Coexistence of Two Electronic Nano-Phases on a CH3NH3PbI3âxClx

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Coexistence of Two Electronic Nano-Phases on a CH3NH3PbI3−xClx Surface Observed in STM Measurements Andrew J. Yost,† Artem Pimachev,† Chun-Chih Ho,‡,§ Seth B. Darling,§,∥ Leeyih Wang,⊥ Wei-Fang Su,‡ Yuri Dahnovsky,† and TeYu Chien*,† †

Department of Physics and Astronomy, University of Wyoming, Laramie, Wyoming 82071, United States Department of Materials Science and Engineering, National Taiwan University, Taipei 10617, Taiwan § Center for Nanoscale Materials, Argonne National Laboratory, Lemont, Illinois 60439, United States ∥ Institute for Molecular Engineering, University of Chicago, Chicago, Illinois 60637, United States ⊥ Center for Condensed Matter Sciences, National Taiwan University, Taipei 10617, Taiwan ‡

S Supporting Information *

ABSTRACT: Scanning tunneling microscopy is utilized to investigate the local density of states of a CH3NH3PbI3−xClx perovskite in cross-sectional geometry. Two electronic phases, 10−20 nm in size, with different electronic properties inside the CH3NH3PbI3−xClx perovskite layer are observed by the dI/ dV mapping and point spectra. A power law dependence of the dI/dV point spectra is revealed. In addition, the distinct electronic phases are found to have preferential orientations close to the normal direction of the film surface. Density functional theory calculations indicate that the observed electronic phases are associated with local deviation of I/Cl ratio, rather than different orientations of the electric dipole moments in the ferroelectric phases. By comparing the calculated results with experimental data we conclude that phase A (lower contrast in dI/dV mapping at −2.0 V bias) contains a lower I/Cl ratio than that in phase B (higher contrast in dI/dV). KEYWORDS: cross-sectional scanning tunneling microscopy, organometallic halide perovskite, electronic phase, mixed halide, density functional theory, nano-phase

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INTRODUCTION Organometallic halide perovskites are drawing substantial attention because of their extraordinary power conversion efficiency (PCE) performance and rapid energy payback time in solar cell applications.1−11 The PCE of organometallic halide perovskite solar cells increased dramatically from 3% in 2009 to 21.1% in 2016.2,12 Although millimeter- to inch-sized single crystals were successfully synthesized and long exciton lifetimes were reported,13−15 details of the underlying mechanism behind this impressive device performance are still unclear. Interestingly, high PCE values have been reported for various types of synthesis methods with various compositions of the mixed halide perovskites, CH3NH3PbI3−xClx,16,17 and other types of mixtures, such as formamidinium lead iodide (FAPbI3) with methylammonium lead iodide (MAPbI3).18 In the mixed materials, phase separation is expected and was commonly reported in manganites.19 Indeed, for CH3NH3PbI3−xClx, it was reported that Cl incorporation in an iodide-based structure is possible only at relatively low concentration levels (3−4%).20 Similar results regarding the mixture of Cl and I phase separation was reported in Sn-based hybrid halide perovskites.21 In addition, structural inhomogeneity at the micrometer scale was reported in single crystal CH3NH3PbI3, which © 2016 American Chemical Society

also affects the optical properties in different locations in the crystal.22 To seek a strategy for better utilization of the organometallic halide perovskite materials, microscopic understanding of the materials is necessary. In particular, it would be informative to know: “Do microscopic phase separations exist in the synthesized thin f ilm or crystals when mixed halogen elements are used? If so, at what length scale?” Here, we utilize crosssectional scanning tunneling microscopy and spectroscopy (XSTM/S) along with density functional theory (DFT) calculations to study CH3NH3PbI3−xClx thin films. Electronic phase separation in the length scale of 10−20 nm is observed and is explained by the local deviation of I/Cl composition ratio.

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RESULTS AND DISCUSSION As shown in Figure 1a, CH3NH3PbI3−xClx thin film thickness is determined by scanning electron microscopy (SEM) to be ∼400 nm. In Figure 1b, the crystallinity of the Received: June 24, 2016 Accepted: October 10, 2016 Published: October 10, 2016 29110

DOI: 10.1021/acsami.6b07721 ACS Appl. Mater. Interfaces 2016, 8, 29110−29116

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ACS Applied Materials & Interfaces

XSTM/S is in a sense a similar method to prepare fresh and clean surfaces for thin film materials, details can be seen in Supporting Information. Figure 2a shows a layered textured topography of the fractured perovskite thin film measured by STM. The rms roughness of the fractured surfaces is determined to be ∼2 nm through an area analysis on the topography image (Figure 2a). This is consistent with the height profile shown in Figure 2c. Interestingly, in dI/dV mapping measured at −2.0 V (Figure 2b), two distinct contrast regions not seen in topography are observed. The contrast seen in dI/dV mapping indicates different electronic local DOS (LDOS) exist at different regions, indicating the existence of the “electronic phase separation”.32 The domain sizes of the electronic phases are determined in the dI/dV mapping (Figure 2b) to be in the range of 10−20 nm. Note that topography-induced artifacts in dI/dV mapping can be ruled out since the topography (Figure 2a) and the dI/dV mapping (Figure 2b) do not show common features. In particular, the electronic domains exhibit preferential orientation along the surface normal, while the topography textures show in-plane features. Note that Figure 2d shows a cross-sectional SEM image of the sample mounted in the XSTM sample holder. The arrows in Figure 2d indicate the surface normal direction and the scan direction, which are also plotted in Figure 2b. The precise angle of the electronic domain orientation is difficult to determine because of the irregular shapes of the individual domains. However, the domain orientations are evidently elongated in the surface normal direction. To understand the origin of the observed electronic phase separation, dI/dV spectra are measured across the domains. Figure 3a shows a zoom-in dI/dV mapping measured at −2.0 V with dimension of 20 nm × 20 nm. The dI/dV point spectra measured at point A and point B are displayed in Figure 3b. Two distinct dI/dV spectra are revealed. It is found that these dI/dV spectra could not be fitted with an exponential growth equation, which is used for typical tunneling spectra. Instead, a power law function, shown below, can describe each spectrum

Figure 1. (a) Cross-sectional view of SEM image showing perovskite thickness ∼400 nm. (b) XRD data indicates comparable quality of the perovskite film reported in literature3,17 with {110} preferred orientation.

CH3NH3PbI3−xClx thin film is revealed by X-ray diffraction (XRD) with a set of peaks at 14.19°, 28.51°, and 43.30°, comparable to the reported thin films with high solar cell performance.3,17 The preferred crystalline orientation of the thin film is in the {110} crystal orientation. It is worth noting that stoichiometric phase separation is evident in XRD, where a small peak at 15.72° associated with pure CH3NH3PbCl3 is visible.3,17 In addition, another small peak at 12.81° associated with PbI2 phase is observed. The question is what are the sizes of those stoichiometric domains? And are there any other types of phase separation in the film due to local I or Cl segregation? To obtain deeper understanding regarding these questions, cross-sectional STM/S (XSTM/S) is utilized.23−27 XSTM/S is used to create a fresh perovskite surface for STM measurements. Other sample preparation methods, such as sputtering/ annealing, are expected to change the highly sensitive perovskite materials. Another popular method to prepare samples for STM/S measurements is the cleaving of single crystal to create a clean surface for STM measurements.28−31

Figure 2. (a) STM topography image, 100 nm × 100 nm, set-point = −2.0 V bias, 100 pA. (b) dI/dV mapping taken simultaneously with the topography. The scan direction (horizontal arrow) and the surface normal direction (angled arrow) are shown as arrows in the bottom of panel b. (c) Height profile of the topography corresponding to the dashed line shown in panel a. (d) SEM image of the top-view of the mounted sample for XSTM/S measurements. The scan and the surface normal directions are determined in SEM image. 29111


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Figure 3. (a) dI/dV mapping at −2.0 V. Image size: 20 nm × 20 nm. (b) dI/dV point spectra measured at low and high contrast regions, corresponding to point A and point B in panel a, respectively. The curves in the positive bias regions are fit with eq 1, resulting in different exponents in the function: 3.9 for A; 5.1 for B. (c) Top: The resulting exponent of the fitting function in positive bias side of the spectrum at each location. The length scale of the change of the exponent is determined to be ∼5 nm (shadowed area). Middle: dI/dV point spectra collected along the line between point A and point B with spatial interval of 0.05 nm. Bottom: The resulting exponent of the fitting function in negative bias side of the spectrum at each location. The length scale of the change of the exponent is determined to be ∼1 nm (shadowed area). (d) The fitting results of V0 for both positive and negative bias regions of each spectrum.

the perovskite system seems to have no relation to either of the reported mechanisms. Further experiments are needed to elucidate the underlying physics of the observed power law dependent dI/dV versus V spectra in these perovskite materials. It is also worth noting that the measured dI/dV spectra are different from those reported on a different but similar material in the literaturevacuum-synthesized thin CH3NH3PbI3 layers on Au(111) surfaces.35 The similarity in the dI/dV spectra indicates that the two domains are similar materials. This implies that the two domains are not related to pure CH3NH3PbCl3 or PbI2, which would exhibit significantly different band gaps and dI/dV spectra. Note that we did not observe any other types of dI/dV spectra in a large area of measurements. This means that the amounts of CH3NH3PbCl3 or PbI2 in the CH3NH3PbI3−xClx thin film are too small for easy detection in the STM measurements or they are clustered somewhere outside of the measuring range. This power law behavior with V0 = 0 implies that the difference among the dI/dV spectra in the two electronic phases is most likely due to a small change in the electronic LDOS or the tunneling barriers. Two possible origins stem from observations and literature reports: (1) ferroelectric domains due to different electric dipole moment orientations; and (2) slight deviation in local halogen element ratio (I/Cl ratio). The former is based on the reported ferroelectric phases,36,37 which may cause a change in the tunneling barrier/ behavior; while the latter is based on the fact that the mixed halogen elements ratio might deviate slightly locally and form varying x in CH3NH3PbI3−xClx with microscopic domains, which may exhibit slightly different LDOS. Note that the existence of ferroelectric phases at room temperature in CH3NH3PbI3−xClx is still controversial, but it is well accepted that the ferroelectric phases exist at low temperature.36−46 For the ferroelectric domain scenario, electric polarization in different domains will produce a local electric field that shifts

at both positive and negative sides of the spectrum; see Supporting Information for more details on the fitting.

dI (V ) = A |V − V0|α dV

(1)

where A is the spectral weight, which is associated with the relative DOS; V is the applied tip−sample bias; V0 is the shift of the power growth law, which is associated with the effective Fermi level of the materials at the measurement location; and α is the exponent in the power law function; where the physical meaning requires further study. The fitted results along with the measured spectra are shown in Figure 3b. The exponents of the fitting function are determined to be 3.9 for spectrum A and 5.1 for spectrum B on the positive bias side. A series of dI/dV point spectra are collected across the domain boundary, along the path shown in Figure 3a. Each dI/dV spectrum is fitted with eq 1 and the resulting values of the exponents of the curves in both the positive and the negative bias regions are plotted in the top and the bottom panels of Figure 3c, respectively. For the positive bias region (the conduction band), the exponents gradually change from ∼4 at point A to ∼5 at point B with a transition length scale of ∼5 nm, which is about 10 unit cells. For the negative bias region (the valence band), the exponents change from ∼8 at point A to ∼10 at around midway toward point B. The dI/dV signal in the negative bias region dropped to zero at a position of ∼4.5 nm, which is where the fitting stopped. The length scale of the change of the exponent in the negative bias region is found to be ∼1 nm, which is about 2 unit cells. As can be seen in Figure 3d, the resulting V0 for both sides of bias are always zero, indicating the energy gaps in the two phases are similar, if not identical. It is worth noting that the power law dependence on the dI/dV versus V in tunneling spectra is rare. To the best of our knowledge, there are only two reported mechanisms that exhibit power law dependence: (1) interacted nanoparticles33 and (2) Luttinger liquid.34 However, 29112


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Figure 4. (a) Schematics of the perovskite structures with the optimized MA molecules orientations. (b) Total energy as a function of the angle between the neighboring MA molecules. (c) Surface DOS of the two MA molecules orientations. (d) Surface DOS of parallel MA molecules orientations of CH3NH3PbI2Cl and CH3NH3PbICl2. (e) Surface DOS of CH3NH3PbI3−xClx with x = 1.0, 1.1, 1.2, 1.3, 1.4, and 1.5, constructed by linear combination of the surface DOS of CH3NH3PbI2Cl and CH3NH3PbICl2.

higher. However, the experimental dI/dV spectra of the two phases show similar band gaps and alternating spectral intensity on the both sides of the Fermi energy. Thus, this eliminates MA molecular orientation as a possible mechanism for the existence of the observed electronic phases. To verify the second scenario where the two phases can be determined by different local I/Cl ratios, we calculate the surface DOS of the relaxed surfaces of CH3NH3PbI2Cl (x = 1) and CH3NH3PbICl2 (x = 2) (Figure 4d), then we construct different linear combinations of these two DOS to describe the trend of the experimental DOS of the two phases. In Figure 4d, the band gap of CH3NH3PbICl2 is ∼1.6 eV, which is larger than that of the CH3NH3PbI2Cl (∼1.3 eV). In Figure 4e, where the linear combinations of the two DOS are presented, the energy gaps remain the same while the surface DOS profile noticeably changes. By comparing experimental data with the calculated results we conclude that phase A contains lower I/Cl ratio than that in phase B. Thus, we believe the observed electronic phases are due to slight deviation of the local I/Cl ratio.

the dI/dV spectra rigidly with respect to the bias, assuming the LDOS is the same for domains with different electric polarization. The experimental dI/dV data (Figure 3) and analysis indicate that the dI/dV spectra in the two phases are not shifting. This observation rules out ferroelectric domains as the possible origin. This conclusion is also confirmed by DFT calculations, which will be described below. The scenario of slight deviation of the local I/Cl ratio is also examined with the DFT calculations. It is believed that the origin of the possible ferroelectric phase is due to the orientation of the ionized methylammonium (MA) molecular ion, (CH3NH3)+, which has an electric dipole moment of 2.0−2.3 D.47,48 To develop a deeper understanding of the ferroelectric phases, the total energy of the bulk structure is calculated as a function of the MA molecular orientation in the crystal. The most stable orientation for the MA molecules is found to be in the x−y plane, defined as the pseudocubic phase of the perovskite crystal, with C−N−x̂ angle ∼50° as shown in Figure 4a. Next, in a two-unit-cell supercell, the total energy of the crystal is calculated as a function of the angle between the neighboring MA molecules. From Figure 4b, it becomes clear that the ferroelectric phase has the lowest total energy, the global minimum. In the same figure, the antiferroelectric phase exhibits a local minimum that is higher than the minimum for the ferroelectric phase, which is consistent with reported theoretical work.48 The calculated orientation energy barrier of the molecular ion in the perovskite is ∼450 meV, which is very large compared to the thermal energy at room temperature (∼25 meV). This indicates ferroelectric phases are very stable at room temperature. In Figure 4c, we present the surface DOS for the two configurations of MA electric dipole orientations. In this calculation, the surface structures are relaxed. First, it is worth noting that the energy gaps of the surfaces for both MA molecules orientations are about 1.2−1.3 eV, which is consist with the STS measurements as shown in Figure 3b. The surface gaps are lower than the measured bulk optical gap (∼1.5 eV).3,20 Second, the antiparallel configuration exhibits a narrower band gap than that of the parallel configuration, and the DOS on both sides of the Fermi energy are distinctly

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CONCLUSION In this work, we have experimentally studied the cross-sectional surface of CH3NH3PbI2Cl thin film using STM/S. We have found the existence of two phases, A and B, with different DOS in 20 nm scale and preferential domain orientations along the surface normal in CH3NH3PbI2Cl. The dI/dV spectra are found to follow a power law explicitly for both positive and negative bias. To find the origin of these two phases, we have checked two scenarios: (1) ferroelectric domains with different electric dipole moment orientations and (2) local deviation of I/Cl ratios. On the basis of DFT solid state calculations, we have found that scenario (1) cannot explain the experimental data because the energy gaps are different for different electric dipole orientations, which is in direct contradiction to the experimental dependences where the gaps are the same for both phases. The calculated results of the scenario of local I/Cl ratio can sufficiently explain the experimental data where the energy gaps remain the same for both phases. In addition, we 29113


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ACS Applied Materials & Interfaces Notes

have determined that phase A with low dI/dV contrast measured at −2.0 V corresponds to lower I/Cl ratio than that of phase B with high dI/dV contrast.

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The authors declare no competing financial interest.

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ACKNOWLEDGMENTS T.-Y.C. and Y.D. acknowledge the U.S. Department of Energy, Office of Basic Energy Sciences, Division of Materials Sciences and Engineering for financial support (DEFG02-10ER46728) of this research. A.J.Y. acknowledges graduate fellowship support from the National Science Foundation and the University of Wyoming EE-Nanotechnology Program (DGE0948027). A.P. acknowledges the School of Energy Resources, University of Wyoming, through its Graduate Assistantship program. Use of the Center for Nanoscale Materials, an Office of Science user facility, was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. C.-C.H., L.-Y.W., and W.-F.S. acknowledge the Ministry of Science and Technology of Taiwan for financial support of this research (MOST 102-2911-I-002-504 and 104-3113-E-002-010).

METHODS

CH3NH3PbI3−xClx films were synthesized by first mixing the perovskite precursors with a 1:3 ratio of PbCl2/CH3NH3I in a solution of dimethylformamide (DMF) followed by heating to 60 °C. This ratio of precursors, which was revealed to produce CH3NH3PbI2Cl (x = 1) via energy dispersive X-ray (EDX) spectroscopy,3 has exhibited great device performance.3,17 The solution was spin-coated onto an As-doped Si(100) substrate at 1500 rpm followed by an annealing process at 90 °C for 2 h in a nitrogen atmosphere. XSTM/S samples were cleaved in situ, along a diced scribe perpendicular to the film/substrate interface to expose fresh layers of the perovskite thin film in cross-sectional geometry. STM/S measurements were performed at a temperature of 50 K on an Omicron variable temperature STM (VT-STM) at −2.0 V bias, 100 pA set-point, unless otherwise noted. A tungsten tip was prepared by chemical etching in a NaOH solution. X-ray diffraction (XRD) data was collected using a Rigaku XRD system at room temperature. Scanning electron microscope (SEM) images were recorded on a FEI Quanta FEG 450 field emission scanning electron microscope with 20 kV accelerating voltage. No additional conductive coating was applied to samples prior to SEM measurements. The solid state geometries of hybrid halide perovskites in cubic and tetragonal configurations were optimized using the DFT formalism with PBE exchange-correlation functional and custom basis sets in the GAUSSIAN09.49 The basis sets assigned to the atoms were SBKJC, 631G(d), and 6-311G for Pb, CH3NH3, and halogens, respectively.50 The electronic properties were studied with the Vienna Ab-initio Simulation Package (VASP)51−54 using the DFT method with the PBE exchange correlation functional.55,56 The projector augmented wave (PAW) pseudopotential57,58 for DFT was used with an energy cutoff of 400 eV for the plane wave basis functions. For the hybrid halide perovskite, the 2 × 1 × 1 and 5 × 1 × 1 supercells were used for the bulk and film structures, respectively. The density of states (DOS) plots were constructed with a 21 × 21 × 21 k-points mesh generated according to the Monkhorst−Pack scheme.59,60 In this work, we used a 4 × 4 × 4 k-points mesh with 0.05 eV Gaussian broadening for the geometry optimization. The single electron configurations 5d106s26p2, 5s25p5, 2s22p3, 2s22p2, and 1s1 were treated as the valence shells for Pb, I, N, C, and H atoms, respectively.

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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/acsami.6b07721.

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REFERENCES

(1) Abate, A.; Saliba, M.; Hollman, D. J.; Stranks, S. D.; Avolio, R.; Petrozza, A.; Snaith, J.; et al. Supramolecular Halogen Bond Passivation of Organometal-Halide Perovskite Solar Cells. Nano Lett. 2014, 14, 3247−3254. (2) 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. (3) Lee, M. M.; Teuscher, J.; Miyasaka, T.; Murakami, T. N.; Snaith, H. J. Efficient Hybrid Solar Cells Based on Meso-Superstructured Organometal Halide Perovskites. Science 2012, 338, 643−647. (4) Burschka, J.; Pellet, N.; Moon, S.-J.; Humphry-Baker, R.; Gao, P.; Nazeeruddin, M. K.; Grätzel, M. Sequential Deposition as a Route to High-Performance Perovskite-Sensitized Solar Cells. Nature 2013, 499, 316−319. (5) You, J.; Hong, Z.; Yang, Y. M.; Chen, Q.; Cai, M.; Song, T.; Chen, C.; Lu, S.; Liu, Y.; Zhou, H.; Yang, Y. Low-Temperature Solution-Processed Perovskite Solar Cells with High Efficiency and Flexibility. ACS Nano 2014, 8, 1674−1680. (6) Noh, J. H.; Im, S. H.; Heo, J. H.; Mandal, T. N.; Seok, S., Il. Chemical Management for Colorful, Efficient, and Stable Inorganic − Organic Hybrid Nanostructured Solar Cells. Nano Lett. 2013, 13 (4), 1764−1769. (7) Chung, I.; Lee, B.; He, J.; Chang, R. P. H.; Kanatzidis, M. G. AllSolid-State Dye-Sensitized Solar Cells with High Efficiency. Nature 2012, 485, 486−489. (8) Crossland, E. J. W.; Noel, N.; Sivaram, V.; Leijtens, T.; Alexander-Webber, J. a; Snaith, H. J. Mesoporous TiO2 Single Crystals Delivering Enhanced Mobility and Optoelectronic Device Performance. Nature 2013, 495, 215−219. (9) Qiu, J.; Qiu, Y.; Yan, K.; Zhong, M.; Mu, C.; Yan, H.; Yang, S. AllSolid-State Hybrid Solar Cells Based on a New Organometal Halide Perovskite Sensitizer and One-Dimensional TiO2 Nanowire Arrays. Nanoscale 2013, 5, 3245−3248. (10) Liu, M.; Johnston, M. B.; Snaith, H. J. Efficient Planar Heterojunction Perovskite Solar Cells by Vapour Deposition. Nature 2013, 501, 395−398. (11) Gong, J.; Darling, S.; You, F. Perovskite Photovoltaics: LifeCycle Assessment of Energy and Environmental Impacts. Energy Environ. Sci. 2015, 8, 1953−1968. (12) Saliba, M.; Matsui, T.; Seo, J.-Y.; Domanski, K.; Correa-Baena, J.-P.; Nazeeruddin, M. K.; Zakeeruddin, S. M.; Tress, W.; Abate, A.; Hagfeldt, A.; Grätzel, M. Cesium-Containing Triple Cation Perovskite Solar Cells: Improved Stability, Reproducibility and High Efficiency. Energy Environ. Sci. 2016, 9, 1989−1997. (13) Dong, Q.; Fang, Y.; Shao, Y.; Mulligan, P.; Qiu, J.; Cao, L.; Huang, J. Electron-Hole Diffusion Lengths > 175 Micro-Meter in

Fracturing method and procedure and fitting method and analysis (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Author Contributions

T.Y.C. developed the experimental concept. A.J.Y. and T.Y.C. conducted the XSTM/S and SEM measurements and analyzed the corresponding data; C-C.H. synthesized the thin film and performed XRD measurements; A.J.Y. and T.Y.C. analyzed the XRD data; A.P. and Y.D. performed and interpreted the DFT calculations. The conclusions were made by discussions among all authors. The manuscript was written by A.J.Y. and T.Y.C. with contributions and suggestions from all authors. 29114


Research Article

ACS Applied Materials & Interfaces Solution-Grown CH3NH3PbI3 Single Crystals. Science 2015, 347, 967−970. (14) Shi, D.; Adinolfi, V.; Comin, R.; Yuan, M.; Alarousu, E.; Buin, A.; Chen, Y.; Hoogland, S.; Rothenberger, A.; Katsiev, K.; Losovyj, Y.; Zhang, X.; Dowben, P. A.; Mohammed, O. F.; Sargent, E. H.; Bakr, O. M. Low Trap-State Density and Long Carrier Diffusion in Organolead Trihalide Perovskite Single Crystals. Science 2015, 347, 519−522. (15) Liu, Y.; Yang, Z.; Cui, D.; Ren, X.; Sun, J.; Liu, X.; Zhang, J.; Wei, Q.; Fan, H.; Yu, F.; Zhang, X.; Zhao, C.; Liu, S. Two-Inch-Sized Perovskite CH 3 NH 3 PbX 3 (X = Cl, Br, I) Crystals: Growth and Characterization. Adv. Mater. 2015, 27, 5176−5183. (16) Nie, W.; Tsai, H.; Asadpour, R.; Blancon, J.; Neukirch, A. J.; Gupta, G.; Crochet, J. J.; Chhowalla, M.; Tretiak, S.; Alam, M. A.; Wang, H.-L.; Mohite, A. D. High-Efficiency Solution-Processed Perovskite Solar Cells with Millimeter-Scale Grains. Science 2015, 347, 522−525. (17) Zhou, H.; Chen, Q.; Li, G.; Luo, S.; Song, T.; Duan, H.-S.; Hong, Z.; You, J.; Liu, Y.; Yang, Y. Interface Engineering of Highly Efficient Perovskite Solar Cells. Science 2014, 345, 542−546. (18) Jeon, N. J.; Noh, J. H.; Yang, W. S.; Kim, Y. C.; Ryu, S.; Seo, J.; Seok, S. Il. Compositional Engineering of Perovskite Materials for High-Performance Solar Cells. Nature 2015, 517, 476−480. (19) Uehara, M.; Mori, S.; Chen, C. H.; Cheong, S. W. Percolative Phase Separation Underlies Colossal Magnetoresistance in MixedValent Manganites. Nature 1999, 399, 560−563. (20) Colella, S.; Mosconi, E.; Fedeli, P.; Listorti, A.; Gazza, F.; Orlandi, F.; Ferro, P.; Besagni, T.; Rizzo, A.; Calestani, G.; Gigli; De Angelis, F.; Mosca, R. MAPbI 3−X Cl X Mixed Halide Perovskite for Hybrid Solar Cells: The Role of Chloride as Dopant on the Transport and Structural Properties. Chem. Mater. 2013, 25 (22), 4613−4618. (21) Yamada, K.; Nakada, K.; Takeuchi, Y.; Nawa, K.; Yamane, Y. Tunable Perovskite Semiconductor CH3NH3SnX3 (X: Cl, Br, or I) Characterized by X-Ray and DTA. Bull. Chem. Soc. Jpn. 2011, 84, 926− 932. (22) Grancini, G.; D’Innocenzo, V.; Dohner, E. R.; Martino, N.; Srimath Kandada, A. R.; Mosconi, E.; De Angelis, F.; Karunadasa, H. I.; Hoke, E. T.; Petrozza, A. CH3NH3PbI3 Perovskite Single Crystals: Surface Photophysics and Their Interaction with the Environment. Chem. Sci. 2015, 6, 7305−7310. (23) Chien, T.; Liu, J.; Yost, A. J.; Chakhalian, J.; Freeland, J. W.; Guisinger, N. P. Built-in Electric Field Induced Mechanical Property Change at the Lanthanum Nickelate/Nb-Doped Strontium Titanate Interfaces. Sci. Rep. 2016, 6, 19017. (24) Chien, T.-Y.; Kourkoutis, L. F.; Chakhalian, J.; Gray, B.; Kareev, M.; Guisinger, N. P.; Muller, D. A.; Freeland, J. W. Visualizing ShortRange Charge Transfer at the Interfaces between Ferromagnetic and Superconducting Oxides. Nat. Commun. 2013, 4, 2336. (25) Chien, T.-Y.; Chakhalian, J.; Freeland, J. W.; Guisinger, N. P. Cross-Sectional Scanning Tunneling Microscopy Applied to Complex Oxide Interfaces. Adv. Funct. Mater. 2013, 23, 2565−2575. (26) Chien, T.-Y.; Guisinger, N. P.; Freeland, J. W. Cross-Sectional Scanning Tunneling Microscopy for Complex Oxide Interfaces. Proc. SPIE 2011, 7940, 79400T. (27) Chien, T.-Y.; Liu, J.; Chakhalian, J.; Guisinger, N.; Freeland, J. Visualizing Nanoscale Electronic Band Alignment at the La2/3Ca1/3MnO3/Nb:SrTiO3 Interface. Phys. Rev. B: Condens. Matter Mater. Phys. 2010, 82, 041101. (28) Edwards, H. L.; Markert, J. T.; de Lozanne, A. L. Energy Gap and Surface Structure of YBa2Cu3O7-x Probed by Scanning Tunneling Microscopy. Phys. Rev. Lett. 1992, 69, 2967−2970. (29) Kallip, S.; Lust, E. In Situ STM Studies of Bi (1 1 1) Electrodes in Aqueous Electrolyte Solutions. Electrochem. Commun. 2005, 7, 863− 867. (30) Massee, F.; de Jong, S.; Huang, Y.; Kaas, J.; Van Heumen, E.; Goedkoop, J. B.; Golden, M. S. Cleavage surfaces of the BaFe2‑xCoxAs2 and FeySe1‑xTex superconductors: A Combined STM plus LEED Study. Phys. Rev. B: Condens. Matter Mater. Phys. 2009, 80, 140507. (31) Hor, Y. S.; Roushan, P.; Beidenkopf, H.; Seo, J.; Qu, D.; Checkelsky, J. G.; Wray, L. A.; Hsieh, D.; Xia, Y.; Xu, S.; Qian, D.;

Hasan, M. Z.; Ong, N. P.; Yazdani, A.; Cava, R. J. Development of Ferromagnetism in the Doped Topological Insulator Bi 2−X Mn X Te 3. Phys. Rev. B: Condens. Matter Mater. Phys. 2010, 81, 195203. (32) Chen, C. J. Introduction to Scanning Tunneling Microscopy, 2nd ed.; Oxford Univ. Press, 2008. (33) Ou, Y. C.; Wu, J. J.; Fang, J.; Jian, W. Bin. Probing Capacitive Coupling and Collective Transport in PbSe Quantum-Dot Arrays Using Scanning Tunneling Spectroscopy. J. Phys. Chem. C 2009, 113, 7887−7891. (34) Hager, J.; Matzdorf, R.; He, J.; Jin, R.; Mandrus, D.; Cazalilla, M. A.; Plummer, E. W. Non-Fermi-Liquid Behavior in Quasi-OneDimensional Li0.9Mo6O17. Phys. Rev. Lett. 2005, 95, 186402. (35) She, L.; Liu, M.; Zhong, D. Atomic Structures of CH3NH3PbI3 (001) Surfaces. ACS Nano 2016, 10, 1126−1131. (36) Kutes, Y.; Ye, L.; Zhou, Y.; Pang, S.; Huey, B. D.; Padture, N. P. Direct Observation of Ferroelectric Domains in Solution-Processed CH3NH3PbI3 Perovskite Thin Films. J. Phys. Chem. Lett. 2014, 5, 3335−3339. (37) Kim, G. Y.; Oh, S. H.; Nguyen, B. P.; Jo, W.; Kim, B. J.; Lee, D. G.; Jung, H. S. Efficient Carrier Separation and Intriguing Switching of Bound Charges in Inorganic-Organic Lead Halide Solar Cells. J. Phys. Chem. Lett. 2015, 6, 2355−2362. (38) Snaith, H. J.; Abate, A.; Ball, J. M.; Eperon, G. E.; Leijtens, T.; Noel, N. K.; Stranks, S. D.; Wang, J. T. W.; Wojciechowski, K.; Zhang, W. Anomalous Hysteresis in Perovskite Solar Cells. J. Phys. Chem. Lett. 2014, 5, 1511−1515. (39) Frost, J. M.; Butler, K. T.; Walsh, A. Molecular Ferroelectric Contributions to Anomalous Hysteresis in Hybrid Perovskite Solar Cells. APL Mater. 2014, 2, 081506. (40) Beilsten-Edmands, J.; Eperon, G. E.; Johnson, R. D.; Snaith, H. J.; Radaelli, P. G. Non-Ferroelectric Nature of the Conductance Hysteresis in CH3NH3PbI3 Perovskite-Based Photovoltaic Devices. Appl. Phys. Lett. 2015, 106, 173502. (41) Eames, C.; Frost, J. M.; Barnes, P. R. F.; O’Regan, B. C.; Walsh, A.; Islam, M. S. Ionic Transport in Hybrid Lead Iodide Perovskite Solar Cells. Nat. Commun. 2015, 6, 7497. (42) Stoumpos, C. C.; Malliakas, C. D.; Kanatzidis, M. G. Semiconducting Tin and Lead Iodide Perovskites with Organic Cations: Phase Transitions, High Mobilities, and Near-Infrared Photoluminescent Properties. Inorg. Chem. 2013, 52 (15), 9019−9038. (43) Fan, Z.; Xiao, J.; Sun, K.; Chen, L.; Hu, Y.; Ouyang, J.; Ong, K. P.; Zeng, K.; Wang, J. Ferroelectricity of CH3NH3PbI3 Perovskite. J. Phys. Chem. Lett. 2015, 6, 1155−1161. (44) Goehry, C.; Nemnes, G. A.; Manolescu, A. Collective Behavior of Molecular Dipoles in CH3NH3PbI3. J. Phys. Chem. C 2015, 119, 19674−19680. (45) Ong, K. P.; Goh, T. W.; Xu, Q.; Huan, A. Mechanical Origin of the Structural Phase Transition in Methylammonium Lead Iodide CH3NH3PbI3. J. Phys. Chem. Lett. 2015, 6, 681−685. (46) Zheng, F.; Takenaka, H.; Wang, F.; Koocher, N. Z.; Rappe, A. M. First-Principles Calculation of the Bulk Photovoltaic Effect in CH3NH3PbI3 and CH3NH3PbI3‑xClx. J. Phys. Chem. Lett. 2015, 6, 31− 37. (47) Frost, J. M.; Butler, K. T.; Brivio, F.; Hendon, C. H.; van Schilfgaarde, M.; Walsh, A. Atomistic Origins of High-Performance in Hybrid Halide Perovskite Solar Cells. Nano Lett. 2014, 14 (5), 2584− 2590. (48) Brivio, F.; Walker, A. B.; Walsh, A. Structural and Electronic Properties of Hybrid Perovskites for High-Efficiency Thin-Film Photovoltaics from First-Principles. APL Mater. 2013, 1, 042111. (49) Frisch, M. J.; Trucks, G. W.; Schlegel, H. B.; Scuseria, G. E.; Robb, M. A.; Cheeseman, J. R.; Scalmani, G.; Barone, V.; Mennucci, B.; Petersson, G. A.; Nakatsuji, H.; Caricato, M.; Li, X.; Hratchian, H. P.; Izmaylov, A. F.; Bloino, J.; Zheng, G.; Sonnenberg, J. L.; Hada, M.; Ehara, M.; Toyota, K.; Fukuda, R.; Hasegawa, J.; Ishida, M.; Nakajima, T.; Honda, Y.; Kitao, O.; Nakai, H.; Vreven, T.; Montgomery, Jr., J. A.; Peralta, J. E.; Ogliaro, F.; Bearpark, M.; Heyd, J. J.; Brothers, E.; Kudin, K. N.; Staroverov, V. N.; Kobayashi, R.; Normand, J.; Raghavachari, K.; Rendell, A.; Burant, J. C.; Iyengar, S. S.; Tomasi, J.; Cossi, M.; Rega, 29115


Research Article

ACS Applied Materials & Interfaces N.; Millam, J. M.; Klene, M.; Knox, J. E.; Cross, J. B.; Bakken, V.; Adamo, C.; Jaramillo, J.; Gomperts, R.; Stratmann, R. E.; Yazyev, O.; Austin, A. J.; Cammi, R.; Pomelli, C.; Ochterski, J. W.; Martin, R. L.; Morokuma, K.; Zakrzewski, V. G.; Voth, G. A.; Salvador, P.; Dannenberg, J. J.; Dapprich, S.; Daniels, A. D.; Farkas, Ö .; Foresman, J. B.; Ortiz, J. V.; Cioslowski, J.; Fox, D. J. Gaussian09, revision D.01and C.01; Gaussian Inc.: Wallingford, CT, 2010. (50) Pimachev, A.; Dahnovsky, Y. Optical and Magnetic Properties of PbS Nanocrystals Doped by Manganese Impurities. J. Phys. Chem. C 2015, 119, 16941−16946. (51) Kresse, G.; Hafner, J. Ab Initio Molecular Dynamics for Liquid Metals. Phys. Rev. B: Condens. Matter Mater. Phys. 1993, 47, 558−561. (52) Kresse, G.; Hafner, J. Ab Initio Molecular-Dynamics Simulation of the Liquid-Metal−amorphous-Semiconductor Transition in Germanium. Phys. Rev. B: Condens. Matter Mater. Phys. 1994, 49, 14251− 14269. (53) Kresse, G.; Furthmüller, J. Efficiency of Ab-Initio Total Energy Calculations for Metals and Semiconductors Using a Plane-Wave Basis Set. Comput. Mater. Sci. 1996, 6, 15−50. (54) Kresse, G.; Furthmüller, J. Efficient Iterative Schemes for Ab Initio Total-Energy Calculations Using a Plane-Wave Basis Set. Phys. Rev. B: Condens. Matter Mater. Phys. 1996, 54, 11169−11186. (55) Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized Gradient Approximation Made Simple. Phys. Rev. Lett. 1996, 77, 3865−3868. (56) Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized Gradient Approximation Made Simple. [Erratum to Document Cited in CA126:51093]. Phys. Rev. Lett. 1997, 78, 1396. (57) Blöchl, P. E. Projector Augmented-Wave Method. Phys. Rev. B: Condens. Matter Mater. Phys. 1994, 50, 17953−17979. (58) Kresse, G.; Joubert, D. From Ultrasoft Pseudopotentials to the Projector Augmented Wave Method. Phys. Rev. B: Condens. Matter Mater. Phys. 1999, 59, 1758−1775. (59) Mosconi, E.; Amat, A.; Nazeeruddin, K.; Grätzel, M.; De Angelis, F. First-Principles Modeling of Mixed Halide Organometal Perovskites for Photovoltaic Applications. J. Phys. Chem. C 2013, 117 (27), 13902−13913. (60) Monkhorst, H. J.; Pack, J. D. Special Points for Brillouin-Zon Integrations. Phys. Rev. B 1976, 13, 5188−5192.

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Coexistence of Two Electronic Nano-Phases on a CH3NH3PbI3âxClx

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