Strong Electron–Phonon Coupling and Self-Trapped Excitons in the

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Strong Electron−Phonon Coupling and Self-Trapped Excitons in the Defect Halide Perovskites A3M2I9 (A = Cs, Rb; M = Bi, Sb) Kyle M. McCall,†,‡ Constantinos C. Stoumpos,† Svetlana S. Kostina,‡ Mercouri G. Kanatzidis,† and Bruce W. Wessels*,‡ †

Department of Chemistry and ‡Department of Materials Science and Engineering, Northwestern University, Evanston, Illinois 60208, United States S Supporting Information *

ABSTRACT: The optical and electronic properties of Bridgman grown single crystals of the wide-bandgap semiconducting defect halide perovskites A3M2I9 (A = Cs, Rb; M = Bi, Sb) have been investigated. Intense Raman scattering was observed at room temperature for each compound, indicating high polarizability and strong electron−phonon coupling. Both low-temperature and room-temperature photoluminescence (PL) were measured for each compound. Cs3Sb2I9 and Rb3Sb2I9 have broad PL emission bands between 1.75 and 2.05 eV with peaks at 1.96 and 1.92 eV, respectively. The Cs3Bi2I9 PL spectra showed broad emission consisting of several overlapping bands in the 1.65−2.2 eV range. Evidence of strong electron−phonon coupling comparable to that of the alkali halides was observed in phonon broadening of the PL emission. Effective phonon energies obtained from temperature-dependent PL measurements were in agreement with the Raman peak energies. A model is proposed whereby electron−phonon interactions in Cs3Sb2I9, Rb3Sb2I9, and Cs3Bi2I9 induce small polarons, resulting in trapping of excitons by the lattice. The recombination of these self-trapped excitons is responsible for the broad PL emission. Rb3Bi2I9, Rb3Sb2I9, and Cs3Bi2I9 exhibit high resistivity and photoconductivity response under laser photoexcitation, indicating that these compounds possess potential as semiconductor hard radiation detector materials.

1. INTRODUCTION Heavy metal−halide semiconductors have long been studied for their optical and electronic properties.1−4 However, interest in halide perovskites intensified in 2009 with the discovery that CH3NH3PbI3 was an efficient, low-cost light absorber for solar cells,5 and since then devices based on these materials have achieved efficiencies as high as 22%.6 Subsequently, a number of organic−inorganic halide perovskite materials have shown potential applications in hard radiation detection,7 LEDs,8,9 lasing,9,10 and photocatalysis.11 These successes have motivated studies of related hybrid materials such as the two-dimensional (2D) organic−inorganic lead iodide analogues for increased stability.12 However, the application of these materials has been limited by the presence of toxic lead13 and stability issues14 related to the organic cation. To bypass the stability issue, we and other groups investigated the inorganic lead halide perovskites CsPbX3 (X = I, Br, Cl), which have shown promise for solar cells15−17 and radiation detectors.18−20 However, the resulting solar cells have not achieved the same performance as the organic−inorganic perovskites and still contain lead. Thus, investigators have turned to the so-called defect perovskites21 as an avenue for eliminating toxic lead and improving device performance.22−24 The halide perovskites share the formula AMX3 where A is a monovalent cation, M is a divalent metal, © 2017 American Chemical Society

and X is a halide anion coordinated to the metal, forming a framework of corner-sharing MX6 octahedra. Defect perovskites are derivative structures that possess the same corner-sharing MX6 octahedra as the AMX3 aristotype but form different structures due to cation deficiencies in order to satisfy charge balance restrictions.21 Tetravalent cations form A2M□X6 structures (□ represents a vacancy) with 1/2 occupancy of the M sites in the A2M2X6 perovskite formula, while trivalent cations form A3M2□X9 structures with 2/3 occupancy of the M sites of the A3M3X9 perovskite formula.25,26 In particular, the A3M2X9 materials have recently attracted a great deal of interest as solar cell materials.23,24,27−31 This group of compounds includes perovskite derivatives containing bismuth iodide, which is a well-studied wide-bandgap semiconductor with potential applications in hard radiation detection32 and as the active layer in photovoltaics.33 Bismuth is an especially attractive candidate for replacing lead due to the isoelectronic nature of the Bi3+ ion to Pb2+ and its uniquely low toxicity relative to the neighboring heavy metals Tl, Hg, and Pb. Similarly, antimony is less toxic than lead while retaining the Received: March 23, 2017 Revised: April 10, 2017 Published: April 26, 2017 4129

DOI: 10.1021/acs.chemmater.7b01184 Chem. Mater. 2017, 29, 4129−4145

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Chemistry of Materials

°C oven, yielding 34 g (95%) of Cs3Sb2I9. The obtained powder was sealed in a quartz ampule (13 mm o.d.) under vacuum (10−3 mbar) and melted in a box furnace at 650 °C to form a dark red ingot. Cs3Bi2I9. Bi2O3 (20 mmol, 9.32 g) was dissolved in 80 mL of 57% aqueous HI to produce a bright red/orange solution in an exothermic reaction. CsI (60 mmol, 15.59 g) was dissolved in 20 mL of water, producing a colorless solution in an endothermic reaction. Addition of the CsI solution to the Bi2O3 solution led to spontaneous precipitation of a bright red solid in an essentially colorless supernatant liquid (pale yellow because of dissolved I2). The addition was done slowly under constant stirring to avoid local concentration of CsI. After cooling to room temperature, the solid was filtered through a fritted-dish funnel under vacuum. The powder was washed with 100 mL of 5:1 HI:H2O solution followed by 300 mL of methanol and dried overnight under vacuum. The solid was further dried in an oven at ∼70 °C in air, with a final yield of 38 g (97%). The obtained powder was sealed in a quartz ampule (13 mm o.d.) under vacuum (10−3 mbar) and melted in a box furnace at 650 °C to form a dark red ingot. SbI3. I2 (150 mmol, 38.07 g, Sigma-Aldrich, 3 N) and Sb shots (100 mmol, 12.18 g, Sigma Alrich, 5 N) were refluxed in 300 mL of toluene overnight, yielding a clear orange solution. The solution was filtered hot and left to cool slowly to room temperature, yielding red flakes. The resulting solid was filtered under vacuum, washed with cold toluene, and dried overnight. The obtained powder was sealed in a quartz ampule (13 mm o.d.) under vacuum (10−3 mbar) and sublimed in a tube furnace with the hot zone set to 350 °C, yielding 45 g (90%) of pure material. This process is required to remove residual Sb2O3 and SbOxIy impurities. BiI3. Bi(NO3)3·5H2O (100 mmol, 48.51 g) was dissolved in 300 mL of warm ultrapure (Millipore) water acidified by 40 mL of concentrated nitric acid, forming a colorless clear solution. In a different beaker, KI (300 mmol, 49.80 g) was dissolved in 50 mL of ultrapure (Millipore) water to form a colorless clear solution. Rapid addition of the KI solution to the Bi(NO3)3 solution immediately resulted in a fine black precipitate and a colorless supernatant solution. The BiI3 powder was filtered under vacuum, washed copiously with ultrapure (Millipore) water, and dried overnight. Subsequently, the BiI3 was washed with a dilute HI solution (1:20 concentrated HI:water) to ensure removal of the nitrate ions and then refiltered and rewashed as above. Yield was 55 g after purification (∼93%). Rb3Sb2I9. SbI3 (12.30 mmol, 6.18 g) and RbI powders (18.46 mmol, 3.92 g, Sigma-Aldrich, 3 N) were crushed together using mortar and pestle, sealed in a quartz ampule (13 mm o.d.) under vacuum and reacted in a box furnace at 650 °C. Rb3Bi2I9. BiI3 (31.37 mmol, 18.5 g) and RbI powders (47.07 mmol, 10 g, Sigma-Aldrich, 3N) were crushed together using a mortar and pestle, sealed in a quartz ampule (13 mm o.d.) under vacuum, and reacted in a box furnace at 650 °C. Crystal Growth. Crystal growth was performed using the powdered raw materials prepared as described above. In a typical procedure 6−20 g of material was sealed in quartz ampules (13 mm o.d.) under vacuum (10−3 mbar). Ingots were grown in a 2-zone Bridgman furnace with the hot zone at 800 °C and the cold zone at 400 °C. The translation speed was 2 mm/h. The resulting ingots were cleaved parallel to the layer plane (parallel to the growth direction for Cs3M2I9, approximately at a 30° angle for Rb3M2I9), coinciding with the van der Waals gap between the layers. The resulting thin wafers (0.2−2 mm) had mirror-like surfaces that required no further processing. Powder X-ray Diffraction. Samples from the as-grown ingots were crushed into fine powder for X-ray diffraction measurements. XRD spectra were measured using a Rigaku Miniflex600 diffractometer with a Cu Kα source (λ = 1.5406 Å) operating at 40 kV and 15 mA with a Kβ foil filter. Single-Crystal X-ray Diffraction. Single-crystal X-ray diffraction experiments were performed at 300 K using a STOE IPDS II diffractometer using Mo Kα radiation (λ = 0.71073 Å) operating at 50 kV and 40 mA. Integration and numerical absorption corrections were performed using the X-AREA, X-RED, and X-SHAPE programs. Using the Olex2 crystallography software package,41 the structure was solved

ns2 lone pair which is partially responsible for the exceptional performance of lead-based compounds.34 Thus, we chose to investigate optical and electrical properties of the bismuth iodide family of A3Bi2I9 compounds and the associated Sbanalogues A3Sb2I9, where A = Cs, Rb is an alkali metal. Detailed single crystal structures have been reported for Cs3Bi2I9,35 Cs3Sb2I9,36 and Rb3Bi2I9,24,37 and very recently for Rb3Sb2I9.38 Optical measurements indicate these compounds are semiconductors with wide band gaps in the 1.89−2.06 eV range and accordingly they exhibit high resistivities (1010−1011 Ω·cm).25,39 The compound Cs3Bi2I9 (Eg = 2.06 eV) does not form in a perovskite-derivative structure, as the BiI6 octahedra share faces to form discrete [Bi2I9]3− molecules, resulting in a zero-dimensional (0D) molecular salt crystal structure (space group P63/mmc). Cs3Sb2I9 can form either a hexagonal [Sb2I9]3− molecular dimer phase or a 2D layered structure depending on the synthesis method. Crystals grown from a concentrated HI solution form the dimer phase isostructural to Cs3Bi2I9, whereas crystals grown from the melt form a 2D layered perovskite derivative structure (space group P3̅m1).36 The single crystals used in the present study were grown by the Bridgman method; therefore, only the layered structure of Cs3Sb2I9 is considered (Eg = 1.89 eV).40,36 The compounds Rb3Bi2I9 (Eg = 1.93 eV) and Rb3Sb2I9 (Eg = 2.03 eV) are isostructural, with both forming distorted monoclinic layers (space group P21/n) where the MI6 octahedra adopt an irregular coordination approximating a rhombohedral elongation.24 Despite the importance of the related Pb-based perovskites, the fundamental properties beyond band gaps and resistivities are not well understood in these materials. In this work, we describe the structural, optical, and electronic properties of these heavy metal iodide perovskite derivatives. Large single crystals of the A3M2I9 (A = Cs, Rb; M = Bi, Sb) materials were grown using the Bridgman method. The crystal structure of Rb3Sb2I9 has been determined by single crystal X-ray diffraction and is isostructural to Rb3Bi2I9. The structural relations of these compounds are used to interpret the strong phonon modes observed via Raman spectroscopy. In situ synchrotron X-ray diffraction over the melt temperature range demonstrated that Cs3Bi2I9 is incongruently melting. Comprehensive low-temperature photoluminescence (PL) measurements are coupled with Raman spectroscopy to demonstrate strong electron−phonon coupling. We propose a model whereby small polarons are induced by the electron− phonon interactions and trap charge carriers, resulting in selftrapped excitons. Phonon-assisted recombination of these selftrapped excitons is responsible for the broad PL emission bands seen in Cs3Bi2I9, Cs3Sb2I9, and Rb3Sb2I9. Photoconductivity measurements under laser illumination indicate a strong photoresponse in Cs3Bi2I9, Rb3Bi2I9, and Rb3Sb2I9, suggesting that these materials are potential candidates for semiconductor hard radiation detectors.

2. EXPERIMENTAL METHODS Synthesis. Cs3Sb2I9. Sb2O3 (20 mmol, 5.83 g) was dissolved in 80 mL of 57% aqueous HI to produce a bright red/orange solution in an exothermic reaction. Cs2CO3 (30 mmol, 9.77 g) was added to 20 mL of 57% aqueous HI to produce a colorless solution in an endothermic reaction. Addition of the CsI solution to the Sb2O3 solution leads to spontaneous precipitation of a bright orange solid in a pale yellow solution. After cooling to room temperature, the solid was removed by vacuum filtration through a fritted-dish funnel. The solid was washed with 100 mL of 5:1 HI:H2O solution followed by 300 mL of methanol and dried overnight under vacuum. The solid was further dried in a 70 4130

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Chemistry of Materials with the ShelXT42 structure solution program using Direct Methods and refined with the ShelXL43 refinement package using Least Squares minimization. Scanning Electron Microscopy and Energy-Dispersive X-ray Spectroscopy (SEM/EDX). SEM images and EDX analyses of the asgrown cleaved crystal surfaces were acquired using a Hitachi S3400NII scanning electron microscope equipped with an Oxford Instruments INCAx-act SDD EDS detector at an accelerating voltage of 25 kV. To avoid excessive sample charging, measurements were taken under a pressure of 40 Pa with the backscattering detector. In Situ Synchrotron X-ray Diffraction. High resolution synchrotron powder X-ray diffraction data were collected at temperatures ranging from 25 to 600 °C using beamline 11-BM at the Advanced Photon Source, Argonne National Laboratory, using an average wavelength of 0.413614 Å. Discrete detectors covering an angular range from −6 to 16° 2θ are scanned over a 28° 2θ range, with data points collected every 0.001° 2θ and scan speed of 0.01°/s. For further details on the sample stage see Stoumpos et al.18 Differential Thermal Analysis Measurements. Differential thermal analysis was performed on powder samples in an evacuated fused silica tube using a Shimadzu DTA-50 thermal analyzer, with alumina as a reference material. UV−vis Diffuse Reflectance Measurements. Optical diffuse reflectance measurements were performed at room temperature using a Shimadzu UV-3600 PC double-beam, double-monochromator spectrophotometer operating from 200 to 2500 nm. BaSO4 was used as a nonabsorbing reflectance reference. The generated reflectance data was transformed to absorbance via the Kubelka− Munk equation44 α/S = (1 − R)2/2R, where R is the reflectance and α and S are the absorption and scattering coefficients, respectively. The absorption edge was estimated by extrapolation of the linear regions. Confocal Raman and Room-Temperature Photoluminescence (PL) Measurements. Raman spectroscopy was conducted using a confocal Horiba LabRAM HR Evolution spectrometer with an excitation wavelength of 785 nm. To reduce laser damage, the laser intensity (100 mW, ∼1 270 000 mW/cm2) was lowered by a series of neutral density filters to 12 700 mW/cm2 for Rb3Bi2I9 and Cs3Bi2I9 or to 40 700 mW/cm2 for Cs3Sb2I9 and Rb3Sb2I9. Room-temperature PL was measured with the confocal Horiba LabRAM HR Evolution spectrometer using an above-bandgap excitation wavelength of 473 nm. To avoid damaging the samples with the laser, the laser power (25 mW, ∼318 000 mW/cm2) was reduced to 3180 mW/cm2 using neutral density filters. Due to the weak PL response at room temperature, further decreases of the laser power led to a loss of signal. At this power the background increases slightly as the measurement is taken from low to high energy, indicating some sample damage is occurring. Low-Temperature Photoluminescence Measurements. Excitation intensity-dependent and temperature-dependent photoluminescence (PL) were measured for each compound. The sample mounting procedure and measurement setup were similar to those described by us previously.45 A continuous wave diode laser (Excelsior One-405 nm, SpectraPhysics) with a constant power of 50 mW and intensity of 3472 mW/ cm2 at 405 nm was used as the excitation source for PL measurements. A set of calibrated neutral optical density filters was used to limit the laser intensity between 35 and 700 mW/cm2 for temperature- and power-dependent measurements. Temperature-dependent measurements were taken at lower intensity to ensure reproducibility. For the chosen laser powers, repeating the measurements yields the same signal, indicating that the samples are undamaged. In the case of Rb3Sb2I9, the power-dependent measurements were performed using an OBIS CW diode laser (OBIS 405 nm LX 100 mW, 15 625 mW/ cm2) as the excitation source. The intensity was adjusted digitally from 78 to 3906 mW/cm2. Electrical Properties Measurements. Devices for electrical measurements were prepared from cleaved samples using several contact materials (Au, Pd, carbon). All contacts were in parallel plate configuration, either parallel or perpendicular to the cleaved surface of the layers. Gold or palladium contacts (50 nm) were deposited

through a shadow mask using an e-beam evaporator (Edwards Auto 306). Cu wires (0.025 mm) were bonded to the sample contacts using carbon paint (Ted Pella graphite), and these wires were attached to Cu tape on a glass substrate. Alternatively, the Cu wires were attached directly to the sample surface using carbon paint as a contact. The high resistivity of these compounds (1010−12 Ω·cm) permits the use of carbon paint contacts, as the contact resistivity (104−5 Ω·cm) is well within experimental error for these measurements. Resistivity and photoconductivity were measured in a setup using a voltage source (Stanford Research Systems PS325) and an ammeter (Keithley 2636A SourceMeter). For photoconductivity measurements, a laser (OBIS 405 nm LX, 15,625 mW/cm2) was used as the excitation source. The beam was operated at 7813 mW/cm2 during the measurement.

3. RESULTS AND DISCUSSION 3.1. Crystal Structure and Thermal Characterization. Single-Crystal XRD for Rb3Sb2I9. Single-crystal X-ray diffraction was conducted on Rb3Sb2I9 at room temperature, and the structure solution parameters can be found in Tables S1 through S4 (Supporting Information). The crystal structure of Rb3Sb2I9 is monoclinic in the space group P21/n (unit cell: a = 14.591(3) Å, b = 8.1879(16) Å, c = 20.584(4) Å, β = 90.36(3)°, V = 2459.1(8) Å3, Z = 4) and is comprised of disordered layers of corner-sharing SbI6 octahedra with Rb cations between them (Figure 1a). This is in good agreement with recent work by Chang et al.38

Figure 1. Rb3Sb2I9 structure as viewed down b axis (a) and down c axis (b); the unit cell is outlined, and SbI6 octahedra are shaded for clarity.

The compound adopts a distorted variant of the layered defect perovskite structure, where every third layer of the AMX3 structure is missing due to the 2/3 Bi occupancy of the M site. The SbI6 octahedra are asymmetric, as each octahedron shares three I atoms with three neighboring octahedra while the other three terminate on the layer surface. The shared iodine atoms have larger Sb−I distances (in the range of 3.15−3.18 Å) than the terminal iodine atoms, which have smaller distances (2.84− 2.87 Å) nearer to the sum of the covalent radii of Sb and I. This gives rise to the honeycomb motif formed by the SbI6 stacks when viewed down the c axis (Figure 1b). Structure Overview in A3M2I9 (A = Rb, Cs; M = Sb, Bi). Powder X-ray diffraction (Figure S1) and SEM-EDS (Figure S2, Table S5) analysis both confirm that the compounds are phase-pure, with no significant deviation from the A3M2I9 stoichiometry observed. Similar to the aristotypical AMX3 perovskite structure, the A3M2I9 crystal structures are dictated by the behavior of the MI6 octahedra. In the case of Cs3Bi2I9, the BiI6 octahedra share faces to form [Bi2I9]3− anions, resulting in a zero-dimensional (0D) molecular salt crystal structure (space group P63/mmc). This structure possesses hexagonal symmetry, as evidenced by the hexagonal channels filled by Cs cations (Figure 2a), but sacrifices the corner-sharing octahedra 4131

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Figure 2. Crystal structure comparison of Cs3Bi2I9 (left), Cs3Sb2I9 (middle), and Rb3Sb2I9 (right) as viewed down the c-axis (a, b, c) and along the baxis (d, e, f).

examine the evolution of the structure on heating (Figure 3). The temperature was raised from 25 to 600 °C, and X-ray

characteristic of perovskite-related structures in favor of disparate [Bi2I9]3− anions (Figure 2d). Despite its apparent structural simplicity, Cs3Bi2I9 exhibits unusual ferroelastic phase transitions and optical properties which are hard to reconcile with its molecular nature.40,46−48 The crystal structures of Cs3Sb2I9, Rb3Sb2I9, and Rb3Bi2I9 are 2D layered derivatives of the perovskite structure possessing the prototypical corner-sharing MI6 octahedra. Melt-grown Cs3Sb2I9 forms a 2D defect perovskite structure (space group P3m ̅ 1) consisting of infinite, hollow layers of corner-sharing [SbI6]3− octahedra. The [SbI6]3− octahedra are rhombohedrally elongated along the crystallographic c-axis (perpendicular to the layer plane), resulting in three short and three long Sb−I bonds which reduce the structure to trigonal symmetry. The layers form hexagonal channels (Figure 2b), bearing a strong resemblance to the parent SbI3 (BiI3-type) crystal structure.49 Contrary to the structure of SbI3, the structural voids in Cs3Sb2I9 are occupied by Cs cations. The cations can occupy two unique positions in the crystal lattice: either in the center of the hexagonal voids or at the edge of the layers bilateral to the Sb atom plane. This configuration relates to the 3D perovskite by cleavage through the (111) crystallographic plane of the cubic perovskite aristotype. In contrast, Rb3Bi2I9 and Rb3Sb2I9 adopt a reduced-symmetry monoclinic (space group P21/n) version of the Cs3Sb2I9 structure. The crystal structure consists of distorted monoclinic layers in which the MI6 octahedra adopt an irregular coordination approximating a rhombohedral elongation. Channels of Rb atoms still form in the voids between MI6 octahedra, but they lie along irregular directions with respect to the layers (Figure 2c). Similarly, the regular I−Sb−I bond angles of Cs3Sb2I9 (Figure 2e) are disrupted, resulting in offset iodine atoms when viewed along the b-axis (Figure 2f). For a thorough analysis of the symmetry relation of these structures to the perovskite aristotype we refer the reader to Chang et al.38 In Situ Synchrotron PXRD Results for Cs3Bi2I9. In situ powder XRD was conducted on Cs3Bi2I9 in a synchrotron beamline (Argonne National Lab 11-BM, λ = 0.413614 Å) to

Figure 3. Synchrotron PXRD of a polycrystalline Cs3Bi2I9 sample on heating from 27 to 600 °C.

diffraction patterns were taken every 6−10 °C. Cs3Bi2I9 Bragg peaks remained from 25 °C until decomposition at 502 °C, whereupon the BiI3 (mp 408 °C) begins to dissociate and after which the only remaining diffraction peaks are from the highermelting CsI (mp = 620 °C). Rietveld refinement was conducted using the program JANA2006 to solve the structure at each temperature,50 allowing the extraction of the thermal expansion parameters (Figure S3, Supporting Information). The a-axis, c-axis, and unit cell volume expansion parameters obtained for Cs3Bi2I9 are αa = 4.79 × 10−5 K−1, αc = 4.83 × 10−5 K−1, and αv = 1.51 × 10−4 K−1, respectively. However, each parameter shows a discontinuity near 410 °C, indicating a phase change. This is likely the dissociation of Cs3Bi2I9 into CsI and BiI3, as the melting and boiling points of BiI3 are 409 and 542 °C, respectively, while those of the ionic CsI are much higher at 632 and 1280 °C, respectively.51 Differential thermal analysis (Figure S4, Supporting Information) provides further evidence for BiI3 dissociation and CsI formation. There is a transition on 4132

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Figure 4. Raman spectra for the A3M2I9 compounds at 300 K using a 785 nm excitation laser.

heating which occurs near the boiling point of BiI3 at 542 °C, well before the final melt at 610 °C that nearly coincides with the melting point of CsI (632 °C). Cs3Bi2I9 contains the same BiI6 octahedra as BiI3, so it is not surprising that BiI3 dissociates from the Cs3Bi2I9 structure at intermediate temperatures while cesium and iodine atoms bond and cause an incongruent melt. Similar results are expected for the other three A3M2I9 compounds, as the low boiling points of SbI3 (bp 400 °C) and BiI3 cause them to dissociate before the alkali salts RbI (mp 656 °C) and CsI reach their melting points.51 This increases the difficulty of growing single crystals and precludes the use of horizontal zone refining and traveling heater methods due to the tendency for phase separation. However, crystallization in the vertical Bridgman configuration is possible because the crystallization point is separated by the melt from the BiI3 (or SbI3) vapor at the top of the tube, so the stoichiometry is maintained at the solid−liquid interface during growth. 3.2. Optical and Electronic Properties. Raman Spectroscopy Results. Raman spectra for the A3M2I9 compounds at room temperature using 785 nm laser excitation reveal structural similarities between the three layered compounds (Cs3Sb2I9, Rb3Sb2I9, and Rb3Bi2I9), with significant differences for Cs3Bi2I9 (Figure 4). Each compound was examined through 2000 cm−1, and no peaks were observed above 200 cm−1. The layered compounds have two major peaks near 150 cm−1 and several overlapping peaks below 85 cm−1. Cs3Bi2I9 shows the same two major peaks above 120 cm−1, along with several strong modes between 120 and 80 cm−1. This difference is likely due to the molecular salt structure of Cs3Bi2I9, as additional Bi−I modes are allowed by the [Bi2I9]3− anions compared to the more uniform MI6 octahedra of the layered compounds. The major peaks of Cs3Bi2I9 are summarized in Table 1, and those of the layered A3M2I9 compounds are listed in Table 2. The Raman spectra of the Bi-containing compounds are in good agreement with previous reports by Vakulenko et al.52 The major peaks for Rb3Bi2I9 are near 145 and 130 cm−1, and the spectrum of Cs3Bi2I9 shows a series of 9 peaks, with the strongest peaks near 150, 105, and 60 cm−1. A detailed analysis of the vibrational modes of the Cs3Bi2I9 space group (P63/mmc) indicates that they derive mainly from the [Bi2I9]3− anion.52 The crystal structure consists of molecular [Bi2I9]3− anions bound together by three Cs+ cations

Table 1. Raman Peaks and Phonon Energies of Cs3Bi2I9 Cs3Bi2I9 vibration

symmetry

v1

A′1

description

terminal Bi−I symmetric stretch E″ terminal Bi−I asymmetric v15 stretch v10 E′ terminal Bi−I asymmetric stretch v2 A11 bridge Bi−I symmetric stretch v7 E″ bridge Bi−I asymmetric stretch Bi−I bending mode not assigned: Cs-[Bi2I9] stretching or other bending mode

freq (cm−1)

energy (meV)

146.6

18.18

127.2

15.77

119.8

14.85

104.5

12.96

90.0

11.16

57.5 48.3

7.13 5.99

per unit. Thus, the dominant vibrations in this crystal arise from the strongly bound [Bi2I9]3− unit and weaker modes from the ionic interactions of that unit with the bridging Cs+ cations. In the [Bi2I9]3− unit, the major vibrational modes are the Bi− I stretching modes associated with the six terminal I atoms and the three bridging I atoms which connect the two Bi atoms together. All notation used to denote Raman modes is consistent with that of Nakamoto.53 There are three Ramanactive stretching modes associated with each set of bonds (the symmetric stretch A′1 and two antisymmetric stretches E″, E′), so there are six distinct stretching modes associated with the [Bi2I9]3− units.54 The force constants for terminal Bi−I bonds are stronger than the bridging Bi−I bonds because the bridging I atoms share their bonds with two Bi atoms. Thus, the higherenergy peaks are attributed to the stretching modes of the terminal bonds. Of these modes, the symmetrical stretching modes have higher degeneracy and thus relatively higher intensity than the antisymmetric modes. Thus, the major peak at 146.6 cm−1 is assigned to the symmetric terminal Bi−I stretch A′1(v1), and its antisymmetric partners are E″(v15) at 127.2 cm−1 and E′(v10) at 119.8 cm−1. The major peak at 104.5 cm−1 is the bridge Bi−I symmetric stretch A′1(v2), and its antisymmetric partner E″(v7) is the 90.0 cm−1 peak. These frequencies are in good agreement with those previously reported,55 as well as those of related compounds.53,56 4133

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Chemistry of Materials Table 2. Raman Peaks and Phonon Energies of 2D Layered A3M2I9 Compounds symmetry and description vibration

*M = Sb, Bi

v1 A1g: M−I sym. stretch v2 Eg: M−I asym. stretch symmetry-breaking M−I Stretch v5 F2g: M−I bending mode not assigned: M−I bending or lattice deformation mode

Cs3Sb2I9

Rb3Sb2I9

Rb3Bi2I9

freq (cm−1)

energy (meV)

freq (cm−1)

energy (meV)

freq (cm−1)

energy (meV)

164.8 146.8

20.43 18.20

17.39 15.71 14.72

10.40 8.38 6.82

20.25 18.29 16.99 16.25 10.19 7.36 6.31

140.3 126.7 118.7

83.9 67.6 55.0

163.3 147.5 137.0 131.1 82.2 59.4 50.9

73.4 64.0 56.7 50.0

9.10 7.93 7.03 6.20

interactions. This low intensity stands in stark contrast to Cs3Bi2I9, which possesses modes of high intensity in this region. The low intensity of the bending modes in Cs3Sb2I9 is due to the constraints imposed by the covalent connections between the octahedra. The directional bonds composing the lattice forbid many of the bending and rotational modes associated with independent MI6 octahedra, and accordingly the lowenergy modes are much weaker than the high-energy stretching modes. In the case of Rb3Sb2I9 and Rb3Bi2I9, the dominant modes are similar to those of Cs3Sb2I9 because the same 2D layers of corner-connected MI6 octahedra are the motifs of the monoclinic structure. Accordingly, the Raman spectra revealed two major high-energy peaks at 140.3 and 126.7 cm−1 for Rb3Bi2I9 and at 163.3 and 147.5 cm−1 for Rb3Sb2I9. These peaks again correspond to the symmetric A1(v1) M−I stretching and E1(v2) antisymmetric Sb−I stretching modes, respectively. In addition, the F2g(v5) bending mode can be observed at 73.4 cm−1 in Rb3Bi2I9 and at 82.2 cm−1 in Rb3Sb2I9. However, it was expected that additional modes would be allowed by the broken symmetries of the monoclinic structure relative to the 2D Cs3Sb2I9 structure. Accordingly, the spectra contain several extra peaks and shoulders not seen in the spectra of Cs3Sb2I9, namely, at 118.7 cm−1 in Rb3Bi2I9 and at 137.0 and 131.1 cm−1 in Rb3Sb2I9. There are also several strongly overlapping lowenergy peaks in the spectra which are largely caused by the various bending modes. These modes are again of low intensity due to the constraints of the covalent connections between the MI6 octahedra comprising the lattice, similar to the case of Cs3Sb2I9. These peak assignments are summarized with labeled Raman spectra in Figure 5. Beyond revealing the structural details, intense Raman scattering at room temperature is an indicator of strong coupling between the excited electrons and Raman active vibrations of the lattice. Raman spectroscopy detects dipole−dipole polarizability, so strong scattering interactions reflect the lattice’s proclivity to local fluctuations. This propensity to form local charges then plays a role in the charge transport, as the free charge carriers will cause the lattice to deform. These displacements travel with the carriers, forming polarons that increase the effective mass of charge carriers.61 Note that this behavior is distinct from that recently reported for CsPbBr3, in which local polar distortions cause a strong central peak in the Raman spectra at room temperature.62 No such peak is seen here, indicating that these compounds are not prone to such fluctuations, possibly due to their lower symmetry. To further probe these interactions and get a complete picture of the optoelectronic properties of the A3M2I9 compounds, the photoluminescence of each compound was studied. As some of these measurements were conducted at

Interestingly, the E′(v11) peak appears absent for the bridging atoms, but the E″(v7) peak is wider than that of the terminal E′ and E″, possibly indicating that both peaks are present but overlap strongly. The lower energy peaks are associated with various bending modes and ionic interactions between the [Bi2I9]3− units and the surrounding Cs+ atoms and cannot be assigned without additional polarization experiments and support from calculations. However, the peak centered at 57.5 cm−1 is unusually strong relative to the higher energy modes. This is a result of the molecular salt structure of Cs3Bi2I9 in which the [Bi2I9]3− anions are not bound directionally by the bridging Cs+ atoms, giving each unit the ability to twist and rotate. This freedom permits the bending modes to occur as often as the stretching modes and is a key difference between 0D compounds and compounds with a covalently bonded lattice which constrain such modes. This peak is likely due to the bending modes of the Bi−I bonds, as the rotational symmetry and Bi−I coupling are sufficiently strong that rotational bending of iodine atoms around the major Bi−Bi axis is of the same order of magnitude as the symmetrical stretching. The Raman spectra of the layered compounds are quite different, with two major peaks above 100 cm−1 and several overlapping peaks of lower energy and much weaker intensity. The spectrum of Cs3Sb2I9 was analyzed using previous studies of the isostructural CH3NH3Bi2Br9 and Cs3Bi2Br9 as a guide.57−59 The point group symmetry of the space group P3̅m1 is D33d, and the relations between the subgroups of site symmetry and the positions of atoms are well-known.57,60 The structure contains six iodine anions as a part of the SbI6 octahedra and three corner-shared I atoms that act as bridges between neighboring octahedra. There are 39 optically active modes, nine of which are Raman-active: the symmetric stretching (4A1g) and asymmetric stretching of bonds (Eg). Considering only the SbI6 octahedra, the intramolecular Raman spectrum consists of five bands (2A1g + 3Eg). Two of these vibrations include Sb atoms, the A1g(v1) symmetric Sb−I stretch, and the antisymmetric Eg(v2) Sb−I planar contraction with stretching in the perpendicular plane.58 These can be correlated with the two major peaks, as the highest-energy modes arise from the Sb−I interactions. Of these, the higher intensity mode in similar systems has been assigned to the A1g(v1) mode due to its higher symmetry.54,58 Therefore, the 164.8 cm−1 band in Cs3Sb2I9 is tentatively attributed to the A1g(v1) symmetric Sb−I stretch and the 146.8 band to the antisymmetric Eg(v2) Sb−I planar contraction. The peak at 83.9 cm−1 likely belongs to the F2g(v5) Sb−I bending mode, which is the other major Raman active mode.53 Below 80 cm−1 lie several peaks of relatively low intensity which strongly overlap, and these are due to various bending modes and Sb−I−Sb 4134

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exciton would dictate that the transition lies below the band edge, implying that the band edge is actually above 2.6 eV. Furthermore, the low electronic dimensionality of the 0D Cs3Bi2I9 supports a higher band gap than the 2D compounds.65 In addition to the n = 1 peak, Machulin et al. observe a second, weaker peak near 2.3 eV in the reflectance spectrum at 4.2 K.64 This peak may be a second exciton with an even higher binding energy, or it may be an exciton related to the degenerate conduction bands calculated by DFT with spin−orbit correction.24,29 The DFT results show several conduction bands with ∼0.3 eV separation,24,29 all of which may contribute to the absorption. The apparent discrepancy between band gap and crystal color is due to the high exciton binding energy (279 meV at 5 K) measured by Machulin et al.64 This binding energy is similar to exciton binding energies of the ionic alkali halide compounds.66 Such high binding energies lead to large Stokes shifts between the absorption and emission spectra,67 which is why no photoluminescence is detected near 2.6 eV (Figure 7).

Figure 5. Waterfall plot with labeled Raman peaks for A3M2I9 compounds.

low temperatures, it should be noted that the shift in Raman mode energy from room temperature to low temperatures is small enough to allow comparison between the two. In particular, the difference is less than 10 cm−1 (1.2 meV) for Cs3Bi2Br9 and Rb3Sb2Br9,57 which are isostructural to the layered compounds Cs3Sb2I9, Rb3Bi2I9, and Rb3Sb2I9, and is within 5 cm−1 (0.6 meV) for Cs3Bi2I9.55 UV−Vis Reflectance Measurements. The optical bandgaps of A3M2I9 were determined with polycrystalline samples at room temperature using UV−vis optical diffuse reflectance spectroscopy (Figure 6). The measured band gaps are 1.89,

Figure 7. Confocal PL of A3M2I9 at 293 K using 473 nm laser excitation.

This exciton binding energy is high enough that the excitonic absorption peaks persist at room temperature where normally they would be absent. These excitonic peaks are responsible for the below-gap absorption, resulting in a significant difference between the optical absorption edge and the electronic band gap. These excitonic bands act as chromophores, resulting in the red hue of the crystal that appears incongruent with the high band gap. A previous report on the isostructural compound (CH3NH3)3Bi2I9 found an analogous absorption feature at 2.51 eV due to an exciton with a binding energy of over 300 meV, along with a band edge near 2.8 eV.68 In this case, the excitonic absorption also causes the crystals to appear red. The absorption spectrum of Cs3Sb2I9 has a small peak near 2.0 eV, which we attribute to an excitonic peak. Such features are not uncommon in the 2D perovskites, as Cs3Bi2Br969 (isostructural to Cs3Sb2I9) and the 2D hybrid Ruddlesden− Popper perovskites70 both exhibit a prominent excitonic absorption peak. In this case, the fundamental (electronic) band edge is the higher energy absorption edge and the lower energy edge is actually the tail of the excitonic absorption.70 This leads to an underestimated band gap if the excitonic edge is used, and it is especially difficult to determine the band gap in the case of lower exciton binding energies. Analogous absorption features have been observed in the 3D perovskites (CH3NH3)3PbX3 (X = I, Br), but the fundamental and excitonic emissions overlap completely due to the relatively low exciton binding energy.71 Similarly, an exciton with low binding energy is likely responsible for the peak in the absorption of Cs3Sb2I9 near 2.0 eV, in which case the true band

Figure 6. Optical absorption spectra of A3M2I9 powder samples at 293 K (derived from diffuse reflectance measurements).

1.93, 2.03, and 2.06 eV for Cs3Sb2I9, Rb3Bi2I9, Rb3Sb2I9, and Cs3Bi2I9, respectively. These results reflect the structural differences of these compounds, such that deviations from the ideal 2D layered structure of Cs3Sb2I9 to the monoclinic Rb3M2I9 and further to the molecular salt structure of Cs3Bi2I9 result in an increased bandgap. The slope of the absorption edges is similar for each compound, with no obvious trend as a function of composition. These values are in good agreement with previous reports with the exception of Cs3Bi2I9, which is reported to have a band gap value ranging from 1.80 eV25 or 1.87 eV63 to 2.857 eV at room temperature.64 Looking to the color of the crystals of the A3M2I9 compounds, Cs3Bi2I9 is the lightest red and Cs3Sb2I9 the darkest red, consistent with the highest and lowest band gaps, respectively. This is difficult to reconcile with the value of 2.857 eV for Cs3Bi2I9, which implies that the compound should be white instead of bright red. However, there is also a prominent peak in the absorption spectrum of Cs3Bi2I9 near 2.6 eV corresponding to the proposed n = 1 excitonic transition previously seen at 2.58 eV by Machulin et al. in reflection measurements on a single crystal.64 The classical definition of 4135

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Figure 8. (a) Broad Cs3Sb2I9 PL emission at 13 K, 5 mW (blue trace) modeled by a modified Gaussian peak (red). (b) Configuration coordinate model where phonons are emitted as the center relaxes from B to C (absorption) or from D to A (luminescence). Here n is the number of phonons involved in the transition, Eph is the phonon energy, and EZPL is the zero-phonon line energy. Nonradiative recombination can occur along the path B to C to X to A via phonon emission and absorption.

carriers to the lattice site (known as self-trapping) and broaden the optical spectra substantially.73 Such broad bands are commonly seen in the alkali halides74,75 and related compounds such as AgBr76 and the organic−inorganic halide perovskites.77,78 A recent report on organic−inorganic perovskites has also shown that the corrugated bilayer structure is more prone to structural deformation than the flat ⟨100⟩ perovskite layer, leading to broad PL.7982 To determine the origin of these emission bands, low-temperature PL was measured for each compound. Analogous to the PL measurements at 293 K, Rb3Bi2I9 showed no PL response in the range of 410 to 900 nm at 15 K. The lack of emission is presumably due to the presence of nonradiative recombination pathways such as native defects and grain boundaries, which can be favored by the indirect bandgap of Rb3Bi2I9.24 An indirect bandgap requires phonons to mediate radiative transitions near the band edge and reduces the probability of radiative recombination, allowing nonradiative recombination to dominate. Low-temperature PL measurements for Cs3Sb2I9, Rb3Sb2I9, and Cs3Bi2I9 revealed one or more broad peaks, but no sharp peaks characteristic of free exciton zero-phonon lines were observed. This is likely due to the temperature (∼15 K) and broadness of the peaks, as the reported intensity of the zerophonon line in Cs3Bi2Br9 (isostructural to Cs3Sb2I9) even at 4.2 K was quite weak in all samples and was not observed for some samples.80 Timmermans et al. estimated an activation energy less than 1.24 meV for quenching of the zero phonon emission. This activation energy is too low in value to attribute to exciton dissociation, as an exciton with lower binding energy than 1.2 meV would not be observed due to thermal energy. This indicates that exciton trapping by defects or impurity centers is responsible for the lack of free-exciton emission. Similarly, the absence of free-exciton peaks in the A3M2I9 compounds is likely due to carrier trapping which precludes direct recombination of free carriers. We begin this section by introducing a configurational coordinate model of phonon-assisted self-trapped exciton (STE) recombination, which depends strongly on the electron−phonon coupling. The low-temperature PL results indicate that this coupling is quite strong and demonstrate that STE recombination is responsible for the broad PL bands observed for the layered materials Cs3Sb2I9 and Rb3Sb2I9. The PL band of the molecular compound Cs3Bi2I9 is much more complex with multiple overlapping peaks which render fitting unreliable. However, the emission shares some similarities with

gap is above the 1.89 eV edge. This agrees with the 1.96 eV emission peak observed in room-temperature photoluminescence measurements (Figure 7), which should not lie above the band gap. A similar feature may be present for Rb3Bi2I9, but the excitonic nature of this peak cannot be confirmed because our Rb3Bi2I9 sample exhibits no photoluminescence at room temperature (Figure 7). No peak or shoulder was observed for Rb3Sb2I9. To better understand these features and check for radiative defect recombination, we performed photoluminescence measurements on these compounds. Photoluminescence of A3M2I9 Compounds. We measured the photoluminescence (PL) properties of A3M2I9 single crystals both at room temperature and at low temperatures (13−15 K). PL measurements at room temperature were collected in the confocal configuration under 473 nm (2.62 eV) excitation laser using 0.25 mW laser power (50 μm spot size). A broad, weak emission in the 1.58−2.2 eV range was observed for Cs3Bi2I9, Cs3Sb2I9, and Rb3Sb2I9 single crystals, while Rb3Bi2I9 showed no response (Figure 7). Each compound exhibited signs of laser damage due to the small spot size, necessitating low laser powers and short exposure times for a reliable measurement. Even under these conditions, the background signal increases slightly as the measurement continues, e.g., the Rb3Bi2I9 black trace in Figure 7, which has no PL response but shows a continuous rise in background as the spectrum is taken from low to high energy. The room-temperature PL of each compound (besides Rb3Bi2I9, which showed no emission) in Figure 7 consists of very broad peaks with full widths at half-maximum (fwhm’s) ranging from 315 to 403 meV. The PL spectrum of Cs3Sb2I9 shows a broad, skewed emission band with a peak at 1.96 eV, the same as the skewed band-edge emission observed previously in Cs3Sb2I9 thin films at room temperature.23 The PL spectra of Rb3Sb2I9 and Cs3Bi2I9 each show an asymmetric, broad emission band with a maximum at 1.92 and 1.93 eV, respectively. The electron−phonon coupling indicated by the intense Raman scattering at room temperature affects the observed luminescence since these interactions broaden the emission bands via scattering.72 The emission and absorption of phonons can alter the energy of transitions and broaden otherwise sharp transitions into wide bands, and this is heavily dependent on the strength of the electron−phonon coupling. The electron− phonon interactions cause local distortions of the lattice, known as polarons, to couple to carriers and increase their effective mass. In the case of strong coupling, the polarons bind the 4136

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Gaussian distribution, which is the convolution of an exponential with a Gaussian:85

that of Cs3Sb2I9 and Rb3Sb2I9, suggesting that recombination of STEs with different binding sites may be the source of these broad peaks. Low-Temperature Photoluminescence of Cs3Sb2I9. The PL spectrum of Cs3Sb2I9 at 13 K shows a broad near band-edge emission at 1.92 eV (Figure 8a), skewed toward lower energy analogous to the room-temperature spectrum (Figure 7). While asymmetry in the line shape at low temperatures can reflect skewness in the impurity band responsible for the emission caused by fluctuations of the local potentials in the crystal,81 here the broadness and skewed shape of the band is attributed to strong electron−phonon coupling which leads to charge localization and emission of phonons that alter the energy of the emitted photon. Such interactions can be modeled using a configuration coordinate diagram (Figure 8b). Here the lower band is the valence band and the upper band is the self-trapped exciton (STE) band, and the conduction band (not pictured) lies above the STE band. The absorption transition upon excitation is from A to B and STE recombination emission is from C to D. The absorption and emission edges begin at point C, as this is lowest point in the excited state energy diagram. This energy EZPL = EC − EA is known as the zero-phonon line (ZPL) of the emission band, as it is the energy absorbed or emitted to go directly via a radiative transition from EC to EA. Higher-energy absorption bands (with energies En = EZPL + Eph) are allowed with the subsequent emission of phonons as the electron relaxes from points B to C. These peaks have equidistant spacing given by the phonon energy Eph, the nth peak involving n phonons has energy En = EZPL + nEph. Similarly, the luminescence spectrum consists of several mirrored radiative transitions from C to D accompanied by n phonons emitted as the center relaxes from point D to point A, each with energy En = EZPL − nEph. This model can also explain low-temperature thermal quenching of luminescence, as efficient quenching can result from nonradiative recombination where the bands cross (point X) provided phonons can make the pathway B to C to X to A viable.83 The resulting luminescence spectrum is the convolution of multiple Gaussians (one for each transition, offset by Eph), the intensities of which follow a Poisson distribution: Int(n) =

e −S S n n!

I (E ) =

⎡ ⎛ ⎞2 E − Emax ⎤⎥ A 1 ω exp⎢ ⎜ ⎟ − ⎢⎣ 2 ⎝ μ ⎠ ⎥⎦ μ μ

z

∫−∞

1 −x 2 /2 e dx 2π (2)

where z is equal to z=

E − Emax ω − ω μ

(2a)

ω is the peak width of the Gaussian component, E and Emax are the photon and peak energies, respectively, and A is the amplitude of the emission. The modification/skewness factor μ determines the direction and amplitude of the asymmetry in the line shape. This expression has been previously used in analysis of the skewed photoluminescence peak of Ge.86 Excellent agreement between the experimental spectrum of Cs3Sb2I9 (sample KM20B-PL) and calculated I(E) was observed (Figure 8a). The dependence of PL emission intensity on photoexcitation intensity was measured at 13 K to determine the recombination kinetics (Figure 9a). An increase in the excitation intensity from

Figure 9. (a) Dependence of Cs3Sb2I9 PL emission on excitation intensity at 13 K. (b) Plot of log(I) vs log(L) for sample KM20B-PL.

1 to 50 mW was accompanied by both an enhancement of emission intensity and a small 0.01 eV redshift in the peak energy. The dependence of the PL intensity I on the laser power L is given by the I ∼ Lκ power law dependence, described by the exponent κ.87 Excitonic (free or bound) emission increases superlinearly with excitation intensity, where the power law coefficient κ is in the 1 < κ < 2 range. For DAP recombination or free-to-bound (FB) transitions the PL intensity increases sublinearly, so the exponent is less than 1. A plot of log(I) vs log(L) is linear with a slope κ of 0.89 ± 0.03, indicating the recombination process is either DAP or free-tobound recombination (Figure 9b). However, the influence of strong electron−phonon coupling can render this recombination model unreliable when the excitons are heavily localized by small polarons in the lattice. Localization imposes severe restrictions on recombination, limiting the rate at which it can occur in the same way that free-to-bound and DAP transitions can only occur at specific locations in the lattice. Such interactions violate the rate equations used to derive the power law model,87 which does not account for localization effects in the case of excitonic emission. Thus, excitonic emission cannot be excluded based on the excitation intensity-dependence data alone. To determine the quenching behavior of the emission, temperature-dependent Cs3Sb2I9 PL spectra were measured using an excitation intensity of 5 mW (Figure 10a). The peak shifts irregularly with temperature, redshifting from 13 to 31 K

(1)

where S is the Huang−Rhys parameter which describes the magnitude of the electron−phonon coupling.72 The superposition of these peaks is known as the phonon wing due to the distinctive shape of the spectrum; see, for example, the green emission of ZnO.84 In this case, the electron−phonon interaction is strong enough that individual phonon replicas are not distinguishable due to the influence of phonon broadening of each Gaussian peak. S represents the mean number of phonons emitted as the excited state relaxes along the path B to C or D to A. The Poisson distribution approaches a Gaussian distribution for sufficiently large values of S (≥100), and below this threshold the distribution resembles a Gaussian peak skewed to higher values of n. Here, the distribution of luminescence energy is skewed to lower energies because high n results in a lower energy En = EZPL − nEph for the transition accompanied by emission of n phonons. This distribution matches the Cs3Sb2I9 PL, which is also skewed to low energy (Figure 8a). To account for this skewness, the measured emission was analyzed using a modified 4137

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Figure 10. (a) Temperature-dependent PL spectra of Cs3Sb2I9 using 5 mW laser power. (b) Full width at half-maximum (fwhm) vs temperature with solid line fit by eq 3. (c) Integrated PL intensity vs 1000/T where the solid line is the least-squares analysis of the data to eq 4.

and then blueshifting to 70 K until the emission quenches at 105 K. The temperature dependence of the full width at halfmaximum [fwhm] of the emission peak, denoted w(T), can be modeled using the theory of Toyozawa.88 This theory models the emission broadening resulting from electron−phonon interactions using a configuration coordinate model where the width of the peak w(T) depends on the average phonon energy:89 1/2 ⎡ ⎛ Eph ⎞⎤ w(T ) = 2.36 S Eph⎢coth⎜ ⎟⎥ ⎢⎣ ⎝ 2kBT ⎠⎥⎦

I (T ) =

I0 1 + ϕ1T

3/2

+ ϕ2T 3/2 exp( −Ei /kBT )

(4)

where I0, ϕ1, and ϕ2 are constants, kB is the Boltzmann constant, and Ei is the activation energy for thermal quenching. The Arrhenius plot was analyzed according to eq 4 (Figure 10c), yielding an activation energy of 14 ± 1 meV. The values of I0, ϕ1, and ϕ2 were (1.59 ± 0.05) × 10−5, 0.005 ± 0.001, and 0.7 ± 0.2, respectively. The temperature-dependent data agrees with the model well, but the lack of emission after 120 K suggests that the calculated activation energy corresponds to a nonradiative recombination process. If the activation energy corresponded to ionization of a donor or acceptor level in DAP recombination, a free-to-bound transition should be observed at higher temperatures as the lower state ionizes,95 but instead the emission quenches completely as the activation energy is approached. This low activation energy indicates that the nonradiative recombination pathway dominates at higher temperatures, as 14 meV corresponds to the thermal energy at 163 K, which is consistent with the quenching of the emission at 120 K. Furthermore, the observed redshift in the peak position with intensity is in contrast to the blueshift characteristic of DAP recombination,72,96 so the emission is not due to DAP recombination. The sublinear increase of the emission intensity with increasing excitation intensity indicates that the emission of Cs3Sb2I9 is due to either free-to-bound or DAP recombination, so it seems that the emission must be a free-to-bound transition. However, there is a peak in the UV−vis reflectance spectrum (Figure 6) similar to the excitonic transition observed in the perovskite halides, including the isostructural Cs3Bi2Br9, indicating that the emission should be excitonic.69 This contradiction arises from the strong electron−phonon interactions, which cause the excited carriers to become highly localized so excitons are immediately trapped by lattice deformations upon formation. The localization imposes severe restrictions on recombination, violating the rate equations used to derive the power law model.87 Thus, excitonic emission cannot be excluded based on the excitation intensity-dependent data alone and must be distinguished from free-to-bound transitions using other observations. Free-to-bound transitions are characterized by a donor or acceptor state responsible for binding a carrier, which at higher temperatures would be ionized and result in band-edge emission between free carriers. This is not consistent with our observation of thermal quenching of the emission which indicates that no such state exists, so the transition should not be from free-to-bound recombination. Furthermore, the excitonic transitions at room temperature suggest a high exciton binding energy that cannot be overcome via thermal activation, ruling out free exciton

(3)

Here S is the Huang−Rhys electron−phonon coupling parameter and Eph is the effective phonon energy. This model applies when there is sufficiently strong electron−phonon coupling; it has been applied to the emission of KCl90 and Cd1−xZnxTe.89 A plot of fwhm (eV) vs temperature (Figure 10b) was analyzed according to eq 3, yielding S = 42.7 ± 3.2 and a phonon energy of Eph = 7.0 ± 0.8 meV. The electron− phonon coupling is quite strong; S = 42.7 is comparable to that observed in the alkali halides, in particular NaCl (S = 42), and is significantly higher than the reported value for CsI (S = 12).91,92 This value is unexpected given the covalency of the Sb−I and Cs−I bonds, but it presumably arises from the layered structure of the material, which tolerates more local polarization than the 3D CsI structure. The layer interfaces are not bonded as in the 3D case. Thus, the SbI6 octahedra have more degrees of freedom since they are bonded to three other octahedra in two directions, as opposed to six bonds in all three directions of the 3D structure. This allows the 2D material to distort substantially in the presence of charge carriers, resulting in stronger electron−phonon coupling even with similar bonding behavior. The measured effective phonon energy of 7.0 meV agrees quite well with the low-energy mode of 6.82 meV (55 cm−1) observed in the Raman spectrum of Cs3Sb2I9 at room temperature. The higher energy modes are frozen out below 50 K, so it is not surprising that the recombination involves a lower-energy mode. The phonon broadening diverges from the fitted line around 55 K (Figure 10b), suggesting that a higher energy phonon is thermally activated and dominates the recombination. However, there is insufficient data to fit this higher energy phonon due to the phase transition at 86 K,93 which shifts the fwhm and after which the emission quickly quenches (105 K). The thermal quenching of the emission intensity with temperature was measured to analyze the recombination process. The Arrhenius plot of integrated PL intensity vs inverse temperature can be analyzed with the DAP equation, which can describe the activation energy of the relevant level or that of a competing nonradiative pathway:94 4138

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Figure 11. (a) Broad PL emission of Rb3Sb2I9 fit by a modified Gaussian peak (dashed blue line) with residual plot (red trace, enlarged 100× for clarity). (b) Dependence of PL emission on excitation intensity for Rb3Sb2I9 at 13 K. (c) Plot of log(I) vs log(L) for sample KM27-PL.

Figure 12. (a) Temperature-dependent PL spectra for Rb3Sb2I9 with 3 mW laser power. (b) fwhm vs temperature, with solid black line denoting the least-squares fit to eq 3 through 105 K and blue line fit to eq 2 offset by 120 K. (c) Arrhenius plot of integrated PL intensity vs 1000/temperature fit with an excitonic model (eq 5).

line broadening attributed to electron−phonon coupling.102 Furthermore, self-trapping is strongly dependent on the dimensionality of the system, with stronger self-trapping in lower-dimensional materials.77 STEs have been shown to cause broad emission covering the visible spectrum in the 2D hybrid halide perovskites (N-MEDA)[PbBr4]103 and α-(DMEN)PbBr4.82 As a corrugated 2D material, the structure of Cs3Sb2I9 provides ample opportunities for trapping as the charges scatter between MI6 octahedra in the bilayers. Thus, the broad band PL emission in Cs 3 Sb 2 I 9 is assigned to recombination involving self-trapped excitons and the emission of multiple phonons. Low-Temperature Photoluminescence of Rb3Sb2I9. PL spectra of Rb3Sb2I9 collected at 13 K for a single crystal reveal a broad, skewed peak at 1.96 eV, with what appears to be a superimposed set of weak phonon replicas (Figure 11a). This peak was successfully modeled using the exponentially modified Gaussian distribution applied to the PL of Cs3Sb2I9 (Figure 11a, blue dashed line). The residual plot (Figure 11a, red trace) shows the difference between the experimental data and the calculated modified Gaussian fitting, revealing several small peaks with a separation of 62 ± 3 meV. These resemble phonon replicas, and the energy spacing is a reasonable value for a phonon. However, 62 meV lies well outside the observed energy range (6−20 meV) for the optical phonons in the Raman spectrum at room temperature so it is unlikely that these peaks are phonon replicas. The intensities of these small peaks are only 1% of the amplitude of the modified Gaussian, so they will not affect the error significantly. The dependence of PL emission intensity on laser power for Rb3Sb2I9 (sample code KM27-PL) was measured at 13 K to probe the nature of the broad band (Figure 11b). An increase in the excitation intensity from 0.5 to 15 mW was accompanied by an increase of the emission intensity but no shift in the peak location. A plot of log(I) vs log(L) was linear with a slope κ of 1.08 ± 0.02 (Figure 11c), indicating an excitonic transition. The excitation intensity-dependence of the emission is consistent with both bound and free-exciton emission; however, the peak

emission. Thus, the observed emission must involve bound exciton recombination. Excitons can be bound by extrinsic defects, native defects, or intrinsic lattice interactions. One such example of the former is excitons bound to isoelectronic impurities, which are responsible for emission in other semiconductors such as GaP:N and AgBr.72,76,97,98 In AgBr, excitons bound to iodine impurity atoms produce a broad emission 106 times stronger than that of the free exciton emission, even though there are ∼107 times more Br atoms than I atoms. However, this process requires the presence of a substantial number of impurities, and there is no indication that such defects exist in Cs3Sb2I9. Our explanation for the PL in Cs3Sb2I9 is recombination between self-trapped excitons (STEs), where small polarons trap wouldbe free excitons.67 The strong electron−phonon interactions induce small polarons in the lattice that locally trap charge carriers, resulting in self-trapped excitons. Self-trapping of excitons often occurs in semiconductors with a large electron− phonon coupling.72 This STE recombination process matches the observed emission behavior in every way, and self-trapped excitons have been observed in many related materials. These PL emission peaks are strongly dependent on electron−phonon interactions, manifesting as narrow lines when the electron−phonon coupling is weak (nearly free excitons, i.e., with a low S parameter), and as broad bands when the coupling is strong (high S parameter).72 For Cs3Sb2I9, both the phonon broadening with a high Huang−Rhys parameter of S = 42.7 and intense Raman scattering at room temperature demonstrate the strength of the electron−phonon coupling. The configuration coordinate model can also explain the lowtemperature quenching of the PL, as efficient quenching can come from nonradiative recombination where the bands cross along the phonon-assisted pathway B to C to X to A (Figure 8b).83 This recombination process has been extensively studied in the alkali halides,75,99,100 and moreover STEs have been observed in both SbI3 and BiI3.101 The 3D hybrid organic− inorganic hybrid halide perovskites also exhibit PL emission 4139

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including radiative recombination of excitons bound to defects such as isoelectronic impurities as described above, but they require a substantial number of impurities to observe emission. In particular, the high intensity of the bound exciton emission supports the notion that the binding mechanism is intrinsic to the structure, as the free exciton peak is entirely absent. Both the intense Raman scattering at room temperature and the large Huang−Rhys factor calculated from the fwhm broadening of the PL emission by phonons at low temperature demonstrate strong exciton−phonon interactions in Rb3Sb2I9. Furthermore, the spectral shape and photoluminescence peak behavior are similar to those of Cs3Sb2I9, supporting the assignment of similar mechanisms to each emission. Thus, we assign the broad band PL in Rb3Sb2I9 to recombination of self-trapped excitons. Low-Temperature Photoluminescence of Cs3Bi2I9. In contrast to the two previous compounds, the low-temperature (13 K) PL spectrum of Cs3Bi2I9 shows several overlapping peaks in the 1.85−2.03 eV range, just below the roomtemperature absorption edge at 2.06 eV (Figure 13). The

is too broad for free exciton emission, which would be very narrow at low temperatures. Thus, the low-T PL emission in Rb3Sb2I9 is excitonic in nature and is due to recombination of bound excitons and multiphonon emission. To further characterize the emission behavior, temperature-dependent PL spectra were taken at an excitation intensity of 3 mW from 13 to 220 K (Figure 12a). The peak shifts irregularly with temperature, redshifting from 13 K to 80 K, saturating until 140 K, and then redshifting again until 220 K after which it can no longer be detected. The temperature dependence of the fwhm can again be modeled using the configuration coordinate model shown in Figure 8b. Analyzing the data in Figure 12b with eq 3 yielded a coupling parameter S1 = 50.4 ± 1.7 and an effective phonon energy of Eph1 = 7.1 ± 0.1 meV. This value corresponds to the low-energy phonon observed at 7.4 meV (59.4 cm−1) in the Raman spectrum of Rb3Sb2I9 at room temperature, as the higher-energy phonon modes are frozen out at low temperature. This behavior, however, holds only to 105 K (black line), after which another phonon dominates. Shifting the zero temperature of eq 3 by 120 K to reflect the thermally activated phonon, this phonon energy can be calculated (blue line, Figure 12b), yielding a coupling parameter S2 = 21.2 ± 0.9 meV and an effective phonon energy of Eph2 = 19.6 ± 0.1 meV for the higher energy phonon. This value likely corresponds to the dominant symmetrical stretching mode A1 (v1) at 20.25 meV (163.3 cm−1), accounting for the thermal shift of Raman peak energy from high to low temperatures. The significant change of the Huang−Rhys parameter S with temperature reflects the change in the effective phonon energy. The parameter S represents the average number of phonons of a single energy emitted as the charge center relaxes from points D to A (Figure 8b), so higher energy phonons reduce the number of phonons required to mediate the same transition. In any case the coupling is quite strong, S1 = 50.4 in particular matches that of KBr (S = 51) and is higher than the value for RbI (S = 40).91,92 This increase, as discussed previously, presumably results from the layered structure, which tolerates more distortion than the rigid 3D RbI lattice. At higher temperatures the coupling weakens to S2 = 21.2, closer to the value of RbBr (S = 25).91 To model the quenching of the PL intensity with increasing temperature, we use the equation of Bimberg et al. for the thermal quenching of exciton emission given by104 I (T ) =

I0 1 + C1 exp( −E1/kT ) + C2 exp( −E2 /kBT )

Figure 13. PL spectra of Cs3Bi2I9 measured as a function of (a) intensity, and (b) temperature.

spectrum is in excellent agreement with the broad emission observed from 1.63 to 2.23 eV at 293 K using the confocal setup (Figure 7). This emission was measured as a function of both excitation intensity and temperature (Figure 13a,b, respectively). The emission shows strong thermal quenching, with the signal vanishing above 105 K. The PL spectrum of Cs3Bi2I9 is comprised of at least four overlapping peaks. This strong overlap renders peak fitting unreliable, as the peaks do not shift significantly with temperature or excitation intensity and thus cannot be resolved. Multiple peak analyses were attempted, but without the precise number of peaks such analysis has limited utility and requires the aid of calculations to assign the different peak energies to specific defects or binding sites. A red emission peak in the related compound Cs3Bi2Br9 was attributed to the presence of oxide impurities from the Bi2O3 precursor,80 which may be one of the peaks seen here. Nevertheless, fitting the emission with a single-peak fit yielded a rough analysis of the overall emission behavior and is included in the Supporting Information. Briefly, this analysis indicates sublinear dependence of PL intensity on excitation intensity, suggesting either a donor−acceptor pair (DAP) recombination or free-to-bound transition. This is in contrast with the observations of excitonic absorption bands via UV−vis reflectance measurements (Figure 6), which indicate that the emission is excitonic. The change in emission peak width (Figure S5b) was attributed to phonon broadening with a phonon of energy Eph = 5.9 ± 0.4 meV. This energy may correspond to the strong bending mode phonon seen in the Raman spectrum at 7.1 meV (57.5 cm−1), since the error in the PL measurement is likely underestimated (see Supporting

(5)

where I0, C1, and C2 are constants and E1 and E2 are activation energies. The Arrhenius plot of integrated PL intensity vs inverse temperature was analyzed with eq 5 (Figure 12c), yielding activation energies of 130 ± 50 meV and 30 ± 3 meV. The values of I0, C1, and C2 were (1.10 ± 0.01) × 10−3, (1 ± 4) × 106, and 90 ± 40, respectively. The errors on C1 and E1 are high because the high-temperature luminescence quenches rapidly and the points cluster together. The observation of two activation energies presumably indicates that there are two nonradiative recombination mechanisms that dominate at different temperatures, together quenching the emission around 220 K. Similarly to the preceding case, the PL emission mechanism in Rb3Sb2I9 is likely to involve recombination of STEs, where small polarons in the lattice bind free excitons. Other possibilities exist which could account for this broad emission, 4140

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Chemistry of Materials Table 3. Photoluminescence Parameters for Cs3Bi2I9, Cs3Sb2I9, and Rb3Sb2I9 compound

absorption edge

peak position (293 K)

peak position (13−15 K)

calc. fwhm (at 0 K)

calc. Ephonon (meV)

observed (Raman) Eph (meV)

Huang−Rhys factor S

Cs3Bi2I9 Cs3Sb2I9 Rb3Sb2I9

2.06 eV 1.89 eV 2.03 eV

1.93 eV 1.96 eV 1.92 eV

1.95 eV 1.92 eV 1.96 eV

125 meV 107 meV 119 meV

5.9 7.0 7.1, 19.6

7.1 6.82 7.4, 20.3

79.5 42.7 50.4, 21.2

Table 4. Conductivity Results for A3M2I9 Compounds compound

sample ID

thickness (cm)

contact material

ρdark(+) (1010 Ω·cm)

σlight/σdark (h+)

ρdark(−) (1010 Ω·cm)

σlight/σdark (e−)

Rb3Sb2I9 Cs3Sb2I9 Cs3Bi2I9 Rb3Bi2I9

KM27-1 KM20B-1 KM8R KM20-1

0.033 0.037 0.094 0.067

Au C Pd C

8.29 52.1 972 31.8

9.39 1.03 10.50 2.95

8.70 52.6 903 32.9

6.15 1.02 9.29 4.00

candidates for radiation detection. Cs3Sb2I9 showed negligible response to laser illumination. The other three materials respond with a photocurrent 2−10× higher than the dark current, summarized in Table 4 as the photoconductivity ratio σlight/σdark. This response was consistent, with all samples studied having response to laser greater than 2× the dark current. While this photoresponse is relatively weak, it is comparable to single crystals of the 3D perovskites CsPbBr3106,107 and MAPbBr3,108 which show photocurrent values roughly an order of magnitude higher than those of the dark current. The lower values measured here could be due to lower mobility or charge carrier generation and are the subject of current investigation. Notably, one piece of the Cs3Bi2I9 ingot exhibited a photocurrent over 100× the dark current, indicating that these materials have higher potential that may be accessible through increased purity or improved syntheses. The consistent photoresponse indicates that self-trapping does not prohibit charge transport and that these compounds merit further study.

Information). The phonon broadening is an indication of strong electron−phonon coupling, with a calculated Huang− Rhys parameter S = 79.5 ± 3.8 which is presumably overestimated due to the presence of multiple peaks. As a 0D molecular structure, Cs3Bi2I9 is expected to have even more efficient self-trapping since charges can be easily bound to the charged Cs+ ions and [Bi2I9]3− units. The Cs3Bi3I9 PL shares many similarities to the broad emission of Cs3Sb2I9 and Rb3Sb2I9, the main difference being the observation of several overlapping peaks. In this case there may be multiple binding sites in the lattice for trapping the excitons, which would give rise to several broad peaks slightly offset by the difference in binding energy. These observations suggest that phononassisted recombination of multiple self-trapped excitons is responsible for the emission of Cs3Bi2I9. The PL properties of the three radiative compounds are summarized in Table 3 below. Resistivity and Photoconductivity. Charge transport was measured to determine how this charge trapping influenced the photoconductivity in the A3M2I9 compounds and to evaluate these materials as solar cell or detector materials. The resistivity and photoconductivity under 405 nm laser excitation were measured at room temperature using a two-probe method. The resistivities are quite high, ranging from 1010 to 1012 Ω·cm for each compound, with the highest and lowest resistivities corresponding to Cs3Bi2I9 and Cs3Sb2I9, respectively (Table 4). While these resistivities are high for solar cell materials, they fall within the desired range for hard radiation detector materials.105 A typical photoresponse to 405 nm laser excitation for Cs3Bi2I9 is shown in Figure 14, which compares the dark conductivity to the photoconductivity. The compounds Cs3Bi2I9, Rb3Bi2I9, and Rb3Sb2I9 are all photoconductive (Figure S6, Supporting Information), making them potential

4. CONCLUSION The optical and electronic properties of Bridgman grown large single crystals of the defect halide perovskites A3M2I9 (A = Cs, Rb; M = Bi, Sb) were examined. In situ powder X-ray diffraction provided evidence for incongruent melting in Cs3Bi2I9 with BiI3 dissociating at 460 °C. Intense Raman emission is observed at room temperature for each compound, indicating the high polarizability and strong electron−phonon coupling of these materials. Photoluminescence spectroscopy reveals broad red emission bands for Cs3Sb2I9, Rb3Sb2I9, and Cs3Bi2I9 below their respective absorption edges. A configuration coordinate model describes the multiphonon line shape and temperature-dependent emission quenching behavior of Cs3Sb2I9 and Rb3Sb2I9. The broadness of the PL emission is attributed to strong electron−phonon coupling comparable to the alkali halides, with calculated phonon energies matching those measured by Raman spectroscopy. These electron− phonon interactions induce small polarons that bind charge carriers and result in self-trapped excitons. The broad luminescence of Cs3Sb2I9 and Rb3Sb2I9 is tentatively attributed to phonon-assisted recombination of these self-trapped excitons. The Cs3Bi2I9 PL spectrum shows several broad overlapping peaks, with similar overall behavior that is consistent with recombination of self-trapped excitons. Rb3Bi2I9, Rb3Sb2I9, and Cs3Bi2I9 exhibited photoconductivity under laser excitation, and all four materials demonstrated high resistivity. The large band gaps, high resistivity, and high

Figure 14. I−V curves in the dark (black trace) and under 50 mW laser (red trace) for Cs3Bi2I9. 4141

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(5) Kojima, A.; Teshima, K.; Shirai, Y.; Miyasaka, T. Organometal Halide Perovskites as Visible-Light Sensitizers for Photovoltaic Cells. J. Am. Chem. Soc. 2009, 131, 6050−6051. (6) NREL Solar Cell Efficiency Chart. http://www.nrel.gov/ncpv/ images/efficiency_chart.jpg (accessed Oct. 25, 2016). (7) Wei, H. T.; Fang, Y. J.; Mulligan, P.; Chuirazzi, W.; Fang, H. H.; Wang, C. C.; Ecker, B. R.; Gao, Y. L.; Loi, M. A.; Cao, L.; Huang, J. S. Sensitive X-ray detectors made of methylammonium lead tribromide perovskite single crystals. Nat. Photonics 2016, 10, 333−339. (8) Tan, Z.-K.; Moghaddam, R. S.; Lai, M. L.; Docampo, P.; Higler, R.; Deschler, F.; Price, M.; Sadhanala, A.; Pazos, L. M.; Credgington, D.; Hanusch, F.; Bein, T.; Snaith, H. J.; Friend, R. H. Bright lightemitting diodes based on organometal halide perovskite. Nat. Nanotechnol. 2014, 9, 687−692. (9) Wong, A. B.; Lai, M.; Eaton, S. W.; Yu, Y.; Lin, E.; Dou, L.; Fu, A.; Yang, P. Growth and Anion Exchange Conversion of CH3NH3PbX3 Nanorod Arrays for Light-Emitting Diodes. Nano Lett. 2015, 15, 5519−5524. (10) Zhu, H.; Fu, Y.; Meng, F.; Wu, X.; Gong, Z.; Ding, Q.; Gustafsson, M. V.; Trinh, M. T.; Jin, S.; Zhu, X. Y. Lead halide perovskite nanowire lasers with low lasing thresholds and high quality factors. Nat. Mater. 2015, 14, 636−642. (11) Chen, Y.-S.; Manser, J. S.; Kamat, P. V. All Solution-Processed Lead Halide Perovskite-BiVO4 Tandem Assembly for Photolytic Solar Fuels Production. J. Am. Chem. Soc. 2015, 137, 974−981. (12) Cao, D. H.; Stoumpos, C. C.; Farha, O. K.; Hupp, J. T.; Kanatzidis, M. G. 2D Homologous Perovskites as Light-Absorbing Materials for Solar Cell Applications. J. Am. Chem. Soc. 2015, 137, 7843−7850. (13) Fabini, D. Quantifying the Potential for Lead Pollution from Halide Perovskite Photovoltaics. J. Phys. Chem. Lett. 2015, 6, 3546− 3548. (14) Niu, G.; Guo, X.; Wang, L. Review of recent progress in chemical stability of perovskite solar cells. J. Mater. Chem. A 2015, 3, 8970−8980. (15) Kulbak, M.; Cahen, D.; Hodes, G. How Important Is the Organic Part of Lead Halide Perovskite Photovoltaic Cells? Efficient CsPbBr3 Cells. J. Phys. Chem. Lett. 2015, 6, 2452−2456. (16) Eperon, G. E.; Paterno, G. M.; Sutton, R. J.; Zampetti, A.; Haghighirad, A. A.; Cacialli, F.; Snaith, H. J. Inorganic caesium lead iodide perovskite solar cells. J. Mater. Chem. A 2015, 3, 19688−19695. (17) Sutton, R. J.; Eperon, G. E.; Miranda, L.; Parrott, E. S.; Kamino, B. A.; Patel, J. B.; Hörantner, M. T.; Johnston, M. B.; Haghighirad, A. A.; Moore, D. T.; Snaith, H. J. Bandgap-Tunable Cesium Lead Halide Perovskites with High Thermal Stability for Efficient Solar Cells. Adv. Energy Mater. 2016, 6, 1502458. (18) Stoumpos, C. C.; Malliakas, C. D.; Peters, J. A.; Liu, Z.; Sebastian, M.; Im, J.; Chasapis, T. C.; Wibowo, A. C.; Chung, D. Y.; Freeman, A. J.; Wessels, B. W.; Kanatzidis, M. G. Crystal Growth of the Perovskite Semiconductor CsPbBr3: A New Material for HighEnergy Radiation Detection. Cryst. Growth Des. 2013, 13, 2722−2727. (19) Sebastian, M.; Peters, J. A.; Stoumpos, C. C.; Im, J.; Kostina, S. S.; Liu, Z.; Kanatzidis, M. G.; Freeman, A. J.; Wessels, B. W. Excitonic emissions and above-band-gap luminescence in the single-crystal perovskite semiconductors CsPbBr3 and CsPbCl3. Phys. Rev. B: Condens. Matter Mater. Phys. 2015, 92, 235210. (20) Dirin, D. N.; Cherniukh, I.; Yakunin, S.; Shynkarenko, Y.; Kovalenko, M. V. Solution-Grown CsPbBr3 Perovskite Single Crystals for Photon Detection. Chem. Mater. 2016, 28, 8470. (21) Aleksandrov, K. S.; Beznosikov, V. V. Hierarchies of perovskitelike crystals (Review). Phys. Solid State 1997, 39, 695−715. (22) Lee, B.; Stoumpos, C. C.; Zhou, N.; Hao, F.; Malliakas, C.; Yeh, C.-Y.; Marks, T. J.; Kanatzidis, M. G.; Chang, R. P. H. Air-Stable Molecular Semiconducting Iodosalts for Solar Cell Applications: Cs2SnI6 as a Hole Conductor. J. Am. Chem. Soc. 2014, 136, 15379− 15385. (23) Saparov, B.; Hong, F.; Sun, J.-P.; Duan, H.-S.; Meng, W.; Cameron, S.; Hill, I. G.; Yan, Y.; Mitzi, D. B. Thin-Film Preparation

photoresponse of these compounds are key characteristics of hard radiation detector materials,105 and future work will involve assessment of these materials for room-temperature detection.

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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.7b01184. Crystallographic refinement tables for Rb3Sb2I9, PXRD, and SEM-EDS analysis of the A3M2I9 compounds, thermal coefficients of expansion for Cs3Bi2I9 derived from synchrotron PXRD data, DTA analysis of Cs3Bi2I9, analysis of low-temperature PL spectra of Cs3Bi2I9, and I−V curves in the dark and under laser for all four compounds (PDF) Crystallographic Information File for Rb3Sb2I9 structure solution (CIF)

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AUTHOR INFORMATION

Corresponding Author

*(B.W.W.) E-mail: [email protected]. ORCID

Kyle M. McCall: 0000-0001-8628-3811 Constantinos C. Stoumpos: 0000-0001-8396-9578 Mercouri G. Kanatzidis: 0000-0003-2037-4168 Bruce W. Wessels: 0000-0002-8957-7097 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work is supported by the Office of Nonproliferation and Verification Research and Development under the National Nuclear Security Administration of the U.S. Department of Energy under Contract DE-NA0002522. Use of the Advanced Photon Source at Argonne National Laboratory was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC0206CH11357. This work made use of the EPIC and SPID facilities of the NUANCE Center at Northwestern University, which has received support from the Soft and Hybrid Nanotechnology Experimental (SHyNE) Resource (NSF NNCI-1542205); the MRSEC program (NSF DMR1121262) at the Materials Research Center; the International Institute for Nanotechnology (IIN); the Keck Foundation; and the State of Illinois, through the IIN.

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REFERENCES

(1) Stoumpos, C. C.; Kanatzidis, M. G. The Renaissance of Halide Perovskites and Their Evolution as Emerging Semiconductors. Acc. Chem. Res. 2015, 48, 2791−2802. (2) Moller, C. K. The structure of perovskite-like caesium plumbo trihalides. Mater. Fys. Medd. Dan. Vid. Selsk 1959, 32, 1−27. (3) Fröhlich, D.; Heidrich, K.; Künzel, H.; Trendel, G.; Treusch, J. Cesium-trihalogen-plumbates a new class of ionic semiconductors. J. Lumin. 1979, 18−19, 385−388. (4) Stoumpos, C. C.; Kanatzidis, M. G. The Renaissance of Halide Perovskites and Their Evolution as Emerging Semiconductors. Acc. Chem. Res. 2015, 48 (10), 2791−2802. 4142

DOI: 10.1021/acs.chemmater.7b01184 Chem. Mater. 2017, 29, 4129−4145

Article

Chemistry of Materials

(41) Dolomanov, O. V.; Bourhis, L. J.; Gildea, R. J.; Howard, J. A. K.; Puschmann, H. OLEX2: a complete structure solution, refinement and analysis program. J. Appl. Crystallogr. 2009, 42, 339−341. (42) Sheldrick, G. SHELXT - Integrated space-group and crystalstructure determination. Acta Crystallogr., Sect. A: Found. Adv. 2015, 71, 3−8. (43) Sheldrick, G. M. Crystal structure refinement with SHELXL. Acta Crystallogr., Sect. A: Found. Adv. 2015, 71, 3−8. (44) Kortüm, G.; Braun, W.; Herzog, G. Prinzip und Meßmethodik der diffusen Reflexionsspektroskopie. Angew. Chem. 1963, 75, 653− 661. (45) Kostina, S. S.; Peters, J. A.; Lin, W.; Chen, P.; Liu, Z.; Wang, P. L.; Kanatzidis, M. G.; Wessels, B. W. Photoluminescence fatigue and inhomogeneous line broadening in semi-insulating Tl6SeI4 single crystals. Semicond. Sci. Technol. 2016, 31, 065009. (46) Aleksandrova, I. P.; Sukhovsky, A. A.; Aleksandrov, K. S. Novel incommensurate phase in Cs3Bi2I9. Solid State Commun. 1998, 105, 323−326. (47) Jorio, A.; Currat, R.; Myles, D.; McIntyre, G.; Aleksandrova, I.; Kiat, J.; Saint-Grégoire, P. Ferroelastic phase transition in Cs3Bi2I9: A neutron diffraction study. Phys. Rev. B: Condens. Matter Mater. Phys. 2000, 61, 3857−3862. (48) Motsnyi, F. V.; Smolanka, O. M.; Sugakov, V. I.; Peresh, E. Y. Giant thermodynamical optical effect near ferroelastic phase transition point in Cs3Bi2I9 layered crystal. Solid State Commun. 2006, 137, 221− 224. (49) Trotter, J.; Zobel, T. The crystal structure of SbI3 and BiI3. Z. Kristallogr. - Cryst. Mater. 1966, 123, 67−72. (50) Petříček, V.; Dušek, M.; Palatinus, L. Crystallographic Computing System JANA2006: General features. Z. Kristallogr. Cryst. Mater. 2014, 229, 345−352. (51) Haynes, W. M. CRC handbook of chemistry and physics: A readyreference book of chemical and physical data, 96th ed.; CRC Press: 2015. (52) Vakulenko, O. V.; Gubanov, V. O.; Kun, S. V.; Motsnyi, F. V.; Peresh, E. Y.; Terekhov, V. A. Raman diagnostics of new types of A3B2X9 layered crystals. Optical Diagnostics of Materials and Devices for Opto-, Micro-, and Quantum Electronics; SPIE Proceedings 2259; SPIE: 1998; pp 351−354. (53) Nakamoto, K. Applications in Inorganic Chemistry. In Infrared and Raman Spectra of Inorganic and Coordination Compounds; John Wiley & Sons, Inc.: 2008; pp 149−354. (54) Laane, J.; Jagodzinski, P. W. Low-frequency vibrational spectra of bromo- and iodobismuthates and the observation of a trans effect. Inorg. Chem. 1980, 19, 44−49. (55) Motsnyi, F. V.; Vuychik, M. V.; Smolanka, O. M.; Peresh, E. Y. Phase Transition in Cs3Bi2I9 Ferroelastic: Investigation by Raman Scattering Technique. Funct. Mater. 2005, 12, 491−496. (56) Sheu, H.-L.; Laane, J. Trans Effect in Halobismuthates and Haloantimonates Revisited. Molecular Structures and Vibrations from Theoretical Calculations. Inorg. Chem. 2013, 52, 4244−4249. (57) Yaremko, A. M.; Valakh, M. Y.; Trylis, O. V.; Baran, J.; Marchewka, M. K.; Ratajczak, H. Anharmonic effects and temperature dependence of the lattice spectra of ferroelectric crystals A3B2C9. J. Mol. Struct. 1999, 511−512, 57−68. (58) Bator, G.; Baran, J.; Jakubas, R.; Karbowiak, M. Raman studies of structural phase transition in Cs3Bi2Br9. Vib. Spectrosc. 1998, 16, 11−20. (59) Miniewicz, A.; Jakubas, R.; Ecolivet, C.; Girard, A. Raman scattering in ferroelectric (CH3NH3)3Bi2Br9 single crystals. J. Raman Spectrosc. 1994, 25, 371−375. (60) Valakh, M. Y.; Lisitsa, M. P.; Trilis, A. V.; Yaremko, A. M.; Peresh, E. Y. Raman scattering in A3B2C9 ferroelectric crystals. Phys. Solid State 1997, 39, 1289−1291. (61) Stoneham, A. M.; Gavartin, J.; Shluger, A. L.; Kimmel, A. V.; Ramo, D. M.; Rønnow, H. M.; Aeppli, G.; Renner, C. Trapping, selftrapping and the polaron family. J. Phys.: Condens. Matter 2007, 19, 255208. (62) Yaffe, O.; Guo, Y.; Tan, L. Z.; Egger, D. A.; Hull, T.; Stoumpos, C. C.; Zheng, F.; Heinz, T. F.; Kronik, L.; Kanatzidis, M. G.; Owen, J.

and Characterization of Cs3Sb2I9: A Lead-Free Layered Perovskite Semiconductor. Chem. Mater. 2015, 27, 5622−5632. (24) Lehner, A. J.; Fabini, D. H.; Evans, H. A.; Hébert, C.-A.; Smock, S. R.; Hu, J.; Wang, H.; Zwanziger, J. W.; Chabinyc, M. L.; Seshadri, R. Crystal and Electronic Structures of Complex Bismuth Iodides A3Bi2I9 (A = K, Rb, Cs) Related to Perovskite: Aiding the Rational Design of Photovoltaics. Chem. Mater. 2015, 27, 7137−7148. (25) Peresh, E. Y.; Sidei, V. I.; Zubaka, O. V.; Stercho, I. P. K2(Rb2,Cs2,Tl2)TeBr6(I6) and Rb3(Cs3)Sb2(Bi2)Br9(I9) perovskite compounds. Inorg. Mater. 2011, 47, 208−212. (26) Giustino, F.; Snaith, H. J. Toward Lead-Free Perovskite Solar Cells. ACS Energy Lett. 2016, 1, 1233−1240. (27) Harikesh, P. C.; Mulmudi, H. K.; Ghosh, B.; Goh, T. W.; Teng, Y. T.; Thirumal, K.; Lockrey, M.; Weber, K.; Koh, T. M.; Li, S.; Mhaisalkar, S.; Mathews, N. Rb as an Alternative Cation for Templating Inorganic Lead-Free Perovskites for Solution Processed Photovoltaics. Chem. Mater. 2016, 28, 7496−7504. (28) Hoye, R. L. Z.; Brandt, R. E.; Osherov, A.; Stevanović, V.; Stranks, S. D.; Wilson, M. W. B.; Kim, H.; Akey, A. J.; Perkins, J. D.; Kurchin, R. C.; Poindexter, J. R.; Wang, E. N.; Bawendi, M. G.; Bulović, V.; Buonassisi, T. Methylammonium Bismuth Iodide as a Lead-Free, Stable Hybrid Organic−Inorganic Solar Absorber. Chem. Eur. J. 2016, 22, 2605−2610. (29) Pazoki, M.; Johansson, M. B.; Zhu, H.; Broqvist, P.; Edvinsson, T.; Boschloo, G.; Johansson, E. M. J. Bismuth Iodide Perovskite Materials for Solar Cell Applications: Electronic Structure, Optical Transitions, and Directional Charge Transport. J. Phys. Chem. C 2016, 120, 29039−29046. (30) Hebig, J.-C.; Kühn, I.; Flohre, J.; Kirchartz, T. Optoelectronic Properties of (CH3 NH 3 ) 3 Sb 2 I 9 Thin Films for Photovoltaic Applications. ACS Energy Lett. 2016, 1, 309−314. (31) Park, B.-W.; Philippe, B.; Zhang, X.; Rensmo, H.; Boschloo, G.; Johansson, E. M. J. Bismuth Based Hybrid Perovskites A3Bi2I9 (A: Methylammonium or Cesium) for Solar Cell Application. Adv. Mater. 2015, 27, 6806−6813. (32) Han, H.; Hong, M.; Gokhale, S. S.; Sinnott, S. B.; Jordan, K.; Baciak, J. E.; Nino, J. C. Defect Engineering of BiI3 Single Crystals: Enhanced Electrical and Radiation Performance for Room Temperature Gamma-Ray Detection. J. Phys. Chem. C 2014, 118, 3244−3250. (33) Lehner, A. J.; Wang, H.; Fabini, D. H.; Liman, C. D.; Hébert, C.A.; Perry, E. E.; Wang, M.; Bazan, G. C.; Chabinyc, M. L.; Seshadri, R. Electronic structure and photovoltaic application of BiI3. Appl. Phys. Lett. 2015, 107, 131109. (34) Ganose, A. M.; Savory, C. N.; Scanlon, D. O. Beyond Methylammonium Lead Iodide: Prospects for the Emergent Field of ns2 Containing Solar Absorbers. Chem. Commun. 2017, 53, 20−44. (35) Lindqvist, O.; Johansson, G.; Sandberg, F.; Norin, T. The Crystal Structure of Caesium Bismuth Iodide, Cs3Bi2I9. Acta Chem. Scand. 1968, 22, 2943−2952. (36) Yamada, K.; Sera, H.; Sawada, S.; Tada, H.; Okuda, T.; Tanaka, H. Reconstructive Phase Transformation and Kinetics of Cs3Sb2I9 by Means of Rietveld Analysis of X-Ray Diffraction and 127I NQR. J. Solid State Chem. 1997, 134, 319−325. (37) Sidey, V. I.; Voroshilov, Y. V.; Kun, S. V.; Peresh, E. Y. Crystal growth and X-ray structure determination of Rb3Bi2I9. J. Alloys Compd. 2000, 296, 53−58. (38) Chang, J.-H.; Doert, T.; Ruck, M. Structural Variety of Defect Perovskite Variants M3E2X9 (M = Rb, Tl, E = Bi, Sb, X = Br, I). Z. Anorg. Allg. Chem. 2016, 642, 736−748. (39) Peresh, E. Y.; Sidei, V. I.; Gaborets, N. I.; Zubaka, O. V.; Stercho, I. P.; Barchii, I. E. Influence of the average atomic number of the A2TeC6 and A3B2C9 (A = K, Rb, Cs, Tl(I); B = Sb, Bi; C = Br, I) compounds on their melting point and band gap. Inorg. Mater. 2014, 50, 101−106. (40) Ivanov, Y. N.; Sukhovskii, A. A.; Lisin, V. V.; Aleksandrova, I. P. Phase Transitions of Cs3Sb2I9, Cs3Bi2I9, and Cs3Bi2Br9 Crystals. Inorg. Mater. 2001, 37, 623−627. 4143

DOI: 10.1021/acs.chemmater.7b01184 Chem. Mater. 2017, 29, 4129−4145

Article

Chemistry of Materials S.; Rappe, A. M.; Pimenta, M. A.; Brus, L. E. Local Polar Fluctuations in Lead Halide Perovskite Crystals. Phys. Rev. Lett. 2017, 118, 136001. (63) Nilă, A.; Baibarac, M.; Matea, A.; Mitran, R.; Baltog, I. Exciton− phonon interactions in the Cs3Bi2I9 crystal structure revealed by Raman spectroscopic studies. Phys. Status Solidi B 2017, 254, 1552805. (64) Machulin, V. F.; Motsnyi, F. V.; Smolanka, O. M.; Svechnikov, G. S.; Peresh, E. Y. Effect of temperature variation on shift and broadening of the exciton band in Cs3Bi2I9 layered crystals. Low Temp. Phys. 2004, 30, 964−967. (65) Xiao, Z.; Meng, W.; Wang, J.; Mitzi, D. B.; Yan, Y. Searching for promising new perovskite-based photovoltaic absorbers: the importance of electronic dimensionality. Mater. Horiz. 2017, 4, 206−216. (66) Kittel, C. Introduction to Solid State Physics, 8th ed.; John Wiley & Sons, Inc.: 2005. (67) Knox, R. S. Introduction to Exciton Physics. In Collective Excitations in Solids; Di Bartolo, B., Ed.; Springer US: Boston, MA, 1983; pp 183−245. (68) Kawai, T.; Ishii, A.; Kitamura, T.; Shimanuki, S.; Iwata, M.; Ishibashi, Y. Optical Absorption in Band-Edge Region of (CH3NH3)3Bi2I9 Single Crystals. J. Phys. Soc. Jpn. 1996, 65, 1464− 1468. (69) Bass, K. K.; Estergreen, L.; Savory, C. N.; Buckeridge, J.; Scanlon, D. O.; Djurovich, P. I.; Bradforth, S. E.; Thompson, M. E.; Melot, B. C. Vibronic Structure in Room Temperature Photoluminescence of the Halide Perovskite Cs3Bi2Br9. Inorg. Chem. 2017, 56, 42−45. (70) Stoumpos, C. C.; Cao, D. H.; Clark, D. J.; Young, J.; Rondinelli, J. M.; Jang, J. I.; Hupp, J. T.; Kanatzidis, M. G. Ruddlesden−Popper Hybrid Lead Iodide Perovskite 2D Homologous Semiconductors. Chem. Mater. 2016, 28, 2852−2867. (71) Sestu, N.; Cadelano, M.; Sarritzu, V.; Chen, F.; Marongiu, D.; Piras, R.; Mainas, M.; Quochi, F.; Saba, M.; Mura, A.; Bongiovanni, G. Absorption F-Sum Rule for the Exciton Binding Energy in Methylammonium Lead Halide Perovskites. J. Phys. Chem. Lett. 2015, 6, 4566−4572. (72) Pelant, I.; Valenta, J. Luminescence Spectroscopy of Semiconductors; Oxford University Press: New York, 2012. (73) Dean, P.; Fischer, B.; Herbert, D.; Lagois, J.; Yu, P.; Cho, K. Excitons; Springer Science & Business Media: 2012; Vol. 14. (74) Mott, N. F.; Stoneham, A. M. The lifetime of electrons, holes and excitons before self-trapping. J. Phys. C: Solid State Phys. 1977, 10, 3391−3398. (75) Purdy, A. E.; Murray, R. B.; Song, K. S.; Stoneham, A. M. Studies of self-trapped exciton luminescence in KCl. Phys. Rev. B 1977, 15, 2170−2176. (76) Czaja, W. Concentration-induced spectral variations of luminescence in AgBr:I. J. Phys. C: Solid State Phys. 1983, 16, 3197− 3213. (77) Song, K. S.; Williams, R. T.; et al. Materials, Extrinsic SelfTrapping, and Low-Dimensional Systems. In Self-Trapped Excitons; Springer Berlin Heidelberg: Berlin, Heidelberg, 1993; pp 351−373. (78) Yangui, A.; Garrot, D.; Lauret, J. S.; Lusson, A.; Bouchez, G.; Deleporte, E.; Pillet, S.; Bendeif, E. E.; Castro, M.; Triki, S.; Abid, Y.; Boukheddaden, K. Optical Investigation of Broadband White-Light Emission in Self-Assembled Organic−Inorganic Perovskite (C6H11NH3)2PbBr4. J. Phys. Chem. C 2015, 119, 23638−23647. (79) Cortecchia, D.; Neutzner, S.; Srimath Kandada, A. R.; Mosconi, E.; Meggiolaro, D.; De Angelis, F.; Soci, C.; Petrozza, A. Broadband Emission in Two-Dimensional Hybrid Perovskites: The Role of Structural Deformation. J. Am. Chem. Soc. 2017, 139, 39−42. (80) Timmermans, C. W. M.; Blasse, G. On the Luminescence of Cs3Bi2Br9 Single Crystals. Phys. Status Solidi B 1981, 106, 647−655. (81) Morgan, T. N. Broadening of Impurity Bands in Heavily Doped Semiconductors. Phys. Rev. 1965, 139, A343−A348. (82) Mao, L.; Wu, Y.; Stoumpos, C. C.; Wasielewski, M. R.; Kanatzidis, M. G. White-Light Emission and Structural Distortion in New Corrugated Two-Dimensional Lead Bromide Perovskites. J. Am. Chem. Soc., 2017, 139, 5210−5215.

(83) Ueta, M.; Kanzaki, H.; Kobayashi, K.; Toyozawa, Y.; Hanamura, E. Theory of Excitons in Phonon Fields. In Excitonic Processes in Solids; Springer Berlin Heidelberg: Berlin, Heidelberg, 1986; pp 203−284. (84) Shi, S. L.; Li, G. Q.; Xu, S. J.; Zhao, Y.; Chen, G. H. Green Luminescence Band in ZnO: Fine Structures, Electron−Phonon Coupling, and Temperature Effect. J. Phys. Chem. B 2006, 110, 10475−10478. (85) Grushka, E. Characterization of exponentially modified Gaussian peaks in chromatography. Anal. Chem. 1972, 44, 1733−1738. (86) Pezzoli, F.; Qing, L.; Giorgioni, A.; Isella, G.; Grilli, E.; Guzzi, M.; Dery, H. Spin and energy relaxation in germanium studied by spinpolarized direct-gap photoluminescence. Phys. Rev. B: Condens. Matter Mater. Phys. 2013, 88, 045204. (87) Schmidt, T.; Lischka, K.; Zulehner, W. Excitation-power dependence of the near-band-edge photoluminescence of semiconductors. Phys. Rev. B: Condens. Matter Mater. Phys. 1992, 45, 8989−8994. (88) Toyozawa, Y. Further Contribution to the Theory of the LineShape of the Exciton Absorption Band. Prog. Theor. Phys. 1962, 27, 89−104. (89) Stadler, W.; Hofmann, D. M.; Alt, H. C.; Muschik, T.; Meyer, B. K.; Weigel, E.; Müller-Vogt, G.; Salk, M.; Rupp, E.; Benz, K. W. Optical investigations of defects in Cd1‑xZnxTe. Phys. Rev. B: Condens. Matter Mater. Phys. 1995, 51, 10619−10630. (90) Morrison, S. R.; Rolfe, J. An anomalous absorption spectrum: The missing zero-phonon line of the F2-centre in KCl crystals. Solid State Commun. 1976, 18, 617−619. (91) Leung, C. H.; Song, K. S. On the luminescence quenching of F centres in alkali halides. Solid State Commun. 1980, 33, 907−910. (92) Dawson, R. K.; Pooley, D. F Band Absorption in Alkali Halides as a Function of Temperature. Phys. Status Solidi B 1969, 35, 95−105. (93) Novikova, M. S.; Makarova, I. P.; Blomberg, M. K.; Bagautdinov, B. S.; Aleksandrova, I. P. Structure and phase transitions in trigonal Cs3Sb2I9 crystals. Crystallogr. Rep. 2001, 46, 26−29. (94) Krustok, J.; Collan, H.; Hjelt, K. Does the low-temperature Arrhenius plot of the photoluminescence intensity in CdTe point towards an erroneous activation energy? J. Appl. Phys. 1997, 81, 1442− 1445. (95) Leroux, M.; Grandjean, N.; Beaumont, B.; Nataf, G.; Semond, F.; Massies, J.; Gibart, P. Temperature quenching of photoluminescence intensities in undoped and doped GaN. J. Appl. Phys. 1999, 86, 3721−3728. (96) Pankove, J. I. Optical Processes in Semiconductors; Dover: New York, 1971. (97) Wight, D. R. Green luminescence efficiency in gallium phosphide. J. Phys. D: Appl. Phys. 1977, 10, 431−454. (98) Thomas, D. G.; Hopfield, J. J. Isoelectronic Traps Due to Nitrogen in Gallium Phosphide. Phys. Rev. 1966, 150, 680−689. (99) Williams, R. T.; Song, K. S. The self-trapped exciton. J. Phys. Chem. Solids 1990, 51, 679−716. (100) Fischbach, J. U.; Fröhlich, D.; Kabler, M. N. Recombination luminescence lifetimes and the self-trapped excition in alkali halides. J. Lumin. 1973, 6, 29−43. (101) Kurisu, H.; Komatsu, T.; Tanaka, T.; Akai, I.; Karasawa, T.; Iida, T. Pressure effects on excitons, phonons and exciton-phonon interactions in BiI3 and SbI3 single crystals. J. Lumin. 1994, 60, 816− 818. (102) Wright, A. D.; Verdi, C.; Milot, R. L.; Eperon, G. E.; PerezOsorio, M. A.; Snaith, H. J.; Giustino, F.; Johnston, M. B.; Herz, L. M. Electron-phonon coupling in hybrid lead halide perovskites. Nat. Commun. 2016, 7, 11755. (103) Hu, T.; Smith, M. D.; Dohner, E. R.; Sher, M.-J.; Wu, X.; Trinh, M. T.; Fisher, A.; Corbett, J.; Zhu, X. Y.; Karunadasa, H. I.; Lindenberg, A. M. Mechanism for Broadband White-Light Emission from Two-Dimensional (110) Hybrid Perovskites. J. Phys. Chem. Lett. 2016, 7, 2258−2263. (104) Bimberg, D.; Sondergeld, M.; Grobe, E. Thermal Dissociation of Excitons Bounds to Neutral Acceptors in High-Purity GaAs. Phys. Rev. B 1971, 4, 3451−3455. 4144

DOI: 10.1021/acs.chemmater.7b01184 Chem. Mater. 2017, 29, 4129−4145

Article

Chemistry of Materials (105) Owens, A.; Peacock, A. Compound semiconductor radiation detectors. Nucl. Instrum. Methods Phys. Res., Sect. A 2004, 531, 18−37. (106) Cha, J.-H.; Han, J. H.; Yin, W.; Park, C.; Park, Y.; Ahn, T. K.; Cho, J. H.; Jung, D.-Y. Photoresponse of CsPbBr3 and Cs4PbBr6 Perovskite Single Crystals. J. Phys. Chem. Lett. 2017, 8, 565−570. (107) Ding, J.; Du, S.; Zuo, Z.; Zhao, Y.; Cui, H.; Zhan, X. High Detectivity and Rapid Response in Perovskite CsPbBr3 Single-Crystal Photodetector. J. Phys. Chem. C 2017, 121, 4917−4923. (108) Fang, Y.; Dong, Q.; Shao, Y.; Yuan, Y.; Huang, J. Highly narrowband perovskite single-crystal photodetectors enabled by surface-charge recombination. Nat. Photonics 2015, 9, 679−686.

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DOI: 10.1021/acs.chemmater.7b01184 Chem. Mater. 2017, 29, 4129−4145