Bromide regulated film formation of CH3NH3PbI3 in low-pressure vapor-assisted deposition for efficient planar-heterojunction perovskite solar cells

Bromide regulated film formation of CH3NH3PbI3 in low-pressure vapor-assisted deposition for efficient planar-heterojunction perovskite solar cells

Solar Energy Materials & Solar Cells 157 (2016) 1026–1037 Contents lists available at ScienceDirect Solar Energy Materials & Solar Cells journal hom...

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Solar Energy Materials & Solar Cells 157 (2016) 1026–1037

Contents lists available at ScienceDirect

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

Bromide regulated film formation of CH3NH3PbI3 in low-pressure vapor-assisted deposition for efficient planar-heterojunction perovskite solar cells Jia Xu a, Jian Yin a, Li Xiao a, Bing Zhang a,c, Jianxi Yao a,b,n, Songyuan Dai a,c,n a

State Key Laboratory of Alternate Electrical Power System With Renewable Energy Sources, North China Electric Power University, Beijing 102206, China Beijing Key Laboratory of Energy Safety and Clean Utilization, North China Electric Power University, Beijing 102206, China c Beijing Key Laboratory of Novel Film Solar Cell, North China Electric Power University, Beijing 102206, China b

art ic l e i nf o

a b s t r a c t

Article history: Received 23 June 2016 Received in revised form 23 August 2016 Accepted 26 August 2016

Obtaining high-quality organic–inorganic halide perovskite films that have smooth and continuous surfaces and large crystal domains and possess photoelectrical properties preferable for photovoltaic applications is of paramount importance to achieve high performance perovskite solar cells (PSCs). The introduction of other halide ions into the synthesis process of methylammonium lead triiodide, CH3NH3PbI3, to fabricate CH3NH3PbI3  xXx (X ¼Cl or Br) (x E0) has been confirmed as an effective approach to optimize optoelectronic properties, thereby enhancing solar cell performance. However, the reported approaches to incorporate chlorine or bromine in perovskite films mainly take place during a liquid-phase synthesis process. Here, we report on bromide regulated CH3NH3PbI3 film formation through a low-pressure vapor-assisted deposition process. PbBr2 was used to either replace or to be mixed with PbI2 to form pre-deposited films that were then reacted with CH3NH3I vapor. Detailed structural, spectroscopic, and morphological characterizations of the perovskite films unambiguously demonstrate the formation of CH3NH3PbI3 films. However, bromine incorporation slows down the perovskite formation process through formation of an intermediate phase, CH3NH3PbBrxIy, which improves the grain domains in the as-fabricated perovskite films. Enhancements in power conversion efficiency (PCE) for the as-fabricated planar-heterojunction structured PSCs were found for samples in which bromine was incorporated. Meanwhile, transient photovoltage decay measurements revealed that carrier recombination was suppressed throughout the entire device. When mixed PbI2/PbBr2 films with a molar ratio of 1:4 were used as re-deposited films, an optimized PCE of 17.40% was obtained. & 2016 Published by Elsevier B.V.

Keywords: Organic–inorganic lead halide perovskites Perovskite solar cells Bromide Vapor-assisted Deposition Planar-heterojunction Power conversion efficiency

1. Introduction Organic–inorganic trihalide perovskites with an ABX3 structure (where A is an organic cation, B a divalent metal ion, and X a halide, i.e., Cl, Br, I, or any mixture thereof) have attracted much attention as novel photovoltaic (PV) materials in the recent years [1–5]. Typically, methylammonium lead triiodide CH3NH3PbI3 or mixed halide variants CH3NH3PbI3-xClx and CH3NH3PbI3-xBrx possess remarkable photoelectrical properties, such as high optical absorption coefficients owing to their suitable band gaps [6], long diffusion lengths of both electrons and holes caused by strong transition dipole moments [7,8], shallow defect levels and absence of additional mid-gap states [9,10], and high mobility of charge n Corresponding authors at: State Key Laboratory of Alternate Electrical Power System With Renewable Energy Sources, North China Electric Power University, Beijing 102206, China. E-mail addresses: [email protected] (J. Yao), [email protected] (S. Dai).

http://dx.doi.org/10.1016/j.solmat.2016.08.027 0927-0248/& 2016 Published by Elsevier B.V.

carriers owing to their low effective masses [11,12]. Moreover, because of the easy operation, low cost, and scalability of the synthesis process of these films, substantial efforts have been invested in optimizing perovskite solar cell (PSC) device designs, material interfaces, and processing techniques. To date, a record certified efficiency of 20.8% has been reported in a mesoporous structured device in which a mesoporous infiltrated anode is combined with a solid perovskite capping layer [13]. Additionally, perovskite materials can also be assembled in a PV cell with a planar heterojunction architecture [14,15]. A planar architecture offers the potential to fabricate both flexible and multi-junction cells [16–18]. Although the performance of PSCs has improved rapidly, the investigation and understanding of the basic properties of organic–inorganic trihalide perovskites is still ongoing and far from complete. It has been broadly observed that the power conversion efficiency (PCE) of PSCs is closely related to the morphology and crystal structure of the perovskite absorbers. Thus, substantial

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efforts have been devoted to obtaining smooth and continuous perovskite films, for example by using additives [19,20] or special solvents [21,22], by changing or mixing the halide salts in the precursors [23,24], through thermal pressure-assisting [25], or by annealing in special atmospheres [26]. Among these approaches, the use of mixed halide salts, e.g., PbCl2 blended with CH3NH3I, has received much attention because of the superior electrical and photophysical properties that have been observed for the mixed halide perovskite CH3NH3PbI3-xClx compared with CH3NH3PbI3 [7,27]. However, both the existence and role of Cl in mixed halide perovskites is still being debated and more investigations are required to validate the results and reconcile all the data presented by multiple groups that all employed different characterization techniques [28–32]. Bromide (in the form of PbBr2 or CH3NH3Br) has been used as a precursor to obtaining CH3NH3PbI3-xBrx with different x values to tailor the bandgap of perovskite films [33,34]. Currently, studies on the effect of bromide as an extrinsic ion on the formation of CH3NH3PbI3-xBrx (x E0) films are scarce. However, it was recently reported that adding r5% PbBr2 into PbI2 via a sequential dip coating process to synthesize perovskite films was successful in achieving CH3NH3PbI3-xBrx (x E0) films on a TiO2 mesoporous anode; solar cells containing such films showed enhanced PCEs compared with similar solar cells fabricated without involving bromide [35]. Moreover, the above cited studies on using mixed halide salts to fabricate mixed halide perovskites all involved a solution approach to synthesize the perovskite films. But besides the solution approach, vapor deposition methods, including dual-source vacuum evaporation [14], low-pressure chemical vapor deposition [36], and vapor-assisted solution process (VASP) [37,38], have also been developed to fabricate high-quality pinhole-free perovskite thin films. Compared with conventional solution processing, vapor deposition has two particular advantages worth noting. First, it avoids solvation and hydration processes as well as undesirable structural transitions that may occur during solution processing [39]. Second, it also effectively reduces the too-fast reaction rate between PbI2 and CH3NH3I [36,38], thereby resulting in an optimized perovskite surface morphology. Because the dynamic and thermodynamic processes of growing films by vapor phase deposition are inevitably different from those during solution processing, it is reasonable to speculate that the effect of PbX2 (X ¼Cl or Br) on the growth and final product of films such as CH3NH3PbI3-xXx (X ¼Cl or Br) would differ compared with solution processed films. In this work, we employ several pure and mixed lead halide salts films in the fabrication process, i.e., pure PbI2, pure PbBr2, and mixed PbI2/PbBr2 with various molar ratios are reacted with CH3NH3I vapor to synthesize perovskite films via low-pressure VASP (LP-VASP). The effect of the bromide content on the as-fabricated perovskite films were investigated through morphological, structural, and spectroscopic characterizations. All the various lead halides used here all formed CH3NH3PbI3 films. Time-dependent structural and morphological characterizations revealed that bromide has a substantial effect on perovskite crystal growth kinetics and film morphology. An intermediate phase, CH3NH3PbBrxIy, was observed during the gas–solid reaction process before the complete exchange of Br by I. In our experiments, all the as-fabricated CH3NH3PbI3 films exhibited similar absorption and photoluminescence spectra regardless of the bromide content involved in the preparation process. Even their fluorescence lifetimes were of the same order of magnitude. However, PSCs based on the various cases we tested showed obvious differences in one regard: there was a remarkable enhancement of open circuit voltage when PbBr2 was involved in the fabrication process. We attribute the open circuit voltage enhancement to a suppression of charge carrier recombination throughout the entire device, which is

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demonstrated by the results of transient photovoltage decay measurements. A maximum PCE of 17.40% under one sun illumination was achieved in planar-heterojunction PSCs with mixed PbI2/PbBr2 films with a molar ratio of 1:4; this is an improvement over performances of similar solar cells fabricated via VASP reported by other groups. Moreover, the results here afford an approach to regulate perovskite film formation through VASP to enhance the PCE of solar cells.

2. Experimental 2.1. Materials PbI2 (99%) and PbBr2 (98%), and 1,2-dichlorobenzene (98%) were purchased from Acros. CH3NH3I (99.5%) and Spiro-OMeTAD (99.7%) were both purchased from Borun Chemicals (Ningbo, China). C60 and DMF were both purchased from Alfa Aesar. Isopropanol was purchased from J&K Scientific Co., Ltd. Tris(2-(1Hpyrazol-1-yl)4-tert-butylpyridine)-cobalt(III) tris(bis(trifluoromethylsulfonyl)imide) (FK209-cobalt(III)-TFSI) was purchased from MaterWin Chemicals (Shanghai, China). Bis(trifluoromethane)sulfonimide lithium salt and tert-butylpyridine were purchased from Sigma-Aldrich. All chemicals were directly used without further purification. Glass substrates with a transparent FTO (thickness 2.2 mm, sheet resistance 15 Ω/square) layer were used for the PSCs. 2.2. Cell fabrication The FTO glass substrates were first etched using Zn powder and 2 M HCl diluted in deionized water to pattern the top and bottom electrodes for the solar cells. Then the etched substrates were sequentially cleaned in a special detergent solution, deionized water, and ethanol. After drying under clean dry air, the FTO glass substrates were annealed at 500 °C for 30 min in air to remove any remaining organic matter. Then, the c-TiO2 underlayer was deposited on the FTO glass substrates by spin coating a precursor solution (titanium isopropoxide in anhydrous ethanol (0.254 M) with the addition of 0.02 M HCl) at 3000 rpm for 30 s, followed by sintering at 500 °C for 30 min. The C60 layer was subsequently spin coated onto the TiO2 layer from a C60 1,2-dichlorobenzene solution (6 mg/ml) at a speed of 1500 rpm, then dried at 60 °C for 2 min 1 M solutions of lead halide were prepared in DMF by constant stirring at 70 °C for 30 min. Six cases of lead halides at different molar ratios were used in our experiments to study their effect on the properties of as-fabricated perovskite films and on PV device performance. The six cases are respectively pure PbI2 (case 1), mixed PbI2 and PbBr2 with different molar ratios of 4:1 (case 2), 3:2 (case 3), 2:3 (case 4), 1:4 (case 5), and pure PbBr2 (case 6). After allowing the FTO/c-TiO2/C60 substrates to cool to room temperature, they were coated with the lead halide solutions by spin coating at 3000 rpm for 30 s, followed by drying at 70 °C for 30 min in a nitrogen-filled glovebox. CH3NH3I powder was evenly spread out around the lead halide coated substrates in a petri dish covered with a lid. The petri dish was placed in a vacuum drying oven (10 kPa) set at 150 °C for 30 min. After this reaction process, the as-prepared perovskite films were first washed with isopropanol then dried and annealed at 150 °C for 5 min in a nitrogen-filled glovebox. For the hole transport layer, 20 ml of a Spiro-OMeTAD mixed solution was spin coated onto each perovskite film at 4000 rpm for 30 s. The Spiro-OMeTAD mixed solution was composed of 1 ml spiroOMeTAD solution (72.3 mg of spiroOMeTAD in 1 ml of chlorobenzene), 28.8 ml tert-butylpyridine, 17.5 ml Li-TFSI solution (520 mg of bis(trifluoromethane)sulfonimide lithium salt in 1 ml of acetonitrile) and 8 ml FK209-cobalt(III)-TFSI solution (300 mg of

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FK209-cobalt(III)-TFSI in 1 ml of acetonitrile). Finally, Au was thermally evaporated on top to form the back electrode (60 nm) at an atmospheric pressure of 4  10  4 Pa. 2.3. Materials and spectroscopic characterization X-ray diffraction (XRD) data were collected on a Bruker X-ray diffractometer with a Cu-Kα radiation source. The 2θ diffraction angle was scanned from 10° to 80°, with a scanning speed of 3.5° per min. Field emission scanning electron microscope (FESEM) images and energy-dispersive X-ray spectroscopy (EDS) spectra were taken with a SU8010 SEM (Hitachi) equipped with an EDS detector. X-ray photoelectron spectroscopy (XPS) measurements were performed using a Thermo ESCALab 250Xi with Al Kα emission at 1486.6 eV. And all the binding energies were referenced to the C 1 s peak at 284.8 eV. The absorption spectra of perovskite films were recorded with a UV–vis spectrophotometer (Shimadzu UV-2450) from 300 to 800 nm in transmission mode. Fluorescence spectra were recorded on a spectrofluorometer (Fluorolog 322). Fluorescence spectra were recorded by exciting the samples with a 450 W Xenon lamp at fixed wavelength 460 nm and scanning the monochromator from 500 to 850 nm. The absorption and photoluminescence spectra measurement samples were multilayered thin films on glass substrates in the form of FTO/TiO2/C60/perovskite. In the measurement, the incident light beam illuminated the sample from the perovskite film sides. Time-resolved photoluminescence (TR-PL) spectra were collected by a transient state spectrophotometer (F900, Edinburgh Instruments). Samples were excited with a 660 nm pulsed diode laser with a repetition rate of 2.5 MHz and an excitation intensity of 14 nJ/cm2. Perovskite film samples for the TR-PL spectra measurement were prepared directly on the glass substrates. 2.4. Photovoltaic device characterization Current density–voltage (J–V) characteristic were measured with a Keithley 2400 source meter together with a solar simulator (XES-300T1, SAN-EI Electric, AM 1.5), which was calibrated using a standard silicon reference cell. The solar cells were masked with a black aperture to define an active area of 0.09 cm2. Incident photon-to-electron conversion efficiency (IPCE) was measured in air using a QE-R measurement system (Enli Technology). Transient photovoltage decay measurements were performed on an electrochemical workstation (Zahner) combined with a module for fast intensity transients. A white LED with an intensity of 1000 W m  2 was used as the light source to illuminate the devices. In the experiments, devices were soaked in light for 2 s before the LED light was turned off.

3. Results and discussion Fig. 1(a) schematically illustrates the LP-VASP approach in our

experiments, where the PbI2/PbBr2 mixed film was intentionally used as the lead source film (for details see the Section 2). Fig. 1 (b) depicts the device configuration of the planar PSCs. Several layers were deposited successively on top of the fluorine-doped tin oxide (FTO) glass substrate as follows: compact TiO2 (c-TiO2), C60, perovskite, 2,2′,7,7′-tetrakis-(N,N-di-p-methoxyphenylamine) 9,9′-spirobifluorene (Spiro-OMeTAD), and Au. In this study, a solution-processed C60 layer was used as an interfacial layer between the c-TiO2 and perovskite layers to facilitate electron injection from the perovskites into c-TiO2 [40,41]. Comparing the PCE of PSCs with and without a C60 layer (Fig. S1) highlights the necessity of the C60 layer in obtaining a solar cell with a high PCE. However, details on the role of the C60 layer are outside the scope of this research article. To ensure the correct molarity of lead in dimethylformamide (DMF) to fabricate lead halide films, six experimental cases were thoroughly researched to understand the effect of bromide content on the structural, spectroscopic, and morphological properties of as-fabricated perovskite films and the PCEs of PSCs fabricated with these films. The six cases were: (case 1) 1.0 M of PbI2; (case 2) 0.8 M of PbI2 plus 0.2 M of PbBr2; (case 3) 0.6 M of PbI2 plus 0.4 M of PbBr2; (case 4) 0.4 M of PbI2 plus 0.6 M of PbBr2; (case 5) 0.2 M of PbI2 plus 0.8 M of PbBr2; and (case 6) 1.0 M of PbBr2. It should be noted that cases 1 and 6 were pure PbI2 and PbBr2, respectively. In this article, these six cases are referred to as PbI2:PbBr2 (5:0), PbI2:PbBr2 (4:1), PbI2:PbBr2 (3:2), PbI2:PbBr2 (2:3), PbI2:PbBr2 (1:4), and PbI2:PbBr2 (0:5), respectively. To correlate the contents of PbBr2 in lead source film to device performance, the as-fabricated perovskite layers discussed in the previous section were employed in PV devices with a standard device structure, shown in Fig. 1(b). 20 nominally identical devices were fabricated for each case. J  V measurements under simulated AM 1.5 G (100 mW cm  2) solar irradiation in a glovebox was used to evaluate the PSC performance; the statistical distribution of the measurement results is shown in Fig. 2. And their mean values summarized in Table 1. Devices based on pure PbI2 case produced short circuit currents (Jsc) between 21.54 and 23.73 mA cm  2, open circuit voltages (Voc) between 956 and 981 mV, fill factors (FF) between 68.63% and 74.08%; the resulting PCEs ranged from 14.45% to 16.15%. In comparison, superior performance was observed for case 2–5 devices, i.e., when PbI2 and PbBr2 coexisted in the films. In those devices Voc ranged from 964 to 1034 mV, Jsc from 21.54 to 23.50 mA cm  2, FF from 66.34% to 75.98%, and the resulting PCEs ranged from 14.74% to 17.40%. The superior PCEs of devices based on cases 2–5 were caused by clearly observable increases in Voc. Conversely, the PCE performance of the devices containing the pure PbBr2 case layers dropped significantly, which was the result of decreases in both Jsc and FF. The improved performance for cases 2–5, as compared with that for cases 1 and 6, reveals a positive effect caused by the contents of PbBr2 in lead source film in the perovskite film production process. The maximum value of the mean PCE was found for case 5 devices, i.e., PbI2:PbBr2 (1:4). Moreover, the individual

Fig. 1. Perovskite film fabrication and device architecture. (a) Schematic illustration of LP-VASP to fabricate the perovskite films using PbI2/PbBr2 mixed film as the lead source film. (b) Schematic device architecture of the planar-heterojunction PSCs.

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Fig. 2. Jsc (a), Voc (b), FF (c) and PCE (d), extracted from J–V measurements of solar cells based on cases 1–6.

Table 1 Summary of average photovoltaic parameters of devices based on cases 1–6. PbI2:PbBr2

Jsc (mA/cm2)

Voc (mV)

FF (%)

PCE (%)

Case Case Case Case Case Case

22.25 22.16 22.88 22.68 22.63 22.12

967 987 989 997 1021 1016

72.05 71.00 72.58 71.93 72.19 70.19

15.50 15.58 16.35 16.36 16.78 15.30

1__5:0 2__4:1 3__3:2 4__2:3 5__1:4 6__0:5

Fig. 3. Performance of best performing PSC fabricated with a molar ratio of PbI2/ PbBr2 of 1:4.

cell maximum PCE was also found for a 1:4 M ratio of PbI2/PbBr2. This device exhibited an outstanding performance with a Jsc of 22.92 mA cm  2, Voc of 1017 mV, FF of 77.50%, and PCE of 17.40%, as shown in Fig. 3. IPCE spectrum, hysteresis, steady-state performance, and repeatability studies of the optimized solar cells are provided in Supplementary Figs. 2–4(a) and (b), respectively. Before characterizing the perovskite films, the different predeposited lead halide films were characterized. Fig. S5 presents the XRD patterns and UV–Vis absorption curves (inset in Fig. S5) of the pre-deposited lead seed films. For the XRD patterns, three points should be noted: 1) Pure PbI2 films crystallize into hexagonal 2 H polytype crystals [1]. The strong diffraction peak at 12.6° corresponds to its (001) lattice plane [1]. 2) The films deposited by spincoating DMF solutions in which PbI2 and PbBr2 were dissolved together in different mole ratios crystallized as PbInBr2-n. And 3) when there was less PbI2 than PbBr2 in the DMF solutions, i.e., as in for cases 4–6, the deposited films no longer had any predominant peaks. As for the UV–Vis absorption curves, only the pure PbI2 film featured an absorption peak at around 515 nm, which is the characteristic bandgap excitation of crystallized PbI2 semiconductors [42]. When the proportion of PbI2 decreased, the absorption onset edges of the as-deposited films showed obvious blue-shifts. The variations in the XRD and UV–Vis absorption patterns evidently showed that the existence of the bromide in the films plays a key role in retarding the crystallization of PbI2. Moreover, the different lead halide films showed different surface morphologies which will be discussed later in this article. Both the changes in crystallinity and morphology of the pre-existing lead seed films played a part in the variations in the morphology of the as-fabricated perovskite films, which will also be discussed later. The crystal structure of each perovskite sample prepared on an FTO/TiO2/C60 substrate (Fig. 4) was characterized via XRD. These

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Fig. 5. XPS characterization of perovskite films fabricated from pure PbBr2 films. Inset: Br 3 s narrow spectrum of the perovskite films.

Fig. 4. XRD patterns of the perovskite films fabricated with different lead halide films (from bottom to up: substrate, cases 1–6).

perovskite films were prepared using mixed lead halide precursors with different molar ratios of PbI2 to PbBr2 as indicated in Fig. 4. We found that all perovskite samples showed exactly the same diffraction peak positions. The dashed lines indicate the peaks attributed to the substrate materials. The other peaks are all in accordance with crystallized CH3NH3PbI3 produced via other methods [34]. The peaks at 14.11°, 23.50°, 24.50°, 28.44°, 31.90°, 35.00°, 40.57°, 42.63°, and 43.22° can therefore be assigned to the (110), (211), (202), (220), (310), (312), (224), (314) and (330) orthorhombic crystal structure peaks, respectively. Furthermore, two points must be noted. 1) In previously reported results, CH3NH3PbI3 films prepared via solution process in the presence of 10% PbBr2 showed a relative shift of 0.08° for the (110) peak [35]. However, in our experiments the XRD patterns for the as-fabricated perovskite films showed the same diffraction peak positions, regardless of the content of bromide incorporation. This means that CH3NH3PbI3 films prepared through different approaches may have different tolerances to the amount of bromide involved in the synthesizing process. The presence of Br in the film was examined further by XPS (Fig. 5). And Br signal cannot be detected even in the sample based on pure PbBr2. Moreover, the presence of Br in the films that were fabricated by exposing pure PbBr2 films to CH3NH3I vapor for 30 min, was also examined by EDS (Fig. S6). There were no Br signals in both the top- and crosssectional-view FESEM characterizations images. These characterization results are very similar to those of CH3NH3PbI3-xClx (x E0) films [28,29]. 2) Different diffraction intensity ratios, for instance (220) versus (310), of the CH3NH3PbI3 films based on each case indicate different final domain orientations in these polycrystalline films. When PbI2 and PbBr2 coexisted in the film, the (220)/ (310) intensity ratios increased significantly. Chloride was confirmed to have an effect in directing the growth of CH3NH3PbI3 structures [30,32]. So far, there are scarcely any experiments to demonstrate the effect of Br on the growth of the final perovskite structure which was fabricated through VASP. Our experimental results show that the content of bromide incorporated in the

CH3NH3PbI3 fabrication process can directly influence the domain orientation of the final CH3NH3PbI3 structure, which is likely caused by different crystal growth or reorganization routes. The morphological characterization of case 1–6 perovskite films are shown in FESEM images (Fig. 6). All perovskite films are densely packed and exhibit polygonal grains of different sizes. It is evident that the average grain size increases with increasing bromide content during preparation. This conclusion is validated by the corresponding cross-sectional FESEM images of these samples, which are shown in Fig. S7. For all samples, the dense TiO2 and C60 layers are too thin to identify from these FESEM images. The thick layer above the FTO in every sample therefore corresponds to the CH3NH3PbI3 layer. With increasing bromide content, a more continuous CH3NH3PbI3 layer was formed. The reason for the increase in the average grain sizes with increasing bromide content during preparation will be discussed later combined with the analysis of the films' morphological evolutions. The absorption and photoluminescence properties of the various perovskite film cases deposited on FTO/TiO2/C60 substrates are shown in Fig. 7(a) and (b), respectively. Although the UV–vis absorption properties differed for the various cases of lead halide films, as shown in inset in Fig. S5, after reacting with CH3NH3I vapor for 30 min, all films displayed similar characteristics regardless of the bromide content in the preparation process. Strong absorption of UV–vis light up to  790 nm was observed for all of the perovskite films, which is consistent with results reported in the literature [7,22]. Additionally, the absorption onset edge did not shift even when the pure PbBr2 film was reacted with CH3NH3I vapor. Photoluminescence spectra (Fig. 7(b)) show a strong emission at 762 nm for all the samples. Halide exchange has been observed in organic–inorganic lead halide perovskites, including the exchange of Cl by I, and the exchange of Br by I [43]. This may explain why there was a blue-shift in the emission peak and absorption onset of CH3NH3PbI3 films prepared through solution processing in the presence of 10% PbBr2 [35]. Based on the experimental results here, it can be suggested further that the exchange process of Br by I is also affected by the method of CH3NH3PbI3 film fabrication. Data from the XRD, UV–vis absorption, XPS and photoluminescence spectra in this study suggest the complete exchange of Br by I in our experiments. The investigation of the exchange process is detailed in the following section. Fig. 8 displays TR-PL spectra of the perovskite films of the case 1–6, which were fabricated on glass substrates. A biexponential

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Fig. 6. Top-view FESEM images of perovskite films fabricated with molar ratios of PbI2 to PbBr2 of 5:0 (a), 4:1 (b), 3:2 (c), 2:3 (d), 1:4 (e), and 0:5 (f), respectively. The scale bars in each image are 2 mm.

decay model was used to fit the fluorescence relaxation spectra to estimate the decay kinetics. We attribute the fast relaxation component, t1, to surface recombination and the longer relaxation component, t2, to recombination occurring in the perovskite bulk [44]. All parameters are summarized in Table 2. Compared with perovskite films fabricated with pure PbI2, films with PbBr2 showed enhanced and decreased surface and bulk recombination, respectively. A calculated average lifetime (tave), which takes into account both t1 and t2 as well as their weighting, was used to evaluate the whole photoluminescence lifetime of the film. The values of tave show little sample variation and there was no discernable pattern for the different bromide ratios in the fabricating process. The comparable carrier lifetimes in all samples indicate that the nonradiative recombination channels in the perovskite films are not affected by the PbBr2 using in the perovskite

fabrication process. Therefore, the carrier lifetimes of the perovskite films do not contribute to the significant variations observed in the device performances. Based on the above experimental results for the XRD patterns, XPS spectra, EDS spectra, steady-state emission spectra, and UV– vis absorption spectra, it can be confirmed all the as-fabricated perovskite films shows properties of bromide-free CH3NH3PbI3, regardless of the molar ratios of PbI2 to PbBr2 in the reaction films. This result can be attributed to halide exchange during the reaction procedure. In order to further understand such the Br-I exchange process in a solid-gas reaction which has rarely been reported on so far, time-dependent structural and morphological characterizations of the perovskites obtained after exposure of the lead halide films to CH3NH3I vapor for different durations were carried out. Three kinds of lead halide film, i.e., pure PbI2, Pure

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Fig. 7. UV–vis absorption spectra (a) and emission spectra (b) of case 1–6 perovskite films.

Fig. 8. Emission dynamics measured at 780 nm upon excitation at 660 nm of case 1–6 perovskite films. Table 2 Summary of time parameters derived from the fitting results of the transient TR-PL decay curves shown in Fig. 8. Sample

5:0

PbI2:PbBr2 4:1

3:2

2:3

1:4

0:5

t1

Value (ns) Rel. (%)

1.181 17.68

1.160 23.13

0.884 47.13

0.958 35.11

0.947 42.27

0.978 39.18

t2

Value (ns) Rel. (%)

31.87 82.32

36.04 76.87

52.00 52.87

33.79 64.89

39.96 57.73

34.98 60.82

tave

Value (ns)

26.44

27.97

27.91

22.26

23.47

21.66

PbBr2, and PbI2:PbBr2 (1:4), were chosen as precursors in this study. The time evolution (0–30 min) of XRD patterns in the 10° to 18° region of pure PbI2, pure PbBr2, and PbI2:PbBr2 (1:4) films exposed to CH3NH3I vapor are illustrated in Fig. 9(a), (b) and (c), respectively. The corresponding XRD patterns over the entire range of angles are shown in Fig. S8. For the pure PbI2 film, diffraction peaks corresponding to PbI2 (2θ ¼12.6°) were absent at exposures times of 10–15 min. Because of the preferential facet of PbI2, diffraction peaks corresponding to CH3NH3PbI3 (2θ ¼14.1°) can only be observed by zooming in on their locality. These features suggest that when pure PbI2 films encounter the CH3NH3I vapor,

CH3NH3PbI3 is formed within 2 min; PbI2 was nearly depleted after 15 min. As shown in Fig. 9(b), the pure PbBr2 film did not have obvious diffraction peaks in this region. Thus, after the pure PbBr2 film was exposed in CH3NH3I vapor for 2 min, bromide perovskites, i.e., CH3NH3PbBrxIy (2θ ¼14.9°) are visible in the XRD pattern. The appearance of a diffraction peak at 12.6° suggests the formation of PbI2. In other words, the transformation of PbBr2 to PbI2 takes place within 5 min in our study. A further increase in the exposure time caused a weakening of the diffraction peaks at 14.9° and they systematically moved toward lower 2θ angles because the gradual substitution of the smaller Br atoms by larger I atoms increased the lattice spacing. When the exposure time was longer than 20 min, a peak typical of CH3NH3PbI3 (2θ ¼ 14.1°) appeared and remained thereafter. The XRD patterns of the PbI2: PbBr2 (1:4) film when reacted with CH3NH3I vapor roughly proceeded somewhere between the above two cases. Although the PbI2 peak at 2θ ¼12.6° was not present in the original PbI2:PbBr2 (1:4) film, it appeared within 2 min of the gas–solid reaction. Similar to the pure PbBr2 film, an intermediate CH3NH3PbBrxIy phase was formed during the reaction process. But the conversion from CH3NH3PbBrxIy to CH3NH3PbI3 was achieved slightly quicker in the PbI2:PbBr2 (1:4) film than in the pure PbBr2 film. Intermediates have been found to play a key role in perovskite films' growth in the wildly adopted 1- and 2-step solution-based approaches [45]. In the original and unoptimized 1-step solution deposition process, an amorphous (CH3NH3)1 þ xPbI3 þ x intermediate phase is formed, which kinetically dominates the initial nucleation, resulting in films with poor crystallinity [30]. Then, many studies focused on changing the intermediate phase with the aim of improving perovskite film deposition; the proposed solutions include using mixed solvents [46], additives [47], and other halogen anions [30]. By incorporating chloride into the precursor solution, the nucleation of perovskites from the amorphous phase was circumvented via creation of kinetically accessible and structurally coherent CH3NH3PbCl3 or CH3NH3PbI3-xClx intermediates [29,30]. These intermediates were also observed in the 2-step solution growth method when chloride was included in the seed phase or immersion solution [48]. Thermally activated CH3NH3Cl sublimation was the primary contributor to chloride loss in the final products in both the 1- and 2-step deposition methods [45]. VASP allows CH3NH3I exposure of the lead seed phase during an annealing process. And the perovskite nucleation occurs at and within the pre-existing lead seed crystallites, which facilitates efficient perovskite nucleation. However, published research on the intermediates during VASP is scarce. The main impact of the intermediate on the final perovskites films' morphologies and

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Fig. 9. XRD patterns of films fabricated by exposing films of (a) pure PbI2, (b) pure PbBr2, and (c) mixed PbI2 and PbBr2 with molar ratios of 1:4 to CH3NH3I vapor for different durations.

Fig. 10. Top-view FESEM images of the films fabricated by exposing pure PbI2 films to CH3NH3I vapor for (a) 0 min, (b) 2 min, (c) 5 min, and (d) 10 min. The scale bars in each image are 2 mm.

crystallinities is the way it alters in the nucleation dynamics of the film [45]. In our experiments, the intermediate CH3NH3PbBrxIy phase can be identified from the XRD patterns (Fig. 9) when bromide was incorporated into the lead seed films. The samples' morphological evolutions (Figs. 10–12 and Figs. S9–S11) further indicate that the impact of the intermediate phase is largely to

guide perovskite nucleation and growth in bromide containing systems. Moreover, the XPS (Fig. 5) and EDS (Fig. S6) results confirm the absence of bromide in the final perovskite films. The transformation from CH3NH3PbBrxIy to CH3NH3PbI3 must be accompanied by the formation and elimination of CH3NH3Br in the films (Reaction (1) in Scheme 1). Similar to the mechanism for

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Fig. 11. Top-view FESEM images of the films fabricated by exposing pure PbBr2 films to CH3NH3I vapor for (a) 0 min, (b) 2 min, (c) 5 min, and (d) 10 min. The scale bars in each image are 2 mm.

chloride loss in the solution depositions, the bromide loss in VASP is probably attributed to the sublimation of CH3NH3Br (Reaction (2) in Scheme 1). Thus, the possible reactions, including the intermediate relevant for the perovskite transformation during the LP-VASP when bromide is included in the pre-existing lead seed films are shown as Reaction (3) in Scheme 1. The different reaction processes of the above three cases can also be observed in the samples' morphological evolutions. The surface morphology of the PbI2 film (Fig. 10(a)) did not change a lot until 5 min in the reaction process. But from 5 min (Fig. 10(c)) to 10 min (Fig. 10(d)), the film underwent significant changes that corresponded with the formation of CH3NH3PbI3; and after 10 min (Fig. S9(a)–(c)) the film morphology no longer changed significantly. In contrast, when the original pure PbBr2 film (Fig. 11(a)) reacted with CH3NH3I for 2 min, some crystals with different shapes were formed on its surface (Fig. 11(b)). This can be assigned to the intermediate CH3NH3PbBrxIy compound forming during this process. Then, these crystals gradually dissolved or alloyed in the growing perovskite film (Fig. 11(c)). The morphology of the perovskite film (Fig. 11(d)) was changed significantly and did not establish itself fully until after 20 min (Fig. S10). Although CH3NH3PbBrxIy exits both in the pure PbBr2 film and PbI2:PbBr2 (1:4) films, when the original PbI2:PbBr2 (1:4) film (Fig. 12(a)) was exposed to CH3NH3I vapor for 2 min (Fig. 12(b)) no scattered crystals were formed; however, its morphology appeared to be similar to that of the PbI2 film, which agrees with the XRD results discussed above. During the reaction process, the whole film appeared flat and densely packed (Fig. 12(c) and (d)). Finally, large scaled grains were formed and composed the whole film (Fig. S11). We conclude that the more PbI2 in the film, the shorter the time required to mutate the film morphology was and the sooner the

CH3NH3PbI3 film established itself; additionally, the more PbBr2 in the film, the more visible the intermediate phase was and the longer the time required for the film to evolve into its final morphology. Therefore, slowing down the perovskite formation process can improve the growth of the crystal domains during annealing. Moreover, the contents of Br in the films which evolved from the original PbI2:PbBr2 (1:4) films exposed to CH3NH3I vapor with different durations were characterized by EDS (Fig. S12). Both the top- and cross sectional-views in the FESEM images showed that the Br signals could be detected within the first 20 min, but vanished when the duration was extended to 30 min. Those results agreed with the conclusions draw from the time evolution XRD patterns (Fig. 9), and also confirmed our speculations regarding the mechanism for the perovskite growth process mechanism. As shown in Figs. 6, 10–12 and S9–S11, when the bromide was incorporated into the lead seed films, the average grain sizes in the final films increased. The possible mechanisms behind such results are provided herein. First, differences in the morphology (Figs. 10 (a), 11(a), and 12(a)) and crystallinity (Fig. S5) of the pre-existing lead seed films with different bromide content guided the different nucleation dynamics of the perovskite films. Second, the final CH3NH3PbI3 film growth process was also directed by the formation of the intermediate CH3NH3PbBrxIy phase in bromide containing systems. Previous experimental results showed that inhibiting the crystallization of PbI2 in the 2-step solution deposition by using strong coordinative dimethylsulfoxide instead of DMF as the solvent for PbI2 produced perovskite films with a small crystal size distribution and flat surface morphology [49]. This is attributed to a slowing down of the reaction dynamic between the lead seed

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Fig. 12. Top-view FESEM images of the films fabricated by exposing mixed PbI2 and PbBr2 with molar ratios of 1:4 films to CH3NH3I vapor for (a) 0 min, (b) 2 min, (c) 5 min, and (d) 10 min. The scale bars in each image are 2 mm.

Scheme 1. Mechanisms for the loss of bromide and possible chemical reactions involved in the perovskite growth process.

film and CH3NH3I. In our experiments, only the pure PbI2 film displayed the typical hexagonal 2 H polytype with an obviously predominant peak of the (001) lattice plane. Such crystallized PbI2 can react with CH3NH3I efficiently as the ordered crystal structure already exits, only requiring the intercalation of CH3NH3I into the lattice to form CH3NH3PbI3 [49]. When bromide was incorporated into the lead seed films, the crystallinity of PbI2 was altered and destroyed to some extent, depending on the proportion of bromide incorporation (Fig. S5). The bromide reacted with CH3NH3I vapor to first form PbI2 and CH3NH3PbBrxIy (Reaction (3) in Scheme 1). Then the reaction between the resultant PbI2 and the CH3NH3I vapor along with the sublimation of CH3NH3Br carried the transformation onward. Thus, the growth dynamic of the CH3NH3PbI3 films must be slowed down in bromide containing systems, which can be confirmed from the samples' morphological evolutions (Figs. 10–12 and S9–S11). Seen from the surface morphological evolution of the pure PbI2 film exposed in the CH3NH3I vapor (Fig. 10), the nucleation sites of the perovskite film are highly correlated with the pre-existing PbI2's surface morphology. Compared with other lead seed films (Figs. 11 and 12), the pure PbI2 film offered more nucleation sites and a higher reaction rate for obtaining the final CH3NH3PbI3 film. The resultant polygonal

grains grew up and crowded the whole film faster than in the films with bromide incorporation. Thus, the average grain size was relatively speaking smaller than in films obtained from lead seed films containing Br. Past research indicates that chlorine incorporation mainly improves carrier transport across heterojunction interfaces [28]. Transient photovoltage decay measurements performed under open circuit conditions are often used to investigate the photoinduced carrier recombination dynamics of PV devices [28,41]. The dynamic process reflected in a transient photovoltage decay measurement is different from the TR-PL measurement in our experiments. Photoinduced carriers in transient photovoltage decay tests cannot be swiped out of the device but recombine. Thus, the carrier lifetime reflects the recombination dynamics of the entire heterojunction across the PV devices. The TR-PL measurements were performed on perovskite films fabricated on glass substrates. Thus, the extracted carrier lifetime can only describe the recombination dynamics within the perovskite films. In Fig. 13 (a), the transient profiles of Voc for cells based on case 1 and 5 show a significantly slower photovoltage decay for the cell based on PbI2:PbBr2 (1:4) than the cell based on pure PbI2. Furthermore, electron lifetime τ can be obtained using the following equation [41]: −1

τ = − kBTe−1( dVoc /dt ) ,

(1)

where e is the elementary charge, T is temperature, and kB is the Boltzmann constant. Fig. 13(b) shows electron lifetime as a function of Voc. The PSC based on the PbI2:PbBr2 (1:4) case shows a much longer electron lifetime than that of the PSC based on the pure PbI2 case, indicating that carrier recombination throughout the entire device has been suppressed in the former cell. Because

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Fig. 13. (a) Transient photovoltage decay curves and (b) the relationship between extracted lifetime of the injected carriers in the devices and Voc of the PSCs based on case 1 (black) and 5 (red). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

the results of the TR-PL measurements (Fig. 8) showed that the perovskite film based on the PbI2:PbBr2 (1:4) case did not exhibit obvious evidence of a significant reduction in carrier recombination, the improvement in the carrier lifetime of the PSC based on the PbI2:PbBr2 (1:4) case may originate from the suppressed interfacial recombination at the interfaces between the perovskites and the adjacent carrier transport layers. Numerical simulations of PSCs have revealed that interfacial recombination loss seriously harms Voc [50]. Thus, less interfacial recombination consequently enhances Voc, which in turn improves the device efficiency; this explains the results presented in Fig. 2. At this point, the incorporation of PbBr2 in the lead seed films in the perovskite film production process through LP-VASP has been confirmed to have a positive effect on the performance of the as-fabricated PSCs. The question of whether other halogen-based lead seed material, such as PbCl2, also have the same regulating effect on perovskite film growth and further on the performance of the as-fabricated PSCs is an interesting one that requires further investigation. We have done some preliminary experiments on using mixed PbCl2/PbI2 as the lead seed films to fabricate CH3NH3PbI3 films through LP-VASP. However, unlike PbI2 and PbBr2, the solubility of PbCl2 in DMF is very poor. When a DMF solution containing 0.5 M PbCl2 and 0.5 M PbI2 was used to fabricate the mixed lead seed films, the efficiencies of the as-fabricated PSCs were lower than those of PSCs prepared using pure PbI2 films, as shown in Fig. S13. Thus, if PbCl2 is to be chosen to regulate the formation of perovskite films through LP-VASP with the aim of fabricating PSCs more efficient than those with the pure PbI2 films, then some special experimental methods will have to be explored to achieve this, which will be the topic of future research.

the bromide content involved during the preparation process. The corresponding fluorescence lifetimes varied little in the samples. The performance of PSCs in which PbBr2 was involved in the fabrication process was improved statistically compared with cells using pure PbI2 during fabrication. The significant open circuit voltage enhancement of PSCs fabricated with PbBr2 can be attributed to carrier recombination throughout the entire device being suppressed, which was demonstrated by transient photovoltage decay measurements. In our experiments, a best-performing PCE of 17.40% was achieved by using mixed PbI2/PbBr2 films with the molar ratio of 1:4. This PCE is superior to those of similar solar cells fabricated via VASP reported by other groups. Our experimental results demonstrate an approach to regulate perovskite film formation through VASP to enhance the PCE of solar cells. Moreover, in order to further boosting device performances, efforts should focus on the investigation of charge transfer behavior across the interfaces in PSCs in which mixed halides were used in the fabrication process.

Acknowledgments This work was supported by the National High Technology Research and Development Program of China (863 Program) (No. 2015AA050602), the National Natural Science Foundation of China (No. 51372083), Jiangsu Province Science and Technology Support Program, China (BE2014147-4), and the Fundamental Research Funds for the Central Universities (Nos. 2014ZZD07 and 2015ZD11).

Appendix A. Supplementary material 4. Conclusions In summary, we investigated the effect of bromide on perovskite thin films fabricated via LP-VASP and the performance of PSCs. Various lead halide films, i.e., pure PbI2, pure PbBr2 and mixed PbI2/PbBr2 with various molar ratios, were exposed to CH3NH3I vapor to synthesize perovskite films. Crystal CH3NH3PbI3 films without any detectable Br signals were confirmed by the XRD and XPS characterization. Time-dependent structural and morphological characterization confirmed that bromide has a substantial effect on perovskite crystal growth kinetics and film morphology. Exchange of Br by I was observed and an intermediate phase, CH3NH3PbBrxIy, was formed during a gas–solid reaction process. All the as-fabricated CH3NH3PbI3 films exhibited similar absorption and photoluminescence spectra regardless of

Supplementary data associated with this article can be found in the online version at http://dx.doi.org/10.1016/j.solmat.2016.08.027.

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