Figures 1A,B illustrate the process of growing MAPb(I_1–xBr_x)₃ thin alloy film using the SPVD technique. 15 × 20 mm glass substrates were trimmed and cleaned sequentially in acetone, isopropanol, and deionized water using an ultrasonic bath. The substrates were dried by blowing with a pressurized stream of pure nitrogen gas. PbI₂ (99.9%) and MAI (98%) powders were obtained from Sigma Aldrich and used as received. Before the deposition, the chamber was cleaned using a solution of extran diluted with deionized water (volume ratio 1:4) and high-pressure cleaner, to remove the effects of residual contaminants. PbI₂ and MAI powders were placed in two separate cylindrical boron nitride boats B₁ and B₂ and inserted into individual heating coils inside an evaporation jar. The coils are connected to a power supply via switches S1 and S2 as shown in Figures 1A,B. The evaporation chamber was evacuated to a pressure of 2.0 × 10^–5 mbar. An in-built quartz crystal monitor, placed at the same level as the substrate inside the evaporation chamber, was used to monitor the thickness of thin-films during the deposition process. The crystal monitor was calibrated for PbI₂ thickness measurement by setting the density to 6.16 g cm^–3 and Z-factor to 1.10, and MAI film thickness was monitored by setting the density to 1.20 g cm^–3 and Z-factor to 2.70. 100 nm thick PbI₂ thin-film was first deposited followed by the deposition of 500 nm thick MAI to form MAPbI₃. The heat generated from the heating crucibles raised the temperature of the substrates to 140°C during the deposition of PbI₂, and to 95°C during the deposit of MAI. Post-annealing of the sample was done in an oven with the temperature maintained at 100°C for 10 min. After the post-annealing process, the prepared MAPbI₃ was loaded in the evaporation chamber, with boats B1 and B2 now containing PbBr₂ and MABr, respectively, as shown in Figure 1B. PbBr₂ thickness was measured by setting the density to 6.66 g cm^–3 and Z-factor to 1.20 and MABr film thickness was monitored by selecting the density to 1.30 g cm^–3 and Z-factor to 2.80. The thickness of PbBr₂ was kept constant at 100 nm while that of MABr was varied from 300 to 500 nm, in steps of 100. Also, the heat generated from heating the crucibles raised the temperature of the substrates to 140°C during the deposition of PbBr₂ and 95°C during the deposit of MABr.

Figure 1C represents the schematic of FTO/MAPb(I_1–xBr_x)₃/Au devices fabricated solely for the determination of charge carrier mobility and trap density using the space-charge-limited current theory. The MAPb(I_1–xBr_x)₃ was grown on FTO as described above and gold (Au) contacts were later deposited on the MHaP through a 0.6 mm diameter shadow mask. Figure 1D shows the schematic diagrams of FTO/c-TiO₂/m-TiO₂/MAPb(I_1–xBr_x)₃/Au devices for dark capacitance-voltage analyses. The precursor solution for deposition titanium dioxide compact (c-TiO₂) layer was prepared by mixing 0.55 ml of titanium diisopropoxide bis(acetylacetonate) and 5 ml of 1-butanol to give 0.3 M TiO_x solution. The c-TiO₂ layer was deposited by spin-coating the 0.3 M TiO_x solution on FTO at 3,000 rpm for 30 s. The film was dried in an oven at 125°C for 5 min. The experiment is repeated once before annealing in a furnace at 500°C for 30 min to produce a compact-TiO₂ layer. TiO₂ paste was prepared by dissolving 1.2 g of TiO₂ nanoparticles in 10 ml of ethanol, 4 ml of terpineol, and 10% ethyl cellulose (0.6 g ethyl cellulose in 10 ml of ethanol). The mixture was stirred overnight at room temperature. The paste was spin-coated at 3,000 rpm for 30 s and dried at 125°C for 5 min before sintering at 500°C for 30 min to form the mesoporous-TiO₂ (m-TiO₂) layer. The MAPb(I_1–xBr_x)₃ thin alloy film was grown on the as described above. Finally, Au contacts were deposited through a 0.6 mm diameter shadow mask.

XRD patterns of the thin-films were measured by a Bruker D2-Phaser X-ray diffractometer using Cu Kα radiation with a wavelength of 1.5405 Å. The angle between the incident and diffracted rays, 2θ was varied from 10 to 50° in steps of 0.05. The measured patterns were used to determine the structures, crystallinity, microstrain, and dislocation densities of the thin-films. The morphological properties were revealed by a Field Emission Scanning Electron Microscope (FE-SEM Zeiss Crossbeam 540), with an accelerating voltage of 2.0 kV. Grain size analysis was performed from the FE-SEM images following the American Standard for Testing Materials (ASTM) using the Image J Software. The optical absorption spectra of the films were recorded using the CARY 100 BIO UV-Vis spectrometer with a wavelength of incident light ranging from 400 to 800 nm.

A 4192A LF Impedance Analyser with measurement frequency ranging from 5 Hz to 13 MHz, was used for the dark current density-voltage (J-V) and capacitance-voltage (C-V) measurements. The C-V measurements were conducted at a constant frequency of 1 MHz so that the defects within the device should not respond to electric field variations (Dharmadasa et al., 2019). The measured values were then used for Mott-Schottky (M-S) analysis to determine the presence of charge accumulation during device operation. Current density-voltage (J-V) characteristics and device performance parameters of the solar cells under illumination were determined using the automatic current-voltage measurement system (Osilla). The light measurements were performed under a solar simulator (Oriel LCS-100^TM Small Area So11A Series, Newport) with simulated solar output conditions of 100 mW/cm² and AM1.5 G reference spectral filtering.

Figure 2A shows the X-ray diffractograms of MAPbI₃ on glass, MAPbBr₃ on glass, and MAPbBr₃ on MAPbI₃ bottom layer. The diffractogram of MAPbI₃ shows sharp planes corresponding to the pure tetragonal crystal structure and I4/mcm space group, according to the literature (Jeon et al., 2014; Saidaminov et al., 2015; Shi et al., 2015). The diffraction patterns of MAPbBr₃ matches the cubic crystal structure which has Pm3̄m space group (Peng et al., 2016; Wang et al., 2017; Tisdale et al., 2018). The thin MAPb(I_1–xBr_x)₃ film diffractograms (Figures 2A,B) show the presence of all the prominent peaks of cubic MAPbBr₃ and the (001) peak corresponding to PbI_2, but do not contain the principal (110) peak characterizing tetragonal MAPbI₃. This points to the fact that the tetragonal phase of MAPbI₃ is transformed to the cubic phase of MAPbBr₃ as MAPb(I_1–xBr_x)₃ is formed, consistent with previous reports (Noh et al., 2013; Gil-escrig et al., 2015; Pistor and Burwig, 2018). The value of bromine mole ratio x was determined using an empirical relationship between the bandgaps of the MAPb(I_1–xBr_x)₃ alloy and those of its MAPbI₃ and MAPbBr₃ components, given in Equation 1,

where b is the bowing parameter that accounts for the curvature of the bandgap energies as a function of composition. Noh et al. (2013) obtained b as 0.33 for Br-I mixing of the MAPbI₃ and MAPbBr₃ components. Similarly, Zhang et al. (2016) obtained b as 0.21 for Br-I mixing upon tuning the bandgap of MAPb(I_1–xBr_x)₃ single crystal. The small values of b for MAPb(I_1–xBr_x)₃ are an indication that MAPbI₃ and MAPbBr₃ are fully miscible in solid solutions. By substituting the bandgaps (from the UV-Vis absorption analysis) and b as 0.33 into Equation 1, we obtained x to be 0.95 for 300 nm and 0.94 for 400 nm MABr which corresponds to structural formulae of MAPb(I₀.₀₅Br_0.95)₃ and MAPb(I₀.₀₆Br_0.94)₃, respectively. Similarly, x is 0.89 when the thickness of MABr equals 500 nm giving the structural formula as MAPb(I₀.₁₁Br₀.₈₉)₃. The values of x for pure MAPbI₃ and pure MAPbBr₃ are 0.0 and 1.0, respectively. Previous studies have shown that MAPb(I_1–xBr_x)₃ with such high amounts of x exists in the cubic phase (Noh et al., 2013; Zhang et al., 2016). Pistor and Burwig (2018) gave an elaborate explanation for the formation of MAPb(I_1–xBr_x)₃ by exposing MAPbI₃ to MABr vapor and MAPbBr₃ to MAI vapor. They proved that it was due to reversible halide exchange at high substrate temperatures and added that it is much easier to replace I with Br than to replaces Br with I. This is because of the smaller radius of Br and small formation enthalpies of MAPbBr₃. Furthermore, they showed that Br replaces I when MAPbI₃ is exposed to MABr vapor while I replace Br when MAPbBr₃ is exposed to MAI vapor. Thus, the high Br content we obtained shows a good correlation with the explanations of Pistor and Burwig. The presence of PbI₂ peak in the diffractograms of MAPb(I_1–xBr_x)₃ shown in Figures 2A,B, is consistent with previous results in the literature (Gil-escrig et al., 2015; Misra et al., 2016). The (001) peak characterizing PbI₂ is not seen in the diffractogram of MAPbI_3. However, it appears in the diffractogram of MAPb(I_1–xBr_x)₃. The shows that PbI₂ is formed with the MAPb(I_1–xBr_x)_3. The formation of PbI₂ alongside MAPb(I_1–xBr_x)₃ may include, amongst other processes, the decomposition of MAPbI₃ as a result of prolonged in-situ annealing and high substrate temperatures within the evaporation chamber as the MAPbBr₃ layer was being prepared. We have shown previously that post-annealing of MAPbI₃ thin-films at 100°C for 40 min leads to the appearance of a sharp PbI₂ peak, which decreases in intensity upon increasing the time to 60 min because the MAPbI₃ is formed again (Fru et al., 2020a). Similarly, annealing MAPbI₃ at temperatures beyond 120°C lead to the decomposition of MAPbI₃ to PbI₂ (Pistor and Burwig, 2018). The intensity of the PbI₂ peak reduces while that of the (100) peak increases as the thickness of MABr is increased from 400 to 500 nm. This may be due to the transformation of the available PbI₂ to MAPb(I_1–xBr_x)₃ as the MABr thickness increased. It could be because of the high temperature (140°C) of the substrate as PbBr₂ was being deposited on MAPbI₃ bottom layer. The exact mechanism of the transformation is not fully understood.

The Williamson-Hall (W-H) plot was used to determine the crystallite size and microstrain from the X-ray line broadening. The W-H plots (Supplementary Figure 1) are based on the Williamson-Hall relation given by Equation 2 (Williamson and Hall, 1953),

where D is the crystallite, β is the full width at half maximum, λ is the wavelength, θ is the Bragg’s diffraction angle, ε is the microstrain and K is the Scherrer’s constant. The Scherrer’s constant is determined by the crystallite shape and is considered as 0.94 for spherical crystallites with cubic symmetry. Figures 2C,D shows how the average crystallite size and microstrain vary as the thickness of MABr is increased from 300 to 500 nm. The crystallite size is observed to range from 28 to 42 nm, in close agreement with the literature (Jang et al., 2015). Besides, the crystallite size first increases to 41 nm when the MABr thickness is 400 nm before decreasing to 28 nm as the thickness is increased to 500 nm. This correlates with the change in intensity of the PbI₂ peak shown in Figure 2B. Thus, the irregular variation in crystallite size could be due to the change in concentration of PbI₂ present in the thin-film. Figure 2D presents the microstrain of the MAPb(I_1–xBr_x)₃ as the MABr is increased from 300 to 500 nm. The microstrain decreases as the thickness MABr rises from 300 to 500 nm. Increasing the thickness of MABr from 300 to 500 nm causes the Br/I mole ratio to change from 24.0 to 7.0. The film containing 500 nm thick MABr with the smallest Br to I ratio has the least microstrain. Thus, the microstrain of MAPb(I_1–xBr_x)₃ decreases as the Br to I ratio decreases.

Figures 3A,B show the absorbance of MAPbI₃, MAPbBr₃ and MAPb(I_1–xBr_x)₃, where x = 0.95, 0.94, and 0.89. As previously discussed, pure MAPbI₃ and MAPbBr₃ have values of x corresponding to 0.0 and 1.0, respectively. The respective absorption onsets of MAPbI₃, MAPbBr₃, and MAPb(I_0.11Br_0.89)₃ are seen at 785, 552, and 580 nm (Figure 3A), corresponding to bandgaps of 1.58, 2.25, and 2.14 eV, respectively. The bandgap of the MAPb(I_0.11Br_0.89)₃ is closer in value to that of MAPbBr₃. This is consistent with MAPb(I_0.11Br_0.89)₃ taking up the cubic crystal structure of MAPbBr₃ as revealed by XRD results above. Incorporating a small amount of I does not change MAPbBr₃ crystal structure (Lehmann et al., 2019). This is because Br-rich compositions have stronger intermolecular forces. Thus re-orientation of the cations is suppressed. Furthermore, pure MAPbI₃ absorbs more than MAPb(I_1–xBr_x)₃ and MAPbBr₃, and MAPb(I_1–xBr_x)₃ more than pure MAPbBr_3. This could be due to the presence of I in MAPbI₃ and MAPb(I_1–xBr_x)₃ thin-films. The greater the mole ratio of I in the structure, the higher its absorbance. Also, an increase in the I mole ratio causes a redshift in the absorption edge, indicative of a decrease in the optical bandgap. This decrease perhaps is because I has a greater ionic radius than Br which causes an expansion of the lattice (Dong et al., 2020). Figure 3B shows that the absorption intensity and onset of MAPb(I_1–xBr_x)₃ depend greatly on the value of x. As x decreases, the onset of absorption increases. A plot of the bandgap against x reveals a highly linear relationship with a positive slope, consistent with previous reports (Gil-escrig et al., 2015). The linear relationship shows that the bandgap of MAPb(I_1–xBr_x)₃ could be tuned for x ranging from 0.89 to 1.0 by increasing the thickness of MABr. This is quite a small range when compared to alloying using other methods but useful to form MAPb(I_1–xBr_x)₃ alloys with small I content which shows improved absorption than MAPbBr₃. Another prominent feature in the absorption spectrum of MAPb(I_1–xBr_x)₃ shown in Figure 3B is the sharp absorption onset, which according to Stranks et al. (2019), is an indication of low levels of energetic disorder. The onsets of absorption were used to calculate the bandgap and the results plotted against the x (Figure 3C). The bandgap is observed to increase linearly with x as expected for HaP with high Br content. The lattice constants of the cubic MAPb(I_1–xBr_x)₃ structures were computed using the same procedure reported in the literature (Fru et al., 2020b), and the results plotted against x as shown in Figure 3D. We observed that the lattice constant is inversely related to x, in agreement with the observation of Kulkarni et al. (2014) who used a solution sequential deposition method. Similarly, the bandgap is inversely related to the lattice constant, as shown in Supplementary Figure 2. Perhaps, the decrease in bandgap with increase MABr thickness is due to increase in the lattice constant. This is because an increase in the lattice constant causes the interatomic distance to increase. Therefore, the binding force between the valence electrons and nucleus is reduced. Thus, the energy required to convert a bound electron to a free or conduction electron reduces.

Figure 4 shows the SEM micrographs and variation of average grain size with the thickness of MABr for MAPb(I_1–xBr_x)₃, where x = 0.95, 0.94, and 0.89. We observed that all the films showed full surface coverage of the substrate. The grains are densely packed, randomly oriented, and with few pinholes. The morphology evolves from a mixture of cuboidal grains and gel-like structures, at 300 nm MABr thickness, to one with only faceted bimodal distributed grains, when the thickness is 500 nm, which is characteristic of the cubic MAPbBr₃ structure (Fru et al., 2020b) shown in Supplementary Figure 3. This correlates well with the increase in the intensity of the (100) distinct peak for cubic MAPbBr₃ as the thickness of MABr is increased from 300 to 500 nm. The gel-like structures seen when the MABr thickness is 300 nm (MAPb(I_0.05Br_0.95)₃ can be linked to the rod-like structures that were seen on faceted grains solution deposited (MAPb(I_0.05Br_0.95)₃ by Kulkarni et al. (2014), and which were associated with high Br content. The average grain size is seen to increase from 150 to 320 nm as the thickness of MABr increases from 300 to 500 nm. The 500 nm MABr thick thin-film, with the most intense (100) characteristic peak corresponding to MAPbBr₃ (Figure 2B), has the highest average grain size. This points to the fact that an increase in average grain size could be due to a rise in the number and size of the MAPbBr₃ grains in the thin alloy films.

Figures 5A,B represents the semi-log and double-log plots of the dark current-voltage (I-V) characteristics of FTO/MAPb(I_1–xBr_x)₃/Au devices, respectively, where x is 0.95, 0.94, and 0.89. The semi-log I-V plot is meant to verify whether the contacts are ohmic, representing an infinite source of charge carriers. This confirms the possibility of using the space-charge-limited current (SCLC) measurement to calculate charge carrier mobility and trap density. The plots are nearly symmetrical, a sign that the contacts of the devices offer a very small barrier to charge injection, and SCLC measurement could be applied. The double-log plot shows two distinct regions for all values of x with slopes equals to 1.0 at low voltages and greater than 2.0 at high voltages, as shown in Figure 5B. This points to the fact that the charge transport mechanism in MAPb(I_1–xBr_x)₃ of the FTO/MAPb(I_1–xBr_x)₃/Au devices is ohmic conduction at low voltages, as is commonly reported in the literature (Röhr et al., 2017, 2018), and space charge limited at high voltages. Two factors are responsible for slopes greater than 2.0; the presence of traps and the existence of a built-in voltage (Röhr et al., 2018). Since the semi-log plot of the forward and reverse current-voltage characteristics are symmetric, then a slope greater than 2.0 is only due to the presence of charge traps. The trap density was calculated using Equation 3 as explained in our earlier report (Fru et al., 2020b).

where N_t is trap density, d is film thickness, q is the electronic charge, V_TFL is the voltage at the point of transition from the ohmic to the space-charge-limited region, ε₀ is the vacuum permittivity, and ε_r the dielectric constant. The dielectric constant for halide perovskites at 295K is typically 25 (Le Corre et al., 2020). The N_t decreased non-uniformly as the thickness of MABr was increased, as shown in Figure 5D. This could be due to the increase in grain size as explained above.

The charge carrier mobility was estimated by SCLC interpretation of the dark I-V characteristics. SCLC mobility measurements for an ideal single-carrier device that contains a trap-free and un-doped semiconductor with no injection barriers are achieved using the Mott-Gunney law (Röhr et al., 2017, 2018). However, as previously explained, there is evidence of the presence of traps. Thus, the charge carrier mobility calculation was performed using the modified Mott-Gunney model shown in Equation 4 (Moiz et al., 2005),

where J is the current density, V is a voltage at the trap-filled limit (TFL); onset of space charge limited region, and θ is the trap factor and is calculated as explained in the literature (Moiz et al., 2005; Fru et al., 2020b). Figure 5C shows that the carrier mobility increases linearly with an increase in the thickness of MABr. This is consistent with the decrease in trap density and an increase in grain size, leading to less charge carrier scattering in the MAPb(I_1–xBr_x)₃ active layer as MABr thickness is increased. The calculated charge carrier mobility values are within the range for inorganic semiconductors (10^–4 − 10³ cm² V^–1 s^–1) (Yu et al., 2015). However, they are two orders of magnitude lower than values of single crystals calculated using the SCLC method (Saidaminov et al., 2015; Herz, 2017; Chen et al., 2018). This is expected since polycrystalline films have grain boundaries that tend to decrease the charge carrier mobility.

Figures 6A,B illustrates the plot of capacitance per unit area (C) against bias voltage (V) and Mott-Schottky (M-S) plots for FTO/c-TiO₂/m-TiO₂/MAPb(I_0.11Br_0.89)₃/Au solar cells, respectively. M-S analysis is a well-established experimental technique to determine the charge carrier density (N_d) and built-in voltage (V_bi) from C-V measurements. It is mostly applied in conventional semiconductor devices with p-n and semiconductor/metal junctions that have a fixed depletion layer and space charge region. The junction capacitance per unit area (C_dl) is given by Equation 5,

where is the vacuum permittivity, ε_r is the dielectric constant of the donor, q is the electronic charge, V is the applied voltage. M-S plot is a graph of C_dl^–2 vs. V from which N_d is obtained from the gradient of the linear part and V_bi from an extrapolation of the linear part to the voltage axis (Figure 6B). The linear region with a negative slope exists on the plot under certain conditions which includes no electrons/holes in the space charge region and a homogenous distribution of the space charges (Shockley, 1948; Fischer et al., 2018). Our results show four distinct regions including an exponential increase in capacitance for voltages ranging from −0.5 to −2.0 V, a constant capacitance from −0.5 to 0.0 V, a linear dependence region from 0.0 to 0.2V, and an exponential increase in capacitance voltages greater than 0.2 V. Bera et al. (2017) obtained a similar deviation from the conventional straight line for other semiconductors. They attributed an exponential increase in capacitance with an increasing negative bias to the accumulation of electrons at the junctions, constant capacitance to a depletion region, and exponential capacitance at a forward bias to the accumulation of holes. Other reasons for constant capacitance may include depletion region exceeds the thickness of perovskite absorber (perovskite are entirely depleted), presence of injection barriers (Yadav et al., 2017), and accumulation of space charges at the interfaces (Hegedus and Shafarman, 2004; Almora et al., 2016; Fischer et al., 2018; Lee and Lee, 2019). Charge accumulation maybe due to a low level of doping, the intrinsic nature of the perovskites, and the presence of mobile ions (Fischer et al., 2018). The exponential capacitance, also known as the diffusion capacitance (Hegedus and Shafarman, 2004), has previously been attributed to the charged ions accumulation occurring at the TiO₂/perovskites and perovskites/Au interfaces in the literature (Almora et al., 2016; Fischer et al., 2018). The presence of capacitances other than the depletion capacitance shows N_d and V_bi cannot be determined from M-S plot. However, it reveals information on charge accumulation at interfaces that are due to ion migration. Thus, we could tell which of the optimum thin-films had a significant level of ion migration. This is useful in perovskite solar cells since a significant level of ion migration will cause hysteresis. There is a difference between V_bi (1.75 eV) from M-S plot and that from the knee of dark current density-voltage (J-V) characteristic (Supplementary Figure 4A). The dark J-V characteristic confirms that the device is highly rectifying with a large V_bi of 1.6 eV. However, the FTO/MAPb(I_1–xBr_x)₃/Au devices (Figure 5A) are only slightly rectifying. This means the rectification is a result of the introduction of TiO₂ and possibly at the TiO₂/MAPb(I_1–xBr_x)₃ junction. Thus, V_bi is the difference between Fermi levels of TiO₂ and MAPb(I_1–xBr_x)₃. The rectification could also be between the TiO₂ and gold electrodes if MAPb(I_1–xBr_x)₃ is intrinsic. The large V_bi at the p-n junction is fundamental for effective charge separation in solar cells leading to less recombination within the depletion layer (Lee and Lee, 2019; Rau and Kirchartz, 2019).

Figures 6C,D shows the C-V and M-S plots for FTO/TiO₂/MAPbI₃/Au solar cells, respectively. The capacitance increases sharply for the first decrease in negative bias, then stays nearly constant as negatives bias decreases to 0.0 and for an increase in the forward bias voltage from 0.0 to 2.0 V. Our C-V curves and Mott-Schottky plots were consistent with those obtained at lower frequencies by Liu and Zhang (2014) and Almora et al. (2016). The constant capacitance for the long-range of biasing voltage may be a sign that the depletion width exceeds the thickness of the perovskites absorber as earlier mentioned with no accumulation of charges at the interfaces. Supplementary Figures 4C,D stands for the J-V and semi-log I-V plots for FTO/TiO₂/MAPbI₃/Au solar cells. The knee in Supplementary Figure 4C shows that the device is rectifying with V_bi of 0.7 V. In our earlier work, we showed that FTO/MAPbI₃/Au device was near ohmic. Thus, the rectification may also be due to the difference between the work functions of the TiO₂ and the Au electrode.

Figure 7 shows the current density-voltage characteristic for FTO/MAPbBr₃/Au, FTO/MAPb(I_0.11Br_0.89)₃/Au, and FTO/MAPbI₃/Au solar cells under illumination. The power conversion efficiency (PCE), fill factor (FF), short circuit current density (J_sc), open-circuit voltage (V_oc), short resistance (R_sh), and series resistance (R_s) of the devices are presented in Table 1. The PCE of FTO/MAPbBr₃/Au, FTO/MAPb(I_0.11Br_0.89)₃/Au and FTO/MAPbI₃/Au solar cells were 0.56, 0.62, and 1.15%. These low PCE values could be due to the absence of charge transport layers and are comparable to those of other charge transport layer-free solar cells in the literature (Chen et al., 2014b; Peng et al., 2016; Duan et al., 2018). However, PCE values show that the halide perovskites are functional. The solar cell performance improves as x decreases. This could be due to the improvement in absorbance, decrease in bandgap and rise in charge carrier mobility as the bromine mole ratio decreased. Besides, the solar cell with x as 1.0 has the highest FF. This is consistent with the high R_sh value and could be because the MAPbBr₃ thin-film is very compact and does not have pinholes. The small V_oc values could be due to the absence of charge transport layers.