# Quasi-2D Ruddlesden-Popper phase induced preferentially oriented FA1-xMAxPbI3 for high-performance sequential perovskite solar cells

Zihao Li, Marie Cresp, Karol Vegso, Nada Mrkyvkova, Peter Siffalovic, Thierry Pauporté

► To cite this version:

Zihao Li, Marie Cresp, Karol Vegso, Nada Mrkyvkova, Peter Siffalovic, et al., Quasi-2D Ruddlesden-Popper phase induced preferentially oriented FA1-xMAxPbI3 for high-performance sequential perovskite solar cells. Chemical Engineering Journal, 2025, 505, pp.159121. 10.1016/j.cej.2024.159121. hal-05378037

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Submitted on 22 Nov 2025

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# Quasi-2D Ruddlesden-Popper phase induced preferentially oriented FA₁-xMAxPbI₃ for high-performance sequential perovskite solar cells

Zihao Li¹, Marie Cresp¹, Karol Vegso²,³, Nada Mrkyvkova²,³, Peter Siffalovic²,³, Thierry Pauporte¹*

¹ Chimie ParisTech, PSL Research University, CNRS, Institut de Recherche de Chimie Paris (IRCP), UMR8247, 11 rue P. et M. Curie, F-75005 Paris, France.

² Institute of Physics, Slovak Academy of Sciences, Dubravska cesta 9, 84511 Bratislava, Slovakia.

³ Center for Advanced Materials and Applications (CEMEA), Slovak Academy of Sciences, Dubravska cesta 5807/9, 84511 Bratislava, Slovakia.

*Author for the correspondence. E-mails: thierry.pauporte@chimieparisrech.psl.eu

## Abstract

Incorporating large alkylammonium cations or crystal seeds into 3D perovskite has been found effective in manipulating crystallographic orientation and increasing the intrinsic stability of polycrystalline perovskite thin films. However, an excess of large-size cations within the bulk layer can detrimentally impact charge transport, resulting in relatively suboptimal device performance. In this study, we substituted a portion of methylammonium chloride (MACl) with an appropriate amount of long-chain n-butylammonium chloride (BACl) in order to modulate the crystallization kinetics of two-step FAPbI₃ polycrystalline films. We show that the addition of BA⁺ cation facilitates the *in-situ* formation of Ruddlesden-Popper (RP) quasi-2D phases (BA₂MAₙ-₁PbₙI₃n+₁, n≥2) at the initial crystallization stage, which, in turn, enables the growth of the preferentially (100) oriented 3D α-FAPbI₃ perovskite films. Furthermore, the mixed alkylammonium system attenuates the elimination rate of dissociation of hydrochloride and alkylamine from deprotonated alkylammonium Cl-salts, yielding films with homogeneous morphology and negligible trap-mediated recombination. Remarkably, only traces (measured at 2 % cation proportion) of BA molecules were detected on the layer surface after complete induction of the oriented growth. The optimal device employing the mixed alkylamine hydrochloride achieved a power conversion efficiency of 23.23% and exhibited exceptional stability, retaining 90% of its initial performance over 2000 hours of storage in a nitrogen-filled glovebox.

**Keywords:** Perovskite solar cell; Directed growth; Crystallization; High performance solar cells

## 1 Introduction

With the rapid development of perovskite solar cells (PSCs) over the past decade, a multitude of methods have been developed to obtain highly crystallized and defect-free perovskite thin films, such as additive engineering[1], designed bulk heterojunction[2] and interfacial modification.[3, 4] Although these techniques have markedly increased power conversion efficiency (PCE) to unprecedented levels[5], the precise manipulation of perovskite crystalline growth to obtain the desired crystallographic orientation [6] and minimize bulk trap-state density remains a formidable
44 challenge.

45 In the realm of the solution deposited organic halide perovskite thin films, crystallinity and
46 surface morphology are highly dependent on the volatilization direction and rate of the solvents and
47 volatile molecules.[7-9] An anisotropic volatilization process can cause non-uniform crystal growth,
48 resulting in undesired film texture, especially when non-(100) oriented perovskite thin films are
49 formed.[10] Yet, the optoelectronic properties of devices are closely related to the preferential
50 orientations of the perovskite layers, an area still relatively under-explored in prior studies. It is well
51 known that the (100), (110), and (111) are dominant textures of α-phase formamidinium lead iodide
52 (FAPbI₃) perovskite thin films. Notably, the (100) texture has attracted attention for its superior
53 photovoltaic performance owing to its more n-type characteristics, reduced defect density and higher
54 carrier mobility.[11] Theoretical analyses additionally indicate that the (100) texture promotes an
55 increased dielectric constant due to its distinct atomic density and symmetry, resulting in diminished
56 exciton binding energies conducive to the generation of free carriers and augmented photocurrent.[12]
57 Regarding the dependence of stability on the preferential orientation, the (111) texture has been
58 shown effective in enhancing the moisture resistance of PSCs, as crystals dominated by the (111)
59 texture exhibit thermodynamic stability against interactions with water molecules.[13] Hence, careful
60 tailoring of favorable preferential film orientation, while limiting undesirable textures or non-
61 isotropic orientations, is crucial for enhancing the optoelectronic properties of perovskite thin layers
62 and elevating the overall performance of PSCs.

63 Thus far, methylammonium chloride (MACl) has been the primary dopant utilized for the
64 deposition of high-purity α-phase FAPbI3 perovskite layers, either through the one-step[5, 14] or two-
65 step method.[6, 15, 16] The addition of MACl facilitates the coordination of Cl- ions at the octahedra
66 terminal corner, which suppresses the formation of face-shared octahedra (δ-phase), resulting in a
67 lower phase transition barrier for α-phase FAPbI3 at room temperature. Optimizing the MACl to
68 FAPbI3 incorporation ratio in the precursor solution has been proven effective in enhancing the
69 preferential crystallographic orientation during perovskite thin film growth.[16] As previous
70 references mentioned,[17, 18] incorporating over 40 mol% of MACl into FAPbI3 precursor
71 effectively induces a predominant α-phase (100) texture in FAPbI3 polycrystalline thin films,[19]
72 thereby significantly enhancing the efficiency and stability of PSCs. Further increase of MACl ratio
73 to 100 mol% could facilitate a 45° lattice tilt and promote uniaxial-oriented growth of (111) textured
74 films of α-FAPbI3 at room temperature.[20] Despite these advances, residual MACl introduced
75 during annealing remains challenging to completely remove from the perovskite layer, granted
76 that chlorine will be entirely removed as hydrochloride (HCl) due to its stronger proton affinity as
77 mentioned in the references.[5, 14, 15] The potential residual methylammonium (MA+) resulting
78 from incomplete deprotonation[21] will cause the increase of the bandgap and compromise
79 stability.[22] Consequently, it is essential to identify viable alternatives to MACl to achieve vertically
80 oriented and phase-pure α-FAPbI3 perovskite thin films.

81 In contrast to methylammonium, long-chain or larger radii deprotonated alkylammoniums
82 possess higher boiling points, suggesting a slower volatilization rate compared to deprotonated
83 methylammonium. This characteristic is anticipated to effectively influence the crystallization
84 kinetics of α-phase FAPbI3.[14] Additionally, larger-size alkylammoniums tend to intercalate within
85 PbI6 octahedron layers, precipitating the formation of low-dimensional perovskites such as one-
86 dimensional (1D), two-dimensional (2D) or quasi-2D lead halide perovskites.[23] In the
87 crystallization process of 3D perovskites, these co-existing low-dimensional perovskites, such as 2D
crystal seeds or bottom template film,[24] can autonomously act as the nucleation centers,[25] fostering the epitaxial growth of 3D perovskite on the 2D perovskite template. This reduces the formation energy and facilitates the nucleation and subsequent crystal growth kinetics.[11] However, excessive incorporation of 2D perovskites may lead to reduced photovoltaic performance owing to their dielectric properties.[26] Hence, controlling the residual quantity of 2D cations in 3D perovskite layer is crucial to realize the efficient charge carrier transport and achieve high-performance photovoltaic devices.

In this work, we utilized a combination of alkylamine hydrochlorides (BACl and MACl) to modulate the crystallization of α-phase FAPbI₃ perovskite layers via a two-step spin-coating method, in which the lead iodide and organo-halogen compounds solutions are deposited sequentially. The presence of synergetic BA⁺ and MA⁺ cations *in-situ* generates Ruddlesden-Popper (RP) quasi-2D phases (BA₂MAₙ-₁PbₙI₃n+₁, n≥2), exhibiting a strong preference for horizontal orientation. This promotes the (100)-oriented preferential growth of 3D FAPbI₃ polycrystalline films with (100) texture. Analysis of the resulting perovskite thin film by NMR and photoluminescence (PL) reveals that remaining BA molecules are predominantly confined to surface layer in the form of quasi-2D phases, comprising only a minor fraction of residual 2D cations (~2.0% cation portion). Glow-discharge optical emission spectroscopy (GD-OES) quantitatively assesses the impact of BACl on the separation rates of chlorine and nitrogen during the annealing process. The resulting perovskite thin films exhibit larger grain sizes, longer charge carrier lifetimes, and lower trap-state densities, all contributing to improved photovoltaic performance. The synergistic incorporation of optimized alkylamine hydrochlorides as dopants into the FAPbI₃ perovskite layers significantly enhances the performance and stability of the PSCs.

## 2 Results and Discussion

### 2.1 Effect of BACl on 3D perovskite crystallinity and orientation
<figure><img src="image_2.png" /><figcaption>**Figure 1.** (a) FWHM and integral intensity of the (100) diffraction peak of FAPbI₃ phase (MACl mol% + BACl mol% = 45 mol%). (b) XRD patterns of fully annealed FAPbI₃ perovskite thin films with different alkylammonium hydrochlorides. (c) NMR spectra of the control and target film (10 thin annealed films dissolved in DMSO-d₆). (d) Estimated organic cation distribution in control and target perovskite thin films based on NMR results. (e) and (f) 2D-GIWAXS patterns of the control and fully annealed (target) perovskite thin films. (g) and (h) zoomed-in 2D-GIWAXS patterns of control and target samples. (i) SEM surface images of FAPbI₃ perovskite thin films (FTO/SnO₂ as the substrate) with different amounts of MACl and BACl.</figcaption></figure>

To investigate the effect of mixed alkylammonium hydrochlorides on the 3D FAPbI₃ perovskite polycrystalline thin films, we measured the conventional X-ray diffraction (XRD) patterns with varying ratio of MACl and BACl in the second-step cation solution treatment, relative to FAI concentrations of 45-0 (control), 40-5, 35-10, 30-15 (target), 25-20, and 20-25 mol% to explore the difference of the film crystallinity before and after annealing treatment (**figure S1 and S2, Supporting Information**). As depicted in **Figure 1a**, when we increased the ratio of BACl while decreasing MACl amount, it resulted in a notable reduction in the full width at half maximum (FWHM) of the (100) diffraction from 0.14° to 0.09° (**Table S1, Supporting Information**). Meanwhile, the integral intensity of the (100) diffraction of α-phase FAPbI₃ exhibited exponential growth, with the maximum intensity of 10⁶ reached for the thin film containing 15mol% BACl, compared to 10⁴ for the control film (**Table S1, Supporting Information**). These observations underscore the beneficial influence of n-butylammonium (BA⁺) on the enhancement of (100)
136 textured growth. However, an intriguing phenomenon was also observed: when the MACl ratio 137 dropped to 20 mol% relative to FAI, even after high-temperature treatment, the α-FAPbI$_3$ phase 138 spontaneously transformed into the δ-FAPbI$_3$ phase (non-optically active) under ambient conditions, 139 suggesting the necessity of a certain amount of MA inside the lattice for stabilizing the α-FAPbI$_3$ 140 phase at room temperature. Based on these findings, we selected the perovskite thin films with the 141 highest preferential orientation and lowest FWHM values, specifically, film containing 30 mol% 142 MACl and 15 mol% BACl, as the target sample for further investigations. 143

143 In Figure 1b, the XRD pattern clearly reveals the (100) diffraction peak of α-phase FAPbI$_3$ at 14.0°, alongside the degraded PbI$_2$ peak at 12.6° induced by humidity during annealing. Notably, in 144 the target perovskite thin film, besides the significantly increased intensity of the (100) diffraction, 145 a new diffraction peak at 8.9° emerged (Figure 1b insert). This peak presumably corresponds to 146 the diffraction peak of residual BA$^+$ induced quasi-2D phase. To determine the type of the organic 147 cations presented in this quasi-2D phase, we deposited a film doped solely with BACl (at 45 mol% 148 relative to FAI) as a control sample to investigate the potential 2D phase between FA$^+$ and BA$^+$ 149 cations. The XRD pattern in Figure 1b from this single BACl-doped film demonstrates that BACl 150 is unable to enhance the intensity of the (100) diffraction of α-phase FAPbI$_3$ film. Additionally, the 151 δ-FAPbI$_3$ phase was observed at 11.7°, and no correlated 2D peak signal was detected. This result 152 indicates that residual BA$^+$ is likely to form a quasi-2D phase with MA$^+$ rather than FA$^+$, as 153 evidenced by the absence of 2D phase peaks in the XRD pattern of the pure BACl-doped 3D 154 perovskite film. Furthermore, due to the high-temperature treatment of the perovskite layer, the 155 residual hydrochloride content left in the deposited film is nearly negligible (as indicated by GD- 156 OES results in Figure 3d). Accordingly, the diffraction peak at 8.9° can be attributed to the 157 Ruddlesden-Popper quasi-2D phase BA$_2$MAPbI$_2$I$_7$ (n=2).[23, 27] 158

159 To further quantify the residual amount of BA⁺ in the 3D perovskite thin film, we conducted 160 ¹H NMR measurements (Figure S3 and S4, Supporting Information) on samples taken from 10 161 fully annealed perovskite thin films deposited on FTO/SnO₂ substrates. The presence of BA⁺ was 162 confirmed in the fully annealed target perovskite thin film, marked with a blue background in 163 Figure 1c, providing clear evidence of BA's existence in the final film. In addition to BA⁺ as the 164 2D cation, we also detected trace amounts of isopropyl-ammonium (IPAm⁺) cations in both the 165 control and target samples, marked with a pink background which is attributed to the reaction 166 between IPA or FA⁺ with MA⁺, respectively, as discussed in a previous reference.[28] Given the 167 coexistence of IPAm⁺ in both control and target perovskite thin films, the possibility of quasi-2D 168 phase formation solely by IPAm⁺ cation can be dismissed. We identified BA⁺ cation as the sole 169 contributor to the quasi-2D phase formation observed in the target perovskite thin film. By 170 integrating the characteristic hydrogen peaks of the organic cations, we quantified the relative 171 proportions of different cations in the perovskite thin films (Figure 1d). The percentages of cations 172 in the control and target perovskite initial precursor solutions and fully annealed films were carefully 173 calculated and compared (Figure 1d and Table S2, Supporting Information). [29] In the control 174 film, the proportion of FA⁺ cation increased from 69.0% to 82.6%, while the proportion of MA⁺ 175 cation decreased from 31.03% to 16.80% after annealing, indicating the volatilization of MA 176 molecules during the crystallization process. For the target film, in addition to FA⁺ (83.0%) and 177 MA⁺ (14.3%), we found that 2.0% of BA⁺ remained in the perovskite film after annealing (initially 178 10.3% BA⁺), indicating that most of the BA molecules volatilized along with MA molecules during 179 the annealing process, while a portion of cations remained in the perovskite thin film in the form of
180 quasi-2D phases.

181 The grazing-incidence wide-angle X-ray scattering (GIWAXS) technique was employed to
182 analyze the preferential orientation of α-FAPbI₃ crystal growth in the presence of various alkylamine
183 hydrochlorides. As illustrated in **Figure 1e**, the fully annealed control perovskite thin film exhibited
184 Debye–Scherrer diffraction rings at qz=1.0 and 0.9 Å⁻¹, corresponding to the (100) diffractions of
185 α-FAPbI₃ and PbI₂ phases, respectively.[5, 6, 30] The azimuthally uniform intensity of the α-FAPbI₃
186 diffraction ring indicates an isotropic orientation of the polycrystals within the perovskite film. In
187 contrast, the GIWAXS pattern of the target perovskite film (**Figure 1f**) shows an increased intensity
188 at qr=0 Å⁻¹, suggesting a textured perovskite film with (100) planes parallel to the substrate, also
189 referred to as face-up orientation.[31] This indicates that the presence of BA molecules promotes
190 this preferential orientation. Furthermore, by focusing on the lower qz range below 0.7 Å⁻¹ (**Figure
191 1h**), the (040) and (020) diffractions of quasi-2D BA₂MAPbI₂I₇ (n=2) phase were clearly observed
192 at qz=0.64 and 0.32 Å⁻¹, respectively.[27] Similarly to the 3D α-FAPbI₃ phase, the quasi-2D phase
193 is highly textured with the (010) planes parallel to the substrate, also referred to as horizontally
194 oriented. The horizontally oriented quasi-2D phase within the perovskite thin film likely promotes
195 the preferential oriented (100) growth of the 3D α-FAPbI₃ phase. To emphasize the significance of
196 BACl, we also examined a sample with only 30 mol% MACl-doped α-FAPbI₃ perovskite film
197 (**Figure S5, Supporting Information**). Compared to the target film, the MACl-doped film exhibits
198 a reduced degree of (100) texture, which highlights the crucial role of BACl in achieving face-up
199 orientation during the crystallization of 3D perovskite.

200 The surface morphology of the perovskite thin films was characterized using scanning electron
201 microscopy (SEM). In **Figure 1i and S6 (Supporting Information)**, the dark grains correspond to
202 the 3D FAPbI3 perovskite crystals, while the white regions are attributed to degraded PbI2, induced
203 by the combined effects of heat and moisture.[32-34] With the incremental addition of BACl dopant
204 in the precursor solution, the grain size of the perovskite layers increased significantly. As shown in
205 **Figure S7 (Supporting Information)**, we calculated surface grain size statistics, revealing that the
206 control perovskite thin film had an average grain size of approximately 1.20 µm, with the largest
207 grains approaching 2 µm. In contrast, the target film exhibited grains exceeding 3 µm in size, with
208 an average grain size of 1.73 µm. The observed increase in grain size is likely due to the slower
209 growth rate caused by the delayed volatilization of both BA and MA molecules compared to the
210 volatilization of MA alone,[5] as discussed in the subsequent section. In contrast, for the films with
211 different amounts of MACl (**Figure S8, Supporting Information**), there is almost no significant
212 change in grain size. The larger grain size in the target film is advantageous as it reduces grain
213 boundaries and defects at the origin of non-radiative recombination, thereby enhancing the open-
214 circuit voltage (Voc) and fill factor (FF). The presence of surface-degraded PbI2 has been reported
215 to be favorable for facilitating charge transport from the perovskite layer to the hole-selective layer.
216 However, excessive PbI2 between the interface will disrupt the charge transfer balance, [35] leading
217 to a lower FF and increased hysteresis.[15] It could be seen that the introduction of BACl effectively
218 mitigates surface PbI2. This is achieved through two mechanisms: first, the formation of two-
219 dimensional phases consumes some of the PbI2, reducing its surface presence; second, the long-
220 chain two-dimensional molecules on the surface act as a barrier, preventing direct contact between
221 grain boundaries and moisture, thereby inhibiting the formation of iodine-deficient phases.
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## 2.2 Effect of quasi-2D phase on the crystallization kinetics of 3D perovskite

<figure><img src="image_3.png" /></figure>

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**Figure 2.** (a) and (b) GIWAXS patterns of the control and target unannealed perovskite thin films. (c) Normalized PL spectra of the unannealed perovskite thin films with different amounts of BACl dopants. (d) and (e) zoomed-in GIWAXS patterns of figures (a) and (b). (f) Normalized PL spectra of the target perovskite thin film in the region of the quasi-2D phase. (g) Target perovskite thin film morphology under optical microscope (scale bar 100 µm). (h) AFM pattern of the target perovskite thin film (scale bar 20 µm). (i)-(k) PL mapping of the target perovskite thin film with the quasi-2D PL peaks at 566nm, 614nm and 662nm (scale bar 10 µm).

To investigate the crystallization kinetics of the 3D perovskite layer in the presence of a quasi-2D phase, we measured the GIWAXS patterns of unannealed perovskite thin films. As depicted in **Figures 2a and 2b**, the perovskite thin films are well crystallized even without thermal annealing. Importantly, no PbI₂ diffraction ring could be detected in either the control or target perovskite thin films, indicating a complete reaction between the PbI₂ layer and organic cations. Moreover, the unannealed target perovskite thin film exhibited a significantly improved (100) texture compared to the control sample (**Figure S9, supporting information**). In the control perovskite thin film, intermediate phases and the non-optically active δ-FAPbI₃ phase were present, unlike in the target sample, exhibiting only the (100) diffraction of α-FAPbI₃ phase. This suggests that BACl lowers the phase transition barrier, facilitating a complete transformation to the α-FAPbI₃ phase at room temperature.[5, 14, 20] In contrast, PbI₂ and δ-FAPbI₃ phases were detected for the film with only 30mol% MACl or 45 mol% BACl (**Figure S10, supporting information**). Additionally, another diffraction spot was observed at $q_z = 0.32 \AA^{-1}$ and $q_z = 0.24 \AA^{-1}$, corresponding to the (020)
diffraction of the quasi-2D phase of BA₂MAPb₂I₇ (n=2) and BA₂MA₂Pb₃I₁₀ (n=3), respectively (Figure 2e). [27] This observation implies that the target perovskite thin films consist of multiple quasi-2D phases and 3D α-phases before annealing. These findings suggest that the quasi-2D phase forms under unannealed conditions and persists through the annealing process. The existence of quasi-2D phases appears to act as a growth template during the initial crystallization stage of the 3D perovskite, promoting the growth of the (100) preferential orientation in the final polycrystalline film.

To further investigate the evolution of the quasi-2D phases during perovskite crystallization, steady-state photoluminescence (PL) measurements were performed. As shown in Figure 3c, the introduction of BACl into the unannealed perovskite thin film resulted in a shift of the normalized 3D perovskite PL peak from 796 nm to 812 nm. This redshift in emission peak wavelength can be attributed to a decrease in MA⁺ cation concentration within the perovskite lattice, consistent with subsequent UV spectrum and EQE measurements presented in the following sections. Furthermore, when the BACl concentration exceeded 15 mol% relative to FAI, quasi-2D phases emerged in the PL spectra, with peaks at 566 nm and 614 nm corresponding to the n=2 and n=3 quasi-2D phases, respectively. These findings align with the GIWAXS results shown in Figure 2e. To determine the spatial distribution of the quasi-2D phases, we measured the PL spectra of the target film using an excitation wavelength of 403 nm. As shown in Figure 2f, the fully annealed target film exhibited n=2 (566 nm), n=3 (614 nm), and n=4 (662 nm) quasi-2D phases when excited from the top of the layer. [26] However, no quasi-2D phases were detected when the excitation direction was reversed, indicating that the quasi-2D phases are predominantly located near the surface of the perovskite layer, with the 3D phase dominating at the bottom.

To investigate the spatial distribution of the quasi-2D phases on the surface of the perovskite thin film, colocalized PL images (Figures 2i-k) with atomic force microscopy (AFM) (Figures 2h) were taken over a selected area (Figures 2g). As shown in Figures 2i-k, the quasi-2D phase is relatively dispersed, with the largest areas exceeding 1 µm². This suggests that the quasi-2D phase not only covers the grain boundaries but also extends across the grain surfaces, which could enhance the interface stability between the perovskite and the ambient environment. To verify this hypothesis, we examined SEM images in Figure S11a and b (Supporting information) of control and target perovskite thin films incorporated into n-i-p devices after 2000 hours of storage and continuous J-V curve measurements. After this period, surface PbI₂ within the perovskite layer was observed not only at the grain boundaries but also on the grain surfaces, indicating degradation caused by external voltage and humidity. Notably, the target perovskite film exhibited limited PbI₂ precipitation on the grain surfaces compared to the control film, confirming the critical role of the BA relevant phases in stabilizing the perovskite surface grains. The XRD pattern depicted in the Figure S11c and d (Supporting information) indicates the same consequence. This improved stability can be attributed to the quasi-2D phase coverage, which shields the perovskite grains from environmental factors leading to degradation.
<figure><img src="image_4.png" /><figcaption>**Figure 3.** Glow-discharge optical emission spectroscopy (GD-OES) profiles of the control and target perovskite thin films before and after annealing resolved by the elements of (a) Sulfur (b) Nitrogen and (d) Chlorine. (c) Normalized integral peak area of nitrogen of control and target perovskite thin films with different annealing times. (e) Normalized integral peak area of chlorine of control and target perovskite thin films with different annealing times. XRD patterns of the (f) control and (g) target perovskite thin films before and after annealing. XRD peak ratio of (h) perovskite (111) / (100) and (i) PbI₂ (100) / perovskite (100) for control and target films with various annealing times.</figcaption></figure>

To elucidate the role of n-butylammonium (BA) molecules in the crystallization kinetics of 3D FAPbI₃ perovskite, we employed real-time glow-discharge optical emission spectroscopy (GD-OES). Initially, we focused on measuring the sulfur (S) element, originating from the solvent dimethyl-sulfoxide (DMSO), present in both control and target perovskite thin films. **Figure 3a** illustrates the sulfur signal compared to the tin (Sn) signal originating from thin conducting oxide, indicating that residual DMSO is mainly on the top of the perovskite thin film with an intensity below 0.1, which could be due to annealing in the first step and exchange with FAI solution in the second step.[36] After a short 5-second annealing of the films at 150°C, the sulfur signal decreases significantly, almost disappearing due to the rapid evaporation, indicating a complete removal of the solvent from the perovskite thin film at the onset of annealing. However, in **Figure S12** (**Supporting information**), the SEM image at this stage confirms that grain sizes are much smaller compared to fully annealed perovskite films, meaning that polycrystalline film growth is not yet
complete.

To investigate the dynamics of crystal growth, we conducted time-dependent GD-OES measurements focusing on nitrogen (N) and chlorine (Cl) elements in the 3D FAPbI₃ perovskite thin films. Before annealing, both N (Figure 3b) and Cl (Figure 3d) elements exhibited more uniform distribution throughout the perovskite film, contrasting with the top-layer presence of sulfur (S) from DMSO. After annealing for 15 minutes at 150°C, a notable decrease in nitrogen content (Figure 3b) was observed, which can be attributed to the volatilization process of deprotonated alkyl-ammonium species. To quantitatively analyze this volatilization process, we integrated the area under the curve of unannealed perovskite films to establish the initial nitrogen content in the bulk film. The remaining nitrogen content was then determined by comparing the integrated area at each time point to the initial integrated area. During the measurement period, we employed 10-second intervals with a cut-off time of 120 seconds, as evidenced by SEM images showing negligible change in grain size beyond this time point (Figure S13). Figure 3c illustrates that during the initial 10 to 70 seconds of annealing, the nitrogen content in the target film exhibited a slower rate of decline compared to the control film. This slower decline indicates a reduced overall volatilization rate of amines, including FA, MA, and BA, attributed to the higher boiling point of BA molecules.[14, 25] This delayed evaporation rate promotes more homogeneous growth of polycrystals, resulting in larger grain sizes and improved morphology in the target perovskite film (Figure S13). Beyond 70 seconds of annealing, both the target and control films showed a similar reduction in nitrogen content, suggesting comparable deprotonation rates. This observation emphasizes that the incorporation of BA molecules during the initial annealing stages slows down the deprotonation process, facilitating the formation of larger grain morphologies in the target perovskite film. The lowered evaporation rate is beneficial for homogeneous growth of the polycrystals, resulting in a larger grain size and better morphology of the target perovskite film.

Regarding chlorine (Cl) in the perovskite thin films, it is almost absent after full annealing (Figure 3d). This phenomenon arises because chlorine ions have a higher affinity for hydrogen compared to iodine or bromine ions. This affinity facilitates their transformation into hydrogen chloride molecules, which then evaporate from the lattice during annealing. When examining the volatilization kinetics (Figure 3e), we observed that both the control and target samples initially exhibited non-uniform rates of Cl volatilization within the first 90 seconds of annealing. Thereafter, however, their volatilization curves converged, suggesting similar overall Cl volatilization rates between the two films. Notably, between 90 and 100 seconds of annealing, the target perovskite thin film exhibited a slower decrease in Cl content compared to the control sample. This observation may be attributed to a slower deprotonation rate in the target film or the incorporation of Cl- ions into certain phases of the lattice, thereby hindering the protonation of chlorine and slowing down its release from the film, which is still ambiguous.

In addition to studying element variations during annealing, we performed real-time XRD measurements to investigate phase transformations and texture changes as annealing time progressed. A comparative analysis revealed notable shifts in the positions of the (100) diffraction peaks between the control sample (Figure 3f) and the target film (Figure 3g) pre- and post-annealing. Initially, the (100) diffraction peak for the unannealed control sample was located at 13.99°, whereas the target film exhibited a slightly lower diffraction angle of 13.95°. This reduced angle can be attributed to a lower concentration of methylammonium (MA⁺) ions in the target film. Upon full annealing, the diffraction angles shifted to higher values for both the control and target
films. This finding contrasts with previous reports on one-step spin-coated perovskite thin films, where the peak position of the (100) diffraction typically shifts to lower angles due to chloride/iodide (Cl⁻/I⁻) ion exchange.[19] The observed shift towards higher angles in our study may suggest that, in the two-step method, some MA⁺ ions initially fail to integrate into the lattice. Annealing likely facilitates the incorporation of additional MA⁺ ions into the lattice, resulting in the higher diffraction angles observed.

Furthermore, we analyzed the intensity variations of specific diffractions during annealing. In particular, the ratio of (111)/(100) diffraction intensities provides a direct metric for quantifying the changes in (100) texture. As can be seen in **Figure 3h**, the (111)/(100) ratio consistently increases for both control and target samples during prolonged annealing. Notably, the intensity ratio in the target sample was lower than that in the control sample, indicating that the incorporation of BACl has a positive influence on the stabilization of the (100) texture. Moreover, as shown in **Figure 3i**, the intensity ratio of lead iodide (PbI₂) decomposition phase on the surface of the target sample was significantly lower than that of the control sample. Besides, the intensity ratio of PbI₂ phase of the target sample remained almost unchanged throughout the annealing process, corroborating previous scanning electron microscopy (SEM) results. This consistency further validates the surface protective effect conferred by the addition of BACl.

## 2.3 Effect of quasi-2D phase on bandgap, carrier transport and defect density

<figure><img src="image_5.png" /><figcaption>**Figure 4.** (a) Variation of bandgap and Urbach energy calculated from UV-Vis spectrum of FAPbI₃ with different amounts of alkylamine hydrochlorides (MACl mol% + BACl mol% = 45 mol%). (b) Steady-state photoluminescence (PL) spectra and (c) time-resolved photoluminescence (TRPL) curves of control and target films deposited on glass substrates. (d) Space charge limited current (SCLC) measurements of control and target perovskite thin films based on the electron-only devices under dark condition. (e) *J-V* curves of control and target devices under dark condition. (f) The linear fits of light-dependence open-circuit voltage (V<sub>oc</sub>) from control and target devices.</figcaption></figure>
The bandgap of FAPbI$_3$ perovskite thin films, deposited on FTO/SnO$_2$ substrates with various alkylamine hydrochlorides, was derived from UV-Vis absorption spectra (Figure S14, Supporting Information) and is presented in Figure 4a. The addition of BACl and the corresponding reduction of MACl led to a narrowing of the bandgap by 10 meV, attributed to the decreased MA$^+$ content in the lattice. Additionally, the Urbach energy ($E_u$) was calculated from the UV-Vis absorption spectra (Figure S15, Supporting Information) using the following equation: $\alpha = \alpha_0 \exp [(hv-E_g)/E_u]$, where $\alpha$ is the absorption coefficient, $\alpha_0$ is a constant, and $hv$ is the photon energy. The $E_u$ of the control perovskite thin film without BACl was 46.9 meV, while the target film exhibited a $E_u$ of 43.8 meV. Under the identical conditions, a smaller $E_u$ value demonstrates a high crystal quality and minimized voltage deficiency.[21]

Figure 4b shows the steady-state PL emission spectra of the fully annealed control and target perovskite thin films deposited on glass substrates. The PL intensity increased with the addition of BACl, which implies the high quality of perovskite crystallinity and a decrease in nonradiative recombination. Besides, the PL peak position is shifted from 800 nm to 805 nm. The slight red shift could be attributed to the reduced MA ion substitution in the perovskite lattice. Time-resolved photoluminescence (TRPL) was performed to quantitatively compare the charge recombination in the perovskite bulk (figure 4c). The bi-exponential model function was employed to measure the decay dynamics reflecting the nonradiative trap-assisted recombination: $Y = A_1\exp(-\tau/\tau_1) + A_2\exp(-\tau/\tau_2)$, where $\tau_1$ and $\tau_2$ indicate the fast and slow decay times corresponding to the radiative and non-radiative recombination, respectively. The average carrier lifetime was calculated according to the following equation: $\tau_{average} = (A_1\tau_1^2 + A_2\tau_2^2)/(A_1\tau_1 + A_2\tau_2)$. Figure S16 (Supporting Information) shows the PL decay curves and fitted curves obtained from the control and target perovskite thin films. Table S4 (Supporting Information) displays the values of $\tau_1$, $\tau_2$ and $\tau_{ave}$ obtained from the control and target thin films. Due to increased $\tau_1$ and $\tau_2$, the average carrier lifetime of the target sample was 2340 ns, much higher than the value (924 ns) of control sample. The increased decay value suggests the larger number of charge carrier in the bulk of perovskite thin film, ascribing to the suppression of non-radiative recombination by the improved Schottky order in the polycrystalline perovskite thin film.[21]

To further study the charge-carrier dynamics in the bulk perovskite films, the trap-state density was analyzed by measuring the space-charge limited current curves. Figure 4d shows the dark I-V curves of the electron-only devices with the structures of FTO/SnO$_2$/Perovskite/PC$_61$BM/Au. The trap-state density ($N_{trap}$) was calculated from the trap-filled limit voltage ($V_{TFL}$) using the following equation: $N_{trap} = 2\varepsilon_0 \times \varepsilon_r \times V_{TFL} / eL^2$, where $\varepsilon_0$ and $\varepsilon_r$ are the vacuum and relative dielectric constants of FAPbI$_3$ ($\varepsilon_r$ = 47),[5] respectively, e is the elementary charge, and L is the thickness of the perovskite thin film. The thickness of the perovskite thin films was measured using cross-section SEM images of the final devices. An average thickness was calculated from five randomly chosen locations within the SEM images (Figure S17, Supporting Information). Other parameters for the calculation are displayed in Table S4 (Supporting Information). The calculated electron-only trap densities were 3.54×10$^{15}$ and 5.85×10$^{15}$ cm$^{-3}$ for the target and control perovskite thin films-based devices, respectively. A similar tendency was also found for the dark I-V curves (Figure 4e). The introduction of the BACl in the perovskite thin film reduced the dark current from 10$^{-7}$ to 10$^{-8}$, approaching one order of magnitude. The quasi-2D induced 3D perovskite thin films exhibited a lower trap-state density, thereby providing a better light absorber capable of facilitating improved charge separation, extraction and transport.
Moreover, the V<sub>oc</sub> response of the control and target devices was evaluated under different light intensities (Figure 4f). The relationship between V<sub>oc</sub> and light intensity can be modeled by the following equation: Voc=n(k<sub>B</sub>T/e) ln(I) + constant, where n is the ideality coefficient, k<sub>B</sub> is the Boltzmann constant, T is temperature, e is the electric elementary charge and I is the light intensity. It can be seen that the control device shows a fitted slope of 1.37k<sub>B</sub>T/q, while the BACl-treated devices show a smaller slope (1.25k<sub>B</sub>T/q). It is widely acknowledged that deviation of the slope from k<sub>B</sub>T/q is indicative of the defect-induced nonradiative recombination in the devices, which further confirmed that the recombination has commendably been suppressed in the target perovskite layer.

## 2.4 Devices performance and stability measurements

<figure><img src="image_6.png" /><figcaption>**Figure 5.** (a) Solar cell structure. (b) The distribution statistics of relevant solar cell parameters for the control (20 devices) and target devices. (c) Typical J-V curve of control and target devices with reverse and forward scan. (d) EQE, differential EQE curves and (e) MPP tracking of the control and target devices. (f) Stability tests of the unencapsulated control and target devices stored in the N<sub>2</sub> filled glovebox. (b-f) Control is 45 mol% MACl + PEAI treatment. Target is 30 mol% MACl + 15 mol% BACl + PEAI treatment. (g) Champion device based on the FTO/SnO<sub>2</sub>/perovskite (RbCl+30/15 MACl/BACl mol%) /PEAI /Spiro-OMeTAD/Au.</figcaption></figure>

The improved properties of the target perovskite thin film prompted us to investigate its use in
photovoltaic solar cells. Accordingly, we assembled 20 n-i-p perovskite solar cells (PSCs) under identical conditions using different amounts of alkylamine hydrochlorides, employing the configuration of FTO/SnO₂/perovskite/PEAI/Spiro-OMeTAD/Au (Figure 5a). Firstly, we determined the optimal ratio of BACl and MACl (Figure S18 and S19, Supporting Information), on the basis of the undoped perovskite film and sole PEAI passivation layer. The results indicated that incorporating 30 mol% MACl and 15 mol% BACl significantly enhanced the efficiency, achieving an average power conversion efficiency (PCE) of 22.1% (device named target), compared to 19.8% for the control devices. This enhancement is primarily attributed to improved open-circuit voltage (Voc) and fill factor (FF), as can be seen in Figure 5b. To further elucidate the device performance, we investigated the hysteresis phenomenon, characterized by the disparity between the J-V curves obtained from reverse and forward scans, in both control and target devices. The hysteresis index (HI) was calculated using the following equation: HI(%)=(PCEreverse-PCEforward)*100/PCEreverse. The results are tabulated in Table S5 (Supporting Information). This table and Figure 5c reveal that the target device exhibits minimal hysteresis (HI~1.0%), while the control device demonstrates a marked difference (HI~5.5%) between the reverse and forward J-V scans, particularly in terms of FF and Voc. This notable reduction in hysteresis can be ascribed to the optimized preferential crystallographic orientation and the reduced presence of PbI₂ at the interface of perovskite films. Excessive PbI₂ at the interface acts as a p-type semiconductor, leading to unbalanced charge transfer dynamics. The optimized BACl doping ameliorates these issues by minimizing PbI₂ residues, thereby promoting more balanced charge transport and enhancing device stability.

External quantum efficiency (EQE) measurements (Figure 5d) were used to further characterize the photoelectric conversion properties in the spectral domain. The tested devices were based on the best-performing control and target perovskite thin films and PEAI passivation layers. Due to the lower concentration of MACl in the target films, the absorption edge of the target devices redshifted from 800 nm to 810 nm, consistent with UV-Vis and PL spectra results. This shift contributed to a higher short-circuit current (Jsc), with the integrated Jsc of target devices reaching 24.25 mA/cm² (Figure S20, Supporting Information), compared to 23.35 mA/cm² for the control devices (Figure S21, Supporting Information). This enhancement can be attributed to the reduced bandgap and enhanced crystallinity of the perovskite thin film.

Further improvements were achieved by using RbX (X=Cl⁻ or I⁻) doping in the initial PbI₂ precursor layer. The RbCl-doped champion perovskite device achieved an impressive power conversion efficiency (PCE) of 23.23% from reverse J-V scanning (Figure 5g), accompanied by a maximum Voc of 1.181V (Figure S22, Supporting Information), which can be assigned to the higher transport capacity by the formation of PbI₂(RbCl).[15] Similar results were achieved with RbI doping (Figure S23, Supporting Information). However, compared with the undoped devices, they exhibited relatively a larger hysteresis index, as evident from Figures S20 and S22 and Table S6, likely due to the accelerated charge transfer between the perovskite and hole transport layers. To address this problem, we did two treatments. We introduced a n-propylamine hydroiodide (PAI) treatment first and then we did a PEAI treatment using a 5 mg/ml in IPA solutions. It led to balance the charge transfer dynamics, as tiny amount of two-dimensional molecules can passivate interfacial defects and balance the charge transport speed.[34] As shown in Figure S24a, this approach effectively mitigated hysteresis while maintaining an efficiency exceeding 23%. The optimized device consistently exhibited over 22.6% steady-state output efficiency (Figure 5g inset),
494 benefiting from superior crystallinity and stable charge transport interfaces within the device.

495 For the stability assessment, we firstly conducted the stability test of perovskite films by XRD
496 measurements (Figure S25, Supporting Information). The result showed that the crystallinity of
497 control and target (both non-doped and PEAI post-treated) perovskite films did almost not change
498 after storing over 150h under the ambient condition. However, under high temperature condition,
499 obvious yellow-phase δ-FAPbI3 and enhanced PbI2 peaks appeared in the control film. For the target
500 film, only the peak of lead iodide was enhanced but in a much less extent than in the control one,
501 and there was no δ-FAPbI3 phases detected, which indicate that the better initial purity and
502 crystallinity of the perovskite film, along with the residual two-dimensional molecules (BA+) on the
503 surface can alleviate the degradation rate of perovskite thin films.

504 To evaluate the short-term and long-term performance of the perovskite solar cells, rigorous
505 tests were conducted. We employed maximum power point tracking (MPPT) to assess short-term
506 stability (Figure 5e). When comparing the control and target devices, the tracking points of the
507 target devices demonstrated greater stability and consistently higher output values. To determine the
508 long-term stability, we subjected the best-performing control and target non-doped, PEAI treated
509 devices to long-term J-V curve measurements. Figure 5f illustrates the results: over the first 100
510 hours of storage under inert atmosphere, both devices initially showed increased efficiency,
511 attributed to enhanced conductivity of the Spiro-OMeTAD layer.[37] However, after 100 h of
512 storage, there was a relatively significant performance attenuation in the control and target devices.
513 By the After 400 h, the efficiency of the control device had decreased to 80% of its initial value,
514 indicating a severe degradation. In contrast, the target device exhibited remarkable stability,
515 maintaining the PCE at 90% of its initial value over an extended period of 2,000 hours. This
516 exceptional stability is a testament to the effectiveness of the optimized doping strategy and interface
517 engineering in the target device, which mitigates degradation pathways and ensures sustained
518 performance under prolonged operational conditions.
519

520

## 3 Conclusion

521 In summary, we successfully substituted a portion of MACl with an appropriate amount of long-
522 chain n-butylammonium hydrochloride (BACl) to optimize the crystallization dynamics of two-step
523 FAPbI3 polycrystalline films. Increasing the BACl concentration to 15 mol% relative to FAI in the
524 precursor solution facilitated the formation of Ruddlesden-Popper (RP) quasi-2D phases (BA2MAn-
525 1PbnI3n+1, n≥2). These RP phases acted as templates, promoting the growth of 3D FAPbI3 perovskite
526 films with preferential (100) orientation. Incorporation of the larger BACl slowed down the
527 dissociation rate of hydrochloride and alkylamine from the Cl-salts, resulting in perovskite films with
528 improved morphology and reduced trap-mediated recombination. BA+ is shown to regulate and
529 decrease MA+ content in the perovskite lattice. We have revealed by 1H NMR measurements that
530 only 2.0% of the cation proportion on the perovskite layer surface consisted of BA molecules after
531 achieving oriented growth. Ultimately, the optimized perovskite film enabled the fabrication of a
532 high-performance device achieving an efficiency of 23.23% after RbCl doping engineering.
533 Furthermore, the target device exhibited exceptional stability, retaining 90% of its initial PCE after
534 2000 hours of storage in a nitrogen-filled glovebox. This study introduces a novel approach to
535 manipulate the crystallographic orientation of perovskite thin films through low-dimensional
536 perovskite engineering, opening new avenues for advancing perovskite solar cell technology.

520
537

## 538 Experimental Section

### 539 Materials

540 Materials: N, N-dimethylformamide (DMF, 99.8%), dimethyl sulfoxide (DMSO, 99.9%), 541 isopropanol (IPA, 99.5%), 4-tert-butylpyridine (96%), acetonitrile (99.8%), chlorobenzene (CB, 542 99.8%) and lithium bis (trifluoromethanesulfonyl) imide (99.95%) were purchased from Sigma-543 Aldrich. Lead iodide (PbI₂, 99.99%), Methylamine hydrochloride (99%), Butylamine hydrochloride 544 (98%) were purchased from TCI Europe N.V. Formamidine hydroiodate (99.9%) and 545 Phenylmethylammonium iodide were obtained from Greatcell Solar Materials. 2,2',7,7'-tetrakis-546 (N,N-di-p-methoxyphenylamine)-9,9'-spirobifluorene (SpiroOMeTAD, ≥99.8%) and PCBM were 547 purchased from Borun New Materials Technology LTD. All chemicals were used as received 548 without further purification. Fluorine-doped SnO₂ (FTO) coated glass substrates is TEC-7 from 549 Pilkington. 550

550

### 551 Device Fabrication

552 Patterned FTO (TEC-7, Pilkington) glass was cleaned with detergent, DI water, acetone, and ethanol 553 in an ultrasonic bath for 15 mins each. Before use, the FTO glass was cleaned by using a dry 554 compressed air gun and treated with UVO for 30 mins to improve its wettability. Then, the SnO₂ 555 solution (1:2.5, diluted with DI water) was spin-coated on the FTO at 4,000 rpm for 30 s with an 556 acceleration of 2,000 rpm/s and annealed on a hotplate in ambient air at 180 °C for 30 min. Before 557 depositing the perovskite, the SnO₂ coated FTO was treated by UVO for another 30 mins to improve 558 the wettability. 559

560 For the two-step perovskite, 1.5 M PbI₂ in DMF: DMSO = 9:1 was stirred at 70° for 12 h. After the 561 solution cooled down to room temperature, 65 µl PbI₂ solution was spin-coated on SnO₂ at 1,500 562 rpm for 30 s and then annealed at 70 °for 1 min in a N₂ filled glovebox. When the film cooled down 563 to room temperature, 100 µl organic salt solution containing FAI: MACl: BACl (90:9-16:0-11 mg 564 in 1 ml IPA) was spin-coated on the PbI₂ films at 1,800-2,400 rpm for 30 s. When the film color 565 was dark grey, the film was quickly transferred from the glovebox to the hotplate in the atmosphere 566 (30%-40% relative humidity) for annealing at 145°C for 15 min. After annealing, the film was 567 quickly transferred into the glovebox to avoid further degradation due to humidity. When the film 568 cooled down to the RT, 70 µl PEAI solution (3-5 mg/ml in IPA) was dropped on the film and spin-569 coated at 5000rpm for 30s, without post-annealing treatment. For the PAI surface treatment, done 570 before the PEAI one, 70 µl PAI solution (3 mg/ml in IPA) was dropped on the film and spin-coated 571 at 5,000rpm for 30 s and annealed at 100°C for 5 mins. The spiro-OMeTAD solution was prepared 572 by mixing 78.2 mg spiro-OMeTAD in 1 mL chlorobenzene, with 30.4 µL 4-tert-butylpyridine 573 solution, and 17.9 µL LiTFSI solution (517 mg in 1 mL acetonitrile) and 18 µL tris(2-1H-pyrazol-574 1-yl)-4-tert-butylpyridine)-cobalt(III) tris (bis (trifluoromethylsulfonyl) imide) (376 mg in 1 mL 575 acetonitrile) in the spiro-OMeTAD precursor solution, respectively, and was spin-casted on the 576 surface of the perovskite films at 2,000 rpm for 30 s. After that, the devices without electrode were 577 oxidized for 12 h in a dry box with a relative humidity of 5%. Finally, gold/silver electrode was 578 thermally evaporated with a thickness of about 60-100 nm on the top of Spiro-OMeTAD layer. 579 For the SCLC devices, the hole transport layer was replaced by the electron transport layer (PCBM). 580 The PCBM powder was dissolved in the CB at a concentration of 20 mg/ml. The fully dissolved 581 solution was spin-coated on the surface of the perovskite layer at 2,000 rpm for 30s. 582

537
581

582 **Materials Characterization**

583 **Device Measurements**

584 The J-V curves were measured by a Keithley 2410 digital source meter, applying a 20mV/s voltage scan rate (from 0-1.2V). The solar cells were illuminated with a solar simulator (Abet Technology Sun 2000) filtered to mimic AM 1.5G conditions (100 mW/cm²). The illuminated surface was delimited through a black mask with an aperture area of 0.076 cm². The power density was calibrated by using a reference silicon solar cell. The maximum power point tracking (MPPT) measurement was performed under ambient conditions. The current was recorded at the voltage of the maximum power. The external quantum efficiency (EQE) spectra were measured by an Oriel QUANTX-300 system.

592

593 **Materials measurements**

594 The structure of the perovskite layers was characterized with a PANalytical X-Pert high-resolution X-ray diffractometer (XRD) operated at 40 kV and 45 mA, emitting the Cu Kα radiation with λ = 1.5406 Å. The absorbance of the perovskite layers was measured by a Cary 5000 UV-Vis-NIR spectrophotometer. A blank glass/FTO/SnO₂ sample was used for the baseline. The photoluminescence spectra were measured by a Cary Eclipse fluorescence spectrophotometer. The co-localized PL and AFM images were measured using a Witec alpha300 system with an excitation laser at 405 nm. The surface and cross-sectional morphology of perovskite thin layers was measured in a field-emission SEM equipment (Zeiss Supra 40) in the in-lens mode.

602

603 The GIWAXS measurements were performed at the SixS beamline at the Soleil synchrotron (Paris-Saclay, France) radiation source and using a laboratory X-ray scattering set-up. Reciprocal space scanning with a moving detector was used to acquire large *q*-space, while additional GIWAXS measurements were performed in low *q*-space in a vacuum in the laboratory to avoid air scattering. The X-ray energy was set to 9.3 keV and the sample-detector distance was set to 509 mm. The X-ray beam size was 500 µm x 50 µm (H x V), and the angle of incidence was 1°. The laboratory GIWAXS measurements were performed at a custom-designed Nanostar system (Bruker, Germany) in an evacuated environment (1 mbar). The X-ray scattering system was equipped with a metal-jet X-ray source generating Ga Kα radiation (E=9.25 keV). The X-ray detector used was a Pilatus 300K (Dectris, Switzerland). The sample-detector distance was set to 157 mm and the angle of incidence was adjusted to 0.3°. The exposure time was set to 600 s.

614

615 NMR spectra were recorded on a Bruker Neo 500 spectrometer, using DMSO-d₆ as solvent. Chemical shifts were reported in ppm downfield from tetramethylsilane and were referenced to the residual hydrogen signal of deuterated solvent (2.50 ppm) for ¹H NMR. ¹H spectra were acquired with a 90° pulse of 10.0 ms, 64K data points over a 10 kHz spectral width, an acquisition time of 3.28 s, a recycle delay of 10 s and 256 scans. The Fourier transform was applied to the FID without prior exponential line-broadening. For each spectrum, the integrations were normalized to the residual DMSO-d₆ proton signal used as an internal reference (arbitrarily noted 100), allowing comparison of the integrations of the different spectra.

623

624 Glow-Discharge Optical Emission Spectrometry (GD-OES) analyses were performed by a
HORIBA Jobin Yvon GD Profiler 2 equipment. This instrument was equipped with a RF-generator (at 13.56 MHz), a standard HORIBA Jobin Yvon glow discharge source with a cylindrical anode of 4 mm inner diameter and two optical spectrometers (a polychromator and a monochromator) for fast-optical detection. The Ar plasma was generated at an Ar pressure of 420 Pa and an applied power of 17 W. The precursor or perovskite film was mounted on an O-ring at one side of the plasma chamber and adopted as a cathode. The interest of the technique was its rapidity and easiness for use. The technique did not require a high vacuum or specific sample preparation. Therefore, the films could be characterized rapidly (in less than 1 min) after each preparation step [38].

633

### Acknowledgements

Z.L. thanks the CSC for funding his Ph.D. scholarship (grant number 202206220024). This work was supported by APVV-20-0111 (P.S.) and IMPULZ no. IM-2023-82 (N.M.). We acknowledge Dr. Alina Vlad and Dr. Alessandro Coati from the SIXS beamline at Soleil synchrotron, where the GIWAXS was measured. The ANR agency is acknowledged for financial support via the ChemSta project ANR-21-CE05-0022.

640

### Conflict of Interest

The authors declare no conflict of interest.

643

644

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