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Efficient bulk heterojunction CH3NH3PbI3-TiO2 solar cells with TiO2 nanoparticles at grain boundaries of perovskite by multi-cycle-coating strategy Jun Shao, Songwang Yang, and Yan Liu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • Publication Date (Web): 25 Apr 2017 Downloaded from http://pubs.acs.org on April 25, 2017

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ACS Applied Materials & Interfaces

Efficient bulk heterojunction CH3NH3PbI3-TiO2 solar cells with TiO2 nanoparticles at grain boundaries of perovskite by multi-cycle-coating strategy Jun Shao†,‡, Songwang Yang*,†, and Yan Liu*,† †

CAS Key Laboratory of Materials for Energy Conversion, Shanghai Institute of Ceramics,

Chinese Academy of Sciences, 588 Heshuo Road, Shanghai, 201899, P. R. China. ‡

University of Chinese Academy of Sciences, Beijing 100039, P. R. China.

KEYWORDS Perovskite solar cells, bulk heterojunction, TiO2 nanoparticle, grain boundary, multi-cyclecoating

ABSTRACT

A novel bulk heterojunction (BHJ) perovskite solar cell (PSC), where the perovskite grains act as donor and the TiO2 nanoparticles act as acceptor, is reported. This efficient BHJ PSC was simply solution-processed from a mixing precursor of CH3NH3PbI3 (MAPbI3) and TiO2 nanoparticles.

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With dissolution and recrystallization by multi-cycle-coating, a unique composite structure ranging from a MAPbI3-TiO2-dominated layer on the substrate side to a pure perovskite layer on the top side is formed, which is beneficial for the blocking of possible contact between TiO2 and hole transport material on the interface. Scanning electron microscopy clearly shows that TiO2 nanoparticles accumulate along the grain boundaries (GBs) of perovskite. The TiO2 nanoparticles at GBs quick extract and reserve photogenerated electron before it transport in the perovskite phase as described in multi-trapping model, which retard the electron-hole recombination and reduce the energy loss, resulting in increased VOC and FF. Moreover, the pinning effect of TiO2 nanoparticles at GBs from the strong bindings between TiO2 and MAPbI3, suppresses massive ion migrations along the GBs, leading to improved operational stability and diminished hysteresis. Photoluminescence (PL) quenching and PL decay confirm the efficient exciton dissociation on the hetero-interface. Electrochemical impedance spectroscopy and opencircuit photovoltage decay measurements show the reduced recombination loss and improved carrier lifetime of BHJ PSCs. This novel strategy of device design effectively combines the benefits of both planar and mesostructured architectures while avoiding their shortcuts, eventually leading to a high PCE of 17.42% under one sun illumination. The newly proposed approach also provides a new way to fabricate TiO2-contained perovskite active layer at a low temperature.

INTRODUCTION Organolead halide perovskites of general formula ABX3 (where A is an organic /inorganic cation, B is a metal cation, and X is a halide anion), have emerged as one of the most appealing


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materials for photovoltaic applications.1-4 Perovskite solar cells (PSCs) have witnessed a certified power conversion efficiency (PCE) of 22.1%5 due to the fascinating characteristics of the material, including high absorption coefficients, long carrier diffusion lengths, and tunable optical band gaps.6-8 The primary material (MAPbI3, where MA=CH3NH3) for highly efficient PSCs is mostly prepared by solution process, where a large number of grain boundaries (GBs) are unavoidably formed.9 The GBs that exhibit in the deposited perovskite active layer are believed to be the most possible causes for energy loss and hysteresis.10, 11 The role of GBs in affecting the photovoltaic properties of PSCs has been widely investigated. Snaith and coworkers have confirmed faster charge recombination at GBs than in interior grains which undermines the device performance.12 Recently, several reports point out that GBs also act as short cut for ion migration that in-turn influence device operation.13, 14 In pursuit of minimizing charge recombination and ion migration within the perovskite films, numerous efforts have been devoted to passivation of surface charge trap states of the perovskite and corresponding GBs to improve the optoelectronic properties of perovskite materials. Multiple passivating agents, such as thiophene and pyridine lewis bases15, graphene16, and even PbI217 and MAI,9 have been studied. However, the delicate control for modification of GBs is still one of the biggest remained challenges for high performance PSCs. Two dominant device architectures, namely mesostructured (MS) PSCs employing a mesoporous scaffold (usually TiO2) and planar heterojunction (PHJ) PSCs sandwiching the perovskite layer between electrode buffer layers, have been developed.18 Although PHJ PSCs without significant hysteresis are obtained,19, 20 the top-performing PSCs have been limited to the mesoporous TiO2-based configuration.21 However, the high-temperature-processed (450 oC to 550 oC) fabrication of mesoporous TiO2 layer (meso-TiO2) makes the manufacture more


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complex and hinders its further application to the plastic substrates.22,

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Moreover, the high

demand of the complete pore-filling of perovskite in MS PSCs makes the perovskite deposition more difficult than the corresponding planar perovskite layers.24 During the preparation of this paper, a TiO2 embedded structure which combined the advantages of MS and PHJ structures was reported.25 The embedded TiO2 nanoparticles could improve the electron extraction and promote the growth of large grains. However, the uniformly dispersed TiO2 nanoparticles would emerge on the surface of the perovskite layer and act as recombination sites due to the possible contact between TiO2 and hole conducting material (HTM), as arising from the vanished capping layer. In this regards, further device design is essential for improving film morphology of active layer and preventing recombination loss on the interface. In this paper, we are motivated to develop a novel device structure of bulk heterojunction (BHJ) PSC where electron acceptor of TiO2 nanoparticles accumulates at the GBs of perovskite, which aims to passivate GB defects and block ion migration. We achieve this target by solution processed deposition from a mixed precursor which contains TiO2 nanoparticles and MAPbI3. With dissolution and recrystallization by multi-cycle-coating of mixed solution and neat MAPbI3 precursor, a unique composite structure ranging from a MAPbI3-TiO2-dominated layer on the substrate side to a pure perovskite layer on the top side is formed, by which the possible contact between TiO2 and HTM is effectively prevented. A complete interpenetration between TiO2 nanoparticles and MAPbI3 grains has been accomplished as compared to normally structured mesoporous TiO2-based devices. The steady-state photoluminescence quenching and photoluminescence decay experiments confirm the efficient charge extraction on the heterointerface of MAPbI3 and TiO2 nanoparticles. Unlike in MS devices, TiO2 nanoparticles in BHJ PSCs function as “electron resevoir” as in the multi-trapping model. Thermalized electrons from


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the TiO2 are then transported through the perovskite phase. This strategy allows a reduction of recombination loss, representing increased VOC and FF. Furthermore, the pinning effect of TiO2 nanoparticle, which comes from the strong bindings between TiO2 and MAPbI3,26, 27 suppresses massive ion migrations along the GBs in the perovskite layer. The incorporation of TiO2 nanoparticles at GBs finally leads to enhanced cell performance, improved operational stability and diminished hysteresis. This unique BHJ PSC demonstrates a tremendously enhanced PCE up to 17.42% that surpasses both fabricated PHJ and MS PSCs. EXPERIMENTAL METHODS TiO2/DMF Solution TiO2 nanoparticles were prepared according to the literature.28 The hydrothermal synthesized TiO2 sol was centrifuged and washed with deionized water, ethanol and anhydrous DMF for several times to remove any residues, respectively. The sediment of TiO2 nanoparticles was then dispersed in anhydrous DMF, followed by stirring and sonication. The TiO2/DMF solution for the coating of low-temperature TiO2 scaffold was 7 wt%, while that for the preparation of mixed precursor solution was 10 wt%. Mixed Precursor Solution PbI2 (461 mg, Sigma-Aldrich), MAI (159 mg, TCI), and DMSO (78 mg, Sigma-Aldrich) were dissolved into anhydrous DMF (600 mg, Sigma-Aldrich), the solution was stirred in a nitrogen filled glovebox at room temperature overnight. Then the TiO2/DMF solution was added to prepare different TiO2 concentration. The mixed precursor solution was stirred for 1 h before use. HTM Solution


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Spiro-OMeTAD (72.3 mg, Merck) was dissolved in chlorobenzene (1 mL, Sigma-Aldrich). 17.5 µL of Li-TFSI (lithium-bis(trifluoromethanesulfonyl)imide) solution (520 mg cm-3 Li-TFSI in 1mL acetonitrile) and 28.5 µL of TBP (4-tert-butylpyridine) were added. Devices Fabrication Fluorine-doped tin oxide (FTO, 14 Ω⁄□, transmittance 85%) glass substrates were ultrasonic bath washed successively in cleaning regent, deionized water, acetone and ethanol for 20 min each. The cleaned FTO substrate was further treated in ultraviolet ozone (UV-O3) for 15 min to remove the organic residues. A 40-nm-thick blocking layer was prepared with sol-gel method. The sol used here was prepared by mixing titanium tetraisopropoxide (TTIP) contained solution A (TTIP, ethanol, acetylacetone) and acid solution B (ethanol, HCl, H2O). The solution was spin coated onto FTO at a spin coating speed of 3000 rpm for 20 s. The substrate was calcined at 510 °С for 30 min in air. After cooling to room temperature, the compact TiO2 films were treated in 40 mM aqueous solution of TiCl4 for 40 min at 70 °С, and then rinsed with deionized water and ethanol. The TiCl4 treated substrates were again calcined at 510 °С for 30 min. The mixed precursor solution was spin-coated on the blocking layer at 5000 rpm for 20 s and diethyl ether (0.5 mL) was dripped on the rotating substrate at the 6th second. After the previous layer had been dried at 100 °С for 2 min, a neat MAPbI3 precursor was then dropped onto the substrate, diethyl ether was again applied during the spin coating. The last mentioned two steps were repeated for two more times to complete the coating of MAPbI3-TiO2 composite layer. The obtained composite film was then annealed at 100 °С for 20 min on a hot plate to form crystalline perovskite. The HTM solution was subsequently deposited on the top of the perovskite layer by spin coating from the spiro-OMeTAD solution at 4000 rpm for 30 s. Finally, 80-nm-thick silver acting as the counter electrode was thermally evaporated on top of the HTM


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layer. For the reference samples, a layer of TiO2 scaffold was deposited on top of the blocking layer by spin coating from the 7 wt% TiO2/DMF solution. The TiO2 film was heated at 100 °С for 30 min on a hot plate to remove DMF. A second layer of neat MAPbI3 was then spin coated on the as-prepared TiO2 scaffold to form a conventional mesostructured PSCs. For VASP method, PbI2 was dissolved in DMF at 85 °С (578 mg mL-1), and mixed with TiO2/DMF solution, where the molar ratio of PbI2 and TiO2 is 5:1. The mixed precursor solution was then spin coated onto the bl-TiO2 at 6500 rpm for 5 s. The substrate was dried on a 100 °С hotplate for 2 min to remove the remaining solvent. The as-prepared PbI2-TiO2 composite film was faced down at a constant distance of around 3 mm against the CH3NH3I powder. The reaction was conducted in a vacuum oven under the pressure of 100 Pa at 110 °С for 6.5 h. After the formation of perovskite, the film was rinsed with isopropanol for 45 s followed by drying with air. Finally, the MAPbI3-TiO2 composite films deposited by VASP method were annealed at 100 °С for 45 min to enhance the crystallinity. Characterization Scanning electron microscopy (SEM) and elemental analysis were performed using a fieldemission scanning electron microscopy (FEI Magellan 400) combined with energy dispersive Xray spectroscopy (EDS). X-ray diffraction (XRD) measurements were performed with Ultima IV X-ray diffractometer using Cu Kα radiation under operation condition of 40 kV and 40 mA at room temperature in the 2θ range of 5-60º, with a scanning speed of 4º min-1. Steady photoluminescence spectroscopy (PL) was recorded on a Shimadzu RF-5301PC device with an excitation wavelength of 470 nm. Time-resolved PL spectra were conducted with fluorescence lifetime spectrometer (Photo Technology International, Inc.). The UV-vis absorption spectra of the perovskite films were characterized on a Shimadzu UV-2550PC spectrometer in the


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wavelength ranging from 400 nm to 900 nm. The current density−voltage (J−V) curves were measured in air at room temperature with a solar simulator equipped with a 450 W xenon lamp and a Keithley-2420 source meter (AM 1.5G, 100 mW cm-2). Light intensity of the measurements was determined using a calibrated silicon solar cell (Oriel-91150) as reference for approximating 1 sun light intensity. The active area of cells was fixed at 0.07 cm2. Steady state current of the PSCs was measured for more than 2 min at a bias voltage equal to voltage corresponding to maximum power point (determined from J-V curve). Impendance spectroscopy was recorder under AM 1.5G illumination by the E4980A Precision LCR Meter at frequencies from 0.01 to 100K Hz. Photovoltage decay measurements were performed by a CHI660C Electrochemical Analyzer. RESULTS AND DISCUSSION The incorporated TiO2 nanoparticles used in the present study was synthesized following the literature28 as described in the experimental section. The as-prepared TiO2 sol was hydrothermally treated at 260 oC for 24 h to obtain well-crystallized nanoparticles with pure anatase phase (as seen in the X-ray diffraction (XRD) pattern of Figure 1a). Figure 1b shows the transmission electron microscopy (TEM) image of the uniformly distributed TiO2 nanoparticles with a diameter of approximately 20 nm. After cooling to room temperature, the resultant colloidal solution was centrifuged and washed with ethanol and anhydrous DMF for several times to remove any residue, respectively. The TiO2 sediment was then dispersed in anhydrous DMF and stirred vigorously over night to achieve a homogeneous suspension with a concentration of 10 wt% (Figure S1a). The resultant TiO2/DMF suspension could be stable for months without noticeable aggregation or precipitation. The mixed solution of MAPbI3/TiO2 was prepared by mixing TiO2/DMF suspension and MAPbI3 precursor of MAI, PbI2 and N,N’-


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dimethylsulfoxide (DMSO) in N,N’-dimethylformamide (DMF). The molar ratio of MAI: TiO2 in the mixing precursor was 1:1.25. As shown in Figure 1c, the translucent yellow color of the neat MAPbI3 precursor turns into opaque upon adding of TiO2 nanoparticles. We noticed that the dispersion of TiO2 nanoparticles in the mixed solution did not remain stable in the presence of environmental cations of MA+, Pb2+ and I- as shown in Figure S1a. TiO2 nanoparticles gradually precipitated and suspended at the lower part of the solution after several weeks of storage, and thus an hour of stirring before the deposition procedure is required. However, it should be noted that the dispersion of mixed solution is stable during the whole spin-coating process. The as-prepared mixed solution of MAPbI3 and TiO2 was employed for deposition of MAPbI3-TiO2 composite films (Figure S1b). The composite film of MAPbI3-TiO2 was firstly spin-coated on the substrate by one-step solution method and annealed at 100 oC for 20 min as illustrated in Figure 1d. The fast rotation speed that suppresses the gravity of TiO2 nanoparticles, along with the quick precipitation of perovskite crystals due to evaporation of solvent, results in a composite structure where the TiO2 nanoparticles are inserted into the matrix of perovskite as shown in Figure 2a. However, the one-step solution process reveals a very poor morphology and coverage. The large voids that penetrate down to the blocking layer are adverse for the transport of photogenerated carriers.29, 30 Besides, a large amount of TiO2 nanoparticles emerged on the surface of the film (Figure 2e) would act as recombination sites because of the direct contact between TiO2 and HTM. Therefore, the corresponding cell performance of PSCs based on spincoated MAPbI3-TiO2 composite film is quite poor (black J-V curve in Figure S2a). In order to block the possible contact between TiO2 and HTM, a drop of neat MAPbI3 precursor was dripped on the center of the composite film as shown in Figure 1d (denoted as 1-cycle coating of mixed solution and neat precursor). The drop spread immediately due to the high solubility of


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PbI2 and MAI in DMF. This step is very important as it creates a TiO2-rich part at the bottom and a neat perovskite part at the top in the whole suspension. After annealing, Figure 2b clearly shows that a top layer of pure perovskite is formed upon the composite film, which indicates that the secondary coating of a neat MAPbI3 precursor is a feasible way to form a quasi-cappinglayer that blocks the possible contact between TiO2 and HTM. Unfortunately, the conventional one-step solution method still produces incomplete coverage upon the MAPbI3-TiO2-1-cycle film (Figure 2f), which is induced by the secondary nucleation dominated process because of high heterogeneous nucleation density in the mixed precursor solution.31 To further improve the crystallinity of perovskite and the coverage of active layer, antisolvent dripping process (AS), in which the antisolvent of ethyl ether was dripped at the 6th second while spinning, was applied.32 However, as high heterogeneous nucleation density in the mixed precursor solution is achieved with existing TiO2 nanoparticles, the secondary nucleation is still prone to take place even with the antisolvent dripping. Large voids which are adverse for carriers transporting are unavoidably formed (Figure 2c). In contrast, with both 1-cycle coating and AS method, a dense layer which contains well-crystallized perovskite grains and incorporated TiO2 nanoparticles is formed for the MAPbI3-TiO2-1-cycle-AS film (Figures 2d). By the drop of a neat precursor, the TiO2 content in the whole suspension along with the heterogeneous nucleation density is reduced. And the primary nucleation due to the salting effect of antisolvent dominates the nucleation process,31 leading to a dense composite film. Undesirably, the device output is still substantially worse than those reported in the literature (blue J-V curve in Figure S2a), which is caused by the recombination sites of TiO2 nanoparticles that emerged on the surface of active layer (white dots in Figure 2h) due to the relatively high molar ratio of TiO2 to MAPbI3 (1.25:1). Although a second layer of neat perovskite was coated, the perovskite grains


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would squeeze the excess TiO2 nanoparticles to the top side of the active layer during the grain growth, which resulted in recombination of photogenerated carriers between TiO2 and HTM. To solve such a problem, we then modulated the molar ratio of MAI: TiO2 in the mixing precursor, which was changed from 10:1 to 5:1 and 1:1 (denoted as T0.1, T0.2 and T1). However, the TiO2 nanoparticles in the composite film for T0.1 is unevenly distributed (Figure 2i). To further fulfill complete interpenetration between MAPbI3 grains and TiO2 nanoparticles, multi-cycle-coating process was used as demonstrated in Figure 1d. After the previous layer had annealed at 100 oC for 2 min, another coating-cycle was initiated. By the drop of another mixed solution and neat precursor in the multi-cycle-coating process, the distribution of TiO2 nanoparticles in the composite films is improved with increased coating-cycles through dissolution and recrystallization.33 3-cycle coating process is sufficient for obtaining a uniform MAPbI3-TiO2 composite film containing well-developed MAPbI3 grains with TiO2 nanoparticles pinned at the GBs (Figures 2j-l). The MAPbI3-TiO2 composite film for T0.2 exhibits the well-interpenetrated morphology and fully coverage (Figure 2o). The BHJ PSC based on 3-cycle-coated composite film achieves decent efficiency above 15% as shown in Figure S2a (red J-V curve in Figure S2a). Figure 3 compares the XRD pattern of MAPbI3-TiO2 composite films prepared by varied printing process and TiO2 content, respectively. In Figure 3a, the directly deposited MAPbI3TiO2 composite film reveals a relatively low intensity of diffraction peaks, which indicates low crystallinity of perovskite. However, dramatically increased intensity ratios of (110)/(310) by the secondary coating of neat precursor have been observed no matter with or without AS method. In particular, the diffraction intensity ratio of (110)/(310) for the MAPbI3-TiO2-1-cycle-AS film is almost four times of the number for the MAPbI3-TiO2 film. This suggests that the secondary coating of neat MAPbI3 precursor provides improved crystallinity and orientation as well as


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blocking of possible contact between TiO2 and HTM. The main diffraction peaks, assigned to the (110), (220), (310) and (330) planes at 14.14º, 28.50º, 31.88º and 43.06º, respectively, are in identical positions for both pristine and MAPbI3-TiO2 composite films (Figure 3b), indicating that both precursor solutions produce the perovskites with tetragonal crystal structure.34 No secondary PbI2 phase has been detected for the composite films as shown in the XRD patterns (Figure 3b), which is consistent with the elemental analysis result in Figure S2b. The EDS mapping in Figure S2b shows that the atomic ratio for lead and iodine is about 1:3 for the MAPbI3-TiO2 composite film. The above mentioned results rule out the possibility that TiO2 nanoparticles affect the crystallite structure of perovskite, which is in agreement with the suggested structure of composite film: a perovskite layer with TiO2 nanoparticles pinned at the GBs as seen in Figure 2. We propose a plausible mechanism for the formation of the uniform and dense MAPbI3-TiO2 composite films. The crystallization of MAPbI3-TiO2 composite films involves two stages: the nucleation and the crystal growth processes of MAPbI3. At the nucleation stage, the film is composed of MAI, PbI2 and TiO2 dissolved in the DMSO/DMF solvent in the initial step during spinning (Figure 1d). TiO2 nanoparticles dispersed in the mixed solution firstly provide many heterogeneous nucleation sites which facilitate the nucleation of perovskite, due to the lower nucleation energy barrier at the heterogeneous interfaces.35 As secondary nucleation are prone to take place,31 composite films with large voids in the active layer are unavoidably obtained, even with ethyl ether dripping. In the second step by the dripping of a neat precursor on the previous layer, the drop spread immediately and provided a relatively lower TiO2 content in the whole suspension. Then the primary nucleation by ethyl ether dripping dominates the nucleation process, leading to a dense perovskite layer. At the grain growth stage by annealing, the


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boundaries of grains will migrate and small grains will grow and infiltrate into each other, while leaving the TiO2 nanoparticles pinned at GBs. By the drop of another mixed solution or neat precursor in the multi-cycle-coating process, the previous film is transformed into an intermediate state which is different from both initial precursor and final film. The TiO2 nanoparticles then suspend and redisperse in such an intermediate state. The distribution of TiO2 nanoparticles in the composite films is improved with increased coating-cycles through dissolution and recrystallization. 3-cycle coating is sufficient for achieving dense MAPbI3-TiO2 composite films with well-interpenetrated TiO2 nanoparticles and perovskite grains. In a word, a novel composite structure ranging from a MAPbI3-TiO2-dominated layer on the substrate side to a pure perovskite layer on the top side is accomplished by dissolution and recrystallization via 3cycle-coating strategy. To verify the elemental distribution in the composite film, energy dispersive X-ray spectroscopy (EDS) was performed on the cross section of the MAPbI3-TiO2 composite film. The EDS maps in Figure 4a clearly show that all the elements, i.e., lead, iodine and titanium are homogeneously distributed. To further confirm the better interpenetration of TiO2 nanoparticles and MAPbI3 grains by this novel method of dissolution and recrystallization, we prepared a control sample which contained the perovskite grains infiltrated into a former coated TiO2 scaffold. The TiO2 scaffold was spin-coated by TiO2/DMF solution (7 wt%) and dried at 100 oC to remove the solvent and form randomly packed pores for the infiltration of perovskite precursor solution. As seen in the cross-sectional SEM images of Figure 4b, voids appear frequently at the bottom of TiO2 scaffold, signifying a low uniformity of perovskite grains in the scaffold. By contrast, dissolution and recrystallization of MAPbI3 by multi-cycle-coating results in a dense layer of well-interpenetrated perovskite grains and TiO2 nanoparticles as shown in


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Figure 4d. Comparison on the elements concentration of Pb and Ti in Figure 4 strongly evidences the above observations. Atomic% of Pb at the mapping spot 1 close to the top of perovskite layer is 16.74, and it dramatically drops to 3.87 at the mapping spot 2 close to the FTO side while Ti concentration greatly increases (Figure 4c). This proves that it is difficult for the effective infiltration of perovskite into a former prepared TiO2 scaffold. By contrast, the interpenetration between TiO2 and MAPbI3 is greatly improved in the MAPbI3-TiO2 composite film. As shown in Figure 4e, atomic% of Pb and Ti at the mapping spot 3 close to the top of composite film and at the mapping spot 4 close to the FTO side are similar. Therefore, the complete interpenetration of TiO2 nanoparticles and MAPbI3 grains is fulfilled by the novel method of dissolution and recrystallization. We subsequently fabricated a new type of BHJ PSC devices using the as-prepared composite films, where the perovskite grains act as light absorber as well as electron donor and the TiO2 nanoparticles act as electron acceptor. The corresponding photovoltaic performances of BHJ PSCs with varied molar ratio of MAI: TiO2 (from 10:1 to 5:1 and 1:1) were examined and compared to those of PHJ and MS PSCs. This unique BHJ PSC has a device structure of FTO/blTiO2/MAPbI3-TiO2/spiro-OMeTAD/Ag, where bl-TiO2 is the TiO2 blocking layer, MAPbI3-TiO2 is the composite film of MAPbI3 and TiO2, spiro-OMeTAD is a HTM of (2,2’,7,7’-tetrakis-N,Ndi-4-methoxyphenylamino)-9,9’-spirobifluorene, and Ag is the silver contact. The device performance was characterized by current density-voltage (J-V) measurements under simulated AM 1.5G solar irradiation (100 mW cm-2). The detailed photovoltaic parameters are given in Table S1. As it was impossible for us to sinter TiO2 nanoparticles in the composite layer to improve their crystallinity, we also fabricated MS PSCs that deposited on low-temperature TiO2 scaffold (denoted as LT) as control samples. The low-temperature TiO2 scaffold (LT-TiO2) was


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spin-coated by a TiO2/DMF solution with a content of 7 wt% and dried at 100 oC for 30 min to remove the solvent. As seen in Figure 5, the prominent enhancement of PCE is observed through dissolution and recrystallization of MAPbI3-TiO2 composite film. This unique BHJ PSCs based on MAPbI3-TiO2 composite film show higher PCE than both conventional PHJ PSCs and MS PSCs, which is mostly attributed to the improved VOC and FF. The MS PSCs show lower VOC than the PHJ devices as more electrons injected to TiO2 lower the Fermi energy level at equilibrium between EF(TiO2) and EF(MAPbI3).36 However the MS PSCs show higher FF than the PHJ devices due to improved charge separation interfacial area between perovskite and TiO2. The lower VOC and higher FF result in equivalent PCE of 13.91% for MS PSCs compared to that of planar counterparts (13.92%). It should be noted that the MS PSCs based on LT-TiO2 show reductions on all the photovoltaic parameters of JSC, VOC, FF and PCE in comparison with the sintered-TiO2-based devices, which is caused by the poor infiltration of perovskite (Figure 4b) and the less efficient electron transport.37 However, by dissolution and recrystallization of MAPbI3, the BHJ PSCs exhibit improvements on both VOC and FF, which leads to an overall PCE of 15.06% that conquers both MS and PHJ devices. The increase of VOC for BHJ PSC compared to MS PSC may be caused by the more sufficient prevention of energy loss on the hetero-interface, which will be discussed later. The remarkable increased FF can be related to the significant enlarged interfacial area by the blend of donor and acceptor. Consequently, the novel BHJ MAPbI3-TiO2 structure demonstrates effective devices. The photovoltaic parameters of BHJ PSCs, are found to be influenced by TiO2 content. The average JSC firstly increases from 21.33 mA cm-2 for T0.1 to 22.17 mA cm-2 for T0.2, but drops to 20.21 mA cm-2 for the higher TiO2 content of T1. The average VOC remains relatively


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unaffected with the increase of TiO2 content from T0.1 to T0.2, then declines slightly to 1.064 V for T1. The average FF decreases slightly from 63.08% for T0.1 to 62.77% for T0.2, and further drops to 56.72% for T1. As a result, the average PCE increases from 14.54% for T0.1 to 15.06% for T0.2, while the high TiO2 content of T1 shows a lower PCE of 12.22%. The optimal molar ratio of MAPbI3-TiO2 for the highest performance is thus determined to be 5:1. The champion cell presents the highest PCE of 15.53% with a JSC of 22.51 mA cm-2, a VOC of 1.085 V and a FF of 63.59%. We further investigate the underlying reason of the changing trend in JSC, VOC and FF with varied TiO2 content. An increase on film thickness with increased TiO2 content in the mixed precursor has been observed in the cross-sectional images of Figure 2j-l. We attribute the improvement in JSC to the enhanced light harvesting properties arising from the thicker composite films, which is also confirmed by the enhanced absorption of the composite films with respect to the pristine perovskite film (Figure S3a). The absorption spectra in Figure S3a show a similar absorption onset at ~760 nm for all the films, exhibiting a little differences in their absorption intensities. The composite films have higher absorptions than those perovskite films that deposited on bl-TiO2 or meso-TiO2 layers, and the absorption intensities increase proportionally to the TiO2 content. IPCE spectra of the BHJ PSCs presented in Figure S3b shows the same trend in the change of the IPCE to that of the JSC obtained from J-V curves. The improvement in VOC and FF is caused by the effective electron injection as proved by the PL quenching results shown in Figure 6a. The PL intensity of the MAPbI3-TiO2 composite films decreases significantly due to the improved interpenetration between perovskite and TiO2. The PL intensity further reduces with increasing TiO2 content. However, the highest TiO2 content of T1 makes the TiO2 nanoparticles emerge on the surface of composite films as


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shown in Figure 2p. This leads to the possible contact between TiO2 and HTM materials and finally causes recombination of photogenerated carriers and thus lower JSC, VOC, FF and PCE. The enhanced photovoltaic property of MAPbI3-TiO2 composite film is a surprise since we would not expect the unsintered TiO2 nanoparticles at the GBs to be able to sustain rapid electron transfer. However, we do observe respectably improved performance of BHJ PSC over that of PHJ and MS devices. We speculate that the working mechanism of BHJ PSC is quite different from both PHJ and MS devices. The TiO2 nanoparticles at GBs are small enough that their deep trap sites are filled under solar illumination, leaving much shallower and more delocalized defects. This means that the deeper states are likely to be full, and mobile carriers are more likely to occupy shallow traps and so are more likely to move.38 The photogenerated electrons in perovskite are quickly extracted by TiO2 nanoparticles, in which the electrons are then “trapped”. Given the suitable energy level alignment, it is feasible that the thermal population of electrons in the TiO2 LUMO levels at the energy of the conduction band of the perovskite are free to transfer back and forth into the perovskite conduction band, such that the TiO2 nanoparticles effectively act as “electron reservoir”.39 The thermal electrons at this level will then be free to be transported through the crystalline perovskite absorber until they trap again in the TiO2, which is described in the multi-trapping model of TiO2.38 The proposed mechanism is consistent with the increased VOC and FF of BHJ PSCs. Since charge transport through the perovskite is more than an order of magnitude faster than through the TiO2,2 the inhibition of electron transfer in the TiO2 moves the EF for electrons closer to the perovskite conduction band and therefore leads to the higher VOC. The improvement of FF, on the other hand, is beneficial from increased interfacial area between TiO2 and perovskite. The interpenetration of TiO2 nanoparticles and perovskite grains is significantly improved as evidenced in Figure 4. Thus the fast extraction and


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reserve of electrons by TiO2 nanoparticles effectively prohibit the recombination of holes and electrons at the GBs, and consequently lead to the higher FF. In this regards, the TiO2 nanoparticles also function in the way of GB defects passivation. To prove the plausibility of the proposed operating physics for BHJ PSCs, we firstly examined the charge extraction on the MAPbI3-TiO2 heterojunction using steady-state and timeresolved photoluminescence spectroscopy. The steady-state photoluminescence (PL) spectra (Figure 6a) demonstrate that the incorporated TiO2 nanoparticles are capable of quenching the PL of perovskite, which implies that excitons are efficiently dissociated at the MAPbI3-TiO2 hetero-interface. With the increase of TiO2 nanoparticles content, the PL is further quenched. Figure 6b shows the PL decays of the perovskite film as well as perovskite composite films that incorporated with TiO2 or ZrO2. The spectra of the perovskite films coated on the TiO2 blocking layer monitored at the peak emission (765 nm) could be fitted to biexponential decays with a fast component and a slow component. The fast decay component τ could be assigned to the interfacial charge separation property40 (the detailed fit data are listed in Table S2). The MAPbI3TiO2 composite film has a faster τ than the pristine perovskite film, while the slower τ for the MAPbI3-ZrO2 composite film corresponds to the insulative behavior of ZrO2 material, which indicates the efficient extraction of the photo-generated electron from perovskite to TiO2. To examine the charge transfer efficiency, the conductivity of various films was measured by putting the films between two highly conducting electrodes and the conductivity was calculated from the I-V curves (Figure S3c).41 The TiO2 films spin-coated by TiO2/DMF solutions and annealed at 100 oC have lower conductivity than the sintered TiO2 film (Table S3), signifying insufficient charge transport. We further tested the conductivity of perovskite films as well as MAPbI3-TiO2 composite films, as shown in Figure 6c and Table S3. The conductivity of pristine


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perovskite film is 0.003 mS cm-1, and it significantly increases to 0.077 mS cm-1 by depositing the perovskite on a sintered mesoporous TiO2 film. The lower conductivity of perovskite film that deposited on LT-TiO2 is caused by the lower conductivity of LT-TiO2 film. The MAPbI3TiO2 composite film prepared by 1-cycle-coating exhibits even lower conductivity than that of pristine perovskite film, which might be caused by the carriers capture of surface trap states on isolated TiO2 nanoparticles. However, with 3-cycle-coating to improve the interpenetration of TiO2 nanoparticles and perovskite grains, the conductivity of MAPbI3-TiO2-3-cycle film is surprisingly improved to 0.093 mS cm-1, indicating sufficient charge transport in the BHJ MAPbI3-TiO2 composite film. The higher conductivity of MAPbI3-TiO2 composite film than that of meso-TiO2/MAPbI3 film is attributed to the reduced interfacial contact resistance which is from the vanishing interface between mesoporous TiO2 layer and perovskite capping layer in the composite film. We further carried out electrochemical impedance spectroscopy (EIS) and open-circuit photovoltage decay (OCVD) measurements to investigate the charge transport and recombination kinetics of the BHJ PSCs. EIS has been widely used in the characterization of PSCs, which is capable of providing useful information on the device working mechanism.42 The Nyquist plots were recorded at 0.80 V under illumination in the frequency range from 0.01 Hz to 100 kHz as depicted in Figure 6d. Generally, the arc at the high frequency region is ascribed to transfer resistance (Rtr), which could be related to the overall charge transport in the cell, including the contribution of bl-TiO2/perovskite interface, perovskite layer and perovskite/HTM interface.43 Given that both the blocking layer and HTM material are all identical in all cases, the Rtr is mainly associated with the modification of perovskite active layer, and its value reflects the electron extraction/transport properties in the perovskite layer. Compared to PHJ PSC, an


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evidently decreased Rtr value has been observed for BHJ PSCs with the incorporation of TiO2 nanoparticles, suggesting the sufficiently improved charge transfer property. The Rtr value for the cell of T0.2 is further reduced with the increasing TiO2 content. However, the increased Rtr value for the cell of T1 is caused by the recombination at the perovskite/HTM interface due to the direct contact between spiro-OMeTAD and TiO2 nanoparticles that exposed on the surface of perovskite layer. Furthermore, OCVD measurement, where the transient VOC was recorded as a function of time upon turning off the illumination, was applied to monitor the electron recombination kinetics of the devices.44 The illumination was kept for 30 seconds before shutting down to achieve a steady-state photovoltage for the devices. As shown in Figure 6e, once the illumination is turned off, the VOC decays sharply due to the photo-carrier recombination, and the VOC decay rate directly indicates the charge recombination rate. Briefly, the electron lifetimes (τn) could be derived from the photovoltage decay curve according to the following formula:

τ௡ =

‫ܭ‬஻ ܶ ܸ݀ை஼ ିଵ ൬ ൰ ‫ݍ‬ ݀‫ݐ‬

where KB is the Boltzmann constant, T is the absolute temperature, and q is the positive elementary charge.44, 45 Figure 6f shows the electron lifetimes as a function of photovoltage. It is observed that the BHJ PSCs present slower photovoltage decays and longer electron lifetimes than PHJ PSC due to the passivation of TiO2 nanoparticles at GBs. The increased electron lifetime of the BHJ PSCs indicates a lower charge recombination rate and a higher charge collection efficiency than those of PHJ PSCs, which is consistent with those observations from EIS measurement. It is noteworthy that the champion PSC with an optimized TiO2 content of T0.2 shows elongated electron lifetime and reduced charge recombination than both PHJ and MS


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PSCs. These elongated carrier lifetime and reduced recombination are consistent with former proposed working mechanism of BHJ PSCs. This novel BHJ PSC, as demonstrated in Figure 6g, has several advantages over those conventional MS or PHJ PSCs. Firstly, the dissolution and recrystallization of MAPbI3 via multicycle-coating process, which avoids the sintering step for mesoporous TiO2 scaffold at high temperature and lowers the fabrication cost, is helpful for the extending applications of PSCs to plastic substrates. And the complete interpenetration between TiO2 and perovskite, along with the prevention of direct contact between TiO2 and HTM, is successfully accomplished by this multi-cycle-coating strategy. Secondly, the grain size of perovskite grains in the MAPbI3-TiO2 composite layer is in the scale of hundreds of nanometers as shown in Figure 2, which conquers the constrained crystallization of perovskite grains within the mesoporous TiO2 scaffold. Perovskite within mesoporous TiO2 scaffold has equiaxed grains of size 10-20 nm, which is governed by the pore size of TiO2 layer, and they have relatively low crystallinity.46 The relatively larger grain size and higher crystallinity of perovskite in the composite layer would exhibit enhanced carriers lifetime and decreased recombination.47 Thirdly, the carriers recombination and ion migrations along the GBs would be diminished by the TiO2 nanoparticles at GBs. The TiO2 nanoparticles function in a GB passivation way as they quickly extract and reserve the electron before it transport in the perovskite, which effectively reduces the recombination loss at the GBs. The “pinning effect” of TiO2 nanoparticles, which comes from the binding between TiO2 and MAPbI3 at the GBs, results in a reduction of ion migration along the GBs. Ion migration has been confirmed to occur in the perovskite material and is found to be more pronounced along the GBs than in interior grains.13, 14, 48 The slow trapping and detrapping process at ionic vacancies leads to operational instability and hysteresis of PHJ PSCs.49 Through


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the pinning of TiO2 nanoparticles at the GBs as illustrated in Figure 6g, the strong binding between iodine atoms and titanium atoms on the major exposed plane of (101) in the anatase TiO2, which has been proved by both experiments and first-principle calculations,26, 27 is induced. In such a situation, the migration channels of the iodine ions can be effectively blocked or interrupted by the pinning of TiO2 nanoparticles. Therefore, efficiency loss and hysteresis behavior, which may be caused by the slow trap-filling process of iodine vacancy (VI) due to ion migration, would be greatly suppressed as demonstrated in Figure 6g. We fabricated BHJ PSCs with ZnO and ZrO2 nanoparticles to prove the benefits of pinning second-phase nanoparticles at the GBs. The cross-sectional SEM images in Figure S4 show that the ZnO and ZrO2 nanoparticles also accumulate along the GBs of perovskite grains, which verifies that this BHJ PSC is widely applicable to other nanomaterials. The improved VOC and FF for MAPbI3-ZnO BHJ structure are attributed to the higher conductivity and faster electron transport of ZnO (Figure S5). In contrast, as ZrO2 nanoparticles are not capable of transporting electrons, reductions on VOC, FF and PCE for MAPbI3-ZrO2 BHJ structure are observed. This comparison also proves that the improved device performance of BHJ MAPbI3-TiO2 PSCs is profited from the electron-extraction property of TiO2 nanoparticles. We further compared operational stabilities of BHJ PSCs with those of PHJ and MS PSCs (Figure S6 and S7). The calculated hysteresis index (HI)50 of BHJ PSCs with varied materials is much lower than that of PHJ PSC (Figure S6). Our former study has pointed out the relationship between hysteresis and ion migration.51 Therefore the greatly diminished HI by incorporation of a second-phase nanoparticles verifies our hypothesis that such pinning of metal oxide nanoparticles at the GBs effectively block the travel path of ions along the GBs and thus the subsequent slow trap-filling process of ion vacancies. Current density and PCE as a function of time in Figure S7 further


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confirm the operational stability of BHJ PSC. PHJ PSC shows the most extreme drop in stabilized power output in comparison with the reverse scan, which is consistent with its fierce hysteresis extent (Figure S6). The BHJ PSC, with the pinning of TiO2 nanoparticles, gives a stabilized output power of 16.02% under 0.85 V constant forward bias, which agrees well with the PCE of 16.75% by J-V measurement. The above-mentioned results have clearly demonstrated that our novel dissolution and recrystallization of MAPbI3-TiO2 is an effective approach to fabricate high performance BHJ PSCs. Based on these results, we applied a two-step vapour assisted solution process (VASP) for device fabrication to further verify the universality of the composite MAPbI3-TiO2 system. We prepared a mixed solution of PbI2 and TiO2 in DMF as the precursor, where the molar ratio of PbI2 and TiO2 is 5:1. A secondary coating of neat PbI2 and 3-cycle coating process were also used to form the PbI2-TiO2 layer. The spin-coated composite layer of PbI2-TiO2 was dried at 100 o

C and then exposed to the CH3NH3I vapour under a reduced pressure (100 Pa) at 110 oC for the

conversion of perovskite.52 A uniform layer with perfectly interpenetrated perovskite grains and TiO2 nanoparticles in a large scale is observed in the low magnification SEM image of Figure S8a. The efficiency deviation of BHJ PSCs as well as PHJ and MS PSCs (detailed photovoltaic parameters for all the cells listed in supplementary Table S4) is shown in Figure S8b. The BHJ PSCs based on MAPbI3-TiO2 composite film show relatively higher PCEs and narrower efficiency deviation. Figure 6h shows the J-V curve for the best-performing solar cell. The JSC, VOC and FF values from the J-V curve of this device are 22.43 mA cm-2, 1.084 V and 71.63%, respectively, these corresponding to a PCE of 17.42% under standard AM 1.5G conditions. Stability of the unencapsulated cells of PHJ, MS and BHJ devices was evaluated in dark at room temperature with a relative humidity of ~1%. As shown in Figure S9, for the PHJ PSCs, the


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average PCE decreases to 65% of its initial value after being stored for 8 days. And the average PCE of MS PSCs decreases to 87% of its initial value for the same period. In comparison, the BHJ PSCs exhibits the best long-term stability, and the average PCE approaches 94% of the initial value over 8 days. In a word, the above results demonstrate that the novel BHJ PSCs show improved PCE, suppressed hysteresis, and enhanced stability compared with those of PHJ and MS architectures. CONCLUSION In conclusion, a novel structure, namely BHJ PSC, has been successfully fabricated via dissolution and recrystallization of MAPbI3 by multi-cycle-coating strategy. The introducing of TiO2 nanoparticles into perovskite precursor can lead to the formation of a BHJ MAPbI3-TiO2 composite film, in which MAPbI3 acts as the electron donor and TiO2 acts as the electron acceptor. A unique structure, which contains a MAPbI3-TiO2-dominated layer on the substrate side and a pure perovskite layer on the top side, is achieved by the multi-cycle-coating process. EDS analysis confirms the complete interpenetration between TiO2 nanoparticles and MAPbI3 in the composite film. Interestingly, the incorporated TiO2 nanoparticles is found to be accumulated along the GBs of perovskite as observed in SEM images. On the one hand, the quick electron extraction and electron reserve properties of TiO2 nanoparticles at GBs efficiently passivate the surface and retard the electron-hole recombination, which results in increased VOC and FF. On the other hand, the pinning effect of TiO2 nanoparticles at the GBs effectively blocks or interrupts the migration channels of iodine ions along the GBs through the binding of perovskite iodine atoms and titanium atoms. It is beneficial to the suppressing of ion migration and diminishing of hysteresis. With reduced recombination loss and enhanced carrier lifetime that have been proved by EIS and OCVD measurements, this novel architecture of BHJ PSCs


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exhibits higher efficiency and better operational stability as compared to its planar or mesostructured counterparts. Finally, a PCE of 17.42% with a JSC of 22.43 mA cm-2, a VOC of 1.084 V and a FF of 71.63% has been achieved. The novel structure of BHJ PSC successfully provides a solution to address two major problems of PSCs: the high temperature sintering process of mesostructured devices, which limits large-scale manufacturing of PSCs; and the fierce hysteresis of planar devices, which hinders the reproducibility of PSCs. This simple solution process by mixing TiO2 nanoparticles in perovskite precursor offers a new and effective way to realize the low-cost and high-output manufacturing of PSCs. ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge on the ACS Publications website (PDF). Images of TiO2 sol, TiO2/DMF solution, MAPbI3-TiO2 mixing solution and perovskite films, J-V curves of BHJ PSCs, EDS analysis of MAPbI3-TiO2 composite film, detailed photovoltaic parameters of PSCs, UV-vis absorption spectra of MAPbI3-TiO2 composite films, IPCE spectra of BHJ PSCs, I-V characteristics of varied TiO2 films, fast component τ of perovskite films, conductivity of varied TiO2 films and perovskite films, SEM images of MAPbI3-ZnO and MAPbI3-ZrO2 composite films, comparison on photovoltaic properties of BHJ PSCs incorporated with ZnO and ZrO2, hysteresis behavior and steady-state output of BHJ PSCs compared with PHJ and MS devices, low magnification SEM and cell performance of BHJ PSCs prepared by VASP method, long-term stability of PHJ, MS and BHJ devices. AUTHOR INFORMATION Corresponding Author


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*E-mail: [email protected] *E-mail: [email protected] Notes The authors declare no competing financial interests. ACKNOWLEDGMENT This work was financially supported by the National High Technology Research and Development Program of China (Grant No.2014AA052002), Shanghai Municipal Natural Science Foundation (Grant No. 16ZR1441000), Shanghai Municipal Sciences and Technology Commission (Grant No. 12DZ1203900) and the Shanghai High & New Technology's Industrialization Major Program (Grant No. 2013-2). REFERENCES (1) 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. (2) Liu, M.; Johnston, M. B.; Snaith, H. J. Efficient Planar Heterojunction Perovskite Solar Cells by Vapour Deposition. Nature 2013, 501, 395-398. (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) 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


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(13) Yun, J. S.; Seidel, J.; Kim, J.; Soufiani, A. M.; Huang, S.; Lau, J.; Jeon, N. J.; Seok, S. I.; Green, M. A.; Ho-Baillie, A. Critical Role of Grain Boundaries for Ion Migration in Formamidinium and Methylammonium Lead Halide Perovskite Solar Cells. Adv. Energy Mater. 2016, 6, 1600330. (14) Shao, Y.; Fang, Y.; Li, T.; Wang, Q.; Dong, Q.; Deng, Y.; Yuan, Y.; Wei, H.; Wang, M.; Gruverman, A.; Shield, J.; Huang, J. Grain Boundary Dominated Ion Migration in Polycrystalline Organic–inorganic Halide Perovskite Films. Energy Environ. Sci. 2016, 9, 1752-1759. (15) Nakita K. Noel.; Antonio Abate.; Samuel D. Stranks.; Elizabeth S. Parrott.; Victor M. Burlakov.; Alain Goriely.; and Henry J. Snaith. Enhanced Photoluminescence and Solar Cell Performance via Lewis Base Passivation of Organic–Inorganic Lead Halide Perovskites. ACS Nano 2014, 8, 9815-9821. (16) Hadadian, M.; Correa-Baena, J.-P.; Goharshadi, E. K.; Ummadisingu, A.; Seo, J.-Y.; Luo, J.; Gholipour, S.; Zakeeruddin, S. M.; Saliba, M.; Abate, A.; Grätzel, M.; Hagfeldt, A. Enhancing Efficiency of Perovskite Solar Cells via N-doped Graphene: Crystal Modification and Surface Passivation. Adv. Mater. 2016, 28, 8681-8686. (17) Chen, Q.; Zhou, H.; Song, T.-B.; Luo, S.; Hong, Z.; Duan, H.-S.; Dou, L.; Liu, Y.; Yang, Y. Controllable Self-Induced Passivation of Hybrid Lead Iodide Perovskites toward High Performance Solar Cells. Nano Lett. 2014, 14, 4158-4163. (18) Park, N.-G.; Grätzel, M.; Miyasaka, T.; Zhu, K.; Emery, K. Towards Stable and Commercially Available Perovskite Solar Cells. Nat. Energy 2016, 1, 16152. (19) Anaraki, E. H.; Kermanpur, A.; Steier, L.; Domanski, K.; Matsui, T.; Tress, W.; Saliba, M.; Abate, A.; Grätzel, M.; Hagfeldt, A.; Correa-Baena, J.-P. Highly Efficient and Stable Planar


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FIGURES


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Figure 1. (a) X-ray diffraction (XRD) pattern of hydrothermally synthesized TiO2 nanoparticles. (b) Corresponding transmission electron microscopy (TEM) image of the TiO2 nanoparticles. (c) Illustration of pristine MAPbI3 precursor and the mixed solution of MAPbI3 and TiO2. (d) Schematic representation of the depositing process used to form the BHJ composite layer of MAPbI3-TiO2.


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Figure 2. Cross-sectional and top view SEM images of varied MAPbI3-TiO2 composite films. (a) Deposited by mixed solution of MAPbI3 and TiO2, (b) Deposited with a secondary coating of neat MAPbI3 precursor, (c) Deposited with antisolvent dripping, (d, i) Deposited with antisolvent dripping and 1-cycle coating of mixed solution and neat precursor, the molar ratio of MAI: TiO2 is varied from (d) 1: 1.25 to (i) 5:1. (j-l) Deposited with antisolvent dripping and 3-cycle coating of mixed solution and neat precursor, the molar ratio of MAI: TiO2 is varied from (j) 10: 1, to (k) 5:1 and (l) 1:1. (e-h, m-p) Corresponding top-view SEM morphologies of the MAPbI3-TiO2 composite films (a-d) and (i-l). The scale bars correspond to 500 nm in length.


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Figure 3. (a) XRD patterns of the MAPbI3-TiO2 composite films prepared by varied deposition procedure. (b) XRD patterns of the 3-cycle-coated MAPbI3-TiO2 composite films with varied TiO2 content compared to pristine MAPbI3 film.


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Figure 4. (a) Cross-sectional SEM image of MAPbI3-TiO2 composite layer and the corresponding EDS maps showing the relative distribution of Ti, Pb and I. (b) Cross-sectional SEM image of perovskite grains infiltrated into a former coated TiO2 scaffold. (c) The corresponding EDS spectrums of the mapping spots 1 and 2. (d) Cross-sectional SEM image of MAPbI3-TiO2 composite layer. (e) The corresponding EDS spectrums of the mapping spots 3 and 4. The inserted tables in the elementary spectrums show the atomic percentages of the corresponding elements.


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Figure 5. Short-circuit current density (JSC), open-circuit voltage (VOC), fill factor (FF) and power conversion efficiency (PCE) of the BHJ PSCs with varied TiO2 content as compared to those of PHJ PSCs and MS PSCs based on sintered TiO2 scaffold or low-temperature TiO2 scaffold.


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Figure 6. (a) Steady state PL spectra of MAPbI3-TiO2 composite films with varied TiO2 content as well as pristine perovskite films that deposited on bl-TiO2 or meso-TiO2. (b) The timeresolved PL decay at 765 nm obtained for the pristine perovskite film and composite films upon excitation at 467 nm. (c) I-V characteristics of the FTO/varied perovskite films/Au devices. (d) Nyquist plot of impedance spectra of BHJ PSCs with varied TiO2 content compared to PHJ PSC and MS PSC. (e) Open-circuit voltage decay and (f) electron lifetimes of BHJ PSCs with varied TiO2 content compared to PHJ PSC and MS PSC. (g) Schematic of BHJ PSC device architectures. (h) J-V curve of the best performing BHJ PSC fabricated by VASP method.


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