RbPbI3 halide perovskite solar cells

Solar Energy Materials and Solar Cells 172 (2017) 44–54

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Solar Energy Materials and Solar Cells journal homepage: www.elsevier.com/locate/solmat

TiO2/RbPbI3 halide perovskite solar cells a,⁎

b

Mi-Hee Jung , Sonny H. Rhim , Dohyun Moon a b c

MARK

c

Department of Nanotechnology and Advanced Materials Engineering, Sejong University, 209, Neungdong-ro, Gwangjin-gu, Seoul 05006, Republic of Korea Department of Physics and Energy Harvest-Storage Research Center, University of Ulsan, Ulsan 680-749, Republic of Korea Pohang Accelerator Laboratory, Beamline Department, 80 jigokro-127-beongil, Nam-Gu, Pohang, Gyeongbuk 37673, Repubic of Korea

A R T I C L E I N F O

A B S T R A C T

Keywords: RbPbI3 Perovskite One dimensional Solar cell

Inorganic-organic halide perovskites hold the great promise for next-generation photovoltaics due to their excellent high performance and low cost. However, major limitation for the commercialization of perovskite solar cell can be attributed to the transformation and degradation of lead halide perovskite during the exposure of environmental humidity and photon irradiation. To solve these problems, herein, we apply the one-dimensional and inorganic RbPbI3 perovskite into the solar cell because it shows the superior stability in environmental conditions. After RbPbI3 perovskite was applied into the solar cell with a FTO/TiO2/RbPbI3/Spiro-MeOTAD/Au configuration, the device exhibits an open circuit voltage of 0.62 V, photocurrent density of 3.75 mA/cm2, fill factor of 44.60%, and 1.04% of PCE by reverse sweeping direction. Even though the performance of RbPbI3 device was still lower than other perovskite solar cells, this approach enabled us to establish the key step to make a highly stable perovskite film, leading to the best photovoltaic performance for real applications.

1. Introduction Organic-inorganic halide perovskite (ABX3, A = CH3NH3, HC(NH2)2; B = Pb, Sn,; X = Cl, Br, I) solar cells have recently received considerable amounts of attention in the field of photovoltaics due to the unprecedented increase in their efficiency from 3.8% [1] to levels exceeding 20% [2]. This is primarily due to the possibility of the formation of defectfree crystalline film at low temperatures with a solution process [3], desirable optical properties such good band gap tenability [4], a long ambipolar diffusion length (> 1 µm for CH3NH3PbI3−xClx and > 175 µm for single-crystal CH3NH3PbI3) [5,6], a high optical absorption coefficient (α = 104–105 cm-1 for hv > 1.7 eV) [7], low exciton binding energy (50 meV) [8] and the availability and use of inexpensive input materials. Since Miyasaka et al. [1] reported a solar cell device based on CH3NH3PbI3(Br3) in 2009, the power conversion efficiency levels of perovskite solar cells have increased rapidly to 10.9% using a P/N heterojunction with a super-structured organometal compound [9]. Later, the efficiency of CH3NH3PbI3-based solar cells was improved to 15% in planar heterojunction solar cells and then boosted to 20% by means of compositional engineering [10] and an intramolecular exchange approach [2]. Recently, as the latest solar cell technology, a two-terminal silicon/perovskite tandem solar cell exhibited a cell efficiency rate of 19.9% with a short-circuit current (Jsc) of 14.0 mAcm-2, a Voc of 1.785 V and a FF of 79.5% [11]. For application to unmanned aerial vehicles, Kaltenbrunner et al. [12] demonstrated an ultrathin (3 µm) and highly flexible perovskite

⁎

Corresponding author. E-mail address: [email protected] (M.-H. Jung).

http://dx.doi.org/10.1016/j.solmat.2017.07.011 Received 14 March 2017; Received in revised form 5 July 2017; Accepted 13 July 2017 0927-0248/ © 2017 Elsevier B.V. All rights reserved.

solar cell using PET foil which showed 12% efficiency and a power-perweight value as high as 23 Wg-1. Tsai et al. [13] prepared a thin singlecrystalline-like film using a two-dimensional Ruddlesden-Popper perovskite. They reported a photovoltaic efficiency rate of 12.52% with no hysteresis and highly improved stability against humidity. The perovskites most studied are organic-inorganic hybrid halides, which are generally composed of a polar organic cation, such as alkyl ammonium or formamidinium, as a space group between inorganic layers. Whilst the use of organic constituents in the perovskite structure offers highly efficient photovoltaic properties, one concern associated with these hybrid perovskites when compared to the conventional thin film compound, is the much lower thermal stability (~ 300 °C for the bulk and 150–200 °C for thin film types) as opposed to that of a conventional inorganic compound. Moreover, given that the organic cation is fully disordered in the structure [14] due to the dipole, perovskite solar cells show hysteresis [15]. To address the issue of the thermodynamic instability of hybrid organic-inorganic perovskites, many efforts have been made to replace the organic cation with an inorganic equivalent alkali metal (Cs or Rb) [15–17], as an inorganic cation is much less volatile and has an anisotropic geometry which can eliminate the phenomenon of hysteresis and accelerate electron movement. Recently, CsSnI3 was reported to have favorable properties as a light absorber due to the high mobility of electrons (~ 2300 cm2 V-1 s-1) and holes (~ 320 cm2 V-1 s-1) [18], a high optical absorption coefficient (104 cm-1) [19], low exciton binding energy (18 meV) [20] and a smaller band gap (1.3 eV) [21]. Chen et al. [22] reported a

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Fig. 1. (a) Crystal structure of RbPbI3 (orthorhombic, space group Pnma) at 298 K, (b) iodine environment for the Rb and Pb2+ cations, (c) X-ray diffraction spectra (XRD) of RbPbI3 single crystal in yellow phase at room temperature and (d) band gaps of RbPbI3 film calculated from the diffuse reflectance spectra.

incorporated into a precursor perovskite, it can decrease surface and bulk defects, which then suppresses the major sites of non-radiative recombinations, resulting in highly stable and efficient solar cells [29]. Moreover, RbPbI3 perovskite exhibits a high absorption coefficient range of 5.21 × 104 cm-1–8.35 × 104 cm-1 from 400 nm to 600 nm and a specific value of 6.93 × 104 cm-1 at 550 nm (Fig. S1), which is comparable to the absorption coefficient of CH3NH3PbI3 perovskite [30]. Investigations of these halide perovskites can lead to a better understanding of how perovskite solar cells degrade [26]. Therefore, there is a need for research on the RbPbI3 perovskite structure at room temperature for application to photovoltaic devices. In this study, we investigated RbPbI3 with the goal of applying it to a stable perovskite solar cell at room temperature. The solar cell structure was FTO/TiO2/RbPbI3/Spiro-MeOTAD/Au, with power conversion efficiency of 1.04% in the reverse sweep direction according to photocurrent-photovoltage (J-V) measurements. The impedance of the photocurrent and photovoltage was measured to investigate the photoexcitation under light illumination (450 nm wavelength). Density functional theory calculations indicated that the electronic structure of RbPbI3 is highly suitable for photovoltaic construction when TiO2 is used as an electron-transport material. However, the one-dimensional connectivity of the RbPbI3 structure results in a flat band structure, which prevents any effective electron transfer into the TiO2 structure. Thus, optimizing these double-sided properties to transfer electrons efficiently into the TiO2 structure is necessary.

planar Schottky cell configuration based on CsSnI3/ITO with power conversion efficiency of nearly 1%. CsSnI3 doped with SnF2 was demonstrated as a light absorber in solid-state perovskite solar cells, exhibiting a high photocurrent density of more than 22 mA cm-2 [23]. However, given that this material can degrade due to structural transitions, issues arise during the actual use of these solar cells. More specifically, the phase of the CsSnI3 perovskite structure changes, with three different black phases possible, abbreviated as B-α (cubic, Pm3m), B-β (tetragonal, P4/mbm), and B-γ (orthorhombic, Pnma) due to the high temperatures [21,24]. Herein, the B-γ phase is the most relevant for solar cell applications because it arises when the temperature is below 352 K. However, when the perovskite structure of B-γ is exposed to air or humidity, it degrades s into a non-perovskite yellow phase (Y-CsSnI3) [21] which does not possess the desirable properties for photovoltaic applications. However, despite the fact that the RbPbI3 analogue is also transformed from a high-temperature cubic structure (Pm3m) to a roomtemperature orthorhombic structure (Pnma) [25], the RbPbI3 perovskite is highly stable at room temperature due to its stabilized 1D orthorhombic configuration (a-a-c+ phase) [26]. In addition, the band structure of the orthorhombic RbPbI3 perovskite was found to be well suited for device integration using an electron-transport material such as TiO2 [25,27]. Recently, Saliba et al. [28] reported that Cs and Rb when added as cation cascades increase the stability of perovskites and produce highly efficient solar cells. It was also reported that when Rb is 45

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Fig. 2. Band structure of (a) RbPbI3, (b) RbSnI3, (c) density of the state plot of RbPbI3; (i) Total DOS, projected DOS for (ii) Rb, (iii) Pb, (iv-vi) I atoms, respectively. In PDOS, contributions from s and p orbital s are denoted in black and red, respectively and (d) density of state plot of RbSnI3; (i) Total DOS, projected DOS for (ii) Rb, (iii) Sn, (iv-vi) I atoms, respectively. In PDOS, contributions from s and p orbital s are denoted in black and red, respectively.

2. Experimental section

2 cm in size. The devices were illuminated through a shadow mask, yielding an active area of 0.175 cm2. To fabricate the solar cell, a 50 nm compact layer of TiO2 was deposited on the FTO substrate by spin coating (2000 rpm 30 s) a mildly acidic solution of titanium isopropoxide (TTIP) in ethanol containing 350 μL TTIP in 5 mL EtOH containing 0.013 M HCl. Then we conducted the annealing process at 550 °C for 30 min. A mesoporous TiO2 layer, which was prepared diluted TiO2 paste (DYESOL-18NRT) with ethanol with a ratio of 1:3.5 w/ w, was spin-coated on the TiO2 compact layer at 500 rpm for 5 s, 3000 rpm for 10 s and 6000 rpm for 30 s and then, sintered at 550 °C for 30 min. To interconnected TiO2 layer, the substrate was further treated with 20 mM TiCl4 aqueous solution at 70 oC for 30 min, rinsed with deionized water and ethanol, and then annealed at 500 °C for 30 min. The equimolar 1 M RbI and PbI2 were dissolved in the DMF solvent. The solution was spin coated on the N type TiO2 layer at 4000 rpm for 30 s followed by the annealing process at 100 oC for 30 min. A solution for spiroMeOTAD coating was prepared by

2.1. RbPbI3 single-crystal preparation The crystal growth of the RbPbI3 perovskite was prepared by a fusion method via available the literature [31]. To grow the RbPbI3 single crystal, we mixed RbI and PbI2 at a stoichiometric ratio and placed these materials in a Pyrex tube which was then sealed in a low vacuum state (1 × 106 Torr). Heating to 800 oC followed, at which the samples were sufficiently melted. After maintaining a temperature of 800 oC for 5 h, the samples were cooled slowly from 800 °C to 40 °C over a time period of 100 h, allowing the formation of crystal nuclei with the subsequent formation of a larger crystal due to the attachment of other crystals on the surface of the crystal nuclei. 2.2. Solar cell fabrication The RbPbI3 solar cell was fabricated with a FTO substrate 2 cm × 46

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2500

2000

Intensity (arb. unit)

deposited at 4000 rpm for 30 s on the perovskite layer. The gold layer of 800 Å for the counter electrode was deposited on the perovskite layer with the thermal evaporation at 1.5 Å/s.

RbPbI3 excitation RbPbI3 emission TiO2/RbPbI3 excitation TiO2/RbPbI3 emission

2.3. Measurement and Characterization

1500

Optical diffuse reflectance measurements were performed at room temperature using a Cary 5000 (Varian) spectrometer operating in a wavelength range of 175–3300 nm. The diffuse reflectance measurement was used to estimate the band gap of the material by converting reflectance to absorption data according to the Kubelka–Munk equation [33]: α/S = (1–R)2(2R)−1, where R is the reflectance and α and S are the absorption and scattering coefficients, respectively. More than 20 solar cells were measured using simulated solar light of AM 1.5 G produced by a 1000 W xenon lamp (Oriel, 91193). Its irradiance was calibrated with a Si reference solar cell (NREL certified KG5 filtered silicon reference diode) to adjust the one sun light intensity (100 mW/ cm2). The emission and exciton spectrum of RbPbI3 were measured with a FluoroMate FS-2 – Fluorescence Spectrometer (Scinco, Republic of Korea). Hall effect measurements were conducted in air using four contacts The Hall bar method utilized DC current flowing through perovskite film applied using a Keithley Model 2400 instrument, and the Hall voltage was recorded using a Keithley Model 4200. Field

1000

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Wavelength (nm) Fig. 3. Fluorescence spectra for the excitation and emission of RbPbI3 and RbPbI3/TiO2 films on glass substrate.

dissolving 72.3 mg spiroMeOTAD in 1 mL of chlorobenzene, to which 28.8 μL of 4-tert- butyl pyridine (TBP) and 17.5 μL of lithium bis(trifl uoromethanesulfonyl) imide solution (520 mg Li-TSFI in 1 mL acetonitrile (Sigma–Aldrich, 99.8%)) were added [32]. Spiro-MeOTAD was

Fig. 4. (a) The UPS spectra of RbPbI3 perovskite layer on FTO and TiO2/FTO substrates, (b) photoemission cut off edge, (c) UPS spectra of the top of occupied states, i.e., HOMO level for the RbPbI3/FTO and RbPbI3/TiO2/FTO substrates and (d) energy level for the FTO/TiO2/RbPbI3/spiro-MeOTAD/Au solar cell configuration.

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Fig. 5. (a) The cross sectional focused ion beam SEM image of the FTO/TiO2/RbPbI3 perovskite/Spiro-OMeTAD/Au solar cell device, (b) SEM image of RbPbI3 perovskite film prepared from a 30 wt% solution in DMF, (c) SEM image and corresponding selected area for EDAX elemental mappings of Ti, O, Rb, Pb and I in the cross sectional RbPbI3/TiO2 layer and (d) The EDS microanalysis on selected area in Fig. 5(c).

frequency of 100 Hz from −1 to 1 V with a 0.1 V voltage step. Ultraviolet photoemission spectroscopy (UPS, model AXIS-NOVA, Kratos) spectra of the ITO surface were studied to determine HOMO level and work function of RbPbI3 in an ultrahigh vacuum (UHV). The base pressure of the chamber was 2 × 10−10 Torr, which rose to 5 × 10−6 Torr during the Ar-ion sputtering due to Ar-gas backfilling. During the UPS measurement the pressure went to the high 10−9 Torr range mainly due to the He. The He I line at 21.2 eV was taken as the excitation source for the UPS. The X-ray source was the Al Kα line at 1478 eV. The optical constant of RbPbI3 thin film was measured at room temperature and low humidity (< 25% RH) by the VASE-ellipsometer (J. A. Woollam Co., Inc).

strength of 0–1.25 T and current of 100 nA were applied. The observed voltage was 0.19–0.25 V. and the calculated resistivity was 8.48 × 104–1.10 × 105 Ω cm. The intensity modulated photovoltage spectroscopy (IMVS) and intensity modulated photocurrent spectroscopy (IMPS) measurements were performed on a ZAHNER CIMPS system (ZAHNER-elektrik GmbH & Co.). The LED was operated with a potentiostatic feedback loop to control the stationary DC voltage and a concurrent sinusoidal modulated AC voltage. The AC amplitude was fixed at 10% of the stationary DC value. The solar cell was placed under the potentiostat unit. IMPS measured the periodic photocurrent response of the cell to the variation of light intensity, while the IMVS experiment was used to measure the periodic modulation of the photovoltage under the change of light intensity modulation. The real and imaginary part of the modulated photovoltage and photocurrent of the IMVS and IMPS were calculated with the Levenberg-Marquardt algorithm program. MottSchottky (C-2– ϕ) analysis (Garmy instrument) was measured at a

2.4. X-ray crystallographic analysis The RbPbI3 crystal was coated with Paratone-N oil (Paratone-N oil was used not only for mounting on a dual-thickness MicroLoop Ld 48

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and (ii) the total energy difference between two successive iteration steps, with the former depending on how we established the charge density criteria. As the systems are all insulators, that would be sufficient. Additional information was already given in the “Computational Methods” section, where we discuss the use of the full-potential linearized augmented plane wave (FLAPW) method. 3. Results and discussion The crystal structure of RbPbI3 was determined by single-crystal diffraction at 100 K, 298 K and 450 K. The refined structural parameters of PbPbI3 are summarized in Table S1 (Supporting Information). RbPbI3 crystallizes in the orthorhombic space group Pnma with four molecules per unit cell. The structural elements of the RbPbI3 were arranged in the double chain of (PbI6)4− along the c-axis, which was held together by Rb ions, and Pb2+ ions were surrounded octahedrally by I- (Fig. 1a). At 298 K, the Pb2+ ions were located inside the distorted iodine octahedra exhibiting two pairs of equatorial Pb-I distances (3.2243(5) Å and 3.2516(5) Å) and two apical distances (3.0490 (6) Å and 3.3811 (7) Å). The Rb ions of RbPbI3 possessed a more irregular nine-fold iodine environment, specifically a tree pair of iodine atoms at distances of 3.8011(8) Å, 3.8048(8) Å and 3.8178(8) Å and three single iodine atoms at distances of 3.8246(10) Å, 4.0550(10) Å and 4.0933(11) Å (Fig. 1b). In this result, one of the Rb-I ions in the ninefold iodine environment was 7.05% shorter than the longest one. This outcome was in the absence of a phase transition in the RbPbI3 up to 634 K [40]. In contrast, the Cs-I distance in the nine-fold iodine environment of CsPbI3 was only 2.8% shorter. Therefore, CsPbI3 underwent a reversible phase transformation of Pm3m (500 K) ↔ P4/mbm (380 K) ↔ Pnma (300 K) with a large volume change due to the high temperature [21]. Refinement of X-ray diffraction data for the yellow phase (Fig. 1c) at room temperature indicated an orthorhombic perovskite structure with the following characteristics: lattice constant a = 10.2976(6) Å, b= 4.7811(2) Å, c = 17.406(1) Å and space group Pnma. The XRD reference of RbPbI3 (ICDD: 98-000-6067) was added into Fig. 1c to confirm the diffraction peaks of RbPbI3. The band gap of RbPbI3 was measured by diffuse reflectance measurements (Fig. 1d) with the band gap estimated from the Kubelka–Munk function f (R) = (1-R)2/(2 R), where R is the reflectance for RbPbI3. The position of the absorption band was determined by the localization of excitons in the PbI2 sublattice. Thus, the top of the valence band of RbPbI3 was formed by Pb 6 s-states mixed with some I 5pstates, and the conduction band was determined by the Pb 6p-states due to the localization of excitons in the PbI2 sublattice of the RbPbI3 compounds. The absorption spectra of RbPbI3 can be treated on the electronic structure of the [PbI6]4− octahedra of the crystal lattices of the compounds [41]. The crystal structure of RbPbI3 is regarded as a quasi-two-dimensional octahedral network with ribbons of edge-connected PbI6 octahedra, resulting in the segmentation of the PbI2 layer [25]. We expect that the band gap of 2.64 eV was due to the edgesharing PbI6 coordination environment, while that of 1.98 eV was caused by the large reduction of the PbI2 band gap. We calculated the band structure of RbPbI3 and compared it to that of RbSnI3. The band structures of RbPbI3 and RbSnI3 are shown in Fig. 2a and b, respectively. At first glance, both exhibit very similar electronic structures with indirect band gaps of 2.27 and 1.84 eV, respectively. Note that these gaps were underestimated with respect to the experiment due to the well-known DFT shortcoming [42–44]. The valence band maximum (VBM) occurred near the Γ point but shifted slightly to the Y point, (0,0.15,0) and (0,1/8,0), for RbPbI3 and RbSnI3, respectively. On the other hand, the conduction band minimum (CBM) was located near X. To investigate this in greater detail, the density-ofstates (DOS) values of RbPbI3 and RbSnI3 are shown in Fig. 2c and d, respectively, where the total and atomic projections are displayed. Owing to the symmetry, I atoms were found to exist in three distinct types, but their DOS contributions were qualitatively almost identical.

Fig. 6. (a) current density-voltage (J-V) curve of the RbPbI3 photovoltaic device and (b) Mott–Schottky (C-2–ϕ) plot for the TiO2/RbPbI3 perovskite/Spiro-OMeTAD interface system. The symbols represent the experimental data and the solid lines represent the fit curves.

assembly (MiTeGen, USA) but also to protect the crystal from a cold nitrogen stream at 100 K. Paratone-N oil is also stable 298 K and 450 K, and protection from the nitrogen stream was accomplished using a Cryojet-5 device (Oxford Instrument Inc.; temperature range 85–500 K)) and the diffraction data was measured at 100 K, 298 K, and 450 K with synchrotron radiation (λ = 0.61000 Å) on a ADSC Quantum-210 detector at BL2D SMC with a silicon (111) double crystal monochromator (DCM) at the Pohang Accelerator Laboratory Korea. The PAL BL2D-SMDC program [34] was used for data collection (detector distance 63 mm, omega scan; Δω = 3°, exposure time 1 s per frame), and the HKL3000sm (Ver. 703r) [35] was used for cell refinement, reduction, and absorption correction. The crystal structure of the RbPbI3 was solved by the direct method with the SHELXT-2014 program [36] and refined by full-matrix least-squares calculations with the SHELXL-2014 program package [37]. 2.5. Computational Methods The calculations of DFT were done using a highly precise full-potential linearized augmented plane wave (FLAPW) method [38] with local density approximation (LDA) [39] for the exchange-correlation potential. Muffin-tin radii of 3.0 a.u. (Rb), 2.50 a.u. (Sn and Pb), and 2.35 a.u. (I) were used with cutoffs for basis expansion and charge, potential representation of 14.82 and 144 Ry, respectively. Summation in reciprocal space was performed with 20 k points in the irreducible wedge. All atomic positions were fully relaxed with a force criteria of 1 × 10-3 eV/Å. The convergence criteria in our calculations are 10-4 htr/ (a.u.)3 in all cases; the difference in the charge density for two successive iterations is less than that value. There are two means of determining the convergence criteria: (i) the charge density difference, 49

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20 18 16 14 12 10 8 6 4 2 0

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Fig. 7. The performance comparison of 10 MAPbI3 cells and 10 RbPbI3 cells for the device stability test before and after exposure to a relative humidity level of 40% at 25 °C for 60 days.

Meanwhile, in the result of the photoluminescence spectra of TiO2/ RbPbI3 films, the energy of the TiO2 band gap (3.2 eV) was very close to the emission band of RbPbI3 at 366 nm. Hence, the RbPbI3 emission band at 366 nm was removed by the absorption of TiO2. It was also observed that the intensities of the 460 nm and 469 nm emission peaks were sharply decreased compared to those of the RbPbI3 film. Moreover, the excitation bands (at 245 nm and 320 nm) of RbPbI3 were linked to the TiO2 absorption region; these peaks disappeared in the TiO2/RbPbI3 film. These results suggest that the charge transfer effectively occurred between the TiO2 and RbPbI3 perovskite layers upon exposure to incident light. The determination of the electronic energy level, the conduction band and the valence band, of RbPbI3 for the adjacent layer allowed us to provide guidelines for the efficient charge transfer from the RbPbI3 perovskite to the TiO2 layer. To clarify these parameters, UPS was carried out in the secondary electron cutoff and HOMO regions for the FTO/RbPbI3 and FTO/TiO2/RbPbI3 electrode structures. Herein, all energies are referenced to an Au Fermi level of 0 eV. Fig. 4a shows the UPS spectra of the FTO/RbPbI3 and FTO/TiO2/RbPbI3 layers. The work function (Φ) was determined from the difference between the photon energy (the He I discharge at hν = 21.2 eV) and the secondary electron cutoff position according to the relationship Φ = hν − [Ecutoff − EFermi]. The Φ value of the RbPbI3/TiO2/FTO substrate was 3.60 eV, while that of RbPbI3/FTO was 4.27 eV due to the conducting electrode effect (Fig. 4b). The HOMO level alignment of the FTO/RbPbI3 and FTO/TiO2/RbPbI3 electrode structures was determined from the onset

Remarkably, VBM was dominated by I p states, whereas CBM was dominated by Pb p states. Previous reports [25,27] used band gap correction in the calculations – such as the GW method or an improved exchange-correlation potential (Tran-Blaha formulation). As such, our method underestimates band gap–theoretical spectra as systematically shifted with respect to the experimental findings. Therefore, the band gap of RbPbI3 as calculated here was 2.27 eV, but the experimental result was 2.64 eV. We are very well aware that the GW method significantly improves band gaps but that it also requires a substantial amount of computational time. The Tran-Blaha formulation is a fairly recent approach, and as such its foundation or validity in terms of many-body physics requires testing. Nonetheless, it has been improved in relation to many materials, with virtually identical computational costs. Furthermore, there are empirical fitting parameters in the TranBlaha formulation, and it underestimates the band gap in some materials. For this reason, we chose a conventional method despite the underestimation of the band gap, as systematic shifts with respect to experiments are well known. Fig. 3 shows each case of the fluorescence spectra for the excitation and emission of RbPbI3 and TiO2/RbPbI3 films on glass substrates. The photogenerated excitation peaks of RbPb3 were observed at 245 nm and 320 nm. This large amount of exciton binding energy is attributable to the exciton motion in the quasi-two-dimensional octahedral network of PbI6. The emission peaks were found to be at 366 nm, 460 nm and 469 nm, dominated by spontaneous radiative emission, because the origin of the photoluminescence was the intrinsic RbPbI3 properties. 50

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0.01 2

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Light Intensity (W/m 2)

Fig. 8. (a) IMVS and (b) IMPS responses for the RbPbI3 perovskite solar cell at 430 nm wavelength illumination. The hole transporting layer was deposited on RbPbI3 layer using a spincoating solution with 72.3 mg spiro-MeOTAD in 1 mL of chlorobenzene containing TBP and Li-TSFI additives. (c) Light intensity (at 430 nm) dependence of fmin and the electron lifetime from the IMVS experiment and (d) electron transit time and diffusion constant for the IMPS results according to the variation of light intensity at 430 nm.

can also be used to determine p-type or n-type semiconductors. The Hall effect arises from the current and consists of charge carriers such as electrons, holes, ions, or all three. The Hall coefficient is defined as RH = Ey/jx Bz = VH t/IB, where Ey is the induced electric field and jx is the current density of the carrier electron. This relationship can be expressed with the carrier density (n), as RH = −1/ne. The mobility (μ) is calculated from the carrier density by the μ = σ/ne equation, where σ is the conductivity. The result of the Hall measurement showed that the electric current in RbPbI3 was carried by moving electrons because RbPbI3 had the features of an ntype semiconductor. The calculated carrier density and mobility of RbPbI3 were 6.24 × 109 cm-3 and 9.11 × 103–1.18 × 104 cm2 V-1 s-1, respectively. Typically, the reported charge mobility of an organic-inorganic hybrid perovskite (CH3NH3PbI3−xClx or CH3NH3PbI3) is less than 1–11.6 cm2 V-1s1 [46]. Bi et al. increased the mobility of perovskite using a low-temperature thermal annealing method, but their result was only 36.0 cm2 V-1s-1 [47]. This suggests that RbPbI3 has more favorable properties compared to the organic–inorganic hybrid perovskite to transport electrons as a light absorber in solar cells. We applied RbPbI3 to a solar cell to determine the performance of the 1D perovskite. A focused ion beam SEM cross-sectional image of the device exhibited a uniform, well-separated layer composed of TiO2, RbPbI3 perovskite and a hole-transport layer (Fig. 5a). The RbPbI3 perovskite layer had a thickness of ~ 560 nm with orthorhombic

values of the UPS spectra in the low-binding-energy region. Fig. 4c shows that the HOMO level of the RbPbI3 layer on the TiO2/FTP substrate was 1.93 eV below the Fermi level (EF), while that of the RbPbI3 layer on the FTO substrate was closer to the EF value (0.92 eV). Based on these results, while the position of the VBM was determined from the sum of the HOMO level and the work function value, the position of the CBM of perovskite was calculated by subtracting the optical band gap (Fig. 1c) from the VBM [45]. The ionization energy (IE) was calculated from the sum of the Φ value and the HOMO level. Both films exhibit an IE of 5.19 eV for the RbPbI3/FTO substrate and 5.52 eV for the RbPbI3/ TiO2/FTO substrate. The electron affinity was calculated from the difference between the Φ value and the CBM position; these values were 1.06 eV for the RbPbI3/FTO film and 0.05 eV for the RbPbI3/TiO2/FTO film. These results indicate that the conducting electrode induced an interface dipole moment in the RbPbI3 layer, which resulted in an increase in the charge extraction. Hence, increased conductivity in the electron transporting layer can cause a decrease in the charge barrier and a loss of the open-circuit voltage. Fig. 4d shows that the band structure of the orthorhombic RbPbI3 perovskite is well suited for device integration when using TiO2 as an electron-transport material. Before we fabricated the solar cell, we measured the carrier density of the RbPbI3 perovskite via Hall measurements. The Hall effect is a very useful means of measuring either the carrier density or a magnetic field. It

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crystals across the film thickness. Fig. 5b shows a SEM image of an RbPbI3 perovskite film sample prepared from a 30 wt% solution in DMF. The RbPbI3 crystal yielded shiny yellow crystalline materials. It was observed that the yellow phase of RbPbI3 upon exposure to ambient conditions remained a thermodynamically stable phase at room temperature. To elucidate the degree of RbPbI3 perovskite penetration into the mesoporous TiO2 layer, energy-dispersive X-ray spectrometry (EDAX) mapping of the elements (Rb, Pb and I), as shown in Fig. 5c, was conducted on the cross-section of the RbPbI3/TiO2 layer. Fig. 5d depicts a typical EDAX spectrum of the yellow rectangle of Fig. 5c. The stoichiometric atomic ratio between Rb, Pb and I is nearly equal to the theoretical composition of RbPbI3 (inset in Fig. 5d). It is important to note that most of the TiO2 layer was covered by the RbPbI3 perovskite layer, indicating the penetration of RbPbI3 into the meso-TiO2 scaffold after the perovskite crystal had formed. Fig. 6a shows the forward-to-reverse sweeping J-V response under 1.5AM solar intensity. In the forward sweeping direction, the device exhibits an open-circuit voltage (Voc) of 0.61 V, photocurrent density (Jsc) of 3.53 mA/cm2, a fill factor (FF) of 44.30%, and PCE of 0.96%. On the other hand, the corresponding values from a reverse direction scan were 0.62 V, 3.75 mA/cm2, 44.60% and 1.04%, respectively. In perovskites that utilize M (cation)-X (anion) frameworks as the inorganic component, the exciton band shifts to a higher energy level as the dimension of the inorganic framework material is reduced from 3D to 1D [48]. Therefore, an inherent limitation of 1D RbPbI3 perovskite is the weak photocurrent arising from their limited optical absorption. Although the band alignment of RbPbI3 is well matched with regard to the CBM and VBM positions of TiO2, the low degree of connectivity in the one-dimensional RbPbI3 crystal structure limits the band structure, resulting in the flat band shown in Fig. 2(a). This yields low Jsc and Voc compared to those of a three-dimensional perovskite such as MAPbI3. However, since the observed Jsc and Voc values from the RbPbI3 perovskite reflected the light harvesting as a light absorber due to the high electron carrier density, we expect that the device performance will be improved by further optimization processes. Kumar et al. [23] demonstrated a photovoltaic device using CsSnI3 as a light absorber. They added SnF2 to CsSnI3 to control the Sn vacancies of CnSnI3, resulting in increased photocurrent densities of more than 22 mA/cm-2. However, in contrast to the excellent photocurrents, the highest obtained opencircuit voltage (Voc) was only 240 mV due to the presence of defects and a high background carrier density within the CsSnI3. They also attempted to increase the Voc by the substituting bromine into the lattice of CsSnI3 [49], but the Voc values remained in the range of 120–410 mV. In addition, as noted in the introduction part, because CsSnI3 is easily degraded due to humidity and oxygen in the atmosphere and becomes a non-perovskite structure (yellow phase), it loses its activity as a light absorber. Similar to CsSnI3, CsPbI3 has a desirable band gap of 1.77 eV for tandem- or single-junction solar cell applications. However, it forms a yellow phase, insulating the orthorhombic non-perovskite phase (denoted as the δo-phase), at room temperature. Photoactive α-CsPbI3 (cubic perovskite structure) with a band gap of 1.77 eV is usually attained at a temperature higher than 300 °C. The high temperature required for the formation of the cubic (perovskite) phase implies the instability of CsPbI3 at room temperature. The αCsPbI3 perovskite solar cells yielded an efficiency rate of 8.4% [50]. However, the cell rapidly degraded in less than 1 h in an ambient atmosphere. Partial substitution of the larger iodide with the smaller bromide in the form of CsPb(BrxI1−x)3 improved the optical and thermal stability at room temperature, leading to an increase in the photoconversion efficiency to 6.5% [17]. However, the efficiency attained with the inorganic perovskite remained limited to less than 10%. For the organic-inorganic perovskite MA(FA)PbI3, despite certified efficiency already exceeding 20%, due to the hygroscopic nature of the organic cations, it is easily degraded by humidity. Therefore, the instability of perovskite materials remains as a bottleneck limiting the

viability of perovskite solar cells on a large scale [51]. In contrast to these perovskites, the RbPbI3 perovskite exhibits high stability in an ambient atmosphere and a high Voc value in solar cell devices despite the fact that they show low photocurrent density due to the 1D structure. Thus, further improvements of the performance of RbPbI3 perovskite solar cells can be expected by the modulation of the 1D RbPbI3 crystal structure. We additionally observed that the device with the RbPbI3 perovskite shows little hysteresis with respect to the J-V scan direction (Fig. 6a). It is generally thought that hysteresis arises due to the compensation of the applied bias during the charge-discharge process [52] or as a result of the formation of polarized domains at the interface [53]. However, the hysteresis in the RbPbI3 solar cell is not likely ferroelectric in other perovskite solar cells given the absence of a polar space group. Therefore, we expect that it is possible to remove the hysteresis of RbPbI3 perovskite solar cells with passivation charge traps at the interface or by improving the crystallized RbPbI3 perovskite morphology. The Mott-Schottky (C-2–ϕ) analysis has commonly been used to determine the dopant density and flat band potential between the two semiconductor junction sides. The C-2–ϕ behavior can be explained in terms of the relative voltage drop at each side of the junction at V = 0, which was VbiRbPbI3/VbiTiO2 = NTiO2·εTiO2/NRbPbI3·εRbPbI3 [54]. The negative-slope region indicates p-type semiconducting properties, whereas the positive-slope region indicates n-type semiconducting behavior. In the FTO/TiO2/RbPbI3/spiro-MeOTAD/Au cell configuration, the electron acceptor density (NA) was calculated from the FTO/TiO2 interface according to the relationship C-2 = 2(Vbi-V)/qNTiO2·εTiO2 (Fig. 6b). The calculated NA value was 5.64 × 1012 cm-3. This information indicated that the electron donor density (ND) was related to the RbPbI3 layer under fully depleted conditions. ND in the RbPbI3 perovskite layer was found to be approximately 5.49 × 1013 cm-3, i.e., less than NA, because carrier extraction was promoted by the spiroMeOTAD layer [55]. We provided data to support the stability of RbPbI3 perovskite solar cells through a comparison with MAPbI3 solar cells. First, we reveal that we created solar cells in air. Generally, because perovskite is very sensitive to the humidity and oxygen in air, perovskite solar cells are created in an inert atmosphere, making them ultimately unsuitable for commercialization. As can be observed in Fig. 7, although the MAPbI3 device showed low efficiency compared to those in the literature, its initial photovoltaic performance appeared to be satisfactory. The device fabricated with MAPbI3 exhibited PCE rates of 8.23–8.34% with an open-circuit voltage (Voc) of 0.74–0.75 V, a short-circuit current density (Jsc) of 17.68–17.69 mA/cm2 and a fill factor of 61–63%. However, when we compared the MAPbI3 cell to the RbPbI3 cell at a relative humidity of 40% at 25 °C for 60 days, better moisture stability was observed for the RbPbI3 cell. In fact, MAPbI3 solar cells show efficiency rates of 0.42–0.62% after 60 days, making it appear that their functionality is mostly lost. However, while the efficiency of the RbPbI3 cell is low, but it can be said that the stability is high because it mostly maintains its initial efficiency, photocurrent density, photovoltage and fill factor values for long periods of time. Previous reports [28,29] suggest that when Rb is incorporated into a precursor perovskite, it can decrease the numbers of surface and bulk defects, thus suppressing the major sites of non-radiative recombination and resulting in a highly stable and efficient solar cell. We expect that although the RbPbI3 devices here are not highly efficient, with further optimization it is likely that they can serve as stable high-efficiency solar cells. To understand the charge transport and electron-hole recombination characteristics of illuminated RbPbI3 perovskite solar cells, the IMVS and IMPS values were measured in light with a wavelength of 430 nm. The photon intensity ranged from 4 W m-2 to 12 W m-2, as displayed in Fig. 8. IMPS measures the photoresponse of a cell to sinusoidal perturbations of the intensity of light. This measure provides information about the dynamics of the charge transport and diffusion constant process under short-circuit conditions, whereas the IMVS 52

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ray crystallography experiment at PLS-II BL2D-SMC beamline was supported in part by MSIP and POSTECH.

measurement determines the electron lifetime under open-circuit conditions because it measures the periodic modulation of the photovoltage between the Fermi level in the dark and the quasi-Fermi level under illumination [56]. Fig. 8 shows typical complex plane plots for the IMVS (Fig. 8a) and IMPS (Fig. 8b) responses at frequencies ranging from 100 mHz to 1 kHz. Given that both the photocurrent and photovoltage had a lag in the responses to illumination, their responses are displayed as a positive real and a negative imaginary part. It was found that the variation of the light intensity did not change the shapes of the IMPS and IMVS plots. If there is a restriction of electrons or carrier diffusion depending on the intensity of light, additional relaxation phenomena appear in the curves of IMPS or IMVS. These phenomena correspond to decaying current transients under pulsed illumination [57]. Therefore, in this experiment, because the shapes of the IMPS and IMVS curves are normal, there is no limitation on the carrier diffusion. However, upon the initial increase in the light intensity, the minimum frequencies in the IMPS and IMVS complex plane increased and then decreased with further increases in the light intensity. From both experiments, the electron lifetime was calculated from the IMVS plot by the equation τn = (2πfmin(IMVS))−1. The mean transit time of photogenerated electrons was obtained from the equation τd = (2πfmin(IMPS))−1. fmin(IMVS) and fmin(IMPS) were determined from the frequency when the phase angles became 45o in Bode plots. Under homogeneous light conditions with a short-circuit condition, the electron diffusion coefficient can be estimated from the equation Dn = d2/ 4τd, where d is the thickness of the TiO2 electron-transport layer [56]. Fig. 8c shows the electron lifetime as obtained from the IMVS results under different light intensities. The electron lifetime decreased with an increase in the light intensity due to the high kinetic energy. The electron transit time increased with the light intensity because it accelerated electron transport into the TiO2 layer at a high light intensity level, whereas the electron diffusion length showed linear dependence on the incident photon flux (Fig. 8d). Under illumination, because the quasi-Fermi level moved toward the conduction band and the deep traps were filled with electrons, the transport of the electrons was not limited and the effective diffusion rate increased. This resulted in an increase in the diffusion coefficient to the light intensity.

Appendix A. Supporting information Supplementary data associated with this article can be found in the online version at http://dx.doi.org/10.1016/j.solmat.2017.07.011. References [1] A. Kojima, K. Teshima, Y. Shirai, T. Miyasaka, Organometal halide perovskites as visible-light sensitizers for photovoltaic cells, J. Am. Chem. Soc. 131 (2009) 6050–6051. [2] W.S. Yang, J.H. Noh, N.J. Jeon, Y.C. Kim, S. Ryu, J. Seo, S.I. Seok, High-performance photovoltaic perovskite layers fabricated through intramolecular exchange, Science 348 (2015) 1234–1237. [3] G. Xing, N. Mathews, S.S. Lim, N. Yantara, X. Liu, D. Sabba, M. Grätzel, S. Mhaisalkar, T.C. Sum, Low-temperature solution-processed wavelength-tunable perovskites for lasing, Nat. Mater. 13 (2014) 476–480. [4] J.H. Noh, S.H. Im, J.H. Heo, T.N. Mandal, S.I. Seok, Chemical management for colorful, efficient, and stable inorganic–organic hybrid nanostructured solar cells, Nano Lett. 13 (2013) 1764–1769. [5] S.D. Stranks, G.E. Eperon, G. Grancini, C. Menelaou, M.J.P. Alcocer, T. Leijtens, L.M. Herz, A. Petrozza, H.J. Snaith, Electron-hole diffusion lengths exceeding 1 μm in an organometal trihalide perovskite absorber, Science 342 (2013) 341–344. [6] G. Xing, N. Mathews, S. Sun, S.S. Lim, Y.M. Lam, M. Grätzel, S. Mhaisalkar, T.C. Sum, Long-Range balanced electron- and hole-transport lengths in organicinorganic CH3NH3PbI3, Science 342 (2013) 344–347. [7] Q. Lin, A. Armin, R.C.R. Nagiri, P.L. Burn, P. Meredith, Electro-optics of perovskite solar cells, Nat. Photonics 9 (2015) 106–112. [8] A. Miyata, A. Mitioglu, P. Plochocka, O. Portugall, J.T.-W. Wang, S.D. Stranks, H.J. Snaith, R.J. Nicholas, Direct measurement of the exciton binding energy and effective masses for charge carriers in organic-inorganic tri-halide perovskites, Nat. Phys. 11 (2015) 582–587. [9] M.M. Lee, J. Teuscher, T. Miyasaka, T.N. Murakami, H.J. Snaith, Efficient hybrid solar cells based on meso-superstructured organometal halide perovskites, Science 338 (2012) 643–647. [10] N.J. Jeon, J.H. Noh, W.S. Yang, Y.C. Kim, S. Ryu, J. Seo, S.I. Seok, Compositional engineering of perovskite materials for high-performance solar cells, Nature 517 (2015) 476–480. [11] S. Albrecht, M. Saliba, J.P. Correa Baena, F. Lang, L. Kegelmann, M. Mews, L. Steier, A. Abate, J. Rappich, L. Korte, R. Schlatmann, M.K. Nazeeruddin, A. Hagfeldt, M. Gratzel, B. Rech, Monolithic perovskite/silicon-heterojunction tandem solar cells processed at low temperature, Energy Environ. Sci. 9 (2016) 81–88. [12] M. Kaltenbrunner, G. Adam, E.D. Glowacki, M. Drack, R. Schwodiauer, L. Leonat, D.H. Apaydin, H. Groiss, M.C. Scharber, M.S. White, N.S. Sariciftci, S. Bauer, Flexible high power-per-weight perovskite solar cells with chromium oxide-metal contacts for improved stability in air, Nat. Mater. 14 (2015) 1032–1039. [13] H. Tsai, W. Nie, J.-C. Blancon, C.C. Stoumpos, R. Asadpour, B. Harutyunyan, A.J. Neukirch, R. Verduzco, J.J. Crochet, S. Tretiak, L. Pedesseau, J. Even, M.A. Alam, G. Gupta, J. Lou, P.M. Ajayan, M.J. Bedzyk, M.G. Kanatzidis, A.D. Mohite, High-efficiency two-dimensional Ruddlesden–Popper perovskite solar cells, Nature 536 (2016) 312–316. [14] H.Y. Ye, W.Q. Liao, C.L. Hu, Y. Zhang, Y.M. You, J.G. Mao, P.F. Li, R.G. Xiong, Band gap engineering of lead-halide perovskite-type ferroelectrics, Adv. Mater. 28 (2016) 2579–2586. [15] G.E. Eperon, G.M. Paterno, R.J. Sutton, A. Zampetti, A.A. Haghighirad, F. Cacialli, H.J. Snaith, Inorganic caesium lead iodide perovskite solar cells, J. Mater. Chem. A. 3 (2015) 19688–19695. [16] S. Dastidar, D.A. Egger, L.Z. Tan, S.B. Cromer, A.D. Dillon, S. Liu, L. Kronik, A.M. Rappe, A.T. Fafarman, High chloride doping levels stabilize the perovskite phase of cesium lead iodide, Nano Lett. 16 (2016) 3563–3570. [17] R.E. Beal, D.J. Slotcavage, T. Leijtens, A.R. Bowring, R.A. Belisle, W.H. Nguyen, G.F. Burkhard, E.T. Hoke, M.D. McGehee, Cesium lead halide perovskites with improved stability for tandem solar cells, J. Phys. Chem. Lett. 7 (2016) 746–751. [18] C.C. Stoumpos, C.D. Malliakas, M.G. Kanatzidis, Semiconducting tin and lead iodide perovskites with organic cations: phase transitions, high mobilities, and near-infrared photoluminescent properties, Inorg. Chem. 52 (2013) 9019–9038. [19] K. Shum, Z. Chen, J. Qureshi, C. Yu, J.J. Wang, W. Pfenninger, N. Vockic, J. Midgley, J.T. Kenney, Synthesis and characterization of CsSnI3 thin films, Appl. Phys. Lett. 96 (2010) 221903. [20] Z. Chen, C. Yu, K. Shum, J.J. Wang, W. Pfenninger, N. Vockic, J. Midgley, J.T. Kenney, Photoluminescence study of polycrystalline CsSnI3 thin films: determination of exciton binding energy, J. Lumin. 132 (2012) 345–349. [21] I. Chung, J.-H. Song, J. Im, J. Androulakis, C.D. Malliakas, H. Li, A.J. Freeman, J.T. Kenney, M.G. Kanatzidis, CsSnI3: semiconductor or metal? High electrical conductivity and strong near-infrared photoluminescence from a single material. high hole mobility and phase-transitions, J. Am. Chem. Soc. 134 (2012) 8579–8587. [22] Z. Chen, J.J. Wang, Y. Ren, C. Yu, K. Shum, Schottky solar cells based on CsSnI3 thin-films, Appl. Phys. Lett. 101 (2012) 093901. [23] M.H. Kumar, S. Dharani, W.L. Leong, P.P. Boix, R.R. Prabhakar, T. Baikie, C. Shi, H. Ding, R. Ramesh, M. Asta, M. Graetzel, S.G. Mhaisalkar, N. Mathews, Lead-free halide perovskite solar cells with high photocurrents realized through vacancy

4. Conclusions In summary, we fabricated 1D inorganic RbPbI3 perovskite solar cells at room temperature under ambient conditions. This study revealed that the band gap alignment of RbPbI3 perovskite is well matched with that of TiO2 for efficient electron transport in a solar cell. We calculated the band structure of RbPbI3 with density functional calculations and found that the valence band maximum occurred near the Γ point, while the conduction band minimum was located near X. Although the RbPbI3 perovskite solar cell does not perform as well as other perovskite solar cells due to its flat band structure, further improvements of the PCE can be expected if the RbPbI3 crystal modulation process is realized. To understand the charge transport and electronhole recombination characteristics of illuminated RbPbI3 perovskite solar cells, IMVS and IMPS measurements were conducted under light at a wavelength of 430 nm. The electron lifetime and transit time decreased with an increase in the light intensity due to the high kinetic energy, whereas the electron diffusion length demonstrated linear dependence on the intensity of the incident light. Acknowledgment This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (2016R1D1A1B04931751). This research was supported by Nano Material Technology Development Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning (2009-0082580). The X53

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M.-H. Jung et al.

[41] O.N. Yunakova, V.K. Miloslavskii, E.N. Kovalenko, E.V. Ksenofontova, The absorption spectra of thin films of ternary compounds in the RbI–PbI2 system, Low. Temp. Phys. 38 (2012) 943–947. [42] R.O. Jones, Density functional theory: its origins, rise to prominence, and future, Rev. Mod. Phys. 87 (2015) 897–923. [43] J. Kohanoff, Electronic structure calculations for solids and molecules: theory and computational methods, 2006. [44] R.M. Martin, Electronic Structure: Basic Theory and Practical Methods, Cambridge University Press, 2008. [45] P. Schulz, E. Edri, S. Kirmayer, G. Hodes, D. Cahen, A. Kahn, Interface energetics in organo-metal halide perovskite-based photovoltaic cells, Energy Environ. Sci. 7 (2014) 1377–1381. [46] C. Wehrenfennig, G.E. Eperon, M.B. Johnston, H.J. Snaith, L.M. Herz, High charge carrier mobilities and lifetimes in organolead trihalide perovskites, Adv. Mater. 26 (2014) 1584–1589. [47] C. Bi, Y. Shao, Y. Yuan, Z. Xiao, C. Wang, Y. Gao, J. Huang, Understanding the formation and evolution of interdiffusion grown organolead halide perovskite thin films by thermal annealing, J. Mater. Chem. A. 2 (2014) 18508–18514. [48] S. González-Carrero, R.E. Galian, J. Pérez-Prieto, Organometal halide perovskites: bulk low-dimension materials and nanoparticles, Part. Part. Syst. Charact. 32 (2015) 709–720. [49] D. Sabba, H.K. Mulmudi, R.R. Prabhakar, T. Krishnamoorthy, T. Baikie, P.P. Boix, S. Mhaisalkar, N. Mathews, Impact of anionic Br– substitution on open circuit voltage in lead free perovskite (CsSnI3−xBrx) solar cells, J. Phys. Chem. C 119 (2015) 1763–1767. [50] Z. Li, M. Yang, J.-S. Park, S.-H. Wei, J.J. Berry, K. Zhu, Stabilizing perovskite structures by tuning tolerance factor: formation of formamidinium and cesium lead iodide, Solid-State Alloy. Chem. Mater. 28 (2016) 284–292. [51] B. Conings, J. Drijkoningen, N. Gauquelin, A. Babayigit, J. D'Haen, L. D'Olieslaeger, A. Ethirajan, J. Verbeeck, J. Manca, E. Mosconi, F. De Angelis, H.G. Boyen, Intrinsic thermal instability of methylammonium lead trihalide perovskite, Adv. Energy Mater. 5 (2015). [52] G.E. Eperon, G.M. Paternò, R.J. Sutton, A. Zampetti, A.A. Haghighirad, F. Cacialli, H.J. Snaith, Inorganic caesium lead iodide perovskite solar cells, J. Mater. Chem. A. 3 (2015) 19688–19695. [53] Y. Hou, W. Chen, D. Baran, T. Stubhan, N.A. Luechinger, B. Hartmeier, M. Richter, J. Min, S. Chen, C.O.R. Quiroz, N. Li, H. Zhang, T. Heumueller, G.J. Matt, A. Osvet, K. Forberich, Z.G. Zhang, Y. Li, B. Winter, P. Schweizer, E. Spiecker, C.J. Brabec, Overcoming the interface losses in planar heterojunction perovskite-based solar cells, Adv. Mater. (2016) 5112–5120. [54] S.M. Sze, Physics of Semiconductor Devices, John Wiley & Sons, Inc, New York, 1981. [55] A. Guerrero, E.J. Juarez-Perez, J. Bisquert, I. Mora-Sero, G. Garcia-Belmonte, Electrical field profile and doping in planar lead halide perovskite solar cells, Appl. Phys. Lett. 105 (2014) 133902. [56] J. Krüger, R. Plass, M. Grätzel, P.J. Cameron, L.M. Peter, Charge transport and back reaction in solid-state dye-sensitized solar cells: a study using intensity-modulated photovoltage and photocurrent spectroscopy, J. Phys. Chem. B. 107 (2003) 7536–7539. [57] M.-H. Jung, Polypyrrole/poly(vinyl alcohol-co-ethylene) quasi-solid gel electrolyte for iodine-free dye-sensitized solar cells, J. Power Sources 268 (2014) 557–564.

modulation, Adv. Mater. 26 (2014) 7122–7127. [24] S.F. Koji Yamada, Hiromi Horimoto, Takashi Matsui, Tsutomu Okuda, Sumio Ichiba, Structural phase transitions of the polymorphs of CsSnI3 by means of Rietveld analysis of the X-ray diffraction, Chem. Lett. (1991) 801–804. [25] J. Brgoch, A.J. Lehner, M. Chabinyc, R. Seshadri, Ab initio calculations of band gaps and absolute band positions of polymorphs of RbPbI3 and CsPbI3: implications for main-group halide perovskite photovoltaics, J. Phys. Chem. C 118 (2014) 27721–27727. [26] J. Young, J.M. Rondinelli, Octahedral rotation preferences in perovskite iodides and bromides, J. Phys. Chem. Lett. 7 (2016) 918–922. [27] R.A. Jishi, O.B. Ta, A.A. Sharif, Modeling of lead halide perovskites for photovoltaic applications, J. Phys. Chem. C 118 (2014) 28344–28349. [28] M. Saliba, T. Matsui, K. Domanski, J.Y. Seo, A. Ummadisingu, S.M. Zakeeruddin, J.P. Correa-Baena, W.R. Tress, A. Abate, A. Hagfeldt, M. Grätzel, Incorporation of rubidium cations into perovskite solar cells improves photovoltaic performance, Science 354 (2016) 206–209. [29] T. Duong, H.K. Mulmudi, H. Shen, Y. Wu, C. Barugkin, Y.O. Mayon, H.T. Nguyen, D. Macdonald, J. Peng, M. Lockrey, W. Li, Y.-B. Cheng, T.P. White, K. Weber, K. Catchpole, Structural engineering using rubidium iodide as a dopant under excess lead iodide conditions for high efficiency and stable perovskites, Nano Energy 30 (2016) 330–340. [30] A.M.A. Leguy, P. Azarhoosh, M.I. Alonso, M. Campoy-Quiles, O.J. Weber, J. Yao, D. Bryant, M.T. Weller, J. Nelson, A. Walsh, M. van Schilfgaarde, P.R.F. Barnes, Experimental and theoretical optical properties of methylammonium lead halide perovskites, Nanoscale 8 (2016) 6317–6327. [31] D.E. Scaife, P.F. Weller, W.G. Fisher, Crystal preparation and properties of cesium tin(II) trihalides, J. Solid State Chem. 9 (1974) 308–314. [32] H.-S. Kim, J.-W. Lee, N. Yantara, P.P. Boix, S.A. Kulkarni, S. Mhaisalkar, M. Grätzel, N.-G. Park, High efficiency solid-state sensitized solar cell-based on submicrometer rutile TiO2 nanorod and CH3NH3PbI3 perovskite sensitizer, Nano Lett. 13 (2013) 2412–2417. [33] L.F. Gate, Comparison of the photon diffusion model and Kubelka-Munk equation with the exact solution of the radiative transport equation, Appl. Opt. 13 (1974) 236–238. [34] J.W. Shin, K. Eom, D. Moon, BL2D-SMC, the supramolecular crystallography beamline at the Pohang Light Source II, Korea, J. Synchrotron Rad. 23 (2016) 369–373. [35] Z. Otwinowski, W. Minor, In Methods in Enzymology, Academic Press, New York, 1997. [36] G.M. Sheldrick, SHELXT – Integrated space-group and crystal-structure determination Acta Crystallogr, Sect. A: Found. Crystallogr. 71 (2015) 3–8. [37] G.M. Sheldrick, Crystal structure refinement with SHELXL Acta Crystallogr, Sect. C: Cryst. Struct. Commun. 71 (2015) 3–8. [38] E. Wimmer, H. Krakauer, M. Weinert, A.J. Freeman, Full-potential self-consistent linearized-augmented-plane-wave method for calculating the electronic structure of molecules and surfaces: O2 molecule, Phys. Rev. B: Condens. Matter 24 (1981) 864–875. [39] L. Hedin, B.I. Lundqvist, Explicit local exchange-correlation potentials, J. Phys. C: Solid State Phys. 4 (1971) 2064–2083. [40] D.M. Trots, S.V. Myagkota, High-temperature structural evolution of caesium and rubidium triiodoplumbates, J. Phys. Chem. Solids 69 (2008) 2520–2526.

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