Gradient Sn-Doped Heteroepitaxial Film of Faceted Rutile TiO2 as an

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Gradient Sn-Doped Heteroepitaxial Film of Faceted Rutile TiO2 as an Electron Selective Layer for Efficient Perovskite Solar Cells Tingting Wu,†,‡ Chao Zhen,† Huaze Zhu,†,‡ Jinbo Wu,†,‡ Chunxu Jia,†,‡ Lianzhou Wang,§ Gang Liu,*,†,‡ Nam-Gyu Park,∥ and Hui-Ming Cheng*,†,⊥

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Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China ‡ School of Materials Science and Engineering, University of Science and Technology of China, 72 Wenhua Road, Shenyang 110016, China § Nanomaterials Centre, School of Chemical Engineering and AIBN, The University of Queensland, St Lucia, Brisbane, Queensland 4072, Australia ∥ School of Chemical Engineering, Sungkyunkwan University, Suwon 440-746, Korea ⊥ Tsinghua-Berkeley Shenzhen Institute, Tsinghua University, 1001 Xueyuan Road, Shenzhen 518055, China S Supporting Information *

ABSTRACT: The high-efficiency photocarrier collection at the interfaces plays an important role in improving the performance of perovskite solar cells (PSCs) because the photocarrier effective diffusion lengths in the lead halide perovskite absorbers usually surpass the incident depths of light. Developing the electron selective layer (ESL) that has good interfaces with photoactive perovskite and current collector layer-like fluorine-doped tin oxide (FTO) is actively pursued. Here, an unusual dense film of faceted rutile TiO2 single crystals with a gradient of the Sn4+ dopant grown heteroepitaxially on the FTO layer is obtained by a hydrothermal route and subsequent thermal treatment. Owing to the global features including low concentration of defects, atomically smooth coherent interface with FTO, and gradient doping-induced built-in electric field to promote the collection of photoelectrons in it, an optimal PSC with such a film as the ESL exhibits an efficiency of 17.2% with an open-circuit voltage of 1.1 V and fill factor of 76.1%, which are among the highest values of the PSCs with rutile TiO2 films as ESLs. KEYWORDS: heteroepitaxial growth, rutile TiO2, gradient Sn doping, electron selective layer, perovskite solar cell

1. INTRODUCTION Converting solar energy into electricity with photovoltaic devices has been regarded a promising route to realize sustainable energy supply.1,2 Metal halide perovskite solar cells (PSCs) as a representative of new generation of photovoltaic devices have attracted extensive attention because of their high efficiency, low cost, and solution processing for the fabrication, and their efficiency has rapidly increased from 9.7% to >23% since the first report on the 500 h-stable solidstate PSC in 2012.3−9 This extraordinary high efficiency is intrinsically associated with the very excellent optoelectronic properties of metal halide perovskite-based materials. Their high absorption coefficient (up to 104 cm−1) requires that the optical active layer in PSCs has only hundreds of nanometers for fully capturing sunlight. Such thickness is apparently shorter than the diffusion length exceeding 1 μm of the photocarriers in the layer,10−14 so that most of the photocarriers can in principle arrive at two interfaces involved with the selective contact layers for the separation. Therefore, controlling the selective contact layers and thus related © XXXX American Chemical Society

interfaces might release room for further improving the performance of PSCs.15−17 Crystalline TiO2 is considered to be a promising electron selective layer (ESL) material and most frequently used in PSCs. Most of the high efficiencies in the development of PSCs were achieved in the devices with TiO2 as the ESL.18−22 However, its intrinsic shortcomings including low charge carrier mobility and abundant defect trap states might restrict the photoelectron collection processes at the TiO2/perovskite interface, TiO2 bulk, and also transparent conducting oxide (TCO)/TiO2 interface. Some strategies such as doping and surface modification have been often used to increase the electron transport ability in TiO2 and transfer efficiency across the TiO2/perovskite interface. Doping can improve the electric conductivity of TiO2 by increasing the carrier concentration (i.e., Nb5+ doping) or the mobility (i.e., Sn4+ doping), and also Received: March 11, 2019 Accepted: May 7, 2019

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DOI: 10.1021/acsami.9b04308 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces

obtained were then transferred on a hot plate and annealed at 100 °C for 90 min to convert them into perovskite films thoroughly. 2.3. Solar Cell Fabrication. Spiro-OMeTAD hole selective layers (HSLs) were subsequently deposited on the perovskite films obtained from the above procedures by spin-coating with a chlorobenzene solution containing 72.3 mg/mL of spiro-OMeTAD, 28.8 μL/mL of tert-butylpyridine and 17.5 μL/mL of bis(trifluoromethane)sulfonimide lithium salt (520 mg/mL in acetonitrile) at 5000 rpm for 30 s. The samples were kept in a dry box overnight before a 70 nm-thick Au layer was casted at the HSL top with thermal evaporation equipment (the base pressure of about 2 × 10−4 Torr and the deposition rate of about 1 Å/s). 2.4. Characterization. J−V measurements of solar cells with an active area of 0.09 cm2 were carried out in ambient air, using a 3A solar simulator (Newport, 91192) as the light source and an EC-Lab (Bio-Logic SP-200) to record the signal. The scan direction for the J− V curve is from 1.3 to 0 V with a scan rate of 200 mV/s. The light intensity was calibrated to 1 sun (100 mW cm−2) with a standard Si solar cell from Newport Company, and all the devices were tested under reverse scanning mode. The external quantum efficiency spectrum was measured with a QTest Station 1000AD system (Crowntech, Inc.). Field-emission scanning electron microscopy (SEM) (FEI NanoSEM 430) was used to acquire SEM images. High-resolution transmission electron microscopy (HRTEM) images were taken on a FEI Tecnai G2 F30 transmission electron microscope operated at 300 kV. X-ray diffraction (XRD) patterns were obtained with a Rigaku diffractometer (D/max 2400) using Cu Kα (λ = 1.54050 Å) as the radiation source. UV−vis absorption spectra were recorded on a JASCO-770 spectrophotometer at room temperature. Transient photocurrent and voltage spectra were measured with a homemade instrument. The device was excited by a laser radiation pulse with a wavelength of 532 nm and pulse width of 6 ns. The solar cell was measured under ambient air.

reduce surface defect trap states (i.e., Al 3+ or Zr 4+ doping).23−25 Alternatively, surface modification with a monolayer of functional molecules (or ions) is also effective to reduce defect trap states at the TiO2/perovskite interface. Two representatives are using fullerene derivatives and small organic molecules with oxygen-containing functional groups such as −COOH and directional groups (i.e., −NH2, −SH, and −Cl) to control the perovskite growth.22,26−30 Particularly, −Cl was reported to be an effective surface modifier of TiO2 ESLs by forming a bridge bonding between TiO2 and perovskite to suppress the formation of defect trap states at the interface, and the fabricated devices with an active area of 0.049 and 1.1 cm2 showed the certified efficiencies of 20.1 and 19.5%, respectively.22 Despite these successes on the basis of tailoring the TiO2/perovskite interface and TiO2 itself, the underlying influence of the TCO/TiO2 interface on photoelectron collection has not been fully explored and deserves to be concerned in the designing of ESLs. The barrier existing at the TCO/TiO2 interface was considered to restrict the photoelectron transfer in dye-sensitized solar cells.31 Engineering the interfaces of planar-structured PSCs, especially the TCO/TiO2 interface by lowering the work function of indium tin oxide to reduce the Schottky barrier led to a record efficiency of 19.3%.32 Developing the strategy of tailoring the properties of the TCO/TiO2 interface remains highly desirable in terms of further releasing space for performance enhancement. In this study, concerning the great importance of the TCO/ TiO2 interface as well as the properties of TiO2 itself and TiO2/perovskite interface, an unusual film of gradient Sndoped rutile TiO2 single crystals with well-defined facets, which are obtained by heteroepitaxially growing on the fluorine-doped tin oxide (FTO) substrate and subsequent heating in air, is constructed as an efficient ESL of the PSC. As a result of the dual functions of such an ESL in both lowering the FTO/TiO2 interface barrier and promoting TiO2 bulk transport, the related PSCs exhibit an efficiency of 17.2% with a short circuit current density (Jsc) of 20.3 mA cm−2, opencircuit voltage (Voc) of 1.11 V, and fill factor (FF) of 76.1%, apparently surpassing that of the devices based on the traditional spin-coated TiO2 ESL.

3. RESULTS AND DISCUSSION 3.1. Design Guidelines of the TiO2 ESL. Two major concerns in the designing of the ESL are good bulk transport of the photocarriers and also their excellent FTO/TiO2 interface transfer. In order to enable good bulk transport of the photocarriers, both reducing the amount of bulk defects and grain boundaries that usually increase electron scattering probability and improving the migration ability of the charge carriers are desirable. For promoting FTO/TiO2 interface transfer of photocarriers, the high-quality interface with good lattice match and low concentration of defects is definitely favorable. A faceted rutile TiO2 crystal film with a gradient of the Sn4+ dopant that is heteroepitaxially grown on the FTO substrate is anticipated to probably have the features to meet all requirements aforementioned based on the following points: (1) the introduction of Sn4+ dopants in TiO2 can increase the electron mobility in TiO2 because more delocalized Sn 5s orbitals are involved in Ti 3d orbitals dominated by the TiO2 conduction band (CB).33 Moreover, a built-in electric field perpendicular to the FTO substrate might be induced by forming a gradient of the Sn4+ dopant with the maximum at the FTO/TiO2 interface, so the electron migration toward FTO can be accelerated. (2) The heteroepitaxial coherent interface between FTO (SnO2) and TiO2, which share the rutile structure, is atomically smooth and has a small potential difference. Such an interface favors the transfer of electrons from TiO2 to FTO. 3.2. Fabrication of the Sn-Doped TiO2 ESL and Microstructures. FTO and TiO2 have been extensively used as the transparent conductive electrode and ESL material, respectively, in normal-structured PSCs. One remarkable feature highlighted here is that rutile TiO2 shares the same

2. EXPERIMENTAL SECTION 2.1. Growing TiO2 Films on the FTO Layer. The rutile TiO2 single-crystal array films were grown heteroepitaxially on the FTO substrates by a modified hydrothermal method as follows. FTO (Pilkington sheet glass, 8 Ω per square)-coated glass was ultrasonically cleaned sequentially in detergent, ethanol, acetone, and isopropanol each for 15 min. The cleaned FTO glass was then treated by oxygen plasma for 15 min. A titanium precursor solution containing 0.0075 M titanium tetrachloride (TiCl4) was prepared by slowly adding TiCl4 into 3.5 M hydrochloric acid (HCl) solution in an ice water bath. FTO substrates were placed in a Teflon-lined autoclave filled with 40 mL of such titanium precursor solution containing 20 mg NaF. The autoclave is kept in an oven at 230 °C for 2 h. After the reaction, the obtained FTO glasses with grown TiO2 single-crystal arrays were washed with deionized water and dried. Sndoped rutile TiO2 films were obtained by heating the hydrothermally grown TiO2 films supported by FTO glasses in a muffle furnace at 500 °C for 30 min. 2.2. Perovskite Film Deposition. The films were deposited on the substrates by spin-coating from a dimethylformamide solution containing 2.64 M methylammonium iodide and 0.88 M PbCl2 at 3000 rpm for 1 min in a glove box filling with Ar gas. The samples B

DOI: 10.1021/acsami.9b04308 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces crystal structure (P42/mnm) with SnO2, which is the matrix material of the FTO substrate, and they have similar crystal parameters. Rutile TiO2 nanostructures heteroepitaxially grown on the FTO substrates have been investigated as photoanodes for photoelectrochemical applications.34−36 Because the SnO2 crystals in the FTO substrate have a lateral size of several hundreds of nanometers, the obtained TiO2 nanostructures usually with a lateral size smaller than 100 nm are unable to fully cover the FTO substrate so that they cannot be used directly as an ESL layer. To solve this limitation, an early attempt was that a dense TiO2 layer as the block layer was introduced to cover the FTO substrate before growing rutile TiO2 nanorods as ESLs in the PSCs with an efficiency of 9.4%.36 Up to date, the heteroepitaxial growth of a dense rutile TiO2 film that can fully cover the FTO substrate that seems to be an ideal ESL has not been realized yet, to our best knowledge. Herein, by using a new synthesis route, such a high-quality film can be directly grown on the FTO substrate and a subsequent thermal treatment process can lead to the formation of a gradient of tin doping in the film, accompanying the quenching of bulk defects, as illustrated in Figure 1.

Figure 2. Top-view SEM images of the (a) commercial FTO layer on the glass substrate and (b) rutile TiO2 single-crystal arrays growing heteroepitaxially on it. (c) Cross-sectional TEM image and (d) HRTEM image of the rutile TiO2 single-crystal array growing on the FTO substrate.

of 0.325 nm is assigned to (110) planes of rutile TiO2. As anticipated, the sharp and atomically continuous interface indicates a coherent crystal structure between TiO2 and FTO as a result of the nature of the heteroepitaxial growth. It is noteworthy that there is an angle of around 14° between (110) planes of rutile TiO2 and SnO2, which arises from the crystal parameter difference between rutile TiO2 and SnO2 as shown in Figure S5. The fact of the smaller spacing (0.325 nm) of TiO2(110) planes than that (0.335 nm) of SnO2(110) planes requires the TiO2 crystal to tilt an angle of θ in order to coherently grow on SnO2(110) planes. The determined value of θ from the relationship of cos(θ) = 0.325/0.335 is around 14.3°, which is in good agreement with the value measured in the HRTEM image. Because of the relatively low temperature and also enclosed space under the hydrothermal condition, many defects always exist in the TiO2 crystals obtained.38,39 The process of thermal treatment at 500 °C in air was carried out with the purpose of reducing the amount of defect trap centers. To evaluate the change of the defect density before and after the thermal treatment, Mott−Schottky electrochemical impedance spectra (EIS) and UV−vis−NIR absorption spectra were collected. The Mott−Schottky curve gives the relationship of surface capacitance to the negative two (C−2) and applied bias (Vbia) in the form of (C−2) = 2(ε0εrNdA2e)−1[(V − Vfb − kBT/e)], where Nd is the donor state density, ε0 is the permittivity of free space, εr is the relative dielectric constant of a semiconductor, A is area of the electrode, kB is the Boltzmann constant, T is temperature, e is the electronic charge, V is applied bias, and Vfb is the flat band potential. The slope of the Mott−Schottky curve is inversely proportional to the defect density. The treated sample has a much larger slope than the pristine one (Figure S6), indicating the greatly lowered concentration of defects in the treated sample. This change can be further confirmed by the comparison of their UV−vis− NIR absorption spectra. The absorption peaks at the long

Figure 1. Schematic of cross sections of the FTO layer supported on glass, rutile TiO2 single-crystal arrays growing heteroepitaxially on the FTO layer, Sn-doped rutile TiO2 single-crystal arrays growing heteroepitaxially on the FTO substrate after thermal treatment. The white points in the third picture represent the Sn atoms in TiO2.

The XRD pattern and top-view and cross-sectional SEM images of a commercial FTO layer supported on the glass substrate were collected to show the features of SnO2 crystals in the FTO layer. The diffraction peaks in the XRD pattern (Figure S1) are assigned to rutile phase SnO2. The enhanced intensity of the diffraction peak at 26.6° (2θ) suggests the preferential growth of the crystals along the crystallographic [100] orientation. The combination of the top-view and crosssectional SEM images (Figures 2a and S2) displays the pillarlike shape of the SnO2 crystals with regular facets exposed. Heteroepitaxial growth of the dense film of rutile TiO2 crystals on the FTO layer, which was carried out in HCl aqueous solution containing the TiCl4 precursor and F− capping agent under the hydrothermal condition,37 leads to the conformal morphology of the crystals as shown in Figure 2b, which is similar to that of SnO2 crystals. Compared to the rough surface and obtuse edges/corners of SnO2 crystals, the TiO2 crystals have a very smooth surface and sharp edges/corners. This is largely caused by the high crystallinity of the TiO2 crystals as indicated by the strong signals of the XRD pattern together with the Raman spectrum (Figures S3 and S4). The crosssectional TEM image of the TiO2/FTO glass in Figure 2c shows that the TiO2 film consists of a single layer of TiO2 crystals, and the layer thickness is around 200 nm. The atomic structures of the regions around the TiO2/FTO interface were investigated on the basis of the HRTEM image in Figure 2d. The region showing the lattice fringes with a spacing of 0.335 nm is assigned to the (110) planes of rutile SnO2, and the region showing the lattice fringes with a spacing C

DOI: 10.1021/acsami.9b04308 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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Figure 3. (a) HAADF image of the as-prepared TiO2 film growing heteroepitaxially on the FTO substrate. (b−e) Sn, Ti, O, and F element energydispersive X-ray spectrometry (EDS) mappings across the as-prepared TiO2/FTO interface region. (f) HAADF image of the TiO2 film on the FTO substrate after thermal treatment. (g−j) Sn, Ti, O, and F element EDS mappings across the TiO2/FTO interface region of the TiO2 film on the FTO substrate. (k,l) Depth-dependent SIMS spectra of Ti and Sn elements in the TiO2 film grown on the FTO substrate before and after the thermal treatment. (m) Schematic of band structure evolution of TiO2 with a gradient of the Sn4+ dopant.

appears near the FTO/TiO2 interface (Figure 3g,h) after the thermal treatment. It means that the thermally driven diffusions of Sn4+ ions from the FTO to TiO2 film and Ti4+ ions from the TiO2 to FTO film occur in the sample. Note that the diffusion length of Sn4+ ions in the TiO2 film is much longer than that of Ti4+ ions in the FTO layer largely because of the lower melting point of SnO2 than TiO2. On the other hand, the O element as a common composition of FTO and TiO2 distributes in the whole image, and its signal intensity is stronger at the TiO2 region due to the partial replacement of O with F in the FTO (Figure 3d). The thermal treatment leads to no obvious change in its distribution (Figure 3i). The lowered signal intensities of F in both FTO and TiO2 regions after the thermal treatment are caused by its thermal evaporation (Figure 3e,j). This is further confirmed by the reduced signal of fluorine in the X-ray photoelectron spectroscopy spectrum of F 1s in the thermally treated TiO2 film (Figure S10). Meanwhile, no Sn signal was detected in the TiO2 films before and after the thermal treatment, suggesting that the diffusion depth of Sn ions in the TiO2 film is shorter than the film thickness. The diffusion behavior of Sn4+ and Ti4+ in the heated TiO2/ FTO sample was further investigated and verified by secondary ion mass spectroscopy (SIMS). With the increase of the sputtering time that corresponds to the increased detection depth in the sample, the signal intensity of Ti4+ second ions for these two samples remains nearly constant before 1500 s and rapidly decays after 1500 s (Figure 3k), and the signal intensity

wavelength region beyond 1500 nm in the spectra are typically originated from the absorption of low-energy photons from some donor defective states in the band gap to the CB of ntype TiO2. The weakened absorption peak in this range for the thermally treated sample means the reduced density of donor states arising from the defects (Figure S7). Meanwhile, the transmittance of the thermally treated sample is apparently higher than that of the pristine one in the visible light range due to the enhanced crystallinity and lowered defect concentration (Figure S8), which is favorable for improving the light absorption efficiency of the perovskite active layer. However, the transmittance of the thermally treated sample is still relatively lower than those of the other reported TiO2 ESLs in the literature due to its larger thickness (∼200 nm). Owing to the high stability of the heteroepitaxial interface and also large crystal size, the rutile TiO2 film after the thermal treatment can well maintain its morphology with sharp edges and corners of the crystals (Figure S9). High-angle annular dark-field (HAADF) images (Figure 3a,f) also demonstrate its original compactness and seamless connection of the crystals. Besides lowering the concentration of defects, the thermal treatment simultaneously causes the diffusion of Sn ions from the FTO layer toward the TiO2 film via their interface. Compared to the sharp edges of the Sn and Ti element mappings that are, respectively, located in the FTO and TiO2 layers of the sample before the thermal treatment (Figure 3b,c), an overlapping zone of Sn and Ti element mappings D

DOI: 10.1021/acsami.9b04308 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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Figure 4. Statistic histograms of (a) open-circuit voltage (Voc) and (b) efficiency of devices based on as-prepared (undoped) TiO2 ETLs (black columns) and thermally treated (Sn doped) ones (red columns). (c) J−V curve of the best performance cell tested under AM 1.5 sunlight simulator irradiation with a power density of 100 mW cm−2 and (d) IPCE spectrum of the best performance device based on the Sn-doped TiO2 ETL.

electric field perpendicular to the FTO substrate (see Figure 3m-III), as experimentally confirmed in the literature.41,42 The built-in electric field may play three positive roles: (1) accelerating electron migration toward the FTO layer; (2) suppressing the recombination of electrons and holes and (3) reducing transfer barrier formed at the FTO/TiO2 interface. Therefore, it is anticipated that the gradient Sn-doped TiO2 film supported by the FTO substrate could have the better ability of collecting electrons than its counterpart without Sn doping. 3.3. Fabrication of Solar Cells and Performance. Two sets of PSCs with the rutile TiO2 single-crystal array films without and with the thermal treatment (Sn doping) as ESLs are fabricated. The perovskite active material used in this study is MAPbI3, which is deposited on the TiO2 ESL by a one-step spin-coating method. Then, the spiro-OMeTAD (HSL) and Au electrode are sequentially deposited by spin-coating and vacuum thermal evaporation methods. The detailed experimental procedures are given in the Experimental Section. The morphologies of the MAPbI3 perovskite films deposited on two kinds of TiO2 films are similar (Figure S12). The structure of the representative device consisting of FTO (∼500 nm)/rutile TiO2 (∼200 nm)/MAPbI3 (∼300 nm)/spiro-OMeTAD (∼100 nm)/Au (∼70 nm) can be viewed from the crosssectional SEM image (Figure S13). Comparison of the performance of two sets of solar cells clearly suggests the much superior properties of the Sn-doped rutile TiO2 film (after heating) as the ESL to the undoped film (before heating). The average Voc of the PSCs with Sn-doped TiO2 ETLs is ∼1.09 V (Figure 4a), which is over 100 mV larger than that (∼0.9 V) of the PSCs with undoped TiO2 ESLs. The best cell with Sn-doped TiO2 ESL has its Voc approaching 1.11 V. Because of the largely improved Voc, the PSCs based on Sn-doped TiO2 ESLs show an average efficiency of 16.5% and a maximum efficiency of 17.2%, outperforming the PSCs based on TiO2 ESLs by 23 and 20%, respectively (Figure 4b). The best performance cell with an

of Sn4+ second ions undergoes an opposite trend with the rapid increase also at around 1500 s (Figure 3l). The rapid intensity change regions starting at around 1500 s corresponds to the interface region. Compared to the sample before heating, the disappearance of the Ti4+ second ion signal is delayed, and the signal of Sn4+ second ions appears in advance in the treated sample. Moreover, the intensity difference of the Ti4+ signal in its decay region before and after heating is apparently smaller than that of the Sn4+ signal in its rise region. These results derived from the SIMS spectra again confirm that the thermally driven diffusions of Sn4+ and Ti4+ ions across the FTO/TiO2 interface do occur, and the diffusion of Sn4+ ions into the TiO2 film is much faster than the diffusion of the Ti4+ ion into the SnO2 layer. We also investigated the case of the commonly used spin-coated TiO2 film supported on the FTO substrate. The similar film depth-dependent SIMS spectra (Figure S11) of the Sn element recorded from the spin-coated films before and after the 500 °C thermal treatment suggest no obvious diffusion of the Sn element from the FTO layer to TiO2 film. All these results indicate that the thermal diffusion induced Sn doping only occurs in the case of the epitaxial TiO2 film/FTO substrate that has a good lattice matching at the interface. SnO2 possesses a high electron mobility of up to 240 cm2 −1 −1 V s , and the incorporation of Sn4+ into TiO2 may improve the electron mobility of TiO2.40 Moreover, the diffusioninduced doping of Sn4+ in the TiO2 film leads to a gradient Sn4+ distribution and the position with the maximum amount of the Sn4+ dopant is located at the FTO/TiO2 interface. The incorporation of Sn4+ dopants causes an upshift of the Fermi level of TiO2 (see Figure 3m-I) and the higher the amount of the Sn4+ dopant, the larger the upshift of the Fermi level is.41 As a consequence of the upshift of the Fermi level in Sn4+doped TiO2 (Sn/TiO2), the band bending occurs at the Sn/ TiO2/TiO2 interface when they contact (see Figure 3m-II). Analogously, the gradient of the Sn4+ dopant formed in TiO2 (gradient Sn/TiO2) shall lead to the formation of a built-in E

DOI: 10.1021/acsami.9b04308 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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Figure 5. (a) Steady-state and (b) TRPL spectra of MAPbI3 films deposited on quartz (blue curve), the as-prepared TiO2 film (black curve), and thermally treated TiO2 film (red curve). (c) Transient photocurrent and (d) transient photovoltage measurements on devices based on the asprepared TiO2 ETL (black curve) and thermally treated TiO2 ETL (red curve). (e) Schematic of energy level diagrams of PSCs based on the asprepared TiO2 ESL and thermally treated TiO2 ESL.

efficiency of 17.2% shows a Jsc of 20.3 mA cm−2, Voc of 1.11 V, and FF of 76.1% (Figure 4c). To verify the Jsc value and analyze wavelength-dependent response behavior, the incident photon-to-current conversion efficiency (IPCE) spectrum has been recorded on the best performance device (Figure 4d). According to the standard AM 1.5 sunlight spectrum, the integrated Jsc value from the recorded IPCE spectrum is 19.4 mA cm−2, which is in accordance with the value obtained from the J−V curve, less than 5% deviation. The IPCE values are higher than 80% at the wavelength range from 400 to 620 nm and maintain over 60% for the whole visible light range. The efficiency achieved here is among the highest values of the PSCs with the (doped) rutile TiO2 film as ESLs.43−45 The dependence of the performance of the devices with the thermally treated TiO2 ESLs on the temperature and duration of thermal treatment is demonstrated in Figure S14. The devices based on the TiO2 ESLs treated at both 400 and 600 °C give the reduced performance compared to that based on the TiO2 ESLs treated at 500 °C. The possible reason is that the low temperature of 400 °C cannot induce the Sn diffusion, and the high temperature of 600 °C may reduce the electronic conductivity of the FTO layer. Increasing the duration at 500 °C from 0.5 to 1 h causes a slight efficiency decrease. A significant performance decrease and poor reproducibility were observed when further increasing the duration to 2 h. The large improvement of Voc can be attributed to the effective suppression of photocarrier recombination at the perovskite/Sn-doped TiO2 ESL interface and TiO2 bulk with favorable features (reduced defect density, improved carrier

migration associated with Sn doping and low barrier at the heteroepitaxial TiO2/FTO interface) revealed above. The dark current density−voltage curves of the devices based on the asprepared and thermally treated TiO2 ESLs further confirm the above analyzations (Figure S15). The device based on the thermally treated TiO2 ESL has a higher dark current onset and lower dark current density compared to that based on the as-prepared TiO2 ESL. This indicates the lowered carrier recombination at the perovskite/thermally treated TiO2 ESL interface. The photoelectrons transferred into the Sn-doped TiO2 ESL from the top perovskite active layer have a greatly reduced probability of being trapped by defective states and can move fast away from the TiO2/perovskite interface into the FTO electrode with the assist of the built-in electric field and smooth electron transfer at the TiO2/perovskite interface. EIS were used to analyze the interface carrier dynamics in PSCs. The Nyquist plots of all the devices were measured at 0.3 V in dark, as shown in Figure S16. Generally, the Nyquist plot of a PSC generally consists of two semiarcs. One appeared at the high frequency range is commonly attributed to the carrier transfer resistance (Rtr) across interfaces and the other at the low frequency range gives the information of the charge recombination resistance (Rrec) at the ESL/perovskite interface. By fitting the Nyquist plots with an equivalent circuit (an Rs, RtrCtr and RrecCrec series circuit), the corresponding resistance values are listed in Table S1. The Nyquist plot recorded on the device based on the Sn-doped TiO2 ESL shows a much bigger Rrec (∼1.15 MΩ) than that (∼0.13 MΩ) based on the undoped TiO2 ESL, suggesting the greatly F

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ACS Applied Materials & Interfaces

mainly caused by the slow photoelectron extraction at the perovskite/TiO2 interface as a result of the Fermi level upshift and the loss of surface defects. The reduced Rs (24.81 Ω) for the thermally treated one is attributed to the gradient Sn4+ doping, which leads to the accelerated migration of electrons in TiO2. The transfer processes of the photogenerated carrier in photovoltaic devices based on TiO2 ESLs before and after the thermal treatment are illustrated in Figure 5e. It is noted that the thermally treated ESL has a relatively low ability in extracting the photoelectrons from the perovskite layer. Thus, the photoelectron accumulation occurs at the perovskite/ESL interface, resulting in a strong hysteresis behavior in the device (Figure S17). Considering the hysteresis formed, the steadystate photocurrent and power conversion efficiency (PCE) output were recorded at a potential of 0.88 V (Figure S18). A stabilized PCE of 14% with a photocurrent density 16 mA cm−2 was achieved. Further study will use a suitable surface modifier to suppress the hysteresis behavior.

reduced carrier recombination at the perovskite/Sn-doped TiO2 ESL interface (Table S1). To further understand the origin of performance improvement, photoluminescence (PL), time-resolved PL (TRPL), transient photocurrent, and photovoltage measurements were carried out. Figure 5a gives the steady PL spectra of three perovskite films that were deposited on the as-prepared TiO2 ESL, thermally treated TiO2 ESL, and quartz, respectively. The film on the quartz exhibits the highest PL intensity because of no photoelectron extraction at the perovskite/quartz interface. The PL intensity of the film on the as-prepared TiO2 ESL is much lower than that on the thermally heated TiO2 ESL, meaning a higher photoelectron extraction efficiency of the asprepared TiO2 ESL. This is further supported by their TRPL spectra (Figure 5b). The lifetimes of the photocarriers in the films deposited on the as-prepared TiO2 ESL, thermally treated TiO2 ESL, and quartz are determined to be 19.40, 34.26, and 58.61 ns, respectively. The shortest lifetime in the film deposited on the as-prepared TiO2 ESL indicates the fastest photoelectron extraction at the perovskite/as-prepared TiO2 ESL interface. The implication from the PL results seems contradictory with the fact of the performance improvement of the devices with the thermally heated TiO2 ESLs. This contradiction can be explained by investigating the carrier dynamics along the entire pathway of the cells. The transient photocurrent decay curves in Figure 5c show that the thermally treated TiO2-based device has faster photocurrent decay than that with the as-prepared TiO2 ESL. This indicates that the photoelectrons travel much faster in the thermally treated TiO2 ESL than in the as-prepared one. This result is consistent with the improved device characteristics but inconsistent with the PL data, where the perovskite on the as-prepared TiO2 ESLs shows the shortest photocarrier decay time. Compared to the PL spectra that were measured in the film component of the devices, the transient photocurrent curves that were measured in the complete devices can fully reflect the photoelectron collection efficiency (Figure 5c). In the complete perovskite devices, the photoelectrons need to transport through the TiO2 ESL and transfer across the TiO2/ FTO interface after being extracted from the perovskite active layer. The more efficient electron extraction ability of the asprepared TiO2 ESL than the thermally treated one is probably associated with the presence of abundant surface defects that act as trapping centers to extract the photoelectrons from the perovskite active layer. However, the bulk transport of the photoelectrons in the as-prepared TiO2 ESL is unfavorable. In contrast, the bulk transport of the photoelectrons in the thermally treated ESL is quite favorable due to the formed gradient band structure and also the built-in electric field as illustrated in Figure 3m-III. It is therefore inferred that the large enhancement of the devices with the thermally treated ESL is dominated by the created bulk band structure. Transient photovoltage measurements provide the insight into carrier recombination rates in the cells. Based on the voltage decay time, the photoelectrons in the device with the thermally treated ESL have a longer lifetime than that in the cell based on the as-prepared ESL (Figure 5d). The prolonged decay time benefits from the substantially reduced density of the trapping centers at the perovskite/thermally treated ESL interface and favorable bulk band structures. This is also consistent with the EIS results (Table S1). Compared to the as-prepared ESL (Rtr = 1.23 × 105 Ω, Rs = 37.82 Ω), the enlarged Rtr (2.09 × 105 Ω) for the thermally treated ESL is

4. CONCLUSIONS By taking advantage of the same crystal structure of rutile TiO2 as that of SnO2 in the FTO layer and also their small lattice mismatch, a film of rutile TiO2 single-crystal arrays with welldefined facets is heteroepitaxially grown on the FTO layer supported on glass by a hydrothermal route. The subsequent thermal treatment in air leads a gradient of the Sn4+ dopant in the film as a result of the diffusion of Sn4+ ions from the FTO layer to TiO2 film. Owing to the single-crystalline nature, coherent FTO/TiO2 interface, and also greatly reduced defect concentration in the resultant Sn-doped TiO2/FTO substrate at the atomic level, such a Sn-doped TiO2 film used as ESLs in PSCs exhibits a much superior performance to its counterpart without Sn doping (the average conversion efficiency of 16.5% vs 13.4%). A highest efficiency of 17.2% is achieved in the optimal cell. The formation of a built-in electric field perpendicular to the FTO substrate associated with a gradient of the Sn4+ dopant is proposed to accelerate electron migration toward the FTO electrode and suppress the recombination of photocarriers, which greatly facilitates electron collection.

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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.9b04308. XRD spectra of FTO and as-prepared and thermally treated TiO2 films; cross-sectional scanning electron microscope images of FTO and a completed photovoltaic device with the thermally treated TiO2 film as the ESL; top-view SEM images of the as-prepared TiO2 film, thermally treated TiO2 film, and perovskite film on these two kinds of substrates; Raman spectra, Mott−Schottky curves, and ultraviolet−visible−near infrared absorption spectra and transmittance spectra of as-prepared and thermally treated TiO2 films; schematic of the slant angle for the (110) crystal plane of TiO2 and SnO2; SIMS spectra of spin-coated TiO2 films; J−V curves including reverse and forward scan directions of the devices based on the thermally treated TiO2 ESL; electronic impedance spectra of PSCs with the asprepared and thermally treated TiO2 as the ESL under 0.3 V bias voltage; fitting parameters of electronic impedance spectra of the perovskite cells with the asG

DOI: 10.1021/acsami.9b04308 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces

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(12) Marchioro, A.; Teuscher, J.; Friedrich, D.; Kunst, M.; van de Krol, R.; Moehl, T.; Grätzel, M.; Moser, J.-E. Unravelling the Mechanism of Photoinduced Charge Transfer Processes in Lead Iodide Perovskite Solar cells. Nat. Photonics 2014, 8, 250−255. (13) 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. (14) Huang, J.; Yuan, Y.; Shao, Y.; Yan, Y. Understanding the Physical Properties of Hybrid Perovskites for Photovoltaic Applications. Nat. Rev. Mater. 2017, 2, 17042. (15) Cho, A.-N.; Park, N.-G. Impact of Interfacial Layers in Perovskite Solar Cells. ChemSusChem 2017, 10, 3687−3704. (16) Bai, Y.; Meng, X.; Yang, S. H. Interface Engineering for Highly Efficient and Stable Planar p-i-n Perovskite Solar Cells. Adv. Energy Mater. 2018, 8, 1701883. (17) Haider, M.; Zhen, C.; Wu, T.; Liu, G.; Cheng, H.-M. Boosting Efficiency and Stability of Perovskite Solar Cells with Nickel Phthalocyanine as a Low-cost Hole Transporting Layer Material. J. Mater. Sci. Technol. 2018, 34, 1474−1480. (18) Leijtens, T.; Eperon, G. E.; Pathak, S.; Abate, A.; Lee, M. M.; Snaith, H. J. Overcoming Ultraviolet Light Instability of Sensitized TiO2 with Meso-superstructured Organometal Tri-halide Perovskite Solar Cells. Nat. Commun. 2013, 4, 2885. (19) 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. (20) Li, Y.; Cooper, J. K.; Liu, W.; Sutter-Fella, C. M.; Amani, M.; Beeman, J. W.; Javey, A.; Ager, J. W.; Liu, Y.; Toma, F. M.; Sharp, I. D. Defective TiO2 with High Photoconductive Gain for Efficient and Stable Planar Heterojunction Perovskite Solar Cells. Nat. Commun. 2016, 7, 12446. (21) Giordano, F.; Abate, A.; Correa Baena, J. P.; Saliba, M.; Matsui, T.; Im, S. H.; Zakeeruddin, S. M.; Nazeeruddin, M. K.; Hagfeldt, A.; Graetzel, M. Enhanced Electronic Properties in Mesoporous TiO2 via Lithium Doping for High-efficiency Perovskite Solar Cells. Nat. Commun. 2016, 7, 10379. (22) Tan, H.; Jain, A.; Voznyy, O.; Lan, X.; García de Arquer, F. P.; Fan, J. Z.; Quintero-Bermudez, R.; Quintero-Bermudez, R.; Yuan, M.; Zhang, B.; Zhao, Y.; Fan, F.; Li, P.; Quan, L. N.; Zhao, Y.; Lu, Z.-H.; Yang, Z.; Hoogland, S. Efficient and Stable Solution-processed Planar Perovskite Solar Cells via Contact Passivation. Science 2017, 355, 722−726. (23) Pathak, S. K.; Abate, A.; Ruckdeschel, P.; Roose, B.; Gödel, K. C.; Vaynzof, Y.; Santhala, A.; Watanabe, S. I.; Hollman, D. J.; Noel, N. Performance and Stability Enhancement of Dye-Sensitized and Perovskite Solar Cells by Al Doping of TiO2. Adv. Funct. Mater. 2014, 24, 6046−6055. (24) Nagaoka, H.; Ma, F.; deQuilettes, D. W.; Vorpahl, S. M.; Glaz, M. S.; Colbert, A. E.; Ziffer, M. E.; Ginger, D. S. Zr Incorporation into TiO2 Electrodes Reduces Hysteresis and Improves Performance in Hybrid Perovskite Solar Cells while Increasing Carrier Lifetimes. J. Phys. Chem. Lett. 2015, 6, 669−675. (25) Chen, B.-X.; Rao, H.-S.; Li, W.-G.; Xu, Y.-F.; Chen, H.-Y.; Kuang, D.-B.; Su, C. Y. Achieving high-performance planar perovskite solar cell with Nb-doped TiO2 compact layer by enhanced electron injection and efficient charge extraction. J. Mater. Chem. A 2016, 4, 5647−5653. (26) Shao, Y.; Xiao, Z.; Bi, C.; Yuan, Y.; Huang, J. Origin and Elimination of Photocurrent Hysteresis by Fullerene Passivation in CH3NH3PbI3 Planar Heterojunction Solar Cells. Nat. Commun. 2014, 5, 5784. (27) Ogomi, Y.; Morita, A.; Tsukamoto, S.; Saitho, T.; Shen, Q.; Toyoda, T.; Yoshino, K.; Pandey, S. S.; Ma, T.; Hayase, S. All-Solid Perovskite Solar Cells with HOCO-R-NH3+I‑ Anchor-Group Inserted

prepared and thermal-treated TiO2 layer as the ESL under 0.3 V bias voltage (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (G.L.). *E-mail: [email protected] (H.-M.C.). ORCID

Lianzhou Wang: 0000-0002-5947-306X Gang Liu: 0000-0002-6946-7552 Nam-Gyu Park: 0000-0003-2368-6300 Hui-Ming Cheng: 0000-0002-5387-4241 Author Contributions

T.W. and C.Z. equally contributed to this work. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank the National Natural Science Foundation of China (nos. 51825204, 51572266, and 51629201) and the Key Research Program of Frontier Sciences CAS (QYZDB-SSWJSC039) for the financial support.

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REFERENCES

(1) Sum, T. C.; Mathews, N. Advancements in Perovskite Solar Cells: Photophysics Behind the Photovoltaics. Energy Environ. Sci. 2014, 7, 2518−2534. (2) Hsiao, Y.-C.; Wu, T.; Li, M.; Liu, Q.; Qin, W.; Hu, B. Fundamental Physics Behind High-Efficiency Organo-Metal Halide Perovskite Solar Cells. J. Mater. Chem. A 2015, 3, 15372−15385. (3) 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. (4) Kim, H.-S.; Lee, C.-R.; Im, J. H.; Lee, K. B.; Moehl, T.; Marchioro, A.; Moon, S. J.; Humphry-Baker, R.; Yum, J. H.; Moser, J. E.; Gratzel, M.; Park, N. G. Lead Iodide Perovskite Sensitized Allsolid-state Submicron Thin Film Mesoscopic Solar Cell with Efficiency Exceeding 9%. Sci. Rep. 2012, 2, 591. (5) 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. (6) 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. (7) Ball, J. M.; Lee, M. M.; Hey, A.; Snaith, H. J. Low-temperature Processed Meso-superstructured to Thin-film Perovskite Solar Cells. Energy Environ. Sci. 2013, 6, 1739−1743. (8) Jeon, N. J.; Noh, J. H.; Yang, W. S.; Kim, Y. C.; Ryu, S.; Seo, J.; Seok, S. I. Compositional Engineering of Perovskite Materials for High-performance Solar Cells. Nature 2015, 517, 476−480. (9) Yang, W. S.; Park, B.-W.; Jung, E. H.; Jeon, N. J.; Kim, Y. C.; Lee, D. U.; Shin, S. S.; Seo, J.; Kim, E. K.; Noh, J. H.; Seok, S. I. Iodide management in formamidinium-lead-halide-based perovskite layers for efficient solar cells. Science 2017, 356, 1376−1379. (10) Stranks, S. D.; Eperon, G. E.; Grancini, G.; Menelaou, C.; Alcocer, M. J. P.; Leijtens, T.; Herz, L. M.; Petrozza, A.; Snaith, H. J. Electron-hole Diffusion Lengths Exceeding 1 Micrometer in an Organometal Trihalide Perovskite Absorber. Science 2013, 342, 341− 344. (11) Xing, G.; Mathews, N.; Sun, S.; Lim, S. S.; Lam, Y. M.; Gratzel, M.; Mhaisalkar, S.; Sum, T. C. Long-Range Balanced Electronand Hole-Transport Lengths in organic-Inorganic CH3NH3PbI3. Science 2013, 342, 344−347. H

DOI: 10.1021/acsami.9b04308 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces between Porous Titania and Perovskite. J. Phys. Chem. C 2014, 118, 16651−16659. (28) Li, Y.; Zhao, Y.; Chen, Q.; Yang, Y.; Liu, Y.; Hong, Z.; Liu, Z.; Hsieh, Y.-T.; Meng, L.; Li, Y.; Yang, Y. Multifunctional Fullerene Derivative for Interface Engineering in Perovskite Solar Cells. J. Am. Chem. Soc. 2015, 137, 15540−15547. (29) Yang, G.; Wang, C. L.; Lei, H.; Zheng, X. L.; Qin, P. L.; Xiong, L. B.; Zhao, X. Z.; Yan, Y. F.; Fang, G. J. Interface Engineering in Planar Perovskite Solar Cells: Energy Level Alignment, Perovskite Morphology Control and High Performance Achievement. J. Mater. Chem. A 2017, 5, 1658−1666. (30) Wu, T.; Zhen, C.; Wu, J.; Jia, C.; Haider, M.; Wang, L.; Liu, G.; Cheng, H.-M. Chlorine capped SnO2 quantum-dots modified TiO2 electron selective layer to enhance the performance of planar perovskite solar cells. Sci. Bull. 2019, 64, 547−552. (31) Snaith, H. J.; Grätzel, M. The Role of a “Schottky Barrier” at an Electron-Collection Electrode in Solid-State Dye-Sensitized Solar Cells. Adv. Mater. 2006, 18, 1910−1914. (32) Zhou, H.; Chen, Q.; Li, G.; Luo, S.; Song, T.-b.; Duan, H.-S.; Hong, Z.; You, J.; Liu, Y.; Yang, Y. Interface Engineering of Highly Efficient Perovskite Solar Cells. Science 2014, 345, 542−546. (33) Long, R.; Dai, Y.; Huang, B. Geometric and Electronic Properties of Sn-Doped TiO2 from First-Principles Calculations. J. Phys. Chem. C 2009, 113, 650−653. (34) Hosono, E.; Fujihara, S.; Kakiuchi, K.; Imai, H. Growth of Submicrometer-scale Rectangular Parallelepiped Rutile TiO2 Films in Aqueous TiCl3 Solutions under Hydrothermal Conditions. J. Am. Chem. Soc. 2004, 126, 7790−7791. (35) Liu, B.; Aydil, E. S. Growth of Oriented Single-Crystalline Rutile TiO2 Nanorods on Transparent Conducting Substrates for Dye-Sensitized Solar Cells. J. Am. Chem. Soc. 2009, 131, 3985−3990. (36) Kim, H.-S.; Lee, J.-W.; Yantara, N.; Boix, P. P.; Kulkarni, S. A.; Mhaisalkar, S.; Grätzel, M.; Park, N.-G. High Efficiency Solid State Sensitized Solar Cell Based on Submicrometer Rutile TiO2 Nanorod and CH3NH3PbI3 Perovskite Sensitizer. Nano Lett. 2013, 13, 2412− 2417. (37) Lai, Z.; Peng, F.; Wang, H.; Yu, H.; Zhang, S.; Zhao, H. A new insight into regulating high energy facets of rutile TiO2. J. Mater. Chem. A 2013, 1, 4182−4185. (38) Guo, W.; Xu, C.; Wang, X.; Wang, S.; Pan, C.; Lin, C.; Wang, Z. L. Rectangular Bunched Rutile TiO2 Nanorod Arrays Grown on Carbon Fiber for Dye-Sensitized Solar Cells. J. Am. Chem. Soc. 2012, 134, 4437−4441. (39) Lv, M.; Zheng, D.; Ye, M.; Xiao, J.; Guo, W.; Lai, Y.; Sun, L.; Lin, C.; Zuo, J. Optimized porous rutile TiO2 nanorod arrays for enhancing the efficiency of dye-sensitized solar cells. Energy Environ. Sci. 2013, 6, 1615−1622. (40) Snaith, H. J.; Ducati, C. SnO2-Based Dye-Sensitized Hybrid Solar Cells Exhibiting Near Unity Absorbed Photon-to-Electron Conversion Efficiency. Nano Lett. 2010, 10, 1259−1265. (41) Duan, Y.; Fu, N.; Liu, Q.; Fang, Y.; Zhou, X.; Zhang, J.; Lin, Y. Sn-Doped TiO2 Photoanode for Dye-Sensitized Solar Cells. J. Phys. Chem. C 2012, 116, 8888−8893. (42) Salvinelli, G.; Drera, G.; Baratto, C.; Braga, A.; Sangaletti, L. Stoichiometry Gradient, Cation Interdiffusion, and Band Alignment between a Nanosized TiO2 Blocking Layer and a Transparent Conductive Oxide in Dye-Sensitized Solar Cell Front Contacts. ACS Appl. Mater. Interfaces 2015, 7, 765−773. (43) Wu, S.; Chen, C.; Wang, J.; Xiao, J.; Peng, T. Controllable Preparation of Rutile TiO2 Nanorod Array for Enhanced Photovoltaic Performance of Perovskite Solar Cells. ACS Appl. Energy Mater. 2018, 1, 1649−1657. (44) Wu, S.; Cheng, C.; Jin, J.; Wang, J.; Peng, T. Low-Temperature Processed Nanostructured Rutile TiO2 Array Films for Perovskite Solar Cells With High Efficiency and Stability. Sol. RRL 2018, 2, 1700164. (45) Liu, X.; Liu, Z.; Sun, B.; Tan, X.; Ye, H.; Tu, Y.; Shi, T.; Tang, Z.; Liao, G. 17.46% Efficient and Highly Stable Carbon-based

PlanarPerovskite Solar Cells Employing Ni-doped Rutile TiO2 as Electron Transport Layer. Nano Energy 2018, 50, 201−211.

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