turned to the ternary compound CuInSe2, and subsequently Cu(In,Ga)Se2 (CIGS). Presently, stable CIGS lab-scale devices are being fabricated with 22% efficiency,6−11 although concerns with availability and high © 2017 American Chemical Society

cost of In and Ga have stimulated efforts to replace these elements. For example, replacing In and Ga with Zn and Sn, respectively, to form Cu2ZnSn(S,Se)4 (CZTSSe) has resulted in stable devices with efficiencies of 12.6%.12 The less-explored I−V−VI 2 chalcogenides, including CuSbS2, also show promise as alternative solar absorbers. In particular, CuSbS2 has basic optoelectronic properties suitable for photovoltaics including a near ideal direct band gap ranging from 1.38 to 1.56 eV, a high optical absorption coefficient (over 104 cm−1), and moderate hole doping (1015 − 1018 cm−3). From an environmental and economic viewpoint, interest in CuSbS2 is driven by the relatively low toxicity of all elements as well as Sb being more abundant and in less demand compared to In and Ga.13−19 Furthermore, unlike CZTS, theoretical calculations indicate that low-energy defects are far from the center of the band gap, meaning that CuSbS2 may exhibit a lower density of recombination-center defects.17 While there are a number of reports on the optical, structural, and electronic properties of CuSbS2 to highlight its potential as a PV absorber, Received: October 13, 2016 Accepted: January 10, 2017 Published: January 24, 2017 4667

DOI: 10.1021/acsami.6b13033 ACS Appl. Mater. Interfaces 2017, 9, 4667−4673

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with an active area of 0.18 cm2. Some devices also included a SpiroOMeTAD hole-transporting layer (HTL) between the CuSbS2 and Au, vide infra. An etchant solution containing HCl and H2O (1:1 ratio by volume), 5 vol % HNO3, and a trace of soap was used to pattern the Sn-doped indium oxide (ITO) coated glass substrates. Kapton tape was used to mask the bottom 2/3 of the substrate. The substrates were immersed in the etchant solution, which was heated to 80 °C, for 10 min or until the ITO was completely removed from the top 1/3 of the substrate. The substrates were then placed in a bath of 1 M KOH to neutralize the acid, followed by removal of the Kapton tape and a DI water rinse. Prior to ALD, the ITO substrates were then cleaned following the same organic solvent cleaning procedure applied for the silicon and fused quartz substrates. TiO 2 thin films were prepared using ALD. Tetrakis(dimethylamido)titanium(IV) (TiTDMA, SigmaAldrich, heated to 75 °C) or titanium(IV) triisopropoxide (TTIP, SigmaAldrich, heated to 80 °C) were used for the titanium precursors and DI water was used as the oxygen source. Films were deposited at 200 °C (TiTDMA) or 225 °C (TTIP) with a timing sequence of 0.15−10− 0.015−10. The film thickness of the TiO2 layer was ∼20 nm. Deposition of the hole-transport layer was done by spin coating a solution containing 500 μL of cholorbenzene, 40 mg of SpiroOMeTAD, 14.4 μL of 4-tertbutylpyridine, and 8.6 μL of lithium trifluoromethanesulfonimide (520 mg in 1 mL acetonitrile) for 30 s at 4000 rpm in the N2-filled glovebox. Thermal evaporation of a gold ingot through a patterned shadow mask produced 100 nm thick top contacts. Finally, Ag paint was used to provide electrical connection to both the exposed ITO bottom contact and the Au top contact. Characterization. The CuSbS2 thin films were characterized by a combination of both in situ and ex situ techniques. In situ quartz crystal microbalance (QCM) analysis was performed using a custom home-built ALD reactor lid equipped with two QCM ports positioned at 5 and 15 cm from the chamber inlet.34 Ex situ analysis included ellipsometry, X-ray diffraction (XRD), Hall measurements, scanning electron microscopy (SEM), X-ray photoelectron spectroscopy (XPS), UV−vis-NIR, and current−voltage measurements. A J.A. Woollam M2000 Ellipsometer was used to determine the thickness of the ALD films when deposited on witness silicon substrates. XRD patterns were collected on a Philips X’Pert Pro MRD diffractometer using Cu Kα radiation (λ = 1.5418 Å) with tube voltage and current set to 30 kV and 30 mA, respectively. The patterns were acquired in continuous mode using a 0.02° step size and measuring for 10.0 s/step. A 60 mm graded parabolic W/Si mirror with a 0.8° acceptance angle and a 1/8° divergence slit were used to condition the incident beam. A 0.27° parallel plate collimator and a flat pyrolytic graphite monochromator was positioned in front of the PW3011/20 sealed proportional point detector to collect the reflected beam. SEM images were collected using a Hitachi S-4700-II SEM equipped with an energy dispersive spectroscopy (EDS) detector for elemental analysis. Images were taken using an accelerating voltage of 5 kV, while elemental analysis with EDS was acquired using a 20 kV accelerating voltage. XPS analysis was performed in a custom built XPS ultrahigh vacuum system using Mg Kα radiation (1253.6 eV) and a hemispherical electron energy analyzer in the fixed absolute resolution mode. Spectra were collected using 44 eV pass energy and either a 0.5 or 0.1 eV step size for survey and high-resolution spectra, respectively. CasaXPS (Copyright 2009 Casa Software, Ltd.) was used to analyze the data. All spectra were shifted to account for sample charging using adventitious carbon as a reference to 284.5 eV. A Varian Cary 5000 with an integrating sphere accessory (DRA-2500) was used for reflectancecorrected UV−vis−NIR spectroscopy. Hall measurements were collected at room temperature with an Ecopia HMS-3000 Hall measurement system both in the N2-filled glovebox and in ambient air. The probe current was selected such that the voltage signal was between 0.1 and 1 V. Dark and illuminated current−voltage measurements of CuSbS2 devices were collected with a CHI 620D potentiostat. A Newport 300 W arc lamp equipped with an AM 1.5 and UV filter was used as the light source, and the output power was calibrated with a Newport silicon photodiode to simulate AM 1.5 conditions (100 mW/cm2). A monochromator coupled to the arc

to date only a few reports have demonstrated its actual use in a photovoltaic device.14,15,17,19−21 Atomic layer deposition (ALD) is emerging as a premier lowtemperature deposition method for thin films for solar energy conversion.22−24 In particular, ALD has been used to deposit thin transparent conducting oxides, window/buffer layers, surface passivation layers, and even the absorber material. It is no surprise then that interest in the ALD of metal sulfide materials has been growing.25 However, only 19 binary sulfides have been prepared to date by ALD, and few reports exist that describe the ALD of mixed-metal sulfides.25−30 In contrast to metal oxides, cation diffusion through many metal sulfides is rapid, which can lead to interdiffusion and even cation exchange when growing mixed metal sulfide stacks by ALD.31 While this makes the fabrication of nanolaminate stacks of sulfides challenging, cation interdiffusion may prove to be beneficial for the low-temperature ALD synthesis of uniform mixed-metal sulfides, as demonstrated for Cu2ZnSnS4.29 Herein we report for the first timea low-temperature ALD route for the synthesis of CuSbS2 thin films by carefully tuning the deposition and precursor ratios. In addition to describing the novel ALD synthesis of CuSbS2, the potential of ALD-grown CuSbS2 thin films as a solar absorber is demonstrated by pairing it with ALD-grown TiO2 in a heterojunction PV device. TiO2 was specifically chosen as the n-type material for two reasons: (1) to ensure a structurally clean (abrupt) interface at the p-n junction given the known interfacial mixing associated with sulfide ALD and (2) to avoid the use of toxic materials like CdS. This work represents one of a very limited number of Cd-free CuSbS2 PV devices and is the first to demonstrate an open-circuit voltage, Voc, comparable to reported CuSbS2 PV devices fabricated with CdS.14,15,17,32

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EXPERIMENTAL SECTION

ALD of CuSbS2 Thin Films. Substratessilicon and fused quartzwere cleaned prior to deposition using a 10 min sonication in acetone, followed by a 10 min sonication in isopropanol. The substrates were dried with a stream of N2 and transferred directly into a N2-filled glovebox. Deposition of the CuSbS2 thin films was performed in an Ultratech-Cambridge Nanotech Savannah 200 ALD reactor modified for use with H2S gas that was coupled to the N2-filled glovebox. The deposition chamber and gas delivery manifold temperatures were 100 and 155 °C, respectively, whereas bis(N,N′disecbutylacetamidinato)dicopper(I) (CuAMD, purchased from DOW Chemical Company) and tris(dimethylamido)antimony(III) (SbTDMA, purchased from SigmaAldrich) were heated to 150 and 55 °C, respectively. Subsequently, 4% H2S (balance N2, purchased from AirGas) was delivered at room temperature through a 500 μm orifice.33 (Safety note: H2S is a poisonous gas but not flammable at 4% concentration.) Films were grown with alternating bilayers that would be expected to produce nanoscale mixing.30 One cycle of Cu−S growth was deposited after every 9 cycles of Sb−S. The timing sequence, reported in seconds, for each deposition followed the typical ALD recipe (t1 − t2 − t3 − t4), where t1 and t3 are the pulse times for the metal and H2S precursors, respectively, and t2 and t4 are the inert purge times. The timing sequence for the Cu−S deposition was 2− 25−0.1−25, and for the Sb−S growth it was 0.5−25−0.1−25. An ∼100 nm thick film was produced with 250 supercycles (2500 ALD cycles total). The mixed sulfide thin films were finally annealed under vacuum at 225 °C in the ALD reaction chamber for 15 min to crystallize the film. During the annealing process, the CuSbS2 thin films were exposed to a 0.1 s 4% H2S pulse introduced into the chamber every 1 min to compensate for any S loss and to maintain the correct stoichiometry. CuSbS2 PV Device Fabrication. CuSbS2 thin-film photovoltaic devices were fabricated using a glass/ITO/TiO2/CuSbS2/Au stack 4668

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Figure 1. In situ QCM analysis of CuAMD exposure to Sb2S3 thin films at a deposition temperature of (a) 160 °C, (b) 140 °C, (c) 125 °C, and (d) 100 °C. The orange regions highlight 10 consecutive CuAMD doses following Sb−S ALD cycling. lamp enabled monochromatic light illumination for external quantum efficiency measurements.

Figure S1. Furthermore, if Sb were to precipitate Cu metal through a complete cation exchange in the sulfur lattice (in analogy to DEZ on Cu2S),31 the mass would increase dramatically with consecutive SbTDMA doses. Therefore, because no significant mass change was observed, we infer that Sb does not readily ion exchange with Cu in a Cu2S film at or below deposition temperatures of 150 °C. The ternary copper sulfide family CuxSbySz consists of four stable compoundsCuSbS 2 , Cu 3 SbS 3 , Cu 3 SbS 4 , and Cu12Sb4S13with varying optical and electronic properties. Therefore, the ALD dosing schedule was tuned to target the CuSbS2 stoichiometry most desirable for use in PV. Table 1

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RESULTS AND DISCUSSION Unlike the ALD of metal oxides, the ALD of metal sulfides does not always result in the simple stack that might be predicted from the dosing schedule. Thimsen et al. found that when a SnSx layer was deposited on top of a Cu2S layer, intermixing of the Sn and Cu cations resulted in a film that was a layered structure of Cu2SnS3 and SnS2 as deposited.28,29 When anticipated and properly controlled, this in situ intermixing may prove to be beneficial for the low-temperature deposition of uniform films of mixed metal sulfides. However, the possibility for complete cation exchange between the film and the gaseous precursor requires that great care must be taken. For example, when an ultrathin film of ZnS was exposed to a common copper ALD precursor (CuAMD), the complete conversion of the ZnS film to Cu2S occurred as a result of the total removal of Zn from the film. Similarly, diethyl zinc was shown to convert a Cu2S film to ZnS, but in this case, Cu metal remained plated out at the substrate surface.31 If unchecked, this facile cation exchange could prevent the growth of mixed metal sulfide thin films by ALD. Therefore, we first investigated the interaction of copper and antimony metal precursors with an ultrathin metal sulfide film containing the other metal. In situ QCM was used to monitor the mass change when an Sb2S3 thin film was exposed to CuAMD precursor and, conversely, when a Cu2S film was exposed to the SbTDMA precursor. A dramatic mass loss was observed when the Sb2S3 film was exposed to 10 doses of CuAMD at 160 °C (Figure 1a). More specifically, ∼250 ng/cm2 were removed after 10 CuAMD pulses, equivalent to the mass of ∼20 Sb2S3 ALD cycles. Given the relative masses of Cu and Sb, if Cu were to effectively remove Sb in the filmanalogous to complete Zn removal from the ZnS films after CuAMD dosing as described above31then observing a mass loss is reasonable. Reducing the deposition temperature to 140 or 125 °C resulted in significantly less mass loss (Figure 1b,c), but even so, the mass loss is not complete after 10 CuAMD pulses at these temperatures. That the mass continues to decline with consecutive CuAMD pulses suggests that cation exchange is still occurring to a small but finite extent under these deposition conditions. Panel (d) in Figure 1 illustrates that cation exchange can be nearly eliminated by lowering the deposition temperature to 100 °C, as indicated by a relatively constant mass during subsequent CuAMD pulsing after the initial CuAMD dose. In contrast, when a Cu2S thin film was exposed to 10 doses of SbTDMA as high as 150 °C, no mass loss was observed,

Table 1. Effect of Deposition and Post-Processing Conditions on the Stoichiometry of Cu−Sb−S Compounds deposition T (°C)

Cu−S:Sb−S cycle ratio

160 140 140 100 100

1:1 1:4 1:9 1:4 1:4

100 100

1:9 1:9

a

postprocessing

compounda/stoichiometryb

n/a n/a n/a n/a annealed at 225 °C for 15 min n/a annealed at 225 °C for 15 min

Cu12Sb4S13a,b Cu12Sb4S13a,b Cu12Sb4S13b CuSbS2b Cu12Sb4S13b CuSbS2b CuSbS2a,b

Determined by XRD. bDetermined by EDX.

(and Table S1 in the Supporting Information) summarizes the CuxSbySz compounds (determined by XRD) and stoichiometries (determined by EDS) observed as a function of the deposition temperature, Cu−S to Sb−S pulse ratio, and postprocessing conditions. At deposition temperatures of 140 and 160 °C, the Cu12Sb4S13 stoichiometry was the only composition obtained. Indeed, the copper-rich Cu12Sb4S13 phase was observed at 140 °C even with additional Sb−S cycles to compensate for the cation exchange process. When the deposition temperature was lowered to 100 °C, CuSbS2 was produced for Cu−S:Sb−S ratios of 1:4 as well as 1:9. However, films grown at 100 °C with a 1:4 Cu−S:Sb−S cycle ratio converted to the copper-rich Cu12Sb4S13 stoichiometry after annealing at 225 °C for 15 min. Therefore, at a deposition temperature of 100 °C, 9 Sb−S cycles were needed for every 1 Cu−S cycle to compensate for the low per-cycle deposition rate of Sb−S and high volatility of Sb when annealed at temperatures above 200 °C. 4669

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a separation of 9.4 eV that suggests Sb3+, consistent with that reported by McCarthy et al.35 The absence of an O 1s peak at 531 eV indicates that no significant surface oxides (e.g., CuxO, Sb2O3) have formed during deposition, postprocess annealing, or after exposure to ambient air. Finally, the S 2p region can be fit with two peaks located at 160.7 (S 2p3/2) and 161.9 (S 2p1/2) eV with a peak splitting of 1.2 eV. These peak positions and separation are consistent with the expected S2− state. Furthermore, the absence of a peak around 168 eV eliminates the possibility of S oxidation to S4+. When taken together, the XPS results authenticate the valence state of the CuSbS2 films to be Cu+Sb3+(S2−)2. Figure 3 (and Figure S4) shows the absorption spectrum of an 80 nm thick CuSbS2 thin film that was deposited at 100 °C

Figure 2 (and Figure S2) shows SEM images and XRD data for CuSbS2 thin films grown at 100 °C using the optimized

Figure 2. (a) Top and (b) side-profile SEM images of a CuSbS2 thin film annealed for 15 min at 225 °C highlight the formation of elongated grain structure. (c) XRD of the annealed CuSbS2 thin film supports preferential orientation of the crystallites within the film.

ALD recipe. While the as-deposited CuSbS2 thin films are amorphous and devoid of any distinct grain structure as deposited, films annealed at 225 °C for 15 min in the ALD chamber show clear grain structure (refer to Supporting Information for details regarding postprocess annealing temperature and time). Representative top and side profile SEM images of the annealed CuSbS2 thin films, Figure 2a and 2b, show columnar grains that are 1−2 μm in length and 0.5 μm in width. The long, thin grains suggest that the crystallites in the thin film have preferential orientation with the long axis (c-axis) of the CuSbS2 orthorhombic unit cell oriented largely in plane, vide infra. The XRD pattern (Figure 2c) exhibits peaks that can all be indexed to the CuSbS2 phase (space group Pnma, ICSD no. 418753).13,15 No peaks ascribed to binary phases Cu2S and Sb2S3 were observed. Not all CuSbS2 peaks from the reference data are visible above the noise, supporting preferential orientation of the crystallites within the film. A conspicuous absence of isolated, 0kl = 0, h0l, and hk0 = 0 Bragg peaks in the XRD pattern, indicating that none of the CuSbS2 principle axes are oriented out-of-plane. In fact, the lattice planes from which the diffraction is observed in Figure 2c are all between 10 and 40° rotated with respect to the [010] or [100] lattice directions, suggesting that the c-axis is tilted between 50 and 80° with respect to the surface normal. This interpretation is supported by the SEM observations, which show what appear to be long, tilted grains within the thin film. The chemical, optical, and electronic properties of the CuSbS2 films were further characterized by XPS, UV−vis−NIR, and Hall measurements. An XPS survey scan indicates the presence of all expected elements (Cu, Sb, and S), Figure S3a. High-resolution scans for the Cu 2p, Sb 3d, and S 2p regions are displayed in Figure S3b−d, respectively. The Cu 2p highresolution scan shows the 2p doublet with binding energies of 931.0 (Cu 2p3/2) and 950.7 (Cu 2p1/2) eV and a separation of 19.7 eV, which is consistent with that previously reported for Cu+ in CuSbS2.17,35 That the Cu 2p peaks are symmetric and narrow, along with the absence of satellite peaks located at 942 (Cu 2p3/2 satellite) and 965 (Cu 2p1/2 satellite) eV, suggest neither Cu2+ or Cu0 were present in measurable quantities in the sample. The 3d core-level spectrum for Sb shows a doublet with binding energies of 528.1 (Sb 3d5/2) and 537.5 (Sb 3d3/2) eV, a spin−orbit splitting ratio near the expected 3:2 ratio, and

Figure 3. Absorption spectrum and Tauc plot (inset) for an 80 nm thick CuSbS2 film annealed at 225 °C for 15 min. A band gap of 1.6 eV was extrapolated from the Tauc plot through a linear regression fit.

and annealed at 225 °C for 15 min. The absorbance rises steeply around 800 nm, reaching an absorption coefficient of >1 × 104 cm−1 by 750 nm. A Tauc plot (Figure 3 inset) shows the relationship between the absorption coefficient, α, and the band gap by plotting αhν2 versus hν (eV), where h is Planck’s constant, and ν is the photon frequency. Fitting the curve in this plot with a linear regression and extrapolating the value at the x-intercept yields a band gap of 1.6 eV. Hall measurements of CuSbS2 films indicate p-type behavior with a hole concentration of 1.4 × 1015 cm−3, a mobility of 3.5 cm2 V−1 s−1, and conductivity of 7.7 × 10−4 Ω−1 cm−1. The measured hole concentration is consistent, albeit on the low end, with what has previously been published for CuSbS2 thin films. The p-type character is attributed to Cu vacancies that give rise to free charge carriers. Through theoretical simulations, Yang et al. estimated a hole concentration from copper vacancies of 1017−1018 cm−3 when CuSbS2 is grown at 300 °C,17 which is higher than that observed in the ALD-grown CuSbS2 films. The conductivity of the film can also be an indicator of stoichiometry and/or impurity phases as demonstrated by Welch et al.14 It was found that Cu-rich films (those approaching Cu12Sb4S13) gave rise to conductivity values exceeding 101 Ω−1 cm−1, those near the ideal CuSbS2 stoichiometry ranged from 10−3 − 101 Ω−1 cm−1, while Sb-rich films showed low conductivity values (10−4 − 10−3 Ω−1 cm−1). The conductivity value measured for the ALD-grown CuSbS2 thin films would suggest that these films might have a slight excess of Sb. It is also plausible that, in addition to an excess of Sb-doses to compensate for loss during annealing, the low growth and postprocess annealing temperatures, leads to fewer 4670

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ACS Applied Materials & Interfaces Table 2. CuSbS2 Thin-Film PV Device Configurations and Performance Metrics

a

device no.

device stack

Voc (V)

Isc (mA)

FF

η (%)

1 2 3 4 5

ITO/TiO2a/CAS/Au ITO/TiO2a/CAS/Au ITO/TiO2a/CAS/Spiro/Au ITO/TiO2b/CAS/Spiro/Au ITO/TiO2b/CAS/Spiro/Au

0.0291 0.025 0.173 0.291 0.313

0.0463 0.0446 0.054 0.031 0.041

0.3022 0.2637 0.2856 0.3495 0.3070

0.0023 0.0016 0.0148 0.017 0.0219

ALD TiO2 with TiTDMA bALD TiO2 with TTIP

to expect the HTL to aid in the extraction of the positive charge carriers (holes) from the CuSbS2 film, reducing the loss of photogenerated charge carriers and increasing the Voc, similar to that in inverted organic PV devices.38 Figure 4c shows the J−V response for the best performing ALD-grown CuSbS2 PV device (ITO/TiO2(TTIP)/CuSbS2/

Cu vacancies and, therefore, lower conductivity values in the ALD-grown CuSbS2 thin films. Finally, photovoltaic devices were assembled to evaluate the applicability of the ALD-grown CuSbS2 thin films. Given the toxicity associated with, as well as the known diffusion/cation exchange with metal sulfides by ALD,28,31 we chose to avoid using CdS as the n-type material for the p-n heterojunction. To the best of our knowledge, only a few reports have demonstrated a CuSbS2-based PV with a Cd-free n-type layer. In 2007, Manolache et al. fabricated a Cd-free CuSbS2 PV device using a dense 100 nm thick TiO2 n-type layer and a 285 nm thick CuSbS2 film both deposited by spray pyrolysis.21 They recorded a Voc, of 0.090 V, a short-circuit current, Isc, of 2.39 × 10−2 mA, and a fill factor, FF, of 0.286. In a more recent publication, Rastogi et al. demonstrated the promise for a PV device using a n-ZnO/CuSbS2 heterojunction, where the 50 nm ZnO layer was formed by sputter deposition, and the 600 nm CuSbS2 layer was deposited electrochemically.20 The device showed diode-like behavior with a turn-on voltage ∼0.5 V; however, the authors did not observe photovoltaic character under illumination. The Cd-free device structure used to assess our ALD-grown CuSbS2 thin films resembles the configuration glass/ITO/TiO2/CuSbS2/optional HTL/Au, as depicted in Figure 4a,b.17,36 (Note: A maximum theoretical Voc = ∼ 0.4 V was estimated for our CuSbS2/TiO2 PV devices. This estimation was based on an Ef = −4.86 eV for the CuSbS2, Ef = −4.5 eV for the TiO2, and the HOMO level of SpiroOMeTAD = −4.9 eV.17,36) The ∼20 nm TiO2 layer was deposited by ALD using two common Ti ALD precursors and water without postdeposition treatments. Table 2 and Figure S5 highlight the CuSbS2 devices fabricated. Device efficiencies were calculated using eq 1, where Isc, Voc, and FF were determined from the I−V plots, and Pin was calculated using the device area (0.18 cm2). η=

IscVoc FF Pin

Figure 4. (a) Schematic and (b) band alignment for CuSbS2 thin-film PV device architecture. (c) J−V curve for the device 5 in the dark (black) and under illumination (maroon). (d) EQE (black), IQE (blue), and LHE (green) plot.

Spiro/Au) in the dark (black) and under illumination (red). The measured Voc, Jsc, FF, and efficiency, η, were 0.313 V, 0.041 mA cm−2, 0.307, and 0.022% respectively. While the overall power efficiency remains well below 1%, this represents an improvement over state-of-the-art Cd-free CuSbS2 devices. The Jsc was further analyzed as a function of wavelength, Figure 4d. The external quantum efficiency, EQE, magnitude is consistent with the low Jsc observed in the J−V scan. The disproportionately strong blue response in the EQE may be attributed to insufficient light absorption or poor charge collection. The internal quantum efficiency (IQE) as derived from the lightharvesting efficiency (LHE) suggests the latter. At 600 nm and below, the CuSbS2 thin film is able to harvest 60% of the incident photons; however, at this wavelength, the IQE is 1019 cm−3. To investigate the role of the HTL, devices were fabricated with and without SpiroOMeTAD. The HTL did not have any apparent effect on the short-circuit current density (Jsc); however, an increase in the Voc by nearly an order-of-magnitude was observed when the Spiro-OMeTAD was included. This increase led to an order-ofmagnitude increase in the solar cell efficiency as well, suggesting the necessity of the HTL in these devices, much like that observed for ultrathin AgBiS2 PV.37 Due to the low conductivity of the ALD-grown CuSbS2 layer, it is reasonable 4671

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ACS Applied Materials & Interfaces the photoexcited carrier dynamics and lifetimes.39 These results suggest that further improvement may come from optimizing the CuSbS2 film thickness as well as the interface chemistry between the CuSbS2 thin film and adjacent layers.

was performed at Argonne National Laboratory, a U.S. Department of Energy Office of Science Laboratory operated under contract No. DE-AC02-06CH11357 by UChicago Argonne, LLC. Use of the Center for Nanoscale Materials, including resources in the Electron Microscopy Center, was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC0206CH11357.

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CONCLUSIONS One of the challenges of metal sulfide ALD is interfacial intermixing and in some cases even complete cation exchange. However, taking advantage of this known phenomenon may provide a low-temperature deposition route for complex multinary sulfides not achievable by other methods. Here, two members of the Cu−Sb−S family, namely, CuSbS2 and Cu12Sb4S13, have now been added to the short list of metal sulfide thin films that may be reproducibly synthesized by ALD. A careful in situ analysis of mass change upon mixed metal dosing as a function of temperature reveals the conditions under which additive and self-limiting growth may be achieved. With this careful control, phase-pure CuSbS2 thin films may be fabricated with optoelectronic properties suitable for photovoltaics as demonstrated for Cd-free devices. Although still far from optimal, this work points one way to earth-abundant, air stable, and flexible PV in which the properties of each layer may be tuned with atomic precision.

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ASSOCIATED CONTENT

S Supporting Information *

REFERENCES

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The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.6b13033. QCM data, EDS data, annealing studies, XPS spectra, J− V curves (PDF)

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ABBREVIATIONS PV, photovoltaic ALD, atomic layer deposition QCM, quartz crystal microbalance XRD, X-ray diffraction SEM, scanning electron microscopy XPS, X-ray photoelectron spectroscopy Voc, open-circuit voltage Jsc, short-circuit current FF, fill factor EQE, external quantum efficiency IQE, internal quantum efficiency LHE, light-harvesting efficiency

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Shannon C. Riha: 0000-0002-7716-6645 Alex B. F. Martinson: 0000-0003-3916-1672 Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS S.C.R. and A.A.K. were supported in part by the U.S. Department of Energy, Office of Science, Office of Workforce Development for Teachers and Scientists (WDTS) under the Visiting Faculty Program (VFP). S.C.R. was supported in part by the Department of Energy (DOE) Office of Energy Efficiency and Renewable Energy (EERE) Postdoctoral Research Awards under the EERE Solar Program administered by the Oak Ridge Institute for Science and Education (ORISE) for the DOE. ORISE is managed by Oak Ridge Associated Universities (ORAU) under DOE contract No. DE-AC0506OR23100. J.D.E., A.B.F.M., and M.J.P. were supported by the Argonne-Northwestern Solar Energy Research (ANSER) Center, an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Science under award No. DE-SC0001059. The research 4672

DOI: 10.1021/acsami.6b13033 ACS Appl. Mater. Interfaces 2017, 9, 4667−4673

Research Article

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