Lead-free solid-state organic-inorganic halide perovskite solar cells

ARTICLES PUBLISHED ONLINE: 4 MAY 2014 | DOI: 10.1038/NPHOTON.2014.82

Lead-free solid-state organic–inorganic halide perovskite solar cells Feng Hao1, Constantinos C. Stoumpos1, Duyen Hanh Cao1, Robert P. H. Chang2 and Mercouri G. Kanatzidis1 * Lead-free solution-processed solid-state photovoltaic devices based on methylammonium tin iodide (CH3NH3SnI3) perovskite semiconductor as the light harvester are reported. Featuring an optical bandgap of 1.3 eV, the CH3NH3SnI3 perovskite material can be incorporated into devices with the organic hole-transport layer spiro-OMeTAD and show an absorption onset at 950 nm, which is significantly redshifted compared with the benchmark CH3NH3PbI3 counterpart (1.55 eV). Bandgap engineering was implemented by chemical substitution in the form of CH3NH3SnI3–xBrx solid solutions, which can be controllably tuned to cover much of the visible spectrum, thus enabling the realization of lead-free solar cells with an initial power conversion efficiency of 5.73% under simulated full sunlight. Further efficiency enhancements are expected following optimization and a better fundamental understanding of the internal electron dynamics and corresponding interfacial engineering. The reported CH3NH3SnI3–xBrx perovskite solar cells represent a step towards the realization of low-cost, environmentally friendly solid-state solar cells.

T

he development of clean alternatives to current power generation methods is immensely important to preserving the global environment and assuring sustained economic growth1. The recent emergence of halide perovskites as light harvesters and hole-transport materials has revolutionized the emerging photovoltaic technologies2–14. Organic–inorganic hybrid perovskite compounds based on metal halides adopt the ABX3 perovskite structure. This structure consists of a network of cornersharing BX6 octahedra, where the B atom is a metal cation (typically Sn2þ or Pb2þ) and X is typically F2, Cl2, Br2 or I2. The A cation is selected to balance the total charge and it can even be a Csþ or a small molecular species15–17. Recent implementation of CH3NH3PbX3 (X ¼ I, Cl, Br) perovskite absorbers with the organic hole conductor 2,2′ ,7,7′ -tetrakis-(N,N-di-p-methoxyphenylamine)9,9′ -spirobifluorene (spiro-OMeTAD) enabled power conversion efficiencies (PCEs) greater than 15% (refs 8,18), and has been recognized as the ‘next big thing in photovoltaics’19–22. A planar heterojunction photovoltaic device incorporating vapour-deposited perovskite (CH3NH3PbI32xClx) as the absorbing layer has demonstrated overall PCEs of over 15% with a high opencircuit voltage of up to 1.07 V, further highlighting the industrial application potential of this configuration in the near future7. Recent studies have indicated that mixed-halide organic–inorganic hybrid perovskites can display electron–hole diffusion lengths of over 1 mm, which is consistent with our reports of very high carrier mobilities in these materials23 and supports our expectations for highly efficient and cheap solar cells using thick absorption layers24,25. However, to realize commercial applications of this technology it is important to reach analogous optical and photovoltaic performance using lead-free organic–inorganic compounds. Here, we report a first attempt using the lead-free perovskite of methylammonium tin iodide (CH3NH3SnI3) as the light-absorbing material to fabricate solution-processed solid-state photovoltaic devices. Featuring an even lower optical bandgap of 1.3 eV than the 1.55 eV achieved with CH3NH3PbI3 , devices with

CH3NH3SnI3 in conjunction with an organic spiro-OMeTAD hole-transport layer showed an absorption onset of 950 nm. Further chemical alloying of iodide with bromide provides efficient energetic tuning of the band structure of the perovskites, leading to a PCE of 5.8% under simulated full sunlight of 100 mW cm22. Building on this very promising initial result and with further reduction of interfacial losses, we believe a substantial increase in efficiency can be achieved. As shown in Fig. 1a,b, CH3NH3SnI3 adopts the perovskite structure type, crystallizing in the pseudocubic space group P4mm at ambient conditions. Unlike CH3NH3PbI3 , which has a lower symmetry at room temperature (b-phase), the Sn analogue adopts its highest symmetry phase (a-phase), even at room temperature. The corner-sharing [SnI6]42 polyhedra form an infinite threedimensional lattice with Sn–I–Sn connecting angles of 177.43(1)8 and 1808 for the a- and c-axes, respectively. The deviation from the ideal cubic (Pm–3m) structure arises from orientational polarization of the CH3NHþ 3 cation along the C–N bond direction, which is imposed on the three-dimensional [SnI3]2 inorganic lattice coinciding with the crystallographic c-axis23. CH3NH3SnI3 is a direct-gap semiconductor with an energy gap of 1.3 eV, as has been shown experimentally and theoretically23,26. The optical bandgap Eg of the CH3NH3SnI3 compound (determined from diffuse reflectance measurements) is shown in Fig. 1c. The optical absorption coefficient (a/S) is calculated using reflectance data according to the Kubelka–Munk equation27, a/S ¼ (1 2 R)2/2R, where R is the percentage of reflected light, and a and S are the absorption and scattering coefficients, respectively. At room temperature, the material displays a strong photoluminescence emission at 950 nm, which corresponds to the onset of the absorption edge (Fig. 1c). The photoluminescence intensity can act as a qualitative measure of the number of photogenerated carriers in semiconductors, as it is proportional to the number of e2hþ pairs produced by the incident light28. As depicted in Fig. 1d, its bulk electrical conductivity s is 5 × 1022 S cm21 at room temperature,

1

Department of Chemistry, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208, USA, 2 Department of Materials Science and Engineering, and Argonne-Northwestern Solar Energy Research (ANSER) Center, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208, USA. * e-mail: [email protected]

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Figure 1 | Crystal structure, XRD pattern, optical absorption and photoluminescence spectra, conductivity and Seebeck coefficient of CH3NH3SnI3 perovskite. a, Perovskite crystal structure of the CH3NH3SnI32xBrx absorber materials. b, Experimental (red) and simulated (black) X-ray diffraction pattern for CH3NH3SnI3. c,d, Optical absorption and photoluminescence spectra (c) and conductivity and Seebeck coefficient (d) as a function of temperature for a sample of CH3NH3SnI3 prepared using the solution method23.

corresponding to a Seebeck coefficient of 2 60 mV K21 (n-type). The compound has a low carrier concentration on the order of 1 × 1014 cm23 and high electron mobilities (me) on the order of 2,000 cm2 V21 s21, which is comparable or even superior to most traditional semiconductors, including Si, CuInSe2 and CdTe, which have comparable bandgap energies. The doping level of CH3NH3SnI3 can be varied greatly depending on the preparation method. Carrier concentrations up to 1 × 1019 cm23 have been reported for CH3NH3SnI3 (ref. 29), showing a strong p-type character and a metallic behaviour suggestive of a heavily doped semiconducting behaviour. We attribute this large difference in the transport properties to Sn4þ impurities that are inherently present in commercial SnI2 , which are readily detectable by a mass loss at 150 8C via thermal gravimetric analysis (Supplementary Fig. 1). Therefore, when assembling the solar cells, care must be taken in depositing films of tin perovskite with low carrier concentration to maximize the carrier mobility within the active perovskite layer. This means that excessive Sn4þ in the sample must be avoided. The valence band maximum (EVB) of the CH3NH3SnI3 compound was determined from ultraviolet photoelectron spectroscopy (UPS) measurements. A representative UPS spectrum for the CH3NH3SnI3 is shown in Supplementary Fig. 2, where the energy is calibrated with respect to the He I photon energy (21.21 eV). The valence band energy EVB is estimated to be 25.47 eV below vacuum level, which is close to the reported value for CH3NH3PbI3 (25.43 eV)5. From the observed optical bandgap, the conduction band energy ECB of CH3NH3SnI3 was determined to be at 24.17 eV, that is, slightly higher than the ECB for the TiO2 anatase electrode (24.26 eV)5. To fabricate the solid-state solar cells, mesoporous anatase TiO2 films were prepared by spin coating a solution of colloidal anatase particles (20 nm in size) onto a 30-nm-thick compact TiO2 underlayer30. The underlayer was deposited by atomic layer deposition on a pre-patterned transparent-conducting-oxide-coated glass 490

substrate acting as the electric front contact of the solar cell. Deposition of the perovskite light-absorbing layer was carried out by spin coating in a nitrogen glove box to avoid hydrolysis and oxidation of the tin perovskite in contact with air. The triarylamine derivative 2,2′ ,7,7′ -tetrakis-(N,N-di-p-methoxyphenylamine)-9,9′ spirobifluorene (spiro-OMeTAD)31 was then applied as a holetransporting material (HTM) on top of the mesoporous TiO2 and perovskite layer. Lithium bis(trifluoromethylsulphonyl)imide and 2,6-lutidine were added in the HTM solution as important dopants to increase the hole mobility31. Figure 2 shows a representative cross-sectional scanning electron microscopy (SEM) image of a typical solar cell device. The mesoporous TiO2 film had an average thickness of 350 nm and was infiltrated with the perovskite nanocrystals using the spin-coating procedure. The HTM penetrates into the remaining pore volume of the perovskite/TiO2 layer and forms a 200-nm-thick capping layer on top of the composite structure. A thin gold layer was thermally evaporated under high vacuum onto the HTM layer, forming the back contact electrode of the device. The solid-state device based on the CH3NH3SnI3 perovskite shows a high mean short-circuit photocurrent density Jsc of 16.30 mA cm22, an open-circuit voltage Voc of 0.68 V and a moderate fill factor (FF) of 0.48 under AM 1.5G solar illumination, corresponding to a PCE of 5.23% (Fig. 3a). This high current density was achieved with submicrometre-thick TiO2 films (that is, 350 nm) because of the large optical absorption cross-section of the perovskite material and the well-developed interfacial pore filling by the hole conductor (Fig. 2). More importantly, the incident photonto-electron conversion efficiency (IPCE) of the CH3NH3SnI3based device covers the entire visible spectrum and reaches a broad absorption maximum of over 60% from 600 nm to 850 nm. It is accompanied with a notable absorption onset up to 950 nm (Fig. 3b), which is in good agreement with the optical bandgap of 1.30 eV. Integrating the overlap of the IPCE spectrum with the AM 1.5G solar photon flux yields a current density of NATURE PHOTONICS | VOL 8 | JUNE 2014 | www.nature.com/naturephotonics


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Figure 2 | Representative cross-sectional SEM view of a completed photovoltaic device with CH3NH3SnI3 perovskite. Individual layers are indicated on the left. Scale bar, 200 nm. Real device images are shown on the right, indicating the colours of the photovoltaic devices made with CH3NH3SnI32xBrx as a function of I/Br ratio.

16.60 mA cm22, which is in excellent agreement with the measured photocurrent density. This confirms that any mismatch between the simulated sunlight and the AM 1.5G standard is negligibly small. It is important to note that, although the obtained Jsc for the CH3NH3SnI3 perovskite device is less efficient than that for the CH3NH3PbI3 device7,8,18,32, the maximum current density that can be generated exceeds 30 mA cm22 when integrating the AM 1.5G solar spectrum below the bandgap of CH3NH3SnI3 perovskite (1.30 eV). To figure out the limiting factor for the fair photocurrent density, devices with thinner active layer thickness (150 nm) were constructed and tested. As shown in Supplementary Fig. 3, a Jsc of 12 mA cm22 was achieved, with a Voc of 0.74 V and a FF of 0.45, thus generating a PCE of 4.44%. Accordingly, the diffusion

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Figure 3 | Photovoltaic and IPCE characteristics for devices with CH3NH3SnI32x Brx perovskites. a,b, Photocurrent density–voltage (J–V) characteristics (a) and corresponding IPCE spectra (b) of devices based on CH3NH3SnI32xBrx (x ¼ 0, 1, 2, 3) perovskites.

length of the tin perovskite might not be a main limiting factor for device performance24,25. The film morphology and quality of the spin-coated CH3NH3SnI3 film were then investigated by SEM (Supplementary Fig. 4). Poor film quality and coverage of the tin perovskite on the mesoporous TiO2 electrodes were observed, which has recently been recognized as an important factor determining perovskite solar cell performance8,13. Future device optimization will therefore focus on the improvement of perovskite film quality and interfacial recombination inhibition. It has recently been observed that the charge accumulates in high density in the perovskite absorber material rather than only in the semiconducting TiO2 electrodes, making this type of photovoltaic device fundamentally different from dye-sensitized solar cells33. Thus, the Voc in a perovskite solar cell is not only related to the energy difference between the HTM potential and the TiO2 conduction bandedge, but could also be correlated with the energy difference between the HTM potential and the conduction bandedge of the perovskite itself. From the abovementioned band alignment it can be inferred that the conduction bandedge ECB of CH3NH3SnI3 is 0.24 eV lower than in CH3NH3PbI3 , thus leading to a lower Voc for the CH3NH3SnI3 perovskite device. Therefore, in an attempt to increase the Voc of these lead-free devices, chemical substitution of the iodide atom with bromide was applied in order to favourably tune the bandgap energetics32. The CH3NH3SnI32xBrx compounds were prepared by mixing stoichiometric amounts of CH3NH3X and SnX2 (X ¼ Br, I), finely homogenized in a mortar in the nitrogen glove box. The resulting solids were sealed in silica ampules under 1 × 1024 mbar vacuum and heated to 200 8C to complete the reaction. As shown by the X-ray diffraction (XRD) patterns in Fig. 4a, this series of compositions forms a continuous solid solution throughout the composition range, without displaying any structural transitions (at room temperature), thus retaining the crystal structure of both end members, crystallizing at the P4mm space group. The properties of the solid solutions are clearly displayed by a continuous contraction of the lattice parameters from the CH3NH3SnI3 to CH3NH3SnBr3 end members, which results in a widening of the bandgap (Table 1). To check the optical properties in the hybrid halide perovskite, ultraviolet–visible diffuse reflectance spectra of CH3NH3SnI32xBrx (x ¼ 0, 1, 2, 3) were measured and transformed to absorption spectra, as mentioned above. Figure 4b shows that the absorption onset of CH3NH3SnI32xBrx (x ¼ 0, 1, 2, 3) hybrid halide perovskites can be tuned from 954 nm (1.30 eV for CH3NH3SnI3) to 577 nm (2.15 eV for CH3NH3SnBr3), thus resulting in significant colour tunability for perovskite photovoltaic devices (as shown in the



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Figure 4 | XRD patterns, absorption spectra and schematic energy-level diagram of CH3NH3SnI32x Brx compounds. a,b, XRD patterns (a) and absorption spectra (b) of the CH3NH3SnI32xBrx (x ¼ 0, 1, 2, 3) perovskites. c, Schematic energy-level diagram of CH3NH3SnI32xBrx with TiO2 and spiro-OMeTAD HTM. The valence band maxima ECB of the methylammonium tin halides were extracted from UPS measurements under high vacuum.

right panel of Fig. 2). The intermediate iodide/bromide hybrid perovskites of CH3NH3SnI2Br and CH3NH3SnIBr2 show absorption onsets of 795 nm (1.56 eV) and 708 nm (1.75 eV), respectively. The valence band energy EVB of the CH3NH3SnI32xBrx compounds was also estimated from the UPS measurements. As illustrated in Fig. 4c, the ECB increased from –4.17 eV below vacuum level for CH3NH3SnI3 to 23.96 eV for CH3NH3SnI2Br and 23.78 eV for CH3NH3SnIBr2 , and finally to 23.39 eV for CH3NH3SnBr3. It is obvious from the band alignment diagram that the change in the bandgap Eg of the CH3NH3SnI32xBrx compounds is mainly due to the conduction band shift to higher energy, with the valence band energy remaining practically unchanged. These changes in energy levels allow for bandgap engineering and the tuning of energetics for more efficient solar cell architectures. Figure 3a,b presents representative photocurrent density–voltage (J–V ) characteristics and IPCE spectra for devices constructed with the CH3NH3SnI32xBrx perovskites as light harvesters. The photovoltaic parameters are summarized in Table 1. As demonstrated in the right panel of Fig. 2, through the chemical compositional control of CH3NH3SnI32xBrx , the corresponding device colour

can be tuned from black for CH3NH3SnI3 to dark brown for CH3NH3SnI2Br and to yellow for CH3NH3SnBr3 with increasing Br content. Notably, Jsc decreased from 16.30 mA cm22 for CH3NH3SnI3 to 8.26 mA cm22 for CH3NH3SnBr3 with increasing Br content, whereas Voc increased from 0.68 V to 0.88 V when switching from the pure iodide to pure bromide perovskite. In addition to the significant improvement in Voc , an increase in FF from 0.48 to 0.59 was also observed upon the incorporation of the Br ions. Amongst the investigated CH3NH3SnI32xBrx perovskites, the device with CH3NH3SnIBr2 delivered the highest PCE of 5.73%, with a Jsc of 12.30 mA cm22, a Voc of 0.82 V and a FF of 0.57. The reduction of Jsc with increasing Br content is directly related to the blueshift of absorption onset, as indicated from the IPCE spectra in Fig. 3b. Consistent with the bandgap tuning, the onset of the IPCE spectra blueshifted from 950 nm for the iodide perovskite to 600 nm for the pure bromide perovskite. Integrating the overlap of these IPCE spectra with the AM 1.5G solar photon flux yields a current density Jcal that is similar to the measured photocurrent density Jsc (Table 1). The improvement in Voc can be attributed to the raised conduction bandedge ECB with

Table 1 | Optical bandgaps and refined lattice parameters of the CH3NH3SnI32xBrx (x 5 0, 1, 2, 3) perovskites and corresponding solar cell performance parameters. Perovskites CH3NH3SnI3 CH3NH3SnI2Br CH3NH3SnIBr2 CH3NH3SnBr3

E g* (eV) 1.30 1.56 1.75 2.15

Lattice parameters (Å) a ¼ 6.169(1) c ¼ 6.173(4) a ¼ 6.041(1) c ¼ 6.053(4) a ¼ 5.948(1) c ¼ 5.953(4) a ¼ 5.837(1) c ¼ 5.853(4)

J †sc (mA cm22) 16.30+0.71 14.38+0.49 12.30+0.47 8.26+0.53

J ‡cal (mA cm22) 16.60 13.96 11.73 7.93

Voc (V)§ 0.68+0.03 0.77+0.02 0.82+0.03 0.88+0.03

FF 0.48+0.03 0.50+0.02 0.57+0.02 0.59+0.02

PCE (%)} 5.23+0.18 5.48+0.15 5.73+0.23 4.27+0.18

R #s (V) 105.00 103.54 65.24 60.98

The photovoltaic parameters were the average of six devices in the same bath. *Optical bandgap determined from the diffuse-reflectance measurements; †short-circuit photocurrent density; ‡calculated photocurrent density from the integration of representative IPCE curves shown in Fig. 3b with the AM 1.5G solar spectrum; §open-circuit photovoltage; fill factor; }power conversion efficiency; #series resistances derived from the J–V curves in Fig. 3a.

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increasing Br content in CH3NH3SnI32xBrx. However, the FF is significantly lower than the reported values for high-efficiency perovskite solar cells7,8,18. It is well recognized that series resistance Rs is one of the major factors influencing the FF of solar cells. This arises mainly from three factors: (1) the active and interfacial layer resistances, (2) electrode resistance and (3) contact resistance. In the present work, variations in Rs among the cells with CH3NH3SnI32xBrx perovskites will be mainly caused by differences in the active layer resistance, because factors (2) and (3) will be similar, being common in all of our devices. We estimated Rs from the slope of the J–V curve at the open-circuit voltage point. As shown in Table 1, the Rs value decreased from 105 V for a device with CH3NH3SnI3 to 60.98 V for CH3NH3SnBr3 , which is in accordance with the observed FF enhancement from 0.48 to 0.59. Another important issue that needs to be specified is that insufficient coverage of the perovskite film (Supplementary Fig. 4) might result in a high frequency of shunt paths allowing direct contact between the HTM and the TiO2 compact layer, which will act as a parallel diode in the device, deteriorating the FF and Voc (ref. 34). Further FF improvements are expected from more efficient interfacial engineering to inhibit back electron recombination. An important issue to note for the tin perovskite solar cell is the well-known poor atmospheric stability of the Sn-based perovskite compared to its lead analogy16,21,35. We carried out a preliminary stability investigation of the CH3NH3SnI3 perovskite solar cell by storing the devices in a nitrogen glove box after sealing with Surlyn films. Encouragingly, the devices retained almost 80% of the initial performance in the first 12 h (Supplementary Table 1). Performance loss arises mainly from the decrease in photocurrent density and FF, primarily due to the p-type doping via Sn2þ oxidation induced during the fabrication process. Far better stability can be expected if more advanced sealing techniques are adopted. In summary, methylammonium tin halide perovskites (CH3NH3SnI32xBrx) have been used as lead-free light harvesters for solar cell applications for the first time. Featuring an ideal optical bandgap of 1.3 eV, devices with CH3NH3SnI3 perovskite together with an organic spiro-OMeTAD hole-transport layer showed a notable absorption onset up to 950 nm, which is significantly redshifted compared with its benchmark CH3NH3PbI3 counterpart (1.55 eV). The bandgap engineering of CH3NH3SnI32xBrx perovskites can be controllably tuned to cover much of the visible spectrum, thus enabling the realization of lead-free, colourful solar cells and leading to a promising initial PCE of 5.73% under simulated full sunlight. Further efficiency enhancements would be expected by the fundamental understanding of the internal electron dynamics and corresponding interfacial engineering. The reported CH3NH3SnI32xBrx perovskites are believed to represent a significant step towards the realization of low-cost, high-efficiency, environmentally benign, next-generation solid-state solar cells.

Methods Materials. Unless stated otherwise, all materials were purchased from Sigma-Aldrich and used as received. Spiro-OMeTAD was purchased from Merck KGaA. CH3NH3I, CH3NH3Br and SnI2 were synthesized and purified according to a reported procedure23. CH3NH3SnI32xBrx compounds were prepared by mixing stoichiometric amounts of CH3NH3X and SnX2 (X ¼ Br, I), finely homogenized in a mortar in a nitrogen glove box. The resulting solids were sealed in quartz ampules under 1 × 1024 mbar vacuum and heated to 200 8C to complete the reaction23. Material characterization. Optical diffuse-reflectance measurements were performed at room temperature using a Shimadzu UV-3101 PC double-beam, double-monochromator spectrophotometer operating from 200 nm to 2,500 nm. BaSO4 was used as a non-absorbing reflectance reference. Photoluminescence spectra were measured with an OmniPV photoluminescence system, equipped with a diode-pumped frequency-doubled Nd:YAG laser (500 mW power output, class 4) emitting at 532 nm coupled with a bundle of eight 400-mm-core optical fibres as an excitation source. Resistivity measurements were made for arbitrary current directions in the a–b plane using a standard point contact geometry. A homemade

resistivity apparatus was used that was equipped with a Keithley 2182A nanovoltometer, Keithley 617 electrometer, Keithley 6220 Precision d.c. source and a high-temperature vacuum chamber controlled by a K-20 MMR system. Seebeck measurements were performed on the same homemade apparatus using Cr/Cr:Ni thermocouples as electric leads, which were attached to the sample surface by means of colloidal graphite isopropanol suspension. The temperature gradient along the crystal was generated by a resistor on the ‘hot’ side of the crystal. The data were corrected for the thermocouple contribution using a copper wire. SEM and energydispersive spectroscopy (EDS) measurements were performed with a Hitachi SU8030 scanning electron microscope equipped with an Oxford X-max 80 SDD EDS detector. Data were acquired with an accelerating voltage of 15 kV. Device fabrication. A fluorine-doped tin oxide-coated glass substrate (Tec15, Hartford Glass) was patterned by etching with Zn metal powder and 2 M HCl diluted in deionized water. The substrates were then cleaned by ultrasonication with detergent, rinsed with deionized water, acetone and ethanol, and dried with clean dry air. A 30-nm-thick TiO2 compact layer was deposited on the substrates by an atomic layer deposition system (Savannah S300, Cambridge Nanotech) using titanium isopropoxide (TTIP) and water as precursors. The mesoporous TiO2 layer composed of 20-nm-sized particles was deposited by spin coating at 4,500 r.p.m. for 30 s using a hydrothermal-synthesized TiO2 paste diluted in ethanol (1:4, weight ratio). After drying at 125 8C, the TiO2 films were gradually heated to 500 8C, baked at this temperature for 15 min, and then cooled to room temperature (25 8C). After cooling to room temperature, the substrates were treated in a 0.02 M aqueous solution of TiCl4 for 30 min at 70 8C, rinsed with deionized water, and dried at 500 8C for 20 min. Before use, the films were again dried at 500 8C for 30 min. CH3NH3SnI32xBrx was dissolved in N,N-dimethylformamide at a weight concentration of 30% while stirring at 70 8C. The solution was kept at 70 8C during the whole procedure. The mesoporous TiO2 films were then infiltrated with CH3NH3SnI32xBrx by spin coating at 4,000 r.p.m. for 45 s and dried at 125 8C for 30 min to remove the solvent. The HTM was then deposited by spin coating at 4,000 r.p.m. for 30 s. The spin-coating formulation was prepared by dissolving 72.3 mg spiro-OMeTAD, 30 ml 2,6-lutidine, 17.5 ml of a stock solution of 520 mg ml21 lithium bis(trifluoromethylsulphonyl)imide in acetonitrile in 1 ml chlorobenzene. Finally, 100 nm of gold was thermally evaporated on top of the device to form the back contact. The devices were sealed in nitrogen using a 30-mmthick hot-melting polymer and a microscope coverslip to prevent oxidation. Device characterization. J–V characteristics were measured under AM 1.5G light (100 mW cm22) using the xenon arc lamp of a Spectra-Nova Class A solar simulator. Light intensity was calibrated using an NREL-certified monocrystalline Si diode coupled to a KG3 filter to bring the spectral mismatch to unity. A Keithley 2400 source meter was used for electrical characterization. The active area of all devices was 10 mm2, and an 8 mm2 aperture mask was placed on top of the cells during all measurements. IPCEs were characterized using an Oriel model QE-PV-SI instrument equipped with a NIST-certified Si diode. Monochromatic light was generated from an Oriel 300 W lamp.

Received 14 January 2014; accepted 18 March 2014; published online 4 May 2014

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Acknowledgements The authors thank T. Marks for use of the solar simulator and IPCE measurement system. Electron microscopy and elemental analysis were carried out at the Electron Probe Instrumentation Center (EPIC) at Northwestern University. This research was supported as part of the ANSER Center, an Energy Frontier Research Center funded by the US Department of Energy, Office of Science, Office of Basic Energy Sciences (award no. DE-SC0001059) and ISEN at Northwestern University.

Author contributions M.G.K. conceived the experiments and directed the study. F.H. and C.C.S. carried out the material synthesis, device fabrication and performance measurements. D.H.C. prepared the TiO2 blocking layer for the electrodes. R.P.H.C. contributed to the revision of the manuscript. All authors discussed the results and commented on the manuscript.

Additional information Supplementary information is available in the online version of the paper. Reprints and permissions information is available online at www.nature.com/reprints. Correspondence and requests for materials should be addressed to M.G.K.

Competing financial interests The authors declare no competing financial interests.

