Fabrication of Porous Lead Bromide Films by Introducing Indium Tribromide for Efficient Inorganic CsPbBr₃ Perovskite Solar Cells

Abstract

In the process of preparing CsPbBr₃ films by two-step or multi-step methods, due to the low solubility of CsBr in organic solvents, the prepared perovskite films often have a large number of holes, which is definitely not conducive to the performance of CsPbBr₃ perovskite solar cells (PSCs). In response to this problem, this article proposed a method of introducing InBr₃ into the PbBr₂ precursor to prepare a porous PbBr₂ film to increase the reaction efficiency between CsBr and PbBr₂ and achieve the purpose of In (Ⅲ) incorporation, which not only optimized the morphology of the produced CsPbBr₃ film but also enhanced the charge extraction and transport capabilities, which was ascribed to the reduction of the trap state density and impurity phases in the perovskite films, improving the performance of CsPbBr₃ PSCs. At the optimal InBr₃ concentration of 0.21 M, the InBr₃:CsPbBr₃ perovskite solar cell exhibited a power conversion efficiency of 6.48%, which was significantly higher than that of the pristine device.

1. Introduction

After more than ten years of rapid development, lead halide-based perovskite solar cells have made remarkable achievements, but they seem to be in a vicious circle where high efficiency and high stability are contradictory to each other. Although the power conversion efficiency (PCE) of organic-inorganic hybrid perovskite solar cells (PSCs) has increased from the initial 3.8% to more than 25% at present [1,2,3,4,5], yet due to the strong volatility of common A-site organic cations, such as organic methylammonium (MA⁺) and formamidinium (FA⁺), the organic components disappear under thermal stress [6,7]. In terms of thermal stability, the all-inorganic cesium-lead halide perovskite CsPbX₃ (X: iodine or bromine), which is formed by using more stable inorganic cesium ions (Cs⁺) to completely replace organic cations, usually performs better stability [8,9,10,11] and is not prone to degradation at temperatures above 400 °C [8,12]. This provides the necessary conditions for the long-term stable use of CsPbX₃ PSCs [13,14,15,16]. The key factor, which influences the stability of CsPbX₃, is the moisture in the air. The presence of humidity changes the phase of the perovskite and reduces the stability of the photovoltaic device, but this does not directly cause the decomposition of CsPbX₃ (mainly I-rich CsPbX₃) and the lack of components [8,17,18]. Of course, this phase change is reversible when heated [19].

Compared with other Cs-based inorganic perovskites, the most prominent advantage of CsPbBr₃ is that it has a highly stable crystal structure. Whether it is the orthorhombic γ-phase at room temperature or tetragonal β-phase and cubic α-phase when heated, the geometric structures of CsPbBr₃ have not changed much, so the electronic structures of different phases are also relatively similar [20]. For this reason, CsPbBr₃ is also regarded as a perovskite material that presents better stability to humidity, heat, and light at ambient temperature [2,8,21,22]. Since Kulbak et al. first prepared the CsPbBr₃ PSCs by a two-step solution-processing method in 2015 [22], in less than ten years, the PCE of the CsPbBr₃ based PSCs have reached more than 10% with an ultrahigh open-circuit voltage (V_OC) of 1.62 V [23], but it still has a large distance compared with CsPbBr₃ PSCs theoretical limit PCE of 16.4% [24] and the maximized PCE of 19.0% for inorganic CsPbI₃ PSCs [10].

In the process of preparing CsPbBr₃ PSCs by the solution-processing method, the solubility of CsBr in commonly used polar solvents is poor, and the concentration differences between CsBr and PbBr₂ solutions are large, which leads to the derivative phases PbBr₂-rich CsPb₂Br₅ and CsBr-rich Cs₄PbBr₆ in the process of the generation of CsPbBr₃ [25]. At the same time, the thickness of the non-optimized prepared perovskite film is low, and the ability to absorb light is inadequate; numerous pinholes appear in the film [26]. Consequently, a decrease appears in the PCE of CsPbBr₃ PSCs. Regardless of whether it is a two-step sequential deposition or a multi-step method to prepare CsPbBr₃, it is necessary to deposit PbBr₂ first and then use CsBr to convert PbBr₂ to CsPbBr₃. Improving the PbBr₂ film preparation process and adjusting the PbBr₂ preparation method can achieve the goals of enhancing the reaction efficiency of the precursor, accurately controlling the subsequent growth of CsPbBr₃ crystals, and reducing the generation of by-products, and finally obtain perovskite film with a high purity phase, large grain size, and high coverage [23,27]. By precisely regulating the film-forming temperature and pore diameter of the PbBr₂ precursor film, Zhao et al. [23] minimized the compressive stress of the perovskite film and prepared CsPbBr₃ grains with a size of up to 1.62 μm, which not only made the PCE of the all-inorganic CsPbBr₃ perovskite solar cell reach 10.7%, the open-circuit voltage (V_OC) as high as 1.6 V, and it also kept the device extremely stable in a high-humidity air environment. Lee et al. [28] introduced CZISSE QDs quantum dots into the PbBr₂ film. CZISSE QDs acted as seeds to promote the crystallization of CsPbBr₃ and, at the same time, penetrated into the m-TiO₂ and CsPbBr₃ perovskite films to increase the electron extraction and transportability of TiO₂, thereby improving the conversion efficiency of the device by 20.6%.

In this work, InBr₃ was introduced into the PbBr₂ precursor solution, so that the multiple ordered crystal orientations of lead bromide grew, and the original rough and extremely uneven grain distribution of the PbBr₂ film evolved into a large uniform-porous film with pores. This morphological change ensured the full diffusion and uniform reaction of CsBr in the PbBr₂ film during the synthesis of CsPbBr₃ and was conducive to the formation of polycrystalline surface growth, high purity phase, and uniform morphology InBr₃: CsPbBr₃ film. The PCE of the small area (0.09 cm²) InBr₃:CsPbBr₃ PSC obtained after conditions optimization was 6.48%, in particular, the V_OC was significantly improved.

2. Experiment Section

2.1. Materials

PbBr₂ (99.99%) and CsBr (99.9%) were purchased from Xi’an Polymer Light Technology Corp. (Xi’an, China) and were not purified. InBr₃ (99.9%) was purchased from Shanghai Macklin Biochemical Co., Ltd. (Shanghai, China). Titanium diisopropoxide bis (acetylacetonate; 75 wt% in 2-propanol) was purchased from Sigma-Aldrich (Louis, MO, USA). Titanium dioxide (TiO₂) paste (18 NR-T) was purchased from Greatcell Solar Limited (Queanbeyan, Australia). N,N-Dimethylformamide (DMF, chromatographic grade, ≥99.9%), methanol (chromatographic grade, ≥99.9%), ethanol (chromatographic grade, ≥99.8%), and isopropanol (≥99.5%) were purchased from Aladdin (Shanghai, China). The fluorine-doped tin oxide coated glass (FTO, 6 Ω/□) and carbon paste were purchased from Opvtech New Energy Co., Ltd. (Yingkou, China) and Shanghai MaterWin New Materials Co., Ltd. (Shanghai, China), respectively.

2.2. Device Fabrication

All the following processes were carried out in a fume hood environment, without artificial control of the temperature, humidity, and airflow rate of the surrounding environment. The fluorine-doped tin-oxide-coated glasses were patterned by laser etching and cleaned by ultrasonic with acetone, isopropanol, ethanol, and deionized water. After being dried by high purity nitrogen, the FTO were further cleaned by an ultrasound treatment for 15 min and washed with ethanol. Afterward, the pre-conditioned FTO were spin-coated with 0.15 M titanium diisopropoxide bis(acetylacetonate) in 1-butanol at 5000 rpm for 20 s and were heated at 125 °C for 5 min. After these substrates returned to room temperature, the above procedure was repeated twice with 0.3 M titanium diisopropoxide bis(acetylacetonate) in 1-butanol and compact TiO₂ (c-TiO₂) was obtained. After that, the mesoporous TiO₂ (m-TiO₂) films were deposited on the above cooling c-TiO₂ by spin-coating at 5000 rpm for 30 s by means of TiO₂ paste diluted with ethanol. Further, the obtained layers were dried at 125 °C for 5 min followed by the muffle furnace at 500 °C for 30 min. After the muffle furnace was lowered to room temperature, pre-coated substrates were acquired.

Perovskite films were synthesized by a multistep solution-processing method. 0.03, 0.09, 0.15, 0.21, and 0.27 mmol InBr₃ were added into 1 mL DMF of PbBr₂ (1 M) and stirred. After the InBr₃ was completely dissolved, the DMF mixed solution was spin-coated on pre-coated substrates at 2000 rpm for 30 s and then heated to 90 °C for 30 min. Afterward, the methanol solution of CsBr (0.07 M) was spin-coated on the InBr₃:PbBr₂ film at 5000 rpm for 30 s and heated to 250 °C for 5 min, and this step was repeated five times. Next, the prepared sample was placed in isopropanol and soaked for 30 min and annealed at 250 °C for 15 min to remove excess CsBr. Finally, the carbon paste was deposited coated on the perovskite films by using the doctor blade coating method and dried at 100 °C for 10 min. The effective area of the back electrode was 3 mm × 3 mm, which defined the active area of each device.

2.3. Characterization

The morphologies of the synthesized films and energy-dispersive X-ray spectroscopy (EDS) mapping images were observed by a scanning electron microscope (SEM, FEI MAGELLAN 400, FEI, Hillsboro, OR, USA). The crystal structure of the synthesized sample was determined by means of X-ray diffraction (XRD, Cu Kα radiation, λ = 1.5418 Å, Rigaku D/max2500, Tokyo, Japan). The steady-state photoluminescence (PL) spectra of perovskite films were performed using a Renishaw InVia micro-Raman spectroscopy system (Renishaw, Wotton-under-Edge, UK) with a 473 nm excitation source. Ultraviolet photoelectron (UPS) and X-ray photoelectron spectroscopy (XPS) were carried out by an X-ray photoelectron spectrometer (EscaLab Xi+, Thermofisher, Waltham, MA, USA). UV-Vis spectrometer (UV-3600, Shimadzu, Kyoto, Japan) was employed to measure the absorption spectrum in the range of 200 nm to 800 nm. The current-voltage (J–V) characteristics and the external quantum efficiency (EQE) of the fabricated solar devices were measured by a solar cells test system (XP3000, Sanyou, Beijing, China) and an EQE measured system (QTest Station 1000A, CROWNTECH, Inc., Macungie, PA, USA), respectively. The impedance was executed at 10⁻¹~10⁷ Hz by using an impedance analyzer in a dark environment (Solartron 1260 coupled to the dielectric interface 1296, Farnborough, UK).

3. Results and Discussion

In the two-step or multi-step method of preparation of perovskite, the quality of the PbBr₂ film determined the morphology of the following perovskite film. Figure 1 shows the top-view SEM images of PbBr₂ films by introducing different concentrations of InBr₃. When there was no InBr₃ in the PbBr₂ precursor solution, as shown in Figure 1a, the surface of the obtained sample was rough, and the PbBr₂ grain distribution was extremely uneven, and a large area of exposed m-TiO₂ could be directly observed. When the PbBr₂ precursor solution was introduced into 0.03 M InBr₃, as shown in Figure 1b, the surface of the PbBr₂ film was flat, and the coverage of the m-TiO₂ film was increased, and the observable exposed m-TiO₂ area was significantly reduced. With the gradual increase in the concentration of introduced InBr₃ (Figure 1c–f), the PbBr₂ film appeared porous, but the number of pores decreased as the concentration of InBr₃ increased. Meanwhile, the porosity volume increased as the concentration of InBr₃ increased. From the cross-sectional view of PbBr₂ shown in Figure S1 (Supplementary Materials), we could clearly see that the pure PbBr₂ film has a flat surface and a uniform thickness of about 50~60 nm, and the m-TiO₂ was filled with PbBr₂. As the concentration of InBr₃ introduced gradually increased, the thickness of the PbBr₂ film also gradually increased (about 70 nm, 90 nm, 100 nm, 120 nm, 160 nm), and the film roughness increased. The above data could clearly demonstrate that the introduction of InBr₃ could effectively affect the morphology of the PbBr₂ film. The increase in the porosity volume, roughness, and thickness of the PbBr₂ film facilitated the diffusion of the subsequent CsBr solution, increased the reaction efficiency with CsBr, and then achieved the full growth of CsPbBr₃ grains [29,30].

To investigate the influence of InBr₃ on the structure of PbBr₂, XRD patterns of InBr₃:PbBr₂ films are shown in Figure 2a. It could be seen that Pure PbBr₂ was in the orthorhombic phase crystal structure (PDF#84-1181) [23]. When InBr₃ was introduced, for all concentrations of InBr₃ used, two new diffraction peaks of (011) and (200) crystallographic planes of PbBr₂ could be found at 2θ of 20.94° and 22.05°, but no diffraction peak belonging to InBr₃ or other protobromides were found. Since InBr₃ did not exist in the form of simple In³⁺ and Br^- in the DMF solution, it was self-ionized and formed various complexes [31,32,33]. Therefore, we speculated that in the process of PbBr₂ crystal growth, In (Ⅲ) could be in the form of free In³⁺ to replace a part of the Pb vacancy or exchange it with Pb, or the In cluster was directly bound to host lattice constituents [33,34,35]. Meanwhile, the PbBr₂ crystal was made to grow along multiple ordered crystal orientations. When using the XPS technique to prove the presence of In in PbBr₂ films, not surprisingly, characteristic peaks belonging to Br 3d, Pb 4f, and In 3d were found in the XPS spectra for the pure PbBr₂ and InBr₃:PbBr₂ films, as shown in Figure 2b. According to Figure 2c–e, the core level In3d_5/2 and 3d_3/2 were located at 445.4 eV and 452.9 eV, respectively, and the Pb 4f_5/2 and 4f_7/2 peaks in Pb 4f spectrum and Br 3d_3/2 and 3d_5/2 peaks in Br 3d spectrum all moved towards higher binding energies, which showed that Pb-Br interactions were enhanced after In³⁺ or In cluster incorporation [36]. Additionally, the EDS mapping was also utilized to confirm the presence of In in the InBr₃:PbBr₂ films. Figure S2 (Supplementary Materials) demonstrated that all elements were uniformly distributed in the corresponding film, especially, there was no aggregation of In elements.

Figure 3 depicts the top-view SEM images of perovskite films without and with InBr₃ with the corresponding cross-section SEM images inserted in the inset. The size of the crystal grain of the pure CsPbBr₃ was quite different, the film uniformity and coverage were also bad, and the bare m-TiO₂ could be clearly seen. As the concentration of introduced InBr₃ gradually increased (0.03~0.21 M), the coverage of m-TiO₂ by CsPbBr₃ films also gradually increased, and the size and number of pores in each film showed a downward trend. This morphological change was conducive to the performance of the perovskite cells. However, when the concentration of InBr₃ was further increased by 0.27 M, there were again obvious holes in the CsPbBr₃ film. This result indicated that the quality and surface CsPbBr₃ film depended on the morphology of the corresponding porous InBr₃:PbBr₂ film greatly that was, the morphology of CsPbBr₃ film could be modified by changing InBr₃ concentration.

The XRD patterns shown in Figure 4a revealed that all CsPbBr₃ films had a cubic structure (PDF#54-0752) [23], and the positions of the diffraction peaks were not significantly shifted to high or low angles, which demonstrated that, although In cluster could promote growth along multiple ordered crystal orientations, it could not change the phase of CsPbBr₃. When the concentration of the introduced InBr₃ was 0.00 M and 0.03 M, there existed two peaks located at 11.7° and 29.4°, respectively, which belonged to (002) and (213) lattice planes of the CsPb₂Br₅ phase [37]. As the concentration of InBr₃ was further increased (0.09~0.27 M), no obvious impurity peak belonging to CsPb₂Br₅ or Cs₄PbBr₆ phase could be observed. In fact, the control of the reaction rate between CsBr and PbBr₂ was a necessary condition for preparing CsPbBr₃ films with a high purity phase and high coverage. Based on Figure 1, the appropriate concentration of InBr₃ could make the PbBr₂ film have higher porosity, which provided more effective diffusion paths for the diffusion of CsBr methanol solution in the PbBr₂ film, and appropriately increased the contact area between CsBr and PbBr₂. That could also ensure the full growth of CsPbBr₃ crystal grains and, at the same time, could prevent the formation of the impurity phase due to excessive PbBr₂ or CsBr. However, if the concentration of the InBr₃ introduced into PbBr₂ was too low or too high, it was not conducive to controlling the reaction rate of CsBr and PbBr₂. In the process of the reaction, due to the incomplete reaction of the precursors or the excessive growth of crystal grains, the morphology of the CsPbBr₃ film was easily deteriorated, accompanied by the formation of byproducts. Further XPS was employed to certify the presence of the incorporated In³⁺ in the InBr₃:CsPbBr₃ film. Figure S3 (Supplementary Materials) exhibited the XPS of Cs 3d, Pb 4f, Br 3d, and In (Ш) 3d for the CsPbBr₃ and the InBr₃:CsPbBr₃ films, respectively. As seen in Figure 4b, compared with CsPbBr₃ film, two In signals corresponding to In 3d_5/2 and 3d_3/2 core levels were detected in InBr₃:CsPbBr₃ film, and Cs 3d, Pb 4f, and Br 3d all moved towards higher values, which means that the chemical state of the [PbBr₆]^4- octahedral was altered and Pb-Br and Cs-Br interactions were enhanced after replacing Pb²⁺ (1.7497 Å) with In³⁺ (1.6590 Å) with a smaller ion radius accompanied by the size of the [PbBr₆]^4- octahedral and the voids decreased [35]. The contraction of lattice and the enhancement of the spatial symmetry of the crystal structure caused by the incorporation of In³⁺ or In cluster could result in an efficient charge transport along with multiple directions, which perhaps was one of the important factors to improve the performance of CsPbBr₃ cells [35,38]. The EDS mapping was used to characterize the cross-sectional of InBr₃:CsPbBr₃, and it was confirmed that In was evenly distributed inside the perovskite, which indicated the successful incorporation of CsPbBr₃ by In (Figure S4, Supplementary Materials).

Subsequently, UV-vis Spectrometer, UPS, PL were used to characterize the cells with the FTO/c-TiO₂/m-TiO₂/CsPbBr₃ structure. Figure S5a (Supplementary Materials) shows the absorption spectra of the CsPbBr₃ with different concentrations of InBr₃. The absorption edge of each perovskite film was at approximately 530 nm within the visible region, which revealed that the concentration change of the introduced InBr₃ did not significantly affect the light absorption range of CsPbBr₃. Correspondingly, the calculated bandgaps (2.34 eV) did not reveal obvious and meaningful changes (Figure S5b, Supplementary Materials). As the concentration of InBr₃ increased, so did the film’s capacity to absorb visible light. This was mainly attributed to the phase-purity of the perovskite film and the full growth of crystal grains, which was beneficial to improve the short current density (J_SC) of the cells. The mechanism of this phenomenon was mainly attributed to the partial substitution of Pb²⁺ by In³⁺ or In cluster [34,35]. Figure 5a,b present the UPS spectra of the pristine and InBr₃ (0.21M):CsPbBr₃ films. By formula valence band maximum E_VBM = 21.22 eV − (E_cutoff − E_onset) [39,40], it could be calculated that the valence band (E_V) of CsPbBr₃ and InBr₃:CsPbBr₃ were −5.60 and −5.28 eV, respectively, which was ascribed to the rearrangement of electrons outside the Cs, Pb, and Br atoms after the incorporation of In³⁺ or In cluster [34,35]. Combined with Figure S5b, the corresponding calculated conduction band (E_C) was −3.26 and −2.94 eV, and the energy band diagram of isolated semiconductors of the PSCs using carbon electrodes is plotted in Figure 5c [39,41]. For HTL-free PSCs, Ev of the perovskite should be deeper than the work function (W_F) of the carbon electrode [39] so as to facilitate the extraction of photogenerated holes and reduce the energy loss of the holes during the transmission process [39]. Obviously, the incorporation of In³⁺ or In cluster effectively reduced the difference in interface energy levels, thereby facilitating the charge extraction and transfer and enhancing the photovoltaic performance of PSCs. In addition, the PL was conducted to analyze the carrier transfer behavior of CsPbBr₃ and InBr₃:CsPbBr₃ films. As shown in Figure 5d, all perovskite films showed the typical emission band around 523 nm. InBr₃ (0.21 M):CsPbBr₃ film showed a strong quenching in contrast with the pristine and other CsPbBr₃ films introduced with InBr₃, which indicated that InBr₃ could effectively inhibit the carrier recombination and enhance the charge extraction ability. The main reason behind this was that the defect density caused by the pinholes of CsPbBr₃ films surface, and the impurity phase of CsPbBr₃ films are obviously improved by adding InBr₃, and the enhancement of the spatial symmetry of the crystal structure caused by partial substitution of Pb²⁺ by In³⁺ or In cluster [34,35,39,42].

The HTL-free PSCs were synthesized based on the standard mesoscopic architecture of c-TiO₂/m-TiO₂/InBr₃:CsPbBr₃/carbon, and the cross-section of the complete device is given in Figure 6a. The current J–V characteristics of relevant devices under reverse scanning are presented in Figure 6b, the corresponding forward scanning curve is shown in Figure S6 (Supplementary Materials), and the key parameters including Jsc, Voc, FF, PCE, and hysteresis index (HI) are summarized in

Table 1. Key J–V parameters of the InBr₃:CsPbBr₃.

Samples	Scan	JSC (mA/cm2)	VOC (V)	FF	PCE (%)	HI
InBr3: 0.00 M	Forward Reverse	4.05 4.21	1.27 1.28	0.48 0.61	2.46 3.29	0.25
InBr3: 0.03 M	Forward Reverse	4.82 4.87	1.27 1.29	0.53 0.62	3.24 3.90	0.17
InBr3: 0.09 M	Forward Reverse	5.14 5.08	1.31 1.32	0.59 0.65	3.97 4.36	0.09
InBr3: 0.15 M	Forward Reverse	5.45 5.49	1.33 1.35	0.64 0.68	4.63 5.04	0.08
InBr3: 0.21 M	Forward Reverse	6.49 6.52	1.37 1.38	0.71 0.72	6.31 6.48	0.03
InBr3: 0.27 M	Forward Reverse	5.95 6.01	1.35 1.37	0.66 0.70	5.30 5.76	0.08

. The PCE of all devices with InBr₃ introduced were better than that of the pristine ones, and all parameters showed a regular trend of first increasing and then decreasing with an increase of the concentration of InBr₃ introduced. When the concentration of InBr₃ was 0.21 M, the corresponding device exhibited the best performance. Compared with the pristine device, the PCE of InBr₃ (0.21 M):CsPbBr₃ device was significantly improved from 3.29% to 6.48% with the continuously increased J_SC of 4.21 and 6.52 mA/cm², Voc of 1.28 and 1.38 V, FF of 0.61 and 0.72, and HI of 0.25 and 0.03. When the InBr₃ concentration was further increased to 0.27 M, the J_SC of the device dropped by about 0.5 mA/cm², while the Voc and FF did not change significantly, which was due to the deterioration of the InBr₃:CsPbBr₃ film morphology. The PEC of PSCs was determined by a variety of complex factors. According to the experimental results of SEM, UV-vis, UPS, and PL, the improvements of Voc and FF were due to the rise in the energy difference between the perovskite conduction band and electron transport layer, thereby reducing the energy loss of the holes in the transmission process. The improved Jsc was not only ascribed to the quality of the InBr₃:CsPbBr₃ film morphology or the increase in film coverage to absorb more photons to generate more electrons but also reduced vacancy defects in the optimized CsPbBr₃ films to improve the charge extraction and transfer process after the incorporation of In³⁺ or In cluster. In addition, the smaller hysteresis of the InBr₃:CsPbBr₃ cells performed than that of the pristine device might be enabled by the passivation function of InBr₃ to diminish the defects of Pb²⁺ and Br⁻ [34,35]. Figure 6c shows the external quantum efficiency (EQE) spectrum. The highest EQE value of 84% was achieved at the InBr₃ concentration of 0.21 M, whereas the reference devices with less or excessive InBr₃ concentration displayed lower EQE responses. This regular change was consistent with the results of J–V characteristics. Additionally, the integrated current density calculated by the EQE curve of each device was very close to the J_SC, and the mismatch was less than 5%. Figure 6d demonstrates the Nyquist plots of pristine CsPbBr₃ and InBr₃ (0.21 M):CsPbBr₃ devices measured at a reverse potential of 1.0 V and the corresponding equivalent circuit model. Table S1 (Supplementary Materials) also provides a list of the fitting values of the series resistance (R_s) and the charge recombination resistance (R_rec). After the introduction of InBr₃, R_rec increased from 765 to 1152 Ω, which showed that the incorporation of In had efficient repression of carrier recombination due to the significantly improved film formation quality of perovskites, thus reducing the trap state density and improving carriers mobility [39].

4. Conclusions

In the process of preparing the CsPbBr₃ film by the multi-step method, we introduced InBr₃ into the PbBr₂ precursor, so that the PbBr₂ film was transformed from a flat membrane to a porous membrane, which was beneficial to improve the reaction efficiency of CsBr and PbBr₂, reduced the impurity in CsPbBr₃, and optimized the surface morphology, and, finally, enabled the performance of CsPbBr₃ PSCs to be significantly improved. When combined with host lattices, the In³⁺ or In cluster could effectively suppress the carrier recombination in the CsPbBr₃ film and shift up the Ev of CsPbBr₃, thereby enhancing the charge extraction and transportation capabilities. When the InBr₃ concentration in the PbBr₂ precursor solution was 0.21 M, the InBr₃:CsPbBr₃ device presented the best photovoltaic performance with a PCE of 6.48% and, especially the V_OC significantly increased by 100 mV compared with the pristine CsPbBr₃. These research results confirmed that InBr₃ has solid potentials for improving the performance of CsPbBr₃ PSCs and also provided a reference for InBr₃ or some other metal bromide applications in the inorganic CsPbI₃ PSCs field and developmental direction.