Article pubs.acs.org/cm
Searching for “Defect-Tolerant” Photovoltaic Materials: Combined Theoretical and Experimental Screening Riley E. Brandt,*,† Jeremy R. Poindexter,† Prashun Gorai,‡,§ Rachel C. Kurchin,† Robert L. Z. Hoye,†,@ Lea Nienhaus,† Mark W. B. Wilson,†,# J. Alexander Polizzotti,† Raimundas Sereika,∥ Raimundas Ž altauskas,∥ Lana C. Lee,⊥ Judith L. MacManus-Driscoll,⊥ Moungi Bawendi,† Vladan Stevanović,‡,§ and Tonio Buonassisi† †
Massachusetts Institute of Technology, Cambridge, Massachusetts 02139, United States Colorado School of Mines, Golden, Colorado 80401, United States § National Renewable Energy Laboratory, Golden, Colorado 80401, United States ∥ Faculty of Science and Technology, Lithuanian University of Educational Sciences, Vilnius 08106, Lithuania ⊥ Department of Materials Science and Metallurgy, University of Cambridge, Cambridge CB3 0FE, United Kingdom ‡
S Supporting Information *
ABSTRACT: Recently, we and others have proposed screening criteria for “defect-tolerant” photovoltaic (PV) absorbers, identifying several classes of semiconducting compounds with electronic structures similar to those of hybrid lead−halide perovskites. In this work, we reflect on the accuracy and prospects of these new design criteria through a combined experimental and theoretical approach. We construct a model to extract photoluminescence lifetimes of six of these candidate PV absorbers, including four (InI, SbSI, SbSeI, and BiOI) for which time-resolved photoluminescence has not been previously reported. The lifetimes of all six candidate materials exceed 1 ns, a threshold for promising early stage PV device performance. However, there are variations between these materials, and none achieve lifetimes as high as those of the hybrid lead−halide perovskites, suggesting that the heuristics for defect-tolerant semiconductors are incomplete. We explore this through firstprinciples point defect calculations and Shockley−Read−Hall recombination models to describe the variation between the measured materials. In light of these insights, we discuss the evolution of screening criteria for defect tolerance and high-performance PV materials.
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INTRODUCTION Thin-film polycrystalline photovoltaic (PV) materials offer the potential for lower-capital intensity manufacturing relative to crystalline silicon PV, but only if they can achieve high PV conversion efficiencies in excess of 20%.1,2 Long minoritycarrier lifetimes are necessary to achieve high efficiency, yet to date, there have been only a few classes of polycrystalline semiconductors that have demonstrated minority-carrier lifetimes in excess of 1 ns.3 Three classes of thin-film PV absorbers have achieved bulk lifetimes in excess of 100 ns, and they [Cu(In,Ga)Se2,4 CdTe,5 and the lead−halide perovskites (LHPs)6] are also the only polycrystalline thin-film materials to achieve PV conversion efficiencies in excess of 20%.7 Given the stability8,9 and scalability10,11 challenges for these leading thin-film materials, it is incumbent upon PV researchers to continue to screen new materials capable of achieving such long minority-carrier lifetimes.12,13 While lifetimes of ≥100 ns are considered a benchmark for commercial success, the available scientific literature suggests that a minority-carrier lifetime of 1 ns is a good benchmark for early stage PV © 2017 American Chemical Society
materials: it allows power conversion efficiencies on the order of 10% in materials with strong optical absorption and suggests the potential for further improvements.3 Including InP and other III−V materials, as well as the Cu2ZnSn(S,Se)4 family of materials14 in addition to CIGS, CdTe, and LHPs, there are several additional classes of thin-film inorganic (or hybrid organic−inorganic) polycrystalline semiconductors that are known to exceed this 1 ns threshold. In prior work, we hypothesized that a promising path to achieving these long lifetimes is through defect-tolerant semiconductors,12 building upon a body of work describing the physical underpinnings of defect tolerance.15−18 Defecttolerant semiconductors are posited to achieve long lifetimes because their intrinsic point defects (and ideally structural defects and extrinsic defects, as well) produce relatively slow electron−hole recombination kinetics. On the basis of a Received: December 30, 2016 Revised: May 8, 2017 Published: May 9, 2017 4667
DOI: 10.1021/acs.chemmater.6b05496 Chem. Mater. 2017, 29, 4667−4674
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Chemistry of Materials
allowed rapid phase identification, allowed determination of whether there was texturing, and reduced the likelihood that diffraction peaks were missed. Phase identification was performed through Le Bail25 or Pawley26 fitting of the diffraction patterns with crystallographic identification files of the tested material space group. Morphology Characterization. Scanning electron microscopy (SEM) imaging of samples was performed on a Zeiss Supra55 instrument with an Everhart−Thornley detector running in secondary electron detection mode, with a 5 kV electron gun voltage. Sections of each sample were prepared for planar and cross sectional imaging by cleaving a small portion of the substrate with a diamond scribe. Special care was taken not to touch the cross sectional area to preserve the asgrown structure, and cleaving was performed within 3 h of the measurements to minimize the environmental exposure of the cross section. Optical Characterization. UV−vis spectrophotometry was performed inside the integrating sphere of a PerkinElmer Lambda 950 UV−vis−NIR spectrophotometer using the method reported previously.27 To calculate the absorption coefficient from the transmittance and reflectance measurements, the thickness needed to be measured. A portion of the thin film was scratched off to create a step edge from the substrate to the top of the film, and a Dektak 150 profilometer was used to measure the height of the step edge. Spectrally resolved photoluminescence (PL) measurements were performed on a Horiba LabRam Evolution multiline PL/Raman spectrometer. All data were collected by using a 532 nm laser, which can access up to 20 mW of power distributed over an ∼500 nm Gaussian spot. Measurements were taken at room temperature in ambient air with 0.01−1% laser fluence (2−200 μW) depending on the measured material to maximize the signal-to-noise ratio. Emitted photons are assumed to be primarily due to radiative band-to-band recombination, and thus, the spectral dependence of the PL emission can help elucidate the electronic band structure of a material. For our purposes, PL emission energies are compared against the expected bandgap of the material and, in conjunction with structural data such as XRD spectra, can be used to confirm the correct material phase. Lifetime Characterization. TCSPC measurements were performed at room temperature in ambient air without background illumination using a 532 nm wavelength excitation laser (PicoQuant LDH-P-FA-530B), similar to a previous reports [a 405 nm wavelength laser was used for (CH3NH3)3Bi2I9].28 A single-photon-sensitive avalanche photodiode (Micro Photon Devices $PD-100-C0C) was used as the detector. The pulse width and detector resolution can be found in ref 28. The instrument response function was measured with the laser scatter from ground glass without filters in front of the detector. The laser power was adjusted such that the instantaneous incident photon fluence remained at 20 nJ/cm2 for BiI3, InI, BiOI, and SbSeI; fluences for (CH3NH3)3Bi2I9 and SbSeI were ∼1 and 62 nJ/cm2, respectively. Although the actual excess carrier concentration will vary depending on the absorption coefficient and charge-carrier lifetime, we anticipate excess carrier concentrations to remain between 1014 and 1016 cm−3 given the range of expected absorption coefficients (104− 105 cm−1) and calculated lifetimes (1−10 ns); see Table S-I. Defect Calculations. To study the physics of intrinsic point defects in InI and BiI3, density functional theory (DFT) calculations were performed using the VASP code.29 The details of these calculations are in the Supporting Information. Supercells containing 72 and 216 atoms were used for InI and BiI3, respectively. Following the methodology described in ref 30, finite-size corrections are applied to obtain reliable defect formation enthalpies: (1) image charge correction for charged defects, (2) potential alignment for charged defects, (3) band filling correction for shallow defects, and (4) band gap correction for shallow donors and/or acceptors. GW calculations were performed to correct for the band gap error in DFT-GGA calculations. The GW-corrected band gaps for InI and BiI3 are 2.1 and 2.15 eV, respectively (including spin−orbit coupling), which agree fairly well with experimental values.31,32 The GGA band structures are provided in the Supporting Information.
Shockley−Read−Hall (SRH) model for electron−hole recombination, we proposed several characteristics of the electronic structure of a material that would reduce the level of recombination in the presence of point defects or lead to defect tolerance.12 These include (i) small capture cross sections through the formation of defects with lower absolute charge states (or ideally neutral), (ii) small capture cross sections of charged defects due to screening by a high dielectric constant (related concepts include degree of ionicity and Born effective charge), and (iii) shallow (near the band edge) defect energy states, in particular through the use of partially oxidized cations such as Pb2+. Other authors have proposed additional criteria, including (iv) anion coordination number and bonding directionality,17 (v) ionicity/covalency of bonding,19−22 and (vi) the formation of paired defects that do not influence the electronic structure.19,23 Further arguments for the long carrier recombination dynamics in LHPs suggest that the material forms relatively few active recombination sites because of a low melting temperature or “self-healing”8,9 or that the radiative recombination is also suppressed by virtue of the material’s electronic structure.24 In this work, we test our original defect tolerance hypotheses by empirically evaluating the carrier lifetimes of several predicted materials. This includes the separate hypotheses that defect-tolerant electronic structures lead to shallow defects and that these shallow defects (and good screening) would then lead to long lifetimes. We also quantify the bulk lifetime by fitting time-resolved photoluminescence (TRPL) data collected using a time-correlated single-photon counting (TCSPC) setup. We evaluate the carrier lifetimes of bismuth triiodide (BiI3), methylammonium bismuth iodide [(CH3NH3)3Bi2I9], indium monoiodide (InI), bismuth oxyiodide (BiOI), antimony sulfoiodide (SbSI), and antimony selenoiodide (SbSeI), presenting the first reported time-resolved photoluminescence on InI, BiOI, SbSI, and SbSeI. After testing these six materials, we discover >1 ns effective lifetimes (accounting for bulk lifetime and surface recombination) in each of them, thus increasing the limited number of preexisting classes of polycrystalline photovoltaic absorbers that exceed this threshold.
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MATERIALS AND METHODS
Synthesis. Materials tested herein are synthesized through several methods: bulk crystal growth through vapor transport (SbSI and SbSeI) or synthesis from the melt (InI), chemical vapor transport (BiOI), physical vapor transport (BiI3), and solution synthesis [(CH3NH3)3Bi2I9]. Details of each synthesis tool are described further in the Supporting Information. All samples for X-ray diffraction (XRD), morphology, photoluminescence (PL), and TCSPC measurements were deposited onto 8 mm × 8 mm Si substrates with a thermal oxide surface layer. All samples for ultraviolet−visible (UV−vis) spectrophotometry measurements were deposited on 1.4 cm × 1.4 cm quartz (Quartz Scientific Inc.) substrates. All substrates were sequentially cleaned ultrasonically in deionized water, acetone twice, ethanol, and isopropanol for 5 min. Subsequently, they were oxygen plasma cleaned for 10 min under vacuum (−90 kPa, gauge). Structural and Phase Characterization. X-ray diffraction measurements were performed using a Bruker GADDS instrument at room temperature in ambient air. The sample was aligned to the center of the goniometer using a laser spot and camera. Twodimensional (2D) diffraction patterns were collected over a period of 15 min from 12° to 74° 2θ using monochromated Cu Kα X-rays. Collecting over a 2D area instead of a one-dimensional line scan 4668
DOI: 10.1021/acs.chemmater.6b05496 Chem. Mater. 2017, 29, 4667−4674
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Figure 1. TRPL decay curves of all six materials measured. Dashed gray lines represent biexponential fits to the data, while solid black lines are fits using a numerical model incorporating both Shockley−Read−Hall and radiative recombination (see the Supporting Information for details and estimated parameters). The fluence used was 20 nJ/cm2 for all traces except for those of (CH3NH3)3Bi2I9 (approximately 1 nJ/cm2) and SbSI (62 nJ/cm2). Data for (CH3NH3)3Bi2I9 were reproduced with permission from ref 33.
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however, all effective lifetimes as well as “slow” exponential time constants exceed 1 ns. Defect Calculations. Figure 2 shows calculated defect formation enthalpies as a function of Fermi level for InI and
RESULTS Photoluminescence Lifetime. TRPL data for all six materials are plotted in Figure 1. There is a clear variation in the TRPL behavior across the materials tested, with (CH3NH3)3Bi2I9 and InI demonstrating the slowest decay, followed by BiI3, SbSeI, SbSI, and BiOI. Biexponential fits to the data from Figure 1 are listed in Table 1, alongside estimates of the effective lifetime, τeff, Table 1. Biexponential Decay Time Constants and Calculated Effective Lifetimes from TRPL on Each Material (traces in Figure 1) biexponential fit
a
material
τfast (ns)
τslow (ns)
calculated effective lifetime (ns)
InIa BiI3 (CH3NH3)3Bi2I9 BiOI SbSeIa SbSIa
1.5 0.6 1.2 0.5 1.0 0.5
6.2 8.1 5.2 2.2 3.6 2.1
7.3 7.3 8.1 2.4 6.0 3.2
Denotes bulk polycrystal samples; others are thin films.
determined by performing numerical fitting incorporating generation, recombination, and diffusion terms and relating them to the change in excess carrier concentration over time. Using the methods of Evenor34 and Rosenwaks,35 the PL intensity was modeled by calculating the integral of the square of the excess carrier concentration into the depth of the film. τeff was calculated from the harmonic sum of the characteristic lifetimes associated with Shockley−Read−Hall, radiative, and surface recombination (see the Supporting Information for details). The traces are not well parametrized with two simple exponential curves given the mix of radiative and nonradiative decay and the intrinsic nature of most materials tested;
Figure 2. Vacancy and antisite defect formation enthalpies for InI and BiI3, as a function of Fermi energy under I-rich conditions. The arrow indicates the approximate expected position of the pinned Fermi level. 4669
DOI: 10.1021/acs.chemmater.6b05496 Chem. Mater. 2017, 29, 4667−4674
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Chemistry of Materials BiI3 under I-rich conditions. We considered intrinsic point defects, including vacancies and antisites. The position of the pinned Fermi level is indicated by the arrow in Figure 2. In each compound, the anion and cation vacancies are the predominant point defects. Antisite IBi is also a relevant defect in BiI3. For InI, both cation and anion vacancies have charge transition levels shallower than those in BiI3, where the charge transition levels are >0.5 eV from the corresponding band edges. Even though the antisites in InI introduce deep levels into the gap, they are unlikely to be present in significant concentrations because of the higher formation enthalpies at the position of the pinned Fermi level. Therefore, InI contains relatively shallow-level intrinsic point defects that are less likely to trap charges.
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Figure 3. Electronic structure of typical III−V, II−VI, or group IV semiconductors (left) compared to the lead halide perovskite crystal structure (right).
DISCUSSION In prior work, we posited that the materials evaluated herein have a higher probability of achieving longer carrier lifetimes, because of the fact that they share electronic structure similar to that of “defect-tolerant” lead−halide perovskites. Our results show that six predicted materials tested exhibit effective lifetimes in excess of 1 ns, which compare favorably to those of other “Earth-abundant” inorganic chalcogenide thin-film semiconductors, which typically show lifetimes of a few hundred picoseconds at this early stage of research.3 However, single nanoseconds are still far from the observed time of hundreds of nanoseconds achieved in lead−halide perovskites. This gap not only is due to the maturity of the lead−halide perovskite field but also may have physical underpinnings. In other words, the selection criteria presented previously appear to be incomplete because of these two considerations: there are strong variations observed among the materials tested, and it must be acknowledged that none of the materials tested to date achieve lifetimes on par with the LHPs (in excess of 1 μs). Indeed, several recent reviews of this emerging ns2 semiconductor class suggest that there are promising candidates, but there have not yet been clear success stories.36−38 Given this result, state-of-the-art predictive theory may only be partially successful in predicting defect-tolerant materials with properties on par with those of lead−halide perovskites. These discrepancies suggest that, as a field, we must direct predictive theory in one of two directions: (1) more accurate and less expensive methods for direct calculation of intrinsic defect energy levels in semiconductors,39,40 to allow a broad search through materials, or (2) additional refinements to our simplified heuristic models for what leads to intrinsically slow recombination kinetics in the presence of defects. Both thrusts should be pursued in parallel and will involve significant effort in both theoretical physics modeling and experimental validation of said models. To provide intuition and a stronger foundation for exploring the questions listed above, we revisit the role of defect-assisted recombination using a SRH model, consistent with the equations presented in ref 12. Initial Heuristic Model for Defect Tolerance. To recap our previous model for achieving shallow defect energy levels in the LHP material family,12 Figure 3 shows the special electronic structure of LHPs, in which intrinsic defects involving vacancies or broken bonds will produce defect energy levels near or within the band edges. Shockley−Read−Hall Recombination. It is important to note that there is some inconsistency in the use of the terms “shallow” and “deep” to describe defect states. Often “shallow”
refers to hydrogenic defects (effective mass-like), while “deep” refers to spatially localized defects. Here, we use the terms to refer to the binding energy of the defect or its energy level within the bandgap, where shallow defects are near a band edge (not as recombination-active) and deep defects are in the middle of the gap. In Figure 4, we aim to quantify the depth of the defect level, such that a deep defect is defined as being detrimental to carrier lifetime and a shallow defect is not. The Shockley−Read−Hall recombination model emphasizes the importance of both defect concentration and defect energy level in controlling recombination kinetics. First, the equilibrium defect concentration is related exponentially to its formation enthalpy as compared to the thermal energy available at a given processing temperature. In Figure 4a, we plot the vacancy concentration in a binary material, such as InI, as a function of the vacancy formation enthalpy, at two temperatures (note that first-principles defect calculations are performed assuming T = 0 K). At 300 K, a formation enthalpy of 0.5 eV leads to ∼1 ppm of the hypothetical vacancy in question, but at 600 K (representative of some deposition or annealing conditions), a 0.5 eV formation enthalpy would lead to ∼1 part per thousand. On the basis of this, we suggest that defects with formation enthalpies above 1 eV may not be particularly relevant for materials with low processing temperatures, but those with lower formation enthalpies are expected to be present in the material. Figure 4b quantifies the SRH carrier lifetime for a range of trap binding energies, Eb. Assuming a low level of injection, and symmetric capture cross sections for holes and electrons (e.g., a neutral defect), we plot the SRH lifetime for three different combinations of doping density and capture cross section. As the doping level increases, the range of trap energy levels that can be considered to be deep grows wider; in other words, shallow defects become worse at higher doping levels. For weakly doped, nearly intrinsic materials (1 ns lifetimes is promising in the context of photovoltaic device performance, as this threshold is consistent with >10% device performance in a wide range of materials.3 This also represents a near doubling of the number of known classes of inorganic polycrystalline thin-film materials with lifetimes in excess of 1 ns. Notwithstanding, longer lifetimes of >100 ns have been observed in all PV materials with efficiencies exceeding 20%. If 100 ns lifetimes are necessary for >20% efficiencies, the gap between the observed lifetimes in this study and those observed in successful thin-film devices invites further reflection on theoretical heuristics for defect tolerance. The first criterion is the symmetry of the crystal structure or electronic structure.48 In previous calculations,12 symmetric crystal structures (rock salts, perovskites, or CsCl-type structures) exhibited a stronger tendency to form direct bandgaps and lower effective masses, while the layered compounds with greater anisotropy in their crystal structure (and electronic structure) exhibited higher effective masses (i.e., mh* > 2) and flatter valence bands.12 None of the materials measured herein exhibit the three-dimensional (3D) structure and crystal symmetry of LHPs, and to date, the only related material to achieve such long lifetimes (>660 ns) is the highersymmetry double perovskite Cs2AgBiBr6.49 Recent work has found more evidence of the correlation between highly symmetric 3D crystal structures or high dimensionality in the electronic structure and improved photovoltaic performance.48
asymmetric rate of carrier capture, given the asymmetry in carrier populations. Given these plots, we can establish what “bad” intrinsic defects look like. First, while neutral defects are preferred, acceptors are much worse when their transition energy is above midgap, and donors similarly when their transition level sits below midgap. A higher injection level and a lower background doping density also reduce the rate of recombination. Finally, defects with formation enthalpies of >1 eV (across all Fermi levels) may be considered irrelevant for processing temperatures of 1 ns in several new PV materials. However, none of the six materials tested herein demonstrate lifetimes on par with those of the LHPs. To explain this, we have performed detailed theoretical studies of two of these materials. We demonstrate that not all materials containing partially oxidized, heavy ns2 cations contain only shallow intrinsic point defects, but that several may be able to achieve effective lifetimes in excess of 1 ns. Building on this initial success, we propose several strategies for revising the models for defect tolerance. In doing so, we hope to increase the predictive power of “defect tolerance” and thereby increase the probability of discovering the next successful polycrystalline PV absorber.
The second criterion is the importance of the choice of anion for a given cation. While cation s-orbital character in the valence band was discussed, it is important to recognize that 70−90% of the VBM partial density of states comes from anion porbitals in these materials. The dispersion of the VBM depends in large part on the hybridization between anion p-orbitals and cation s-orbitals. For BiI3, deep s-orbitals (Bi) and higherenergy p-orbitals (I) lead to very little interaction and minimal contributions from Bi to the valence band edge. Furthermore, the oxidation state of the cation will determine the charge state of the defects that tend to form. Given that Bi is in the +3 charge state, VBi can take on multiple charge states, including −1 and −3. Our calculations suggest that the −1/−3 charge transition level for VBi is quite deep in the bandgap. These deep, highly charged acceptors would be very detrimental to lifetime. Moving toward cations with a +1 and +2 oxidation state may decrease the likelihood of forming deep hole traps. Finally, while the original defect tolerance theories related primarily to cation vacancy defects, recent theory has offered some potential lessons regarding the anion vacancies.17,50 Shi and Du suggest that shallow anion vacancies arise when the cation dangling bonds that surround the vacancy have minimal overlap or interaction once the anion is removed. This occurs when the anion has a low coordination number (e.g., only 2 in the case of perovskites), and when cation−cation distances are long. This occurs more readily for larger anions (and larger lattice constants). They demonstrate that several selenides, iodides, and selenoiodides with low anion coordination all show shallow anion vacancies. Meanwhile, sulfides such as SnS and PbS demonstrate deep sulfur vacancies.51,52 This work may point to larger, less coordinated anions like iodine in the perovskite crystal structure. What is clear is that the MAPbI3 perovskite offers a very special combination of cation and anion, coupled with crystal structure, that lead to its shallow intrinsic point defects. There are likely very few compounds that hit this exact recipe correctly, but with careful design, we are hopeful another successful material may be discovered. For shallow cation vacancies, a partially oxidized cation in a low-valence state would be a good start: In+, Sn2+, Tl+, Pb2+, Cu+, and Ag+ have all demonstrated this property in multiple compounds. For shallow anion vacancies, one might look for a disperse CBM formed from heavy cation p-orbitals or to form the conduction band from nonbonding states as in TiO2.53 Alternatively, one may think about minimizing the potential cation−cation hybridization that occurs when an anion vacancy is formed,17 by reducing the coordination of the anion (using crystal structures like perovskite), increasing the cation−cation spacing (with large anions like I− or Br−), and/or ensuring that the orientation of the cations’ orbitals produces minimal overlap (for example, avoid 90° bond angles for materials with cation pderived CBMs). In fact, these suggestions for shallow anion vacancies also translate to shallow cation vacancies: the large cations listed above produce larger anion−anion separation, and if they do interact to form a deep state, the presence of cation s or d antibonding orbitals at the VBM ensures that this state ends up being resonant with the VB, not in the bandgap. Satisfying all or some of these features may require more clever searches or design based on crystal structure, bond lengths, coordination numbers, or the actual orbital geometries themselves.
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ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b05496. Full phase and optical characterization of all materials studied here, the approach and equations for TRPL fitting, and details on the first-principles calculations, including GGA-corrected band structures and phase diagrams (PDF)
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. ORCID
Riley E. Brandt: 0000-0003-2785-552X Jeremy R. Poindexter: 0000-0002-6616-9867 Lea Nienhaus: 0000-0003-1412-412X Present Addresses @
R.L.Z.H.: Cavendish Laboratory, University of Cambridge, J. J. Thomson Avenue, Cambridge CB3 0HE, U.K. # M.W.B.W.: Department of Chemistry, University of Toronto, Toronto, ON M5S 3H6, Canada. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS The authors acknowledge Dr. Vera Steinmann and Dr. Rupak Chakraborty [Massachusetts Institute of Technology (MIT)] for helpful discussions regarding materials synthesis and Dr. Rafael Jaramillo (MIT) for assistance with TRPL modeling. The theory and synthesis portions of this work were supported primarily through the Center for Next Generation Materials by Design (CNGMD), an Energy Frontier Research Center funded by the Department of Energy (DOE) Office of Science, Basic Energy Sciences; the characterization portion of this work was supported primarily through a TOTAL research grant funded through MITei. M.W.B.W., L.N., and M.B. acknowledge support from the Center for Excitonics, an Energy Frontier 4672
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Research Center funded by the DOE BES under Award DESC0001088 (MIT) for photoluminescence measurements. L.C.L. and J.L.M.-D. thank the EPRSC Centre for Doctoral Training: New and Sustainable Photovoltaics, and the Cambridge Winton Programme for the Physics of Sustainability for funding. Research on bismuth compounds was funded under NSF Grant CBET-1605495. This work made use of the Shared Experimental Facilities supported in part by the MRSEC Program of the National Science Foundation under Grant DMR-1419807, and the Center for Nanoscale Systems (CNS, at Harvard University), a member of the National Nanotechnology Infrastructure Network (NNIN), supported by the National Science Foundation (Grant ECS-0335765).
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