A-Site Cation in Inorganic A3Sb2I9 Perovskite Influences Structural

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Cite This: Chem. Mater. 2018, 30, 3734−3742

A‑Site Cation in Inorganic A3Sb2I9 Perovskite Influences Structural Dimensionality, Exciton Binding Energy, and Solar Cell Performance Juan-Pablo Correa-Baena,* Lea Nienhaus, Rachel C. Kurchin, Seong Sik Shin, Sarah Wieghold, Noor Titan Putri Hartono, Mariya Layurova, Nathan D. Klein, Jeremy R. Poindexter, Alexander Polizzotti, Shijing Sun, Moungi G. Bawendi, and Tonio Buonassisi* Massachusetts Institute of Technology, 77 Massachusetts Avenue, Cambridge, Massachusetts 02139, United States

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ABSTRACT: Inspired by the rapid rise in efficiencies of lead halide perovskite (LHP) solar cells, lead-free alternatives are attracting increasing attention. In this work, we study the photovoltaic potential of solution-processed antimony (Sb)based compounds with the formula A3Sb2I9 (A = Cs, Rb, and K). We experimentally determine bandgap magnitude and type, structure, carrier lifetime, exciton binding energy, film morphology, and photovoltaic device performance. We use density functional theory to compute the equilibrium structures, band structures, carrier effective masses, and phase stability diagrams. We find the A-site cation governs the structural and optoelectronic properties of these compounds. Cs3Sb2I9 has a 0D structure, the largest exciton binding energy (175 ± 9 meV), an indirect bandgap, and, in a solar cell, low photocurrent (0.13 mA cm−2). Rb3Sb2I9 has a 2D structure, a direct bandgap, and, among the materials investigated, the lowest exciton binding energy (101 ± 6 meV) and highest photocurrent (1.67 mA cm−2). K3Sb2I9 has a 2D structure, intermediate exciton binding energies (129 ± 9 meV), and intermediate photocurrents (0.41 mA cm−2). Despite remarkably long lifetimes in all compounds (54, 9, and 30 ns for Cs-, Rb-, and K-based materials, respectively), low photocurrents limit performance of all devices. We conclude that carrier collection is limited by large exciton binding energies (experimentally observed) and large carrier effective masses (calculated from density functional theory). The highest photocurrent and efficiency (0.76%) were observed in the Rb-based compound with a direct bandgap, relatively lower exciton binding energy, and lower calculated electron effective mass. To reliably screen for candidate lead-free photovoltaic absorbers, we advise that faster and more accurate computational tools are needed to calculate exciton binding energies and effective masses.

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INTRODUCTION Lead-based hybrid organic−inorganic halide perovskite solar cells (LHPs) have attracted much attention for photovoltaic (PV) applications because of their low fabrication cost and high power conversion efficiency (PCE).1,2 PSCs have made impressive progress in just a few years with PCEs of 3.8%3 in 2009 to a certified 22.1%4 in 2016. These leadcontaining absorbers present a concern regarding the water solubility of the toxic PbI25 and therefore challenging their wide-scale use. This has inspired researchers to develop nontoxic perovskite alternatives. To find alternative Pb-free materials with similar optoelectronic properties to the Pb-containing analogues, selection criteria need to be established. In particular, the properties that make lead-halide perovskites “defect-tolerant” need to be better understood. Research points at energetically shallow intrinsic point defects,6,7 which may account for the long charge carrier lifetimes8,9 allowing for some of the lowest open-circuit voltage (VOC) deficit of any photovoltaic material.10,11 Shallow defect states are linked to the formation of antibonding orbitals at band edges,7 and this antibonding character can occur in materials other than Pb halide perovskites, such as those containing other heavy ns2 metal cations. In particular, Sb- and Bi-based materials have been identified as such alternatives.7,12 © 2018 American Chemical Society

Materials based on low-dimensional bismuth-containing inspired on LHPs, with a structure A3Bi2I9, where A is methylammonium (MA), formamidinium (FA), Cs, or Rb, have recently attracted attention as PV absorbers.13−17 As predicted,7 these materials exhibit charge carrier lifetimes above 1 ns,14,18 a metric often used to assess the suitability of materials for PV applications.19 However, the low-dimensional nature of these compounds and their large strongly bound electron−hole pairs (excitons), with binding energies above 300 meV, make them undesirable for applications as thin-film planar absorbers due to difficult charge extraction. Materials based on 0D and 2D antimony-containing perovskites with a structure A3Sb2I9, analogous to Bi perovskites, have more recently attracted the attention of researchers looking for the next lead-free material. Theory and experimental work have shown that evaporated Cs3Sb2I9, with a layered 2D perovskite structure, can be suitable for PV applications due to its direct bandgap of 2 eV. However, the solution-processed, room-temperature-stable Cs3Sb2I9 polymorph consists of 0D dimers.20 Similarly, solution-processed Received: February 13, 2018 Revised: May 21, 2018 Published: May 22, 2018 3734

DOI: 10.1021/acs.chemmater.8b00676 Chem. Mater. 2018, 30, 3734−3742

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Chemistry of Materials MA3Sb2I9 produces 0D structures.21 Because dimensionality is an important factor affecting PV metrics,22 and 2D and 3D materials tend to have more preferred electronic transport symmetry than 0D,7 we believe A3Sb2I9 with 2D structures are more desirable. Exciton dissociation into free carriers (quantified by the material’s exciton binding energy) is another important parameter that leads to high photocurrents and therefore PV performance. Controlling dimensionality and exciton binding energy is therefore vital in the quest for highefficiency Pb-free PSCs. Perovskites are well-known for the ease with which their optoelectronic properties can be tuned for tailored functionalities. By simple halide or A-site cation substitution, Pb halide perovskites are able to exhibit bandgaps that span the visible spectrum.1 Therefore, studying the role of the A-site cation in A3Sb2I9 on its optoelectronic properties is desired. In this work, we modify the A3Sb2I9 perovskite structure and the material’s properties by exchanging the A-site inorganic cations, including Cs, Rb, and K. We show that the A-site cation dramatically affects the structural properties of the material. Cs3Sb2I9 takes on a 0D (dimer) structure consisting of face-sharing SbI6 octahedra.23 On the other hand, Rb3Sb2I924 and K3Sb2I9 take on a 2D structure, with the corning-sharing octahedra similar to the layered LHP analogues.25 This leads to changes in the absorption spectra and optical bandgap. Thin films of the Cs and K compounds exhibit indirect bandgaps of 2.43 and 2.03 eV and exciton binding energies of 166 and 120 meV, respectively. Thin films of Rb showed a direct bandgap of 2.02 eV and exciton binding energy of 107 meV. When employed as solar cell absorbers in a TiO2/A3Sb2I9/Spiro-OMeTAD/Au device configuration, the Cs and K materials performed poorly, yielding PCEs of 0.03% and 0.17%, respectively. The Rb material yielded a higher efficiency above 0.76%. The solar cells are mostly limited by their low photocurrents, although these are also highly dependent on the perovskite composition. The Cs perovskite yielded 0.13 mA cm−2 due to its large exciton binding energy and dimer structure. The Rb perovskite shows over an order of magnitude higher photocurrent of 1.84 mA cm−2 due to improved charge extraction of 2D materials and to its lower exciton binding energy. The Rb3Sb2I9 compound presents a very promising material and would benefit greatly from a lower exciton binding energy and effective mass. The lessons from this work will help the photovoltaic community design new materials with the potential for higher efficiencies.

Figure 1. Optical properties of A3Sb2I9. (A) Film absorptance calculated from 100 − T − R, where T and R are transmittance and reflectance, respectively, with 2 nm wavelength resolution. Photographs of films deposited by spin-coating solutions of A3Sb2I9 where the A site is (B) Cs, (C) Rb and (D) K; scale bar is 1 cm. Calculated (from DFT) band structure of (E) Cs3Sb2I9, (F) Rb3Sb2I9, and (G) K3Sb2I9. The calculated band gaps are indirect between M−K in the valence band and Γ in the conduction band for Cs3Sb2I9, direct at Γ for Rb3Sb2I9, and slightly indirect between Γ in the valence band and along Γ−Y in the conduction band for K3Sb2I9.

bandgaps21 undesirable for PV applications. The cations with the smallest cationic radii, Na and Li, form white and yellow films, respectively. Bandgap Estimation. Given the appropriate range of bandgaps as solar cell absorbers (specially for tandem configurations), the Cs, Rb, and K materials are investigated in-depth. The absorptance of these Sb perovskite-inspired materials is shown in Figure 1A, as measured by UV−vis spectroscopy. The corresponding photographs of the measured films are shown in Figure 1B−D. Clearly, exchanging the A-site cations can significantly change the absorption onset. The extracted bandgaps were derived from Tauc plots and are presented in Figure S2. The Cs material forms a yellow film and has an absorption onset at 2.43 eV. The Rb analogue, by contrast, has a sharper onset at around 2.03 eV. Similarly, the K perovskite has a sharp onset at 2.02 eV. The density function theory (DFT)-calculated bandgap for the Cs compound is underestimated at 1.88 eV (indirect), as compared to the experimental value of 2.43 eV, in contrast to the work by Saparov et al., which calculates the bandgap by DFT at 2.4 eV for the dimer compound.20 The other two compounds are close to the experimental values at 1.99 eV (direct) and 2.03 eV (slightly indirect) for the Cs, Rb, and K materials (Figure 1E− G and Figure S2). Both the Rb and K materials show relatively smooth curves with no pronounced features. The Cs perovskite sample shows a small feature at 570 nm, which can be attributed to an exciton forming at room temperature. Interestingly, processing plays a role on the structure and therefore on the optoelectronic properties of Cs3Sb2I9. In a previous study, samples prepared by evaporation showed a

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RESULTS AND DISCUSSION Thin-Film Preparation. We prepared thin films of the Sbcontaining perovskites with both organic and inorganic cations, including FA, MA, Cs, Rb, K, Na, and Li (Figure S1). Solutions containing AI and SbI3 were spin-coated on glass substrates at room temperature, where A is the A-site cation. The AI:SbI3 (with a molar ratio of 3:2) salts are dissolved in dimethylformamide (DMF) and prepared fresh before each deposition. The sulfur-containing solvent, dimethyl sulfoxide (DMSO), tends to coordinate with Sb, forming a dark solution over the course of a few hours. Therefore, with the exception of CsI (which does not dissolve in DMF), DMSO was not used in order to avoid changes in morphology due to such interactions, as has been suggested for Pb-based perovskites.26 The films produced with this method are smooth, as will be discussed below. As seen in the photographs in Figure 1, changing the cation produces a relatively wide range of colors. The large organic cations form yellow films (FA and MA), with wide 3735

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Figure 2. Structural properties of A3Sb2I9. X-ray diffraction patterns (A, B, C) and corresponding crystal structures (D, E, F) with phase stability diagrams (G, H, I) for (A, D, G) Cs3Sb2I9, (B, E, H) Rb3Sb2I9, and (C, F, I) K3Sb2I9 from powders and films synthesized from solution. Powders were used to extract all peak information from the A3Sb2I9 compounds. Ritvield refinement was performed on the powder materials, and the calculated patterns for AI and A3Sb2I9, where A is Cs, Rb, or K, are presented for each material. The asterisk denotes a minority impurity phase (unconverted KI) in the K3Sb2I9 film. Films of Cs3Sb2I9 and Rb3Sb2I9 did not show any minority impurity phases.

stability of K3Sb2I9 relative to KI or SbI3 is much narrower compared to the Rb and Cs analogues (Figure 2G−I), as predicted by our calculations. Rietveld refinement of the powders (Figure 2A−C) shows that both phases are present; the calculated patterns for the pure phase (A3Sb2I9) and minority impurities (AI) are shown. The structures are shown in Figure 2D−F, as 0D, 2D, and 2D for Cs-, Rb-, and K-based materials, respectively. On the basis of the refined XRD patterns, and taking into account only the pure A3Sb2I9, we construct the structures of the studied compounds in order to better visualize them.20 Cs3Sb2I9 crystallizes in space group P63/mmc (no. 194); the Sb2I93− dimers share their triangular faces, making isolated structures or 0D (Figure 2D). In contrast, Rb3Sb2I9 and K3Sb2I9, form 2D structures with K or Rb cations acting as spacers between the corner-sharing octahedra layers.20,27,28 The Rb and K compounds form lead-free ⟨111⟩-stacked layered perovskite architecture isostructural to (NH4)3Bi2I9. Both these structures crystallize in the space group P3̅m1; these were refined starting from Cs3Sb2I9, and the peak profiles are consistent with a 2D layered perovskite structure with corning-sharing SbI6− octohedra as described elsewhere.20 The thin films of these compounds show the same characteristics of 0D, 2D, and 2D structures for the Cs, Rb, and K materials (see discussion in Figure S4). Thin films were also deposited on preheated substrates (otherwise known as “hot casting”) as described in other reports29 dealing with 2D materials to see the effect on structure and photovoltaic performance, which will be evaluated below. Figure S5 presents the XRD patterns of the three spin-coated materials, comparing room temperature versus the hot casting approach with preheated substrates at

smaller absorption onset at 1.89 eV in a 2D structure. The same work reported dimer structures with larger bandgaps when processed from solution (similar to our work). Without the need to evaporate, by exchanging the A-site cation, it is possible to control the dimensionality of the compound and its excitonic characteristics. The density of states (DOS) is plotted in Figure S3 for all three compounds. Structural Properties. To understand the structural properties of these compounds, we performed one-dimensional X-ray diffraction (XRD) on powders and films of A3Sb2I9 perovskites based on Cs, Rb, and K, as seen in Figure 2A−C, respectively. Looking at the films of these compounds alone, it is clear that due to high preferential orientation of the peaks between 20° and 30° (showing that the films have preferred orientation of (003), similar to the previous reports20), measurements on randomly oriented powders are needed in order to extract all peak information. We therefore prepared powders of these materials using the same precursor solutions used for films, as described above. Precursor solutions Cs, Rb, and K of AI:SbI3 (3 mL of AI and 2 mL of SbI3 for 1 M precursor solutions) in DMF were stirred under constant heat (80 °C) for 3 h. Powders were made by chlorobenzene addition; details can be found in the Experimental Section. Interestingly, powders produced from this method show some degree of minority phase impurity formation or unreacted AI. All samples show a majority phase of the A3Sb2I9, with at least 80% of the powder estimated to be of the desired perovskite phase. On the basis of the deconvolution, we see that thin films of K3Sb2I9 show small peaks that belong to KI; however, the Rb and Cs analogues do not show any traces of the unreacted AI and show only the pure phase (A3Sb2I9). This is not surprising, given that the chemical potential range for thermodynamic 3736

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Figure 3. Luminescence properties of A3Sb2I9 thin films. (A) Time-resolved photoluminescence at room temperature. The laser power was 12.5 μW, and the excitation wavelength was 532 nm. The data were fitted using a triexponential decay function. (B) Extracted photoluminescence counts (integrated from the time-resolved photoluminescence) as a function of temperature. The photoluminescence counts are divided by the laser power (ranging from 1.3 to 67 nW) and corrected with the power law coefficient.24,30 (C) Extracted constant decay of the long-tail photoluminescence as a function of temperature for Cs3Sb2I9, Rb3Sb2I9, and K3Sb2I9, from (D), (E), and (F), respectively. Temperature-dependent time-resolved photoluminescence of (D) Cs3Sb2I9, (E) Rb3Sb2I9, and (F) K3Sb2I9 thin films. The temperature ranges from 80 to 320 K as labeled to the right of (E).

at 774 nm, despite having a much wider bandgap (Figure S2), which may indicate trap-assisted emission. All samples emitted poorly at room temperature, and it was therefore necessary to cool down the films to measure them. The TRPL signal for all three compounds is relatively slow with an initial fast decay followed by a slow component. We fitted the TRPL signal to a triexponential function; here we report the longest time constant (see Table S1 for a summary of the fits). The Cs material exhibits a slow component of 54 ns, which is expected for 0D structures where charges are confined within the dimer.31 As the structure changes from 0D to 2D for the Rb and K samples, the lifetimes become shorter. Similar to Cs samples, the Rb analogues show a fast initial decay followed by long tail of 9 ns. The K samples show a similar fast initial decay and a long tail which is slower than that of the Rb, and shorter than Cs samples, of 30 ns. The details of all time constants extracted from the fits of the TRPL data can be found in Table S1. From the TRPL measurements it is clear that all materials are suitable for solar cell applications due to their long-lived charge carriers, and this should not be a limiting factor. Defect tolerance has also been suggested by theory for this class of materials, where deep traps tend not to form in the presence of many types of vacancies.7 Because charge carrier lifetime alone is not the only feature useful to understand the viability of a material as a photoabsorber in solar cell, we investigate PL counts

50°C for 5 min. Films of the Cs materials (Figure S5A) do not show pronounced differences between the deposition methods. In contrast, upon hot casting, Rb films (Figure S5B) show an increased preferential orientation along the (006) plane, similar to other reports,28 rendering it detrimental for charge transport in PV devices due to their planes being parallel to the substrate. Because these layered structures tend to exhibit anisotropic charge transport, orientation along this plane should be avoided, and planes perpendicular to the substrate are desirable. Interestingly, the hot casting approach seems to have the opposite of the intended effects and, in contrast to the organic− inorganic lead halide 2D perovskites, does not aid in orienting the 2D planes perpendicular to the substrate. This should translate into lower photocurrents in solar cell and will be investigated below. The effects of hot casting are also detrimental for K3Sb2I9 (Figure S5C) where the peaks of KI become more prominent. However, the hot casting method does seem to have a positive effect in the orientation where the film becomes less oriented on the 25.9° peak (003) and other peaks, such as the one at 44.4° (006), become more prominent. Photoluminescence Properties. We measured the timeresolved photoluminescence (TRPL) properties of A3Sb2I9 thin films based on Cs, Rb, and K (Figure 3A). The steady-state photoluminescence (Figure S6) shows sharp emission at around 635 nm for the Rb and K samples upon cooling the thin films to 200 K or lower. Interestingly, the Cs samples emit 3737

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Chemistry of Materials Table 1. Characteristics of the A3Sb2I9 Compounds Studieda chemical formula

m*h

m*e

EG/eV (calcd from DFT)

Cs3Sb2I9 Rb3Sb2I9 K3Sb2I9

1.55 1.55 1.83

1.40 1.11 3.15b

1.89 1.99 2.03

bandgap type

EG/eV (from absorptance in Figure S2)

VOC/ mV

JSC/ mA cm−2

fill factor/%

power conv efficiency/%

2.43 2.03 2.02

404 660 338

0.13 1.84 0.413

58 63 50

0.03 0.76 0.07

indirect direct slightly indirect

a DFT-computed effective masses, bandgaps, and types; bandgap obtained from absorptance, and solar cell parameters for the best performing device of each composition: open-circuit voltage, short-circuit photocurrent density, fill factor, and efficiency. bThis fit did not fully converge; therefore, the value might be underestimated.

mass (Table 1), we can predict that the Rb compounds will perform better in a solar cell than the K and Cs counterparts. Thin-Film Morphology. To further understand the morphology change induced by the A-site cations, we prepared thin films of A3Sb2I9 on substrates for solar cell applications. A sprayed pyrolyzed TiO2 layer followed by a mesoporous TiO2 electron selective layer serves as the foundation for the photoabsorber crystal growth. In Figure 4, scanning electron

(integrated from the TRPL signal and devided by the laser power and corrected with the power law coefficient24,30) as a function of temperature. As shown in Figure 3B, the PL intensity increases as temperature is decreased for all three materials. This is indicative of increasing emission from the excitonic state, and a reduction of charge generation caused by thermally splitting the exciton, a result of less thermal energy available at lower temperatures. To extract the exciton binding energy, we investigate the thermal quenching of the excitonic emission fitted using24,32 I (T ) =

I(0 K) 1 + A exp( −E B /kT )

where I and A are constants and EB is the exciton binding energy. Since the emission quenches very readily at higher temperatures (emission is mostly quenched above roughly 105 K),24 the excitation power is increased to obtain a higher count rate. To account for the slight power law dependence of these materials, the total integrated counts are corrected by the excitation power with the material-specific power law coefficient.30 We use 0.87 for Cs24 and 1.08 for Rb24 and estimate 1.022 for K based on our PL data. To account for the dark counts on the detector, we estimate the dark counts from the prepulse background and subtract 1500 cps from all total counts. Because of the nonmonotonic behavior of the photoluminescence quantum yield (PLQY) of the Cs and K perovskites (a result of the shallow traps, from which the exciton cannot thermalize up into the emissive state), the lowest temperatures at which the fast component appears in the lifetime are omitted in the fit. Fitting the resulting data using the above equation yields exciton binding energies of 166 ± 12, 107 ± 10, and 120 ± 6 meV for the Cs, Rb, and K compounds, respectively (Table 1). Looking at the TRPL as a function of temperature (summarized in Figure 3C, lifetimes can be seen in Figure 3D−F), we can further understand the excitonic behavior of these compounds. The three materials studied exhibit longer lifetimes as temperature is reduced, and the long tails at 80 K are on the order of microseconds. For the Cs and K materials, an initial fast component emerges, which can be attributed to trapping of the exciton into shallow traps, out of which the exciton cannot escape at low temperature. Extracting the lifetimes from the long tail of the traces in Figure 3D−F, can also give us insights into the exciton binding energy. Fitting the resulting data (Figure 3C) using the same equation as above yields exciton binding energies of 184 ± 46, 95 ± 5, and 137 ± 24 meV for the Cs, Rb, and K compounds, respectively. Larger errors for the Cs and K films are due to the nonmonoexponential decay of the PL. On the basis of the lower exciton binding energy and its 2D nature with lower effective

Figure 4. Morphology of A3Sb2I9 thin films. Images of (A, D) Cs3Sb2I9, (B, E) Rb3Sb2I9, and (C, F) K3Sb2I9 thin films grown on fluorine-doped tin oxide glass/mesoporous TiO2 (300 nm) substrates. (A, B, C) are scanning electron micrographs, whereas (D, E, F) are atomic force micrographs.

micrographs (SEM) show the different morphology induced by the change in A-site cation. Cs3Sb2I9 films grown on mesoporous TiO2 (Figure 4A) show films with little to no features, composed of small grains, with sizes in the tens of nanometers. Rb3Sb2I9 forms larger grains in the hundreds of nanometers and some as large as 1 μm (Figure 4B). K3Sb2I9 forms films with very large grains in the micrometer range. Small crystallites are visible, and it is possible that these large grains are composed of smaller crystals in the compact film. We performed atomic force microscopy (AFM) to further understand the morphological features of these films (Figure 4D,E). Roughness is a very important feature in the preparation of solution-processed solar cells because the layers are spincoated and therefore sensitive to morphological features. In general, low roughness is desired for planar junction solar cells for homogenous charge extraction. We calculated a root-meansquare (RMS) value of 49.1, 43.1, and 156.7 nm for the Cs, Rb, and K films, respectively. The roughness of Cs3Sb2I9 and Rb3Sb2I9 is relatively low and would not represent a major obstacle for solar cell preparation. Ideally, the roughness of this layer is much lower than the layer that will be subsequently deposited, which in this case is a 200 nm Spiro-OMeTAD film. The roughness of the K 3 Sb 2 I 9 film might present a 3738

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Figure 5. Thin-film solar cells employing A3Sb2I9 thin films as photoabsorbers. (A) Schematic of the cross section of a completed solar cell, followed by cross-sectional scanning electron micrographs of solar cells based on (B) Cs3Sb2I9, (C) Rb3Sb2I9, and (D) K3Sb2I9 phtoabsorbers, with the device configuration glass/fluorine-doped tin oxide/mesoporous TiO2/photoabsorber/Spiro OMeTAD/gold. Scale bar is 200 nm. (E) Current−voltage plots for all compounds and (F) for the champion Rb3Sb2I9 device. (G) Short-circuit current (JSC) and (H) power conversion efficiency distribution for the three different photoasborbers in full solar cells as measured at AM1.5 at room temperature with a scan rate of 10 mV s−1.

(95−107 meV, Figure 3B), are structured in 2D, and are more thermodynamically stable than the K analogues (Figure 2 and Figure S3). This means that excitons are more likely to dissociate into free carriers at room temperature than the Cs analogues. Those dissociated charges are then able to travel more freely in a 2D structure to dissociate charges at the contacts. This combination of higher dimensionality and lower exciton binding energy makes Rb3Sb2I9 the more promising candidate as a photoabsorber. The highest efficiency achieved with the Rb3Sb2I9 is 0.76% (Figure 5F), and if the 2D planes were oriented perpendicular to the contacts, we might achieve much higher values. Several devices were prepared to understand the reproducibility of these devices, and the results for Short-circuit current (JSC) and PCE are summarized in Figure 5G,H. As expected, all devices fabricated with Cs3Sb2I9 and K3Sb2I9 are limited by the JSC, whereas Rb3Sb2I9 devices show around an order of magnitude higher values. As discussed above, we attempted an approach of hot casting, which has been shown to yield orientation of 2D planes perpendicular to the contacts,29 facilitating charge extraction. From the XRD analysis (Figure S5) we see no clear changes in the orientation of the planes in the films processed on heated substrates (50 °C preheating). We made solar cells with all three compounds using the hot casting approach (Figure S7) and found that Cs and Rb samples perform slightly worse, with photocurrents lower than the room temperature analogues. The samples using the K photoabsorber by hot casting showed slightly higher performance than the room temperature samples. This is surprising as these samples show large amounts of impurity phase in the XRD (Figure S5) patterns. LHPs have been shown to be tolerant to remnants of unconverted PbI2, showing a beneficial morphological effect and no drop in performance.36 It is possible that similar to LHPs, these corner-sharing Sb-based octahedra are tolerant to large amounts of KI impurities, while improving morphology, which can lead to the slightly higher performance. To understand whether excitons in these materials can be dissociated into free carrier, translating into larger photocurrents, we measured JSC as a function of temperature (0−60

morphological obstacle for obtaining efficient solar cells due to the possibility of shunt creation. Solar Cell Performance. To further prove our hypothesis that the Rb compound will perform better than the K and Cs counterparts, we fabricated full devices. The device stack (Figure 5A) is composed of a fluorine-doped tin oxide-coated glass, an A3Sb2I9 photoabsorber-infiltrated mesoporous TiO2 layer (ca. 200 nm), a thin layer of compact A3Sb2I9 (50−100 nm), Spiro-OMeTAD (ca. 200 nm), and a top gold electrode (100 nm). Cross-sectional SEM images of three representative devices are shown for all photoabsorbers (Figure 5B−D). The Cs3Sb2I9 (Figure 5B) device shows a compact photoabsorber atop the mesoporous TiO2, with a thickness of roughly 50 nm, whereas for the Rb3Sb2I9 (Figure 5C) and K3Sb2I9 (Figure 5D) devices this layer is 100 and ca. 20 nm, respectively. Current−voltage curves for the different photoabsorbers are presented in Figure 5E, where all materials show photovoltaic properties. The Cs3Sb2I9 devices show the lowest photocurrents of around 0.1 mA cm−2, as expected due to the large exciton binding energy (166−184 meV, Figure 3B) and low dimensionality (0D, Figure 2D). The former makes it difficult to dissociate into free carrier at room temperature which would be otherwise thermally activated and affects the effective mass (and therefore mobilities; Table 1) of the materials. 3D lead halide perovskites were originally believed to exhibit exciton binding energies of around 50 meV;33 however, recent reports show lower energy of around 16 meV,34,35 allowing for easy dissociation into free carriers. Dimensionality in these devices also plays an important role, as 0D materials tend to have confined charges (similar to quantum dots), making charge transport challenging from dimer to dimer.31 Devices with Rb3Sb2I9 show an order of magnitude increase in photocurrents to >1 mA cm−2. Part of this may be explained by the slightly narrower bandgap of this compound compared the Cs analogue as well as the 50 nm thicker photoactive layer. However, an order of magnitude is very large difference, and most of it cannot be explained by those two factors. The Rb photoabsorbers have three additional advantages over the other two compounds: they have lower exciton binding energies 3739

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Chemistry of Materials °C), as shown in Figure S8. The devices were measured at 60 °C, and the temperature was progressively modulated to 0 °C and back to 60 °C. While there were changes in JSC with temperature modulation, these were rather small, suggesting either that the exciton cannot be thermalized at temperatures as low as 60 °C or that the photocurrents are limited mainly by the electronic dimensionality of these materials. Nonetheless, for the Rb and K compounds an optimum JSC is found at around room temperature, whereas for Cs the optimum is at either 60 or 0 °C. Since the thermalization of the exciton is less likely to happen at 0 °C, it is unlikely this behavior is related to excitonic behavior.

components. The A3Sb2I9 solutions were spin-coated onto the mesoTiO2/bl-TiO2/FTO substrate by consecutive two step spin-coating method as described above. A solution of Spiro-OMeTAD (Lumtec) in chlorobenzene (100 mg/mL) as a hole transporting material (HTM) was mixed with bis(trifluoromethylsulfonyl)imide lithium salt (Li-TFSI, Sigma-Aldrich), tris(2-(1H-pyrazol-1-yl)-4-tertbutylpyridine)cobalt(III) tris(bis(trifluoromethylsulfonyl)imide) (FK209, Lumtec), and 4-tert-butylpyridine (tBP, Sigma-Aldrich). The molar ratio of additives for spiro-OMeTAD was 0.5, 0.03, and 3.3 for Li-TFSI, FK209, and tBP, respectively. This solution was deposited on top of the A3Sb2I9 film by spin-coating at 4000 rpm for 20 s. Finally, a 100 nm Au top electrode was deposited by thermal evaporation. The active area of this electrode was fixed at 0.16 cm2. General Characterization. The crystal structure and phase of the materials were characterized using an XRD (Rigaku SmartLab and Bruker D8 DISCOVER X-ray diffractometer). Rietveld refinements of all A3Sb2I9 compounds were performed using Highscore Plus version 4. The morphology was investigated by a Zeiss Merlin field-emission scanning electron microscope (FESEM, Zeiss). J−V curves were measured using a solar simulator (Newport, Oriel Class AAA, 91195A) with a source meter (Keithley 2420) at 100 mW cm−2, AM 1.5 G illumination, and a calibrated Si-reference cell certified by the NREL. The J−V curves were measured at 10 mV/s. J−V curves for all devices were measured by masking the active area with a metal mask of 0.094 cm2. AFM measurements were performed with an Asylum Research Cypher S AFM (Oxford Instruments). Images were collected with an uncoated silicon tip (300 kHz, 26 N/m, OPUS) in tapping mode. AFM images were analyzed with the free Gwyddion software.37 All images are shown with line-wise flattening to remove tilting effect of the substrate plane. Root-mean-square (RMS) values were determined by 20 × 20 μm2 images. Photoluminescence Study. Temperature-dependent time-resolved photoluminescence measurements were obtained by timecorrelated single photon counting (TCSPC). The sample was mounted in a liquid nitrogen cooled coldfinger cryostat (Janis) and evacuated to a pressure below 5 × 10−4 Torr. The sample was excited by a pulsed 532 nm wavelength laser (PicoQuant LDH-P-FA-530-B), and the laser power was adjusted by a neutral density filter wheel to obtain a