Performance-limiting metastable intragrain impurity nanoclusters in metal halide perovskites

Domain: physics
Published: 2025-12-18
Editorial outcome: major_revision

Abstract
Impurity control in metal halide perovskite thin films is crucial for advancing the efficiency and long-term stability of perovskite solar cells. While previous efforts have mainly addressed impurities at grain boundaries and hetero-interfaces, efforts have rarely been made to reveal and modulate (sub-) nanoscale impurities within grains. In this work, using low-dose scanning transmission electron microscopy, we observed the nontrivial existence of intragrain impurity nanoclusters, identifying a previously unknown, metastable orthorhombic non-perovskite phase with a parallel chain-like structure, in nominally processed Cs-incorporated mixed-cation perovskite films. First-principles calculations and quasi-insitu observations reveal that these intragrain nanoclusters adversely affect the optoelectronic properties and chemical stability of the perovskite. Guided by these insights, we incorporate inner salts during the solution processing, effectively reducing nanocluster density and enhancing device performance. Clarifying unconventional impurities hidden in perovskite crystal grains opens a research avenue in fundamental understanding and nano-engineering for improved perovskite solar cells and optoelectronics.

Introduction
Metal halide perovskites represent a broad group of semiconductor materials characterized by flexible compositions and diverse crystal structures. In typical three-dimensional perovskites, the chemical formula follows ABX3, where A refers to either cesium or an organic cation, B is commonly Pb, and X denotes a halide anion1. ABX3 perovskites have been proven to be excellent optoelectronic materials with long carrier diffusion lengths and tunable bandgaps2–5. They can also be processed under mild conditions based on solution methods, unleashing their potential to innovate solar cells at a low cost6–15. Perovskite solar cells (PSCs) have now become a major photovoltaic (PV) technology that can compete with the existing silicon PVs. The past decade has witnessed the rapid development of PSCs toward promising power conversion efficiencies (PCEs) and stability16–19.

Impurity control in perovskites has been recognized as one prominent approach because impurities cannot only affect the transport of photo-generated carriers but also usually initiate the degradation of materials20–23. In the literature, most efforts have been devoted to engineering the impurity phases at grain boundaries (GBs) and device heterointerfaces, which are relatively straightforward to characterize, while the grain interiors have usually been assumed to be phase pure in these studies8,23–26. Nevertheless, owing to the soft crystal lattice, perovskite may accommodate various nanoscale heterogeneous impurity phases within the grain interiors27–30. One challenge of studying these intragrain impurities is that they are fully buried in the crystalline perovskite matrix and have no apparent surface features. As a result, they are undetectable using conventional surface characterization methods such as scanning electron microscopy (SEM), atomic-force microscopy (AFM), or X-ray diffraction (XRD)28. To overcome this limitation, low-dose scanning transmission electron microscopy (STEM) has emerged as a powerful tool with extremely high spatial resolution. This technique has recently been applied to directly imaging the atomic-scale structure of perovskite grains in real space while minimizing beam-induced damage27,29,31–35. It is intriguing to leverage this prominent characterization development to unravel the microstructure and performance-limiting effects of intragrain impurity phases in perovskites, which remains an uncharted territory.

Based on low-dose STEM, this work resolves the atomic structure of a previously unknown metastable A2BX4 impurity phase buried within the crystalline grains of nominally processed Cs-incorporated mixed-cation perovskite thin films. It exhibits a 1D chain-like crystal structure, imparting unfavored large bandgaps and metastability, as we reveal using density functional theory (DFT) calculations. The A2BX4 impurity phase exhibits a nanocluster morphology, mostly with dimensions of less than 20 nm. Based on quasi-insitu STEM observation, we also demonstrate that these intragrain impurity nanoclusters (IGINs) can trigger the degradation of perovskite grains, causing detrimental effects on the optoelectronic processes in perovskites. This work is linked to our previous study, which has revealed the intriguing healing behavior of intragrain impurities30. However, the mechanisms behind the formation and healing of intragrain impurities remained unclear at that time. This work now clarifies the atomic-level structure and the role of the A2PbX4 impurity phase, leading to a plausible chemical mechanism for the formation and phase-healing of intragrain impurities. Specifically, we find that the formation of the PbI2 and A2PbX4 impurity pair results from intragrain cation segregation, as described by the dissociation reaction: 2APbX3 → PbI2 + A2PbX4. The phase-healing process is the combination reaction of the impurity pair: PbI2 + A2PbX4 → 2APbX3, leading to an intragrain cation homogenization. For a proof-of-concept demonstration, we develop an IGIN-mitigation method that features the use of 3-(N,N-dimethylpalmitylammonio)-propane-sulfonate (NDPS) inner salt to homogenize solution cations and manipulate intragrain microstructures in the resultant films. Performance and stability improvements of PSCs are shown, featuring an improved PCE of 26.03% in the n-i-p device setting and a T95 operational stability of over 1000 h (T95 means the time for a PSC to decrease to 95% of its initial PCE). This work, in conjunction with our previous work30, provides atomic insights into understudied intragrain microstructural phenomena and unravels their performance-limiting mechanisms, pointing to a pathway in the nanoengineering of thin films for improved PSCs.

Results
Lattice-resolved microstructure of intragrain impurities

Mono-cation (FAPbI3, MAPbI3, CsPbI3), FA-Cs based (FAxCs1-xPbI3) (x=0.5, 0.8, 0.9) films, and triple-cation ((FA0.95Cs0.05)PbI3)0.975(MAPbBr3)0.025 compositions (FA, formamidinium; MA, methylammonium) are selected for low-dose STEM intragrain structure investigations since they are representative compositions for achieving PSCs with high PCEs and stability. The FA0.5Cs0.5PbI3 samples were used for accurate characterization of the intracrystalline microstructure for their relatively good resistance to electron irradiation. We would like to point out that our research results based on FA0.5Cs0.5PbI3 samples are applicable to other Cs-incorporated mixed-cation perovskite films with different Cs contents, as confirmed by our low-dose STEM observations on multiple compositions. We focused on the PSC device cross-sections, which can facilitate linking microscopic observations to the resultant device performance. The focused ion beam (FIB) fabricated PSC device specimens are encapsulated by ultrathin amorphous carbon layers as demonstrated in our previous studies. This mitigates the electron-dose sensitivity issue of perovskite samples during STEM observations29,30. The sample geometry and STEM experimental settings are schematically illustrated in Fig. 1a, b.

Figure 1a and Supplementary Fig. 1 present representative atomic-resolution high-angle dark-field STEM (HAADF-STEM) images obtained from a fresh FA0.5Cs0.5PbI3 PSC device. The beam current of the electron probe used in the present study was reduced to 1 pA to minimize electron-beam damage. All atomic positions in the perovskite grain match the projection along [010] zone axis of a perovskite phase with an orthorhombic structure (space group Pbnm), coinciding with previously reported results29,30,36. A high density of irregularly shaped IGINs, with projected area less than 20 nm2, was observed within the perovskite grains. The enlarged HAADF-STEM images for these IGINs are shown in Supplementary Fig. 2. We then performed the imaging over, in, and under focus conditions, as shown in Fig. 1b-c and Supplementary Fig. 3. The images matched well with the [010] projection of perovskite at over and under focus conditions. However, once the electron beam is in focus, the atomic lattice refers to the non-perovskite impurity phase, confirming that IGINs are buried entirely inside the perovskite grain.

Two types of nano-scale intragrain impurities, PbI2 and non-PbI2, and their intriguing healing behavior were resolved in FA-Cs based films30. However, the mechanism behind the formation and healing of these intragrain impurities was unclear. A careful comparison revealed that the previously reported non-PbI2 phase, which was speculated to have a zero-dimensional crystal structure, has the same projected structure as the IGINs shown in Fig. 1 and Supplementary Fig. 2. Uncovering the atomic structure of IGINs, thus, is the key to revealing the plausible chemical mechanism for their formation and annihilation. However, only atomic-resolution projected images in two or more different directions can accurately reconstruct the three-dimensional atomic structure of a crystal phase. To elucidate the accurate atomic structure of this non-perovskite impurity phase, we obtained the atomic-resolution HAADF-STEM images of IGINs in three projections ([010]o, [001]o and [100]o) (Fig. 1d-f). The observed IGIN-phase fits a C-face centered orthorhombic structure (space group Cmmm) with the lattice parameters a = 0.926 nm, b = 1.309 nm and c = 0.463 nm (Supplementary Fig. 4). The chemical formula of the unit cell is A2BX4, where A is FA+ or Cs+, B is Pb2+, and X- is I-. In an FAxCs1-xPbI3 film, the IGIN phase can be expressed as (FAxCs1-x)2PbI4. The positions of atoms within the unit cell are listed in Supplementary Table 1. We then used image simulations to determine the ratio of FA+ and Cs+, as detailed in Supplementary Fig. 5, showing the simulated atomic images of (FA0.4Cs0.6)2PbI4 match well with experimental HAADF-STEM images (Supplementary Fig. 5). The simulated HAADF-STEM images and electron-diffraction patterns along the [001]IGIN, [010]IGIN and [100]IGIN directions (Fig. 1g-i) also match well with the experimental data (Fig. 1d-f). As further revealed by the reconstructed atomic structure shown in Supplementary Fig. 4, the A2BX4 phase exhibits a one-dimensional crystallographic feature formed by edge-shared BX6-octahedral chains, and the A cations are located between these one-dimensional chains.

We then calculated the XRD pattern of the A2BX4 phase using our proposed impurity structure. As shown in Supplementary Fig. 6, the main diffraction peaks of A2BX4 overlap with those of PbI2, indicating that XRD is unreliable for distinguishing A2BX4 from PbI2. In particular, the nanoscale A2BX4 phase is encapsulated in the crystalline perovskite matrix, which can hardly be detected in perovskite films where PbI2 is usually a native impurity. Furthermore, these A2BX4 IGINs generally exist in FA0.5Cs0.5PbI3, FA0.8Cs0.2PbI3, FA0.9Cs0.1PbI3, and triple-cation ((FA0.95Cs0.05)PbI3)0.975(MAPbBr3)0.025 Cs-incorporated perovskite films (Supplementary Figs. 7-9). No such IGINs was observed in mono-cation (FAPbI3, MAPbI3, CsPbI3) samples as shown in Supplementary Figs. 10-12. Since Cs doping is currently common in high-performance perovskites, it is important to reveal the impact of IGINs on perovskites. Notably, the Cs content in the A2BX4 phase is higher than that in the perovskite grain, demonstrating that the cation segregation may trigger IGIN formation. Therefore, suppressing the cation segregation can potentially avoid the formation of IGINs, which will be discussed later.

Performance-limiting effects of intragrain impurities

We then conducted quasi-insitu and ex situ low-dose STEM observations to assess the impact of IGINs on the stability of the FA-Cs perovskite grains. For quasi-insitu observations, the samples were removed from the electron microscope chamber immediately after observations to minimize electron beam-induced artifacts, and exposed to ambient conditions (25 °C, 50% relative humidity (RH)). The structural evolutions were continuously tracked through low-dose STEM imaging. The samples for ex situ observations were transferred to vacuum packages immediately after fabrication and then investigated after one week of storage. Figure 2a and b illustrate structural changes in the FA0.5Cs0.5PbI3 film upon exposure to ambient conditions. According to the atomic resolution image and the corresponding fast Fourier transformation (FFT) pattern (inset in Fig. 2a), only IGINs exist in the pristine sample, confirmed by the FFT spot of (020) plane of IGINs with 0.656 nm interplanar spacing indicated by yellow circle in the inset in Fig. 2a. After the exposure, PbI2 phases evolve at the original locations of (FA0.4Cs0.6)2PbI4 impurities, as shown in Fig. 2b. The emergence of a new spot with an interplanar spacing of 1.457 nm (corresponding to the (001) plane of 2H-PbI2), indicated by white circle in the corresponding FFT pattern in the inset in Fig. 2b, confirmed the formation of PbI2. Conversely, the GB regions without IGINs showed no notable structural alterations. Moreover, as shown in Supplementary Fig. 13, after the FA0.5Cs0.5PbI3 device was stored in a vacuum package for one week, the co-existence of (FA0.4Cs0.6)2PbI4 and high-density PbI2 can be observed, which is consistent with our previous observations30. The average PbI2 particle size exceeds that of pristine (FA0.4Cs0.6)2PbI4 impurities (Fig. 1 and Supplementary Fig. 1), aligning with quasi-insitu observations. Our quasi-insitu and ex situ results collectively demonstrate that the one-dimensional (FA0.4Cs0.6)2PbI4 formed during the fabrication is meta-stable, prone to decompose into PbI2 under environmental stress. The IGIN-derived PbI2 will further trigger the degradation of the surrounding perovskite phase.

We then performed DFT calculations to study the effect of A2BX4 IGINs on the structural stability and optoelectronic properties of perovskite grains (details in Methods). It was proposed that the valence band maximum (VBM) and the conduction band minimum (CBM) of perovskites and the relevant phases composed of octahedral units are determined by the length of Pb-I and the angle of Pb-I-Pb37. The cations, located at gaps between octahedral units, such as FA+, MA+, and Cs+, have a limited effect on their electronic structure38. For the sake of clarity and to reduce computational costs, we use CsPbI3 and Cs2PbI4 as models to calculate the impact of IGIN formation on the stability and electronic structure of the perovskite system. Structural optimization for the IGIN phase utilizes the experimentally determined orthorhombic structure as the initial configuration, which exhibits a quasi-face-centered cubic (FCC) sublattice of I- ions. This feature is commensurate with that of the perovskite phase, where the I- anions, along with the FA+/Cs+ cations, also form a quasi-FCC sublattice. The lattice parameters of CsPbI3 and Cs2PbI4 obtained from DFT calculations are shown in Supplementary Table 2. We then calculated the formation energies of CsPbI3 and Cs2PbI4 with reference to the stable binary compounds of CsI and PbI2. According to the computed phase diagram, Cs2PbI4 is a metastable phase at the tie line between CsPbI3 and Cs4PbI6, as shown in Supplementary Fig. 14a. The DFT calculations reveal that CsPbI3 and Cs2PbI4 have formation energies of -0.048 and -0.016 eV per formula unit (f.u.), respectively (Fig. 2c). The less negative formation energy of Cs2PbI4 indicates its lower thermodynamic stability and higher tendency to decompose into CsI and PbI2 (Supplementary Fig. 14b), the latter of which has been confirmed through quasi-insitu investigations. We also calculated the electronic density of states (DOS) for CsPbI3 and Cs2PbI4 (Fig. 2d). It is worth noting that Cs2PbI4 exhibits a considerably larger bandgap (2.84 eV) than CsPbI3 (1.83 eV). After aligning the energy bands according to the vacuum level, a type I heterojunction distribution is formed as illustrated in Supplementary Fig. 15. The CBM of the perovskite phase is 0.24 eV lower than that of the IGIN phase, while the VBM is 0.77 eV higher. Therefore, extra energy barriers and charge-blocking regions within perovskite grains can be introduced by the formation of IGINs. The charge carriers generated in the perovskite phase can hardly be transferred to the IGIN phase. Instead, they will bypass the IGIN phase along the perovskite domain during their transfer. This causes the electron transfer pathway to become longer, resulting in more consumption and electron-hole recombination.

In addition, the formation of IGINs also brings about numerous additional intragrain interfaces. Three types of perovskite-IGIN interfaces were resolved at the atomic level, based on which atomic models were built for calculating the corresponding electronic structures (Fig. 2e–s). One half of each atomic model contains CsPbI3, while the other half contains Cs2PbI4. This configuration enables us to study the targeted interface between the two phases directly, without introducing other interfaces. The structural similarity of both phases, sharing the same structure type (quasi-FCC) of specific sublattices, leads to a low mismatch between specific lattice planes of CsPbI3 and Cs2PbI4, resulting in the semi-coherent interfaces as revealed by STEM imaging (Fig. 2e, j, o). As shown in Fig. 2e–g, the interface (denoted as interface I) between the (101) lattice plane of CsPbI3 and the (110) plane of Cs2PbI4 could generate multiple defect states near the CBM of CsPbI3. These defect states are primarily localized on specific PbI6 polyhedra at the interface, mainly contributed by the p orbitals of Pb atoms. In comparison to the PbI2-perovskite interface, which was previously demonstrated to be benign when there are no point defects at the interface30, the localized trap states exhibited at a pristine perovskite-IGIN interface suggest that the IGIN phase is intrinsically detrimental to the device PCEs. Even without other defects, the IGINs could potentially increase the recombination of the photo-excited charge carriers. On the contrary, the VBM of the interfacial region is contributed by the perovskite domain (Supplementary Fig. 16), exhibiting a delocalized character similar to that in the bulk phase.

The interfaces can serve as sinks for migrating point defects. Therefore, the effect of point defects on the optoelectronic properties of a device with IGINs has also been examined, as displayed in Fig. 2h and i for interface I. Given the relatively low formation energy of the iodine vacancies (VI) both in the bulk perovskite phase and at the perovskite-IGIN interfaces (Supplementary Fig. 17), it is expected that VI defects will accumulate over time and interact with the perovskite-IGIN interfaces during device operation. The calculation result shows that VI could generate localized trapped states distributed at the interface, increasing carrier recombination rate, similar to the case described in our previous study30. In addition, such analyses can be extended to other interface configurations, denoted as interface II and III as shown in Fig. 2j and o. Notably, defect states can emerge for each configuration, regardless of whether a point defect exists or not. The defect states are located near the CBM and primarily localized on the Pb atoms at the interface (Fig. 2k-n and p-s), which could significantly enhance recombination between electrons and holes as localized trap states. On the other hand, no empty states emerge near the VBM, and the highest occupied state is delocalized (Supplementary Figs. 18 and 19), meaning that the interfacial structures will only contribute to defect states near the CBM.

Synthetic control of intragrain impurity nanoclusters

Our previous study has revealed the general existence of nanoscale Cs+ cation segregation and the resultant intragrain uneven strain distribution in FA-Cs perovskites through geometrical phase analysis (GPA)29. Here, as IGINs may form concurrently with cation segregation, we utilized GPA to identify the mismatch of lattice between perovskites and IGINs, as well as the residual strain in perovskite grains induced by cation segregation (Fig. 3a-b and Supplementary Fig. 20). In perovskite grains containing a large density of IGINs, pronounced strain concentration appears not only around the IGINs, but also in adjacent perovskite regions. However, when it comes to regions without IGINs, the strain concentration becomes subtle, suggesting the critical role of Cs cation segregation in promoting the formation of such impurities.

To eliminate IGINs, we introduced an amphiphilic surfactant additive of NDPS inner salt into the precursor solution. Inner salts are known to tailor the perovskite crystallization and passivate the defects39–41. DFT calculations indicate that NDPS features an electron-deficient quaternary ammonium group and a negatively charged sulfonic group42,43 (Supplementary Fig. 21), making it well-suited for interacting with both cations and anions. Functioning as a surfactant additive, NDPS molecules process both hydrophilic and hydrophobic moieties. These molecules can spontaneously organize on the surface of the precursor solution, orienting their hydrophobic ends toward the air. This behavior lowers the surface tension of solution and reduces the critical radius of the nucleus, that is, the number of solute particles required to form a stable nucleus39–41. The decrease in surface tension is corroborated by a measurable reduction in the contact angle of perovskite solution on the substrate upon surfactant addition, as shown in Supplementary Fig. 22.

Top-view SEM imaging showed that incorporating NDPS into FA0.9Cs0.1PbI3 precursor solutions does not significantly affect the perovskite grain size compared to additive-free controls (Supplementary Figs. 23 and 24). XRD analysis indicated that the NDPS with a large molecular size does not incorporate into the perovskite lattice or form large-scale impurities in films, as shown in Supplementary Fig. 25. In addition, XRD results also indicated that the peaks of impurities gradually decreased with an increase in the concentration of NDPS, suggesting strong coordination between Pb2+, I- and NDPS molecules. We then performed X-ray photoelectron spectroscopy (XPS) analysis on perovskite films with and without NDPS. As shown in Supplementary Fig. 26, the emergence of S 2p peak and a new peak around 402.5 eV in N 1s spectra obtained from films with additives clearly show the successful incorporation of NDPS into perovskite films. Furthermore, the Pb 4f and I 3d peaks of films with 0.5 mg mL-1 and 1 mg mL-1 NDPS additives shifted to lower binding energy as compared with the control film, indicating the coordination between Pb2+, I-, and NDPS, as shown in Fig. 3c and d.

To evaluate the residual strain in FA0.9Cs0.1PbI3 films with and without NDPS, we performed grazing incidence x-ray diffraction (GIXRD) analyses using the 2θ-sin2(ψ) method44,45. The (012) reflection at about 31.5°, based on the cubic symmetry of FA0.9Cs0.1PbI3 (Pm3m, a=0.6362 nm), was selected for strain analysis due to its multiplication and more significant diffraction intensity, which can provide the most reliable structural symmetry information. The setup details for the X-ray diffractometer are listed in Supplementary Table 3. As shown in Fig. 3e, f and Supplementary Fig. 27, when increasing ψ from 0° to 50°, the (012) peak gradually shifts to lower angles in both the control and the 0.5 mg mL-1 NDPS perovskite films, as indicated by the dot lines, confirming the presence of residual tensile strain. However, a smaller angular shift in the 0.5 mg mL-1 NDPS sample compared with the control film suggests partial strain alleviation. At NDPS concentrations above 1 mg mL-1, the peak shift vanished entirely, demonstrating the complete release of residual tensile strain (Fig. 3f).

To confirm the inhibitory effect of additives on IGINs, we further conducted low-dose atomic-resolution HAADF-STEM imaging and GPA analysis of FA0.5Cs0.5PbI3, FA0.8Cs0.2PbI3, FA0.9Cs0.1PbI3, and (FA0.95Cs0.05)0.975(MAPbBr3)0.025 with 1 mg mL-1 NDPS (target) samples, as shown in Supplementary Fig. 28. No IGINs were observed in the perovskite grains, consistent with the released strain concentration due to the homogenization of cations. To reveal the distribution of NDPS in the perovskite film, we further performed electron energy loss spectroscopy (EELS) measurements by exploiting the signal provided by sulfur in angular dark-field (ADF) STEM mode, which unequivocally identifies the additive from the perovskite (Fig. 3g and Supplementary Fig. 29, details in Methods). Low-magnification ADF-STEM imaging of the target films revealed no detectable impurities. Elemental mapping of S confirmed that the NDPS was concentrated at the GBs and the upper surface of the perovskite layer, consistent with XRD analysis. This distribution suggests that, due to its relatively large molecular size, NDPS is spontaneously excluded from the growing perovskite lattice and accumulates at surfaces and GBs during film formation (Fig. 3g). The long alkyl chain of NDPS enables it to self-assemble on both the perovskite grains and solution surfaces, acting as a reservoir that promotes solute ion intercalation and functions as a molecular sieve to moderate the evaporation of solvents. Furthermore, the highly polar sulfonyl (S=O) group facilitates the formation of strong hydrogen bonding networks among polar solvent molecules, thereby decelerating the crystallization process39–41. The polar hydrophilic head groups anchor to charged ions and crystal surfaces, while the hydrophobic alkyl tails aggregate around the perovskite grains, enhancing structural organization during film formation. The strong coordination between Pb2+, I-, and NDPS around GBs can passivate GBs, improving the stability of perovskite films. More importantly, as demonstrated by our quasi-insitu characterization, the grain interiors with IGINs begin to degrade prior to GBs (Fig. 2a, b), implying the importance of eliminating IGINs rather than passivating GBs.

Our results clearly demonstrate that NDPS can uniformly distribute the cations in resulting perovskite films, removing the A2BX4 phase in the final films. It is reported that the O atom in tetrabutylammonium anion ([TFSI]) can act as a monodentate and bidentate ligand to form dative bonds with alkali metal ions, though chelation effect46,47. This prompted us to compare the ability of NDPS and [TFSI] to interact with Cs+ through DFT calculations to reveal the intrinsic mechanism of NDPS in inhibiting cationic segregation. Solvent effects of DMF were incorporated using the SMD polarizable continuum model. The calculated results show that the cations (Cs⁺) form bidentate or polydentate complexes by coordinating with the oxygen atoms in both the DNSP and the [TFSI]. As indicated in the Gibbs free energy presented in Supplementary Table 4, this coordination is an exothermic process, which is thermodynamically favorable and contributes significantly to the overall energy stabilization of the system. Furthermore, our analysis reveals that the DNSP molecule exhibits a stronger coordination ability compared to [TFSI], aligning with our microstructural analysis. These results clearly indicate that the chelation effect occurs between DNSP and Cs+ in the precursor solution and can effectively disperse Cs+. In the subsequent spin coating process, NDPS, as a surfactant, can reduce the size of the crystal nucleus and control grain growth, there by further benefiting cation homogeneity. We then proposed the possible mechanisms for the NDPS-induced IGINs chemo-elimination, as schematically illustrated in Fig. 3h. The amphiphilic nature of NDPS, featuring both negative and positive active groups (Supplementary Fig. 21), enables interactions with solute species, including the cationic clusters in the precursor solution. These interactions enhance miscibility and dispersion by forming kinetically stabilized cationic clusters with reduced sizes. The NDPS coordination also slows the overall crystallization, allowing sufficient time for cation intermixing. This prevents the formation of cation-segregated phases such as A2BX4.

Effect of intragrain impurities on photovoltaic properties

We further performed time-resolved photoluminescence (TRPL) spectroscopy on the control and target films with the sample structure of glass/ITO/ETL/perovskite, with results shown in Fig. 4a. Based on the fitted parameters (Supplementary Table 5), the average PL lifetime decreases from 61.81 to 9.79 ns, suggesting effective suppression of non-radiative recombination due to the elimination of IGINs. To further investigate the carrier dynamics, transient absorption (TA) spectroscopy was conducted on both films (Fig. 4b, c). Following photoexcitation, both bandgap renormalization and hot-carrier cooling were observed to occur within the initial 2 ps for all sample48. Notably, the target film exhibited accelerated delay kinetics with a narrower full-width at half-maximum (FWHM) in the absorption profile at 0.84 ps, measuring 35.3 nm for the control and 32.8 nm for the target samples. This spectral narrowing indicates reduced energy level broadening in the target film, which can be attributed to the elimination of IGINs.

We further evaluated the ambient stability of both films under controlled environmental conditions (22 °C, 40% RH). After 40 days of storage, the absorbance of the target film remained nearly unchanged, indicating significantly improved environmental stability compared to the control film (Supplementary Figs. 30 and 31). The current density-voltage (J-V) characteristics of the best-performing PSCs are presented in Fig. 4d. The target device, free of IGINs, achieved an enhanced PCE of 26.03%, with a short-circuit current (Jsc) of 26.29 mA cm-2, an open-circuit voltage (Voc) of 1.206 V, and a fill factor (FF) of 0.821. In contrast, the control device exhibited a PEC of 24.50%, with a Jsc of 25.95 mA cm-2, a Voc of 1.165 V, and an FF of 0.810. A direct comparison of these parameters is provided in Supplementary Table 6. The notable enhancement in PCE is primarily attributed to the increased Voc. To further validate the improvement in device performance, we measured steady-state photocurrent outputs at the maximum power point (MPP) for both devices, also observing a faster transition from dark to illuminated conditions in the target PSC (Fig. 4e). Additionally, the target devices demonstrated better reproducibility, as supported by a statistical analysis of PCE across 30 devices (Supplementary Fig. 32). The PSCs without IGINs also demonstrated enhanced operational stability. As shown in Fig. 4f, the target device demonstrated an impressive 95.21% PCE retention stability upon 1000 h of continuous operation under one-sun-intensity illumination (ISOS-L-1l protocol). In contrast, the PCE of the control device degraded rapidly in the initial 100 h. We then compared the storage stability of the two cases. The cross-sectional STEM images in Supplementary Fig. 33 show that PbI2 can be observed in perovskite intragrain regions and grain boundaries in the pristine one after being stored in a glove box for two weeks. For comparison, no degradation products can be observed in the target one. Meanwhile, atomic-resolution imaging revealed that residual strain-free perovskite maintained the integrity of the crystal structure. Still, in the control sample, the perovskite partially degraded to PbI2 nanoparticles, confirming that the device without IGINs showed better stability and durability.

Methods
Raw chemicals
Lead bromide (PbBr2, ≥99%), 4-tert-butylpyridine (TBP, 96%), bis(trifluoromethane)sulfonimide lithium salt (99.95%), dimethylformamide (DMF, 99.8%), dimethyl sulfoxide (DMSO, 99.9%), chlorobenzene (CB, 99.8%), toluene (TB, 99.8%), acetonitrile (CH3CN, >99.9%), diethyl ether (DE, 99%) and ethyl acetate (EA, 99.8%) were supplied by Merck (USA). Lead iodide (PbI2, >99%) and 3-(N,N-dimethylpalmitylammonio)-propane-sulfonate inner salt (>98%) were obtained from TCI (Japan). Formamidinium bromide (FABr, >99.99%), formamidinium iodide (FAI, >99.99%), methylammonium chloride (MACl, >99.99%), methylammonium bromide (MABr, >99.99%), and methylammonium iodide (MAI, >99.99%) were sourced from Greatcell Solar (Australia). Tin(IV) oxide colloidal dispersion (15 wt.% in H2O), poly[bis(4-phenyl)(2,4,6-trimethylphenyl)amine] (PTAA, Mn 15,000-20,000), tetrakis(pentafluorophenyl)borate (TPFB, 99%), and cesium iodide (CsI, 99.998%) were provided by Alfa Aesar (USA). Spiro-OMeTAD (99.8%) was purchased from Borun Chemical Co., Ltd. (China). All chemicals were employed directly without additional purification steps.

Preparation of solutions and thin film synthesis
SnO2 nanoparticles were diluted with distilled water (v/v 1:5) to form the ETL solution. The suspension was spin-coated onto pre-patterned ITO substrates at 3000 rpm for 30 s, followed by annealing at 150 °C for 10 min.

To prepare the HTL, 72.3 mg of Spiro-OMeTAD, 28.8 μL of 4-tert-butylpyridine, and 17.5 μL of lithium bis(trifluoromethane)sulfonimide in acetonitrile solution (520 mg mL-1) were dissolved into 1 mL CB. A 50 μL portion of this solution was spin-coated onto the perovskite surface at 3000 rpm for 30 s. For stability assessments, PTAA served as the HTL. The PTAA solution was prepared by dissolving 30 mg PTAA and 3 mg TPFB in 1 mL CB, and 50 μL was spin-coated at 2000 rpm for 30 s. Subsequently, an 80 nm Au was deposited by thermal evaporation.

FA0.9Cs0.1PbI3, FA0.8Cs0.2PbI3, FA0.5Cs0.5PbI3, FAPbI3, MAPbI3, and CsPbI3 perovskite films were used for the fundamental investigation of intragrain impurities. For FA0.9Cs0.1PbI3, a 1 M precursor solution was prepared by dissolving 154.8 mg FAI, 26.0 mg CsI, 461.0 mg PbI2, and 20.3 mg MACl in DMF/DMSO (v/v, 7:3). Similarly, the FA0.8Cs0.2PbI3 precursor (1 M) was obtained by dissolving 137.6 mg FAI, 52.0 mg CsI, 461.0 mg PbI2, 20.3 mg MACl in the same solvent mixture. Film deposition for both compositions followed an identical three-step spin-coating process. 40 μL of precursor was applied to SnO2-coated ITO substrates at 500 rpm for 5 s, 1000 rpm for 10 s, and 6000 rpm for 30 s. During the final stage, 400 μL of EA was rapidly drop-cast at 10 s of the second step. The resulting films were annealed at 150 °C for 20 min. FA0.5Cs0.5PbI3 thin films were prepared from a 1 M precursor solution obtained by dissolving 86.0 mg FAI, 129.9 mg CsI, 461.0 mg PbI2, and 40.3 mg MACl in 1 mL DMSO. Film deposition was carried out by spin-coating at 1000 rpm for 10 s, followed by 5000 rpm for 60 s. During the second step, two aliquots of 380 μL EA were rapidly drop-cast at 20 and 40 s. The resulting films were annealed at 100 °C for 2 min inside a glove box, and subsequently at 140 °C for 60 min under ambient conditions (30-40% RH). For FAPbI3 perovskite thin films, a perovskite precursor solution with 1 M concentration was obtained by dissolving 633.0 mg FAPbI3 single crystal in DMF/DMSO (v/v, 9:1). Film deposition was carried out by a spin coating process of 1000 rpm for 5 s, followed by 5000 rpm for 120 s. After 10 s of the high-speed step, 150 μL CB was rapidly dripped onto the center of the substrate. The films were then annealed at 150 °C for 20 min. MAPbI3 thin films were prepared from a 1 M precursor solution containing 159.05 mg MAI and 461.0 mg PbI2 dissolved in DMF/DMSO (v/v, 7:3). The MAPbI3 perovskite films were fabricated at 500 rpm for 3 s, followed by a 4000 rpm 30 s process. During the second step, 200 μL of CB was drop-casted quickly at 20 s, followed by 150 °C 20 min annealing. For preparing CsPbI3 perovskite thin films, a 1 M perovskite precursor solution was prepared by dissolving 259.81 mg CsI and 461.0 mg PbI2 in DMF/DMSO (v/v, 4:6). The CsPbI3 perovskite films were fabricated at 1000 rpm for 10 s, followed by a 3000 rpm 30 s step. The films were annealed at 250 °C for 10 min after a resting period of 20 min.

For the device demonstration, ((FA0.95Cs0.05)PbI3)0.975(MAPbBr3)0.025 perovskite films were employed. The precursor solution (1.44 M) was prepared by dissolving 705.3 mg PbI2, 228.8 mg FAI, 18.2 mg CsI, 33.67 mg MACl, and 18.2 mg MAPbBr3 single crystal in the mixture of 890 µL DMF and 110 µL DMSO. The MAPbBr3 single crystal was synthesized separately by reacting MABr and PbBr2 in a 1:1 molar ratio at 0.2 M concentration in DMF, followed by crystallization using dichloromethane as anti-solvent. Films deposition was carried out by spin-coating 50 μL of precursor at 1000 rpm for 10 s and then at 4000 rpm for 30 s. During the final stage, 1 mL of DE was quickly dropped onto the substrate center within 1 s in the last 20 s of the second step. The resulting films were annealed at 100 °C for 40 min. Additives, when required, were incorporated directly into the precursor solution at the corresponding concentrations.

Materials characterization
XRD and GIXRD measurements were performed using a high-resolution diffraction (D8 Advance, Bruker) equipped with Cu Kα radiation (λ=1.5406 Å), employing a step size of 0.01° and a scan rate of 10° min-1. Top-view SEM images were acquired with a Helios 5CX (Thermo Fisher) operated at 5 kV. XPS analysis was conducted using an XR6 (Thermo Fisher) with monochromated Al Kα radiation and a pass energy of 20 eV. The TRPL measurements were carried out on a spectrograph (Ando Kymera 328i) with excitation from a 375 nm picosecond laser (LDH-DC-375, PicoQuant). The emitted signal was collected by a signal-photon avalanche photodiode. The TRPL spectra were analyzed by fitted to a bi-exponential delay function of f(t)=A1exp(-t/τ1)+A2exp(-t/τ2), where A1 and A2 represent the decay amplitudes, and τ1 and τ2 correspond to the decay time constants. TA measurements were performed using infrared laser pulses generated by an amplifier (Pharos, Light Conversion) with a pulse duration of 170 fs, a central wavelength of 1030 nm, and a repetition rate of 1 kHz. The pump beam was produced by directing the output into an optical parametric amplifier (Orpheus-F, Light Conversion), yielding a 640 nm pulse. Pump intensity was controlled with a neutral density filter wheel and a neutral density filter, resulting in a sample fluence of 2.31 μJ cm−2. A probe beam with 0.1 mJ of pulse energy was attenuated and focused onto an yttrium aluminum garnet (YAG) crystal to generate the supercontinuum probe spectrum. The probe was then directed onto the sample, spatially overlapping with the pump spot. Temporal delay between the pump and probe pulses was adjusted using a mechanical translation stage, and the transmitted signal was recorded with a USB spectrometer. UV-vis spectra were recorded using a commercial UV-visible spectrophotometer (Cary 300, Agilent, USA).

Solar cell fabrication and testing
ITO substrates were cleaned sequentially by ultrasonic cleaning in cleaning fluid, deionized water, acetone, and isopropanol. A SnO2 layer was then deposited by spin-coating at 3000 rpm for 30 s, followed by annealing at 150 °C for 30 min in ambient air. The resulting ETL-coated substrates was subsequently transferred into a nitrogen-filled glovebox (H2O < 0.01 ppm, O2 < 0.01 ppm). The perovskite layer was deposited as described previously, after which the HTL layer was spin-coated at 4000 rpm for 30 s. Finally, an 80-nm Au electrode was thermally evaporated onto the HTL. J-V characteristics of the PSCs were recorded using a source meter (2162, Keithley) under simulated AM1.5G spectrum (one-sun illumination; 100 mW cm-2) provided by an AAA-class solar simulator (Sirius-ss, Zolix) inside a nitrogen-filled glovebox. Measurements were conducted in reverse scan mode with a delay time of 10 ms. The device active area was defined as 0.069 cm2 using a shadow mask. Light intensity calibration was performed with an Oriel reference solar cell accredited by the National Institute of Standards and Technology to the ISO-17025 standard. Since the spiro-OMeTAD hole-transporting layer depends on molecular packing that can deteriorate under extended operation, it may compromise device durability49,50. To evaluate stability, PSCs were instead fabricated using PTAA for device stability tests. For light-soaking tests, unencapsulated devices were placed in a sealed quartz holder inside a N2-filled glovebox and exposed to a one-sun-intensity illumination. For the MPP tracking, PSCs were positioned in test chambers with a continuous flow of N2 gas to maintain a temperature of 50 °C. A bias at MPPs was applied on the devices and data points were collected at regular intervals following the ISOS-L-1l protocol.

Low-dose STEM characterization
Cross-sectional PSC specimens were prepared using a dual-beam focused ion beam (FIB) platform (Helios 5CX, Thermo Fisher). Sequential Pt protection layers of 30 nm (electron deposition) and 300 nm (ion deposition) were applied to the top electrode of the device. Lamella preparation was carried out with a gallium ion beam operated at 30 kV and 0.1-24 nA. To reduce ion-beam-induced damage, the surface amorphous layer was removed using a low-energy cleaning step at 1 kV and 39 pA. The prepared lamellas were subsequently transferred to a high-vacuum sputter coater (Leica ACE 600) for deposition of a protective carbon layer. Low-dose STEM was performed on an aberration-corrected microscope (Spectra 300, Thermo Fisher) equipped with a field emission gun and operated at 300 kV. The probe convergence half-angle was set to 29.9 mrad, and the HAADF-STEM detector collected electrons within a half-angle range of 57 to 200 mrad. To minimize beam-induced damage, the electron probe current was reduced to 1 pA, and the dwell time per pixel in atomic-resolution imaging was limited to 3 μs. For quasi-insitu investigations, the pixel dwell time was reduced to 2 µs. Atomic-resolution images were acquired with an average frame size of 34.5×34.5 nm2 at 2048×2048 pixels. Each image acquired in quasi-in situ experiments were recorded with a larger field of view of 100.8×100.8 nm2 at same pixel resolution. Under these conditions, the average dose rate was estimated to be 53 e Å-2 s-1 for atomic-resolution imaging and further reduced to 6 e Å-2 s-1 for quasi-insitu experiments. The EELS mapping was conducted in ADF-STEM mode using a Gatan Image Filter K3 camera, with a dispersion of 0.18 eV channel-1, a 5 mm entrance aperture, and a beam current of 10 pA.

Calculation methods
Density functional theory (DFT) calculations were conducted using the Vienna Ab initio Simulation Package (VASP), with the projector-augmented wave approach employed to describe the electron-ion interactions51. Generalized gradient approximation of Perdew-Burke-Ernzerhof (PBE)52 was used and a kinetic energy cutoff of 450 eV was set, with a Gaussian smearing width of 0.05 eV. The lattice parameters and atomic positions were fully relaxed until the interatomic forces were converged to a threshold of 0.02 eV Å−1. For the calculations of formation energy and density of states, Γ-centered 5×5×4 and 4×3×8 k-meshes were used for CsPbI3 and Cs2PbI4, respectively. A slab model was employed for band alignment calculations of CsPbI3 and Cs2PbI4, in which a vacuum space of 15 Å was placed above the slab. The formation energies (ΔEf) were calculated according to the following equations:

ΔEf(CsPbI3) = Efu(CsPbI3) - Efu(CsI) - Efu(PbI2) (1)

ΔEf(Cs2PbI4) = Efu(Cs2PbI4) - 2Efu(CsI) - Efu(PbI2) (2)

where Efu represents the total energy per formula unit of the substance in the brackets.

The formation energy of an I vacancy was calculated from the following expression:

ΔEf(VI) = E(VI) - E(pristine) + E(I) (3)

where E(VI) and E(pristine) are the total energies of the configurations with and without an I vacancy, and E(I) is the total energy per atom of bulk iodine in the ground state.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The Authors focus on a class of nanoscale features of perovskite films for solar cells that are not often studied - they refer to them as intra-grain impurity nanoclusters and, unlike the commonly studied features at the grain boundaries, exist within individual perovskite grains, generally with some level of epitaxial matching. They are generally harmful to solar cell performance, and their direct study is very challenging because their lattice features are very similar to the surrounding material and they evolve under the electron beam more than the perfect halide perovskite lattice.\nThe STEM experiments described are very challenging, as well as the data interpretation. The Authors have demonstrated in the past the ability to conduct careful experiments on these materials. The study of nano-domains in halide perovskites is a very interesting and active direction - for example there is a very recent Nature Nanotech paper (https://www.nature.com/articles/s41565-025-01917-0) which looks into impurity domains with different techniques and without the same direct imaging resolution as this manuscript.\nSpecifically, this manuscript covers a set of FA/Cs perovskites and proposes a mitigation strategy for these impurities via the use of NDPS, which results in better PCE and stability.\nThis work combines very good EM investigation with complex sample preparation and an extensive set of experiments and simulations.\nSomething that should be pointed out more clearly in the manuscript (the paper is referenced, but not elaborated upon much) is that this builds on a previous nature comms paper (ref 30 - Cai, S. et al. Intragrain impurity annihilation for highly efficient and stable perovskite solar 638 cells. Nat. Commun. 15, 2329 (2024)). In the previous paper a subset of the Authors had identified intragrain impurities in these systems and concluded that they were in some cases identifiable as PbI2 and in others as a 0D phase. These impurities had been observed as a function of electron dose, the process resulting in annealing/healing, and a laser treatment was suggested to improve performance.\nI think it is important that the Authors here elaborate (if they desire, in a rebuttal letter, but also this should go in the main manuscript) on how these findings relate to previous information. I fully understand that the identification process for these IGINs is complex and current findings might overrule previous hypotheses (the previous paper is not very conclusive on the nature of the second class of IGINs), but for this there should be a frank discussion, perhaps including motivations on why the current measurements are more representative. This should apply to both the identification of the phases and the evolution of the impurities under the electron beam.\nThe proposed mechanism of cation segregation for formation of IGINs is rather interesting and a good approach to rationally designing an additive for improving overall performance.\nSome minor comments:\nLine 190 - the definition of in situ and ex situ in this context is unclear. I assume by in situ the Authors mean evolution under the electron beam, and ex situ exposure to air and moisture? This should be made clearer in the text.\nIn page four, in reference to the emergence of the PbI2 phase, the FTT pattern before and after degradation is brought as evidence; the concept is clear in the image but much less so in the text. It would be necessary to better explain the reasoning used (so as to understand the concept without images) and quantify the lengths to verify that the phase that emerges after degradation is PbI2.\nIn the GIXRD experiments section, the reflection (012) at 31.5° is attributed using a cubic symmetry. A reference of the .cif file used should be added, or at least the size of one side of the cubic cell should be defined, so that the accuracy of the deductions can be verified.\nFig. 1A - the GPA colourscale should have units - I assume it is % strain. Same for Supp Fig 15, 22, 27\nSupp Fig 9 - \"squires\" should obviously be \"squares\"\nLine 320 - Regarding the use of sulfur as \"tracer element\" - normally tracers are something that is added for the purpose of tracking. Here my understanding is that the compound already contains sulfur, and it is a convenient atomic species to map because it gives a good signal (EELS or EDS), unlike lighter elements. I think this should be clarified in the text. For example this could be changed to \"[...] (EELS) measurements by exploiting the signal provided by sulfur, which unequivocally identifies the additive from the perovskite.\"", "reviewer_2": "This manuscript by Wang et al. reports on the investigation of intragrain impurity nanoclusters (IGINs) in metal halide perovskites, specifically FA-Cs perovskite films. The authors employ low-dose scanning transmission electron microscopy (STEM) and first-principles calculations to identify and characterize these IGINs, which are revealed to be metastable orthorhombic non-perovskite A₂BX₄ phases. These impurity phases are shown to negatively impact the optoelectronic properties and chemical stability of the perovskite films. To mitigate this issue, the authors propose a strategy using a salt additive, NDPS, which results in notable improvements in both power conversion efficiency (PCE) and operational stability of perovskite solar cells (PSCs). The STEM-based findings are indeed interesting and contribute to the fundamental understanding and nano-engineering of perovskite materials. However, I find serious inconsistencies between the STEM analyses, DFT calculations, and device-level demonstrations, which hinder the overall coherence of the study.\n1. Overall, the manuscript gives the impression of merging three separate studies into a single story, making it difficult to derive firm conclusions. For instance, the STEM investigations focus on FA₀.₅Cs₀.₅PbI₃ to study A₂BX₄-type impurity phases (Figures 1 and 2), while DFT calculations are conducted using CsPbI₃. In contrast, the device performance improvements are demonstrated using a composition of (FA₀.₉₅Cs₀.₀₅PbI₃)₀.₉₇₅(MAPbBr₃)₀.₀₂₅, which is entirely disconnected from the compositions studied in STEM and DFT. This mismatch undermines the credibility of the proposed mitigation strategy. Additionally, the hole transport layer (HTL) was changed from Spiro-OMeTAD to PATT purely to showcase enhanced stability, further complicating the narrative. It appears the authors are generalizing insights obtained from one system to another with limited justification.\n2. The authors claim in lines 145–147 that “Cs content in the A₂BX₄ phase is higher than that in the perovskite grain, demonstrating that cation segregation may trigger IGIN formation.” However, the STEM data is based on FA₀.₅Cs₀.₅PbI₃, a composition already rich in Cs. Conversely, the device structure used for evaluating NDPS contains a much lower Cs content. This inconsistency raises doubts about whether the same IGIN formation mechanism applies across the different compositions studied.\n3. Additional confusion arises from the inclusion of other impurity compositions such as Cs₂PbI₄ and FA₀.₄Cs₀.₆PbI₄ (line 193). Moreover, the optimal NDPS concentration differs between the device study (1 mg/mL) and STEM analysis (2 mg/mL). It is essential to use consistent concentrations across all experimental sections to allow for meaningful correlation of results. Otherwise, it becomes unclear whether observations from one part of the study are applicable to the others.\n4. The DFT calculations shown in Figures 2c and 2d lack sufficient detail regarding the perovskite composition and model parameters. Was FA₀.₅Cs₀.₅PbI₃ used as the host matrix? How was the Cs₂PbI₄ impurity structurally incorporated into the simulation cell? How was the concentration of impurities determined, and how do these compare with the impurity densities observed via STEM?\n5. The identification of a specific A₂BX₄ orthorhombic non-perovskite phase as the dominant IGIN is a key finding. However, the authors should elaborate on whether other impurity phases might be expected in perovskites with different FA/Cs ratios or other compositions (e.g., those used in the final device structure). Furthermore, on page 5, lines 140–141, the authors acknowledge that XRD cannot distinguish A₂BX₄ from PbI₂. Given that STEM is a local probe, how can the presence and suppression of the A₂BX₄ phase be validated on a statistically significant scale? In line 292–294, the authors claim that XRD peaks of impurities decrease with increasing NDPS concentration, contradicting their earlier statement.\n6. In lines 219–221, the authors state that Cs₂PbI₄ has a larger bandgap (2.84 eV) compared to CsPbI₃ (1.83 eV), implying that IGINs introduce charge-blocking regions. However, if Cs₂PbI₄ is embedded within the grain, further clarification is needed on how these impurity phases spatially impede charge transport.\n7. In lines 196–198, the authors describe that after one week of vacuum storage, FA₀.₅Cs₀.₅PbI₃ develops a high density of PbI₂ and FA₀.₄Cs₀.₆PbI₄ IGINs. However, reference 30 (J. Am. Chem. Soc. 2022, 144, 4, 1910–1920) reports excellent shelf-life stability for FA₀.₅Cs₀.₅PbI₃ (Figure S5). The discrepancy between these findings requires a thorough explanation.\n8. The rationale for selecting NDPS as an additive is not sufficiently justified. The authors should provide a mechanistic explanation of how NDPS enhances miscibility, forms kinetically stabilized cationic clusters, and suppresses impurity formation. Moreover, the mechanism by which NDPS slows crystallization should be elaborated upon with supporting evidence.\n9. In Supplementary Figures 24–25, the authors show that perovskite films without IGINs are more stable under ambient conditions. However, given that IGINs are embedded within the grains and that moisture-related degradation typically initiates at the surface, it remains unclear how IGINs accelerate the degradation process. Further discussion on the degradation mechanism is necessary.", "reviewer_3": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts.", "reviewer_4": "This manuscript investigates the nature of intragrain impurity nanoclusters (IGINs) in metal halide perovskite films and proposes a chemical additive strategy for their mitigation. The authors employ a combination of advanced characterization techniques, notably low-dose scanning transmission electron microscopy (STEM), and density functional theory (DFT) calculations to identify a metastable A₂BX₄-type impurity phase. While the quality of the experimental data, particularly the atomic-resolution STEM imaging, is commendable, I have significant concerns regarding the manuscript's novelty and conceptual advance, especially in light of the authors' own recent publication on a highly similar topic (Intragrain impurity annihilation for highly efficient and stable perovskite solar cells, Nature Communications, 2024, 15, 2329).\nMy primary concern is the substantial overlap with their previous work in terms of the research subject, core methodology, and ultimate scientific goal. Both studies focus on IGINs, rely heavily on STEM as the central characterization tool, and aim to improve device performance by addressing these impurities. While the current work identifies a specific A₂BX₄ phase, this feels more like an incremental progression rather than the fundamentally new discovery that would warrant publication in a top-tier journal. The conceptual framework remains largely the same: identifying and eliminating intragrain impurities.\nFurthermore, several of the key elements presented as innovative in this manuscript are, in fact, well-established in the field. The use of low-dose STEM for characterizing beam-sensitive perovskites has been reported in numerous prior studies and can no longer be considered a novel technique in itself. Similarly, the use of inner salt additives to passivate defects and control crystallization in perovskites is a widely explored strategy. While the authors apply it to mitigate IGINs, the underlying chemical approach lacks fundamental novelty.\nGiven these significant concerns about the manuscript's independent conceptual advance, I cannot recommend publication in its current form. The authors must convincingly differentiate this work from their prior publication and from the broader literature.\nSome major questions the authors:\n1. The authors' previous Nature Communications paper already established the existence of IGINs and a method for their “annihilation”. The core claim of this new manuscript is the identification of a specific A₂BX₄ phase. Could the authors please clarify why the identification of this specific stoichiometry constitutes a fundamental conceptual leap, rather than a more detailed characterization of the previously reported phenomenon? How does knowing the phase is A₂BX₄, rather than simply PbI₂ or another impurity, fundamentally change our understanding or the mitigation strategies in a way that is not incremental?\n2. The manuscript repeatedly highlights the use of low-dose STEM. However, this is a well-established technique for perovskite characterization (e.g., Science 370, 548 (2020), Nature Communications, 2023 14:1940, ACS Energy Lett. 2023, 8, 2620−2629, and a review paper on low dose TEM: Energychem, 2023,5,100105). Could the authors please specify what is uniquely innovative about their application of this technique in the current study that has not been demonstrated before?\n3. The use of zwitterionic additives, such as the NDPS reported here, for defect passivation and crystallization control is a common strategy in the perovskite field. The proposed mechanism that interaction with precursor clusters to homogenize the solution and slow crystallization is also a well-understood concept for many additives. What is the specific novelty of using NDPS for IGINs that distinguishes it from the vast body of work on additives for grain boundary passivation and bulk defect reduction? Is there a chemical interaction that is uniquely effective against the formation of the A₂BX₄ phase specifically?\n4. The manuscript needs to more explicitly and convincingly articulate its unique contribution relative to the authors' own Nature Communications 2024 paper. Beyond identifying a different impurity composition (A₂BX₄ vs. PbI₂) and proposing a different solution (chemical suppression vs. physical annihilation), what is the overarching, new scientific message? Without a clear distinction, this work appears to be a follow-up study that is too closely related to the previous one for a separate high-impact publication.\n5. The study focuses on FA-Cs based perovskites. How prevalent is this newly identified A₂BX₄ phase in other important perovskite systems (e.g., FAPbI3, MAPbI₃, triple-cation, all-inorganic)? Is this a universal, previously overlooked impurity, or a phenomenon specific to the composition studied? The claimed broad impact hinges on this point."}, "author_response": "General Response: We sincerely thank all the reviewers for their insightful and constructive\ncomments on our manuscript. We have thoroughly examined each comment and addressed\nthem with detailed, point-by-point response. We have significantly revised manuscript to\nclearly differentiate our present work from our prior publication and the broader literature.\nAdditional theoretical calculations and mechanistic characterizations were performed to\nimprove the accuracy and consistency of our conclusions. Through more in-depth studies, we\nhave further clarified the novelty, significance, and rigor of our findings.We are confident that\nthe carefully revised manuscript based on the detail responses is now suitable for its publication\nin Nature Communications.\nResponse to Reviewer #1:\nThe Authors focus on a class of nanoscale features of perovskite films for solar cells that are\nnot often studied - they refer to them as intra-grain impurity nanoclusters and, unlike the\ncommonly studied features at the grain boundaries, exist within individual perovskite grains,\ngenerally with some level of epitaxial matching. They are generally harmful to solar cell\nperformance, and their direct study is very challenging because their lattice features are very\nsimilar to the surrounding material and they evolve under the electron beam more than the\nperfect halide perovskite lattice.\nThe STEM experiments described are very challenging, as well as the data interpretation. The\nAuthors have demonstrated in the past the ability to conduct careful experiments on these\nmaterials. The study of nano-domains in halide perovskites is a very interesting and active\ndirection-for example there is a very recent Nature Nanotech paper\n(https://www.nature.com/articles/s41565-025-01917-0) which looks into impurity domains\nwith different techniques and without the same direct imaging resolution as this manuscript.\nSpecifically, this manuscript covers a set of FA/Cs perovskites and proposes a mitigation\nstrategy for these impurities via the use of NDPS, which results in better PCE and stability.\nThis work combines very good EM investigation with complex sample preparation and an\nextensive set of experiments and simulations.\nSomething that should be pointed out more clearly in the manuscript (the paper is referenced,\nbut not elaborated upon much) is that this builds on a previous nature comms paper (ref 30 -\nCai, S. et al. Intragrain impurity annihilation for highly efficient and stable perovskite solar\n638 cells. Nat. Commun. 15, 2329 (2024)). In the previous paper a subset of the Authors had\nidentified intragrain impurities in these systems and concluded that they were in some cases\nidentifiable as PbI2 and in others as a 0D phase. These impurities had been observed as a\nfunction of electron dose, the process resulting in annealing/healing, and a laser treatment was\nsuggested to improve performance.\nI think it is important that the Authors here elaborate (if they desire, in a rebuttal letter, but\nalso this should go in the main manuscript) on how these findings relate to previous information.\nI fully understand that the identification process for these IGINs is complex and current\nfindings might overrule previous hypotheses (the previous paper is not very conclusive on the\nnature of the second class of IGINs), but for this there should be a frank discussion, perhaps\nincluding motivations on why the current measurements are more representative. This should\napply to both the identification of the phases and the evolution of the impurities under the\nelectron beam.\nThe proposed mechanism of cation segregation for formation of IGINs is rather interesting\nand a good approach to rationally designing an additive for improving overall performance.\nWe thank the reviewer for the excellent summary of our work and the overall positive\ncomments on its novelty, impact and quality. We agree that it is necessary to clarify the\nconnection and difference between present work and our prior publication. Our previous report\nrevealed the existence of PbI₂ and non-PbI₂ intragrain impurities and found that they can\nundergo a healing reaction under appropriate electron irradiation or laser intensity to transform\ninto perovskites, thereby improving the photoelectric properties of perovskites. However, due\nto the unknown atomic configuration of the non-PbI₂ phases, the formation mechanism of these\nintracrystalline precipitates and the mechanism of the healing reaction remained unclear.\nBuilding on this previous work, we have precisely reconstructed the atomic structure of the\nnon-PbI intracrystalline impurities according to the atomic-resolution images obtained from\n2\nrepresentative directions. We then carried out image simulation to determine their chemical\ncomposition to be (FA Cs ) PbI , and proposed possible reaction of their formation and\n0.4 0.6 2 4\nhealing process:2(FA,Cs)PbI → PbI + (FA,Cs) PbI .\n3 2 2 4\nFurthermore, we revealed that the presence of A BX impurity adversely affects the\n2 4\noptoelectronic properties and stability of the perovskite by combining DFT calculations with\natomically resolved quasi-in situ experiments. Therefore, the present work and our previous\nwork provide an atomic-scale understanding of how the intragrain microstructure affects the\nperformance of perovskites together, and the two works complement each other. We appreciate\nthe reviewer to provide constructive comments to help us future boost the quality of this\nmanuscript, and we have incorporated the relevant discussions into the revised manuscript\n(lines 74-82, 160-168).\n1. Line 190 - the definition of in situ and ex situ in this context is unclear. I assume by in situ\nthe Authors mean evolution under the electron beam, and ex situ exposure to air and moisture?\nThis should be made clearer in the text.\nThank you for your suggestions. To avoid confusion and express more accurately, we revised\n\"in situ\" to \"quasi-in situ\" and explained the experimental process in detail. For quasi-in situ\nobservations, the samples were removed from the electron microscope chamber immediately\nafter observations to minimize electron beam induced artifacts, and exposed to ambient\nconditions (atmosphere, 25 ℃ and 50% relative humidity (RH)). The structural evolutions were\ncontinuously tracked though low-dose STEM imaging. The samples for ex situ observations\nwere transferred to vacuum packages immediately after fabrication and then investigated after\none week storage.\n2. In page four, in reference to the emergence of the PbI phase, the FTT pattern before and\n2\nafter degradation is brought as evidence; the concept is clear in the image but much less so in\nthe text. It would be necessary to better explain the reasoning used (so as to understand the\nconcept without images) and quantify the lengths to verify that the phase that emerges after\ndegradation is PbI .\nchanges should be explained in detail in the text to facilitate understanding. Therefore, we\nrevised the relevant content to “According to the atomic resolution image and the\ncorresponding fast Fourier transformation (FFT) pattern (inset in Fig. 2a), only IGINs exist in\nthe pristine sample, confirmed by the FFT spot of (020) plane of IGINs with 0.656 nm\ninterplanar spacing indicated by yellow circle in the inset in Fig. 2a. After the exposure, PbI\n2\nphases evolve at the original locations of (FA Cs ) PbI impurities, as shown in Fig. 2b. The\n0.4 0.6 2 4\nemergence of a new spot with an interplanar spacing of 1.457 nm (corresponding to the (001)\nplane of 2H PbI ), indicated by white circle in the inset in Fig.2b, in the corresponding FFT\n2\npattern (inset in Fig.2b) confirmed the formation of PbI .” (lines 228-242).\n2\n3. In the GIXRD experiments section, the reflection (012) at 31.5° is attributed using a cubic\nsymmetry. A reference of the .cif file used should be added, or at least the size of one side of\nthe cubic cell should be defined, so that the accuracy of the deductions can be verified.\nThank you for the constructive suggestions. We have added a description of the structures used,\nincluding the space groups and lattice parameters, to the revised version (lines 387-388).\n4. Fig. 1A - the GPA colourscale should have units - I assume it is % strain. Same for Supp Fig\n15, 22, 27\nThank you for your suggestions. We have modified the corresponding figures after carefully\nchecking the calculation results of GPA.\n5. Supp Fig 9 - \"squires\" should obviously be \"squares\"\nThank you for pointing out typos in our manuscript. We apologize for any inconvenience\ncaused by the format errors and typos. We havecorrected it accordingly and carefully reviewed\nthe entire manuscript to enhance the language quality.\n6. Line 320 - Regarding the use of sulfur as \"tracer element\" - normally tracers are something\nthat is added for the purpose of tracking. Here my understanding is that the compound already\ncontains sulfur, and it is a convenient atomic species to map because it gives a good signal\n(EELS or EDS), unlike lighter elements. I think this should be clarified in the text. For example\nthis could be changed to \"[...] (EELS) measurements by exploiting the signal provided by sulfur,\nwhich unequivocally identifies the additive from the perovskite.\"\nThank you for reading carefully and helping us clarify the concepts. We have modified the\nrelevant content in the revised manuscript (lines 403-406).\nResponse to Reviewer #2:\nThis manuscript by Wang et al. reports on the investigation of intragrain impurity nanoclusters\n(IGINs) in metal halide perovskites, specifically FA-Cs perovskite films. The authors employ\nlow-dose scanning transmission electron microscopy (STEM) and first-principles calculations\nto identify and characterize these IGINs, which are revealed to be metastable orthorhombic\nnon-perovskite A BX phases. These impurity phases are shown to negatively impact the\n2 4\noptoelectronic properties and chemical stability of the perovskite films. To mitigate this issue,\nthe authors propose a strategy using a salt additive, NDPS, which results in notable\nimprovements in both power conversion efficiency (PCE) and operational stability of\nperovskite solar cells (PSCs). The STEM-based findings are indeed interesting and contribute\nto the fundamental understanding and nano-engineering of perovskite materials. However, I\nfind serious inconsistencies between the STEM analyses, DFT calculations, and device-level\ndemonstrations, which hinder the overall coherence of the study.\nWe thank the reviewer for the evaluation of our work and the constructive comments. We\nrespectfully clarify that the STEM analyses, DFT calculations, and device-level demonstrations\nare in fact logically consistent. The perceived “inconsistencies” mainly arose from insufficient\nexplanation in the original manuscript, rather than actual discrepancies among these parts.\nSpecifically, the STEM analyses identify the presence and structural features of intragrain\nimpurity nanoclusters (IGINs), the DFT calculations reveal their metastable energetic nature\nand electronic impacts, and the device-level studies demonstrate the corresponding\nperformance and stability improvements upon removing these IGINs. These three components\ntogether form a coherent narrative from atomic-scale characterization to theoretical\nunderstanding, and finally to device engineering. To address the reviewer’s concern, we have\nrevised the manuscript to clarify the rationale behind the use of different perovskite\ncompositions, and we explicitly explain how the experimental observations and theoretical\nresults are connected. We have also added supplementary experimental data and DFT\ncalculations to further reinforce this logical flow. We believe the revised version now presents\na clear and coherent connection among all parts of the study.\n1. Overall, the manuscript gives the impression of merging three separate studies into a single\nstory, making it difficult to derive firm conclusions. For instance, the STEM investigations\nfocus on FA Cs PbI to study A BX -type impurity phases (Figures 1 and 2), while DFT\n0.5 0.5 3 2 4\ncalculations are conducted using CsPbI . In contrast, the device performance improvements\n3\nare demonstrated using a composition of (FA Cs PbI ) (MAPbBr ) , which is\n0.95 0.05 3 0.975 3 0.025\nentirely disconnected from the compositions studied in STEM and DFT. This mismatch\nundermines the credibility of the proposed mitigation strategy. Additionally, the hole transport\nlayer (HTL) was changed from Spiro-OMeTAD to PTAA purely to showcase enhanced stability,\nfurther complicating the narrative. It appears the authors are generalizing insights obtained\nfrom one system to another with limited justification.\nThank you for your valuable comments. We would like to clarify that IGINs are ubiquitously\npresent in Cs-incorporated perovskite films with different Cs contents, as confirmed by our\nlow-dose STEM observations on multiple compositions (Figure R1). Based on this\nuniversality, we deliberately selected different Cs ratios for different types of measurements to\naccommodate the technical requirements and limitations of each characterization method,\nrather than to study unrelated systems. This experimental design does not affect the scientific\nconclusions regarding the impact of IGINs, which remain consistent across all compositions\nexamined.Moreover, the IGIN removal strategy we developed is universal and applicable to\nall FA-Cs perovskite systems.\nFor STEM analyses, we used FA Cs PbI films because increasing the Cs content slightly\n0.5 0.5 3\nimproves the electron beam tolerance of FA-Cs perovskites. Low-Cs compositions, such as\n(FA Cs )PbI and (FA Cs PbI ) (MAPbBr ) , can withstand a maximum\n0.9 0.1 3 0.95 0.05 3 0.975 3 0.025\ncumulative electron dose below 100 e·Å-2. Exceeding this threshold generates many vacancies,\ninduces structural changes, and results in artifacts. These effects make it impossible to\naccurately determine the spatial distribution of impurities (inside grains vs. on surfaces),\nanalyze their three-dimensional atomic structure, or perform quasi-in situ experiments.\nTherefore, to ensure high-fidelity STEM observations without beam-induced artifacts, we\ntherefore selected a composition with higher Cs incorporation.\nFor device fabrication, we used the composition (FA Cs PbI ) (MAPbBr ) .A small\n0.95 0.05 3 0.975 3 0.025\namount (2.5%) of MAPbBr was incorporated to optimize the photovoltaic performance within\n3\nthe FA-Cs perovskite system, which has been used in many high-impact studies and high-\nefficiency perovskite solar cells. As shown in Figure R1d, IGINs are still clearly observable\nin this composition, indicating that the presence of IGINs remains relevant even in optimized\ndevice configurations. This allows us to evaluate the effect of IGIN removal under conditions\nwhere device performance is already near optimal, ensuring that any observed improvements\nare directly attributable to the mitigation strategy. We have added a brief description of the\nreason in the revised manuscript (lines 131-138).\nFigure R1 a-d Atomic-resolution HAADF-STEM images obtained from FA Cs PbI ,\n0.5 0.5 3\nFA Cs PbI , FA Cs PbI and (FA Cs PbI ) (MAPbBr ) samples.\n0.8 0.2 3 0.9 0.1 3 0.95 0.05 3 0.975 3 0.025\nFor DFT calculations,\nThe valence band maximum (VBM) and the conduction band minimum (CBM) of perovskite\nand the relevant phase composed of octahedral units are determined by the length of Pb-I and\nthe angle of Pb-I-Pb. The atoms, located at gaps between octahedral units, such as FA, MA and\nCs, have limited effect on their electronic structures (e.g., Chem. Rev. 120, 7867-7918 (2020),\nScience 375, eabj1186 (2022)). However, DFT calculations for compounds with complex\ncompositions, such as (FA Cs PbI ) (MAPbBr ) , require much more computing cost\n0.95 0.05 3 0.975 3 0.025\nand time than that with sample compositions, such as CsPbI . We compared the calculated the\n3\nelectronic structures of perovskite and IGINs with different compositions, as shown in Fig. 2\nin the revised manuscript and Figure R2. Figure R2a, 2b are the calculated DOS of\n(FA Cs )PbI and (FA Cs ) PbI under Cs/FA mixed occupation, respectively. It can be\n0.5 0.5 3 0.4 0.6 2 4\nseen that their electronic structures are very close to those of CsPbI and Cs PbI . The formation\n3 2 4\nenergy of (FA Cs )PbI is relatively lower than that of (FA Cs ) PbI under Cs/FA mixed\n0.5 0.5 3 0.4 0.6 2 4\noccupation (Figure R3). We further calculated the electronic structures of the interfaces\nbetween perovskite and IGINs with different compositions, as shown in Fig. 2 and Figure R4.\nThe electronic structure and defect state properties of the CsPbI /Cs PbI interfaces are\n3 2 4\nconsistent with those of (FA Cs )PbI /(FA Cs ) PbI . Our calculated results for systems\n0.5 0.5 3 0.4 0.6 2 4\nwith different compositions confirmed previous results that the electronic structures of\nperovskite relevant systems are dominated by Pb and I. For the sake of clarity and to reduce\nthe cost of computation, in manuscript we use CsPbI and Cs PbI as models to calculate the\n3 2 4\neffect of IGIN formation on the stability and electronic structure of the perovskite system.\nDue to the inadequate description, some confusion has arisen. Therefore, we have revised the\nrelevant content as shown below and displayed it in red font in the reversed manuscript.\nFigure R2 DOS of (FA Cs )PbI a and (FA Cs ) PbI b under Cs/FA mixed occupancy,\n0.5 0.5 3 0.4 0.6 2 4\nrespectively.\nFigure R3 Formation energies of (FA Cs )PbI and (FA Cs ) PbI under Cs/FA mixed\n0.5 0.5 3 0.4 0.6 2 4\noccupancy, respectively.\nFigure R4 a, c and e DOS of interface I, II and III between (FA Cs )PbI and\n0.5 0.5 3\n(FA Cs ) PbI without vacancies. b, d and f Spin-polarized DOS of interfaces I, II and III\n0.4 0.6 2 4\nwith an V at the interfacial region.\nI\nFor HTL selection, we used PTAA to avoid device degradation dominated by spiro-OMeTAD\ninstability before perovskite layer degradation. Here we want to compare the intrinsic stability\nof perovskite layer before and after the mitigation of IGINs.\nDoped spiro-OMeTAD is known to crystallize under thermal stress at ~100 °C, disturbing\ninterface contact, and devices employing spiro-OMeTAD HTLs typically exhibit ≥20 %\nefficiency loss after ISOS-L-1 testing at 85 °C, even though most perovskites remain stable in\nthis temperature range. Moreover, LiTFSI/tBP dopants are hygroscopic and can migrate under\nbias, inducing morphological changes and side reactions before the perovskite layer begins to\ndegrade. In contrast, PTAA is completely amorphous, shows no glass transition or melting\nfeature up to 300 °C, and forms smooth and intimate contact with the perovskite absorber. Its\nsuperior thermal and environmental robustness enable the observed device degradation to be\nattributed primarily to the effect of IGIN elimination on the intrinsic stability of the perovskite\nlayer, rather than HTL-related failure. This experimental approach is consistent with previous\nhigh-impact studies that replaced spiro-OMeTAD with PTAA or other more robust HTLs to\nbetter probe intrinsic perovskite stability (e.g., Science 377, 531-534 (2022); Science 385, 161-\n167 (2024); Nat. Photon. 19, 701-708 (2025)).\n2. The authors claim in lines 145-147 that “Cs content in the A BX phase is higher than that\n2 4\nin the perovskite grain, demonstrating that cation segregation may trigger IGIN formation.”\nHowever, the STEM data is based on FA Cs PbI , a composition already rich in Cs.\n0.5 0.5 3\nConversely, the device structure used for evaluating NDPS contains a much lower Cs content.\nThis inconsistency raises doubts about whether the same IGIN formation mechanism applies\nacross the different compositions studied.\nThank you for this question. We would like to clarify that IGINs are ubiquitously present in\nCs-incorporated perovskite films with different Cs contents, as confirmed by our low-dose\nSTEM observations on multiple compositions (Figure R5). This suggests that the underlying\nmechanism of IGIN formation-namely, cation segregation leading to the nucleation of A₂BX₄-\ntype phases-is general and not limited to Cs-rich systems. To further validate this, we examined\nthe effect of NDPS treatment on samples with different Cs ratios and found that NDPS\nconsistently suppressed IGIN formation, as shown in Figure R6. This result supports the notion\nthat the formation and elimination mechanisms of IGINs are applicable across a wide\ncompositional range in FA-Cs systems.\nGiven the IGIN in the (FA Cs )PbI sample has a higher density and larger size, and the\n0.5 0.5 3\n(FA Cs )PbI sample also has good resistance to electron irradiation, as described in our\n0.5 0.5 3\nresponse to your first question, we choose this composition to conduct quasi-in situ and ex situ\nSTEM observations. Therefore, our conclusions based on the microstructure of (FA Cs )PbI\n0.5 0.5 3\nare applicable to other FA-Cs-based perovskite films. To avoid any confusion, we have clarified\nthis point in the revised manuscript (lines 131-138) to avoid any potential confusion.\nFigure R5 (same as Figure R1) a-d Atomic-resolution HAADF-STEM images obtained from\nFA Cs PbI , FA Cs PbI , FA Cs PbI and (FA Cs PbI ) (MAPbBr ) samples.\n0.5 0.5 3 0.8 0.2 3 0.9 0.1 3 0.95 0.05 3 0.975 3 0.025\nFigure R6 a-d Atomic-resolution HAADF-STEM images obtained from FA Cs PbI ,\n0.5 0.5 3\nFA Cs PbI , FA Cs PbI and (FA Cs PbI ) (MAPbBr ) samples with 1 mg/mL\n0.8 0.2 3 0.9 0.1 3 0.95 0.05 3 0.975 3 0.025\nNDPS.e-h Corresponding GPA analysis for images shown in a-d.\n3. Additional confusion arises from the inclusion of other impurity compositions such as\nCs PbI and (FA Cs ) PbI (line 193). Moreover, the optimal NDPS concentration differs\n2 4 0.4 0.6 2 4\nbetween the device study (1 mg/mL) and STEM analysis (2 mg/mL). It is essential to use\nconsistent concentrations across all experimental sections to allow for meaningful correlation\nof results. Otherwise, it becomes unclear whether observations from one part of the study are\napplicable to the others.\nThank you for the constructive comments and suggestions. Based on our atomic-resolution\nSTEM imaging and image simulation results, the composition of IGINs in various FA-Cs based\nperovskite films is determined as (FA Cs ) PbI₄. For the sake of clarity and to reduce the\n0.4 0.6 2\ncost of computation, in manuscript we use CsPbI and Cs PbI as models to calculate the effect\n3 2 4\nof IGIN formation on the stability and electronic structure of the perovskite system. Due to the\ninadequate description, some confusion has arisen. Therefore, we have revised the relevant\ncontent and displayed it in red font in the reversed manuscript.\nWe agree that keeping the same concentrations of NDPS in devices and samples for STEM\nanalysis is crucial for the meaningful results. We thank the reviewer again for your suggestion.\nWe measured the distribution of S element in samples with 1 mg/mL NDPS addition using low-\ndose STEM EELS method and updated the results in Fig. 3 in the revised manuscript.\nElemental mapping results of S showed that the NDPS was enriched at top surface and GBs,\ninstead of entering the intragrain, aligned with the XRD results. Based on these results, we\nproposed the possible mechanisms for the NDPS-induced IGINs chemo-elimination. This is\nelaborated in detail in the response to question 8.\n4. The DFT calculations shown in Figures 2c and 2d lack sufficient detail regarding the\nperovskite composition and model parameters. Was FA Cs PbI used as the host matrix?\n0.5 0.5 3\nHow was the Cs PbI impurity structurally incorporated into the simulation cell? How was the\n2 4\nconcentration of impurities determined, and how do these compare with the impurity densities\nobserved via STEM?\nThank you for your suggestions. We describe in detail the models and parameters used in the\nDFT calculations in the revised version (lines 286-288). As described in our response to\nquestion 1, the effect of Cs content on the electronic structure and physical properties of the\nperovskite and A BX phases, as well as the interface electronic structure, is very limited, but\n2 4\nit significantly increases the computational cost. We also compared the\nFA Cs PbI /(FA Cs ) PbI and CsPbI /Cs PbI systems, confirming this conclusion.\n0.5 0.5 3 0.4 0.6 2 4 3 2 4\nTherefore, we only present the results based on the CsPbI /Cs PbI system in Fig. 2 in the\n3 2 4\nrevised manuscript. To avoid confusion, we have also provided explanations in the revised\nmanuscript.\nNeither CsPbI nor Cs PbI was used as the host matrix in the simulation cells shown in Fig.\n3 2 4\n2g, i, l, n, q, and s in the revised manuscript. In fact, one half of each simulation cell contains\nCsPbI while the other half contains Cs PbI . This configuration enables us to directly study\n3 2 4\nthe targeted interface between both phases without the introduction of other interfaces (those\nwith different indices of lattice planes). If we construct a simulation cell with one phase as the\nmatrix and the other embedded inside the former one, there will be multiple interfaces, thus\npreventing us from investigating the electronic properties of the targeted one. In this case, the\nconcentration of the IGIN phase in the simulation cell deviates significantly from the\nexperimental value. Nevertheless, the electronic structure of an interface is a robust property\n(related with the corresponding lattice planes rather than the concentration of phases) and it\ncan therefore be used as a critical factor for evaluating the influence of IGINs on the\nphotovoltaic properties.\nMoreover, the computational complexity of DFT calculations increases super-linearly with the\nnumber of atoms, and calculations for the embedded configurations will need much larger\nsimulation cell, thus bringing huge computational cost challenges. The number of atoms in the\nsystem calculated by DFT is generally limited to the range of 100-500. In present study, we\nused DFT to calculate the stability and electronic structure of the A BX phase, as well as the\n2 4\nelectronic structure of the perovskite/IGIN interfaces. For these structures, the computational\ncost remains acceptable. But for more realistic and sophisticated structures, the computational\ndemand exceeds our capabilities. We revealed that the presence of IGINs leads to degradation\nof the perovskite film and negatively impacts its optoelectronic properties. However,\ncomputational limitations prevented us from scaling our DFT calculations to the same scale as\nour experimental samples. Our atomic-resolution low-dose investigations show that the IGIN\nstructure and interface structure in perovskite films of different compositions are consistent\nwith the model used in our DFT calculations, as shown in Figure R7, confirming the\nuniversality of our calculation results and the reliability of our performance predictions.\nFigure R7 a-d Atomic-resolution HAADF-STEM images obtained from FA Cs PbI ,\n0.5 0.5 3\nFA Cs PbI , FA Cs PbI and (FA Cs PbI ) (MAPbBr ) samples. Insets in a-d\n0.8 0.2 3 0.9 0.1 3 0.95 0.05 3 0.975 3 0.025\nare the atomic models of the perovskite-IGIN interfaces.\n5. The identification of a specific A BX orthorhombic non-perovskite phase as the dominant\n2 4\nIGIN is a key finding. However, the authors should elaborate on whether other impurity phases\nmight be expected in perovskites with different FA/Cs ratios or other compositions (e.g., those\nused in the final device structure). Furthermore, on page 5, lines 140–141, the authors\nacknowledge that XRD cannot distinguish A BX from PbI . Given that STEM is a local probe,\n2 4 2\nhow can the presence and suppression of the A BX phase be validated on a statistically\n2 4\nsignificant scale? In line 292-294, the authors claim that XRD peaks of impurities decrease\nwith increasing NDPS concentration, contradicting their earlier statement.\nWe thank the reviewer for the constructive comments and suggestions. To investigate whether\nthere are other impurities in perovskites with different FA/Cs ratios, we further carried out SEM\nand STEM investigations on phase compositions in the FA Cs PbI and\n0.8 0.2 3\n(FA Cs PbI ) (MAPbBr ) films, in addition to our original analyses on\n0.95 0.05 3 0.975 3 0.025\nFA Cs PbI and FA Cs PbI composition films. As shown in Figure R8, bright contrast\n0.5 0.5 3 0.9 0.1 3\nprecipitates at grain boundaries were occasionally observed across all compositions. Atomic-\nresolution STEM confirmed these precipitates to be PbI , in agreement with XRD results\n2\n(Figure R9). Importantly, within the perovskite grains, STEM observations across multiple\nbatches consistently revealed that A₂BX₄ phase was the main intragrain impurity phase in\nfreshly prepared samples, with only occasional PbI and no other phases (such as δ-CsPbI and\n2 3\nδ-FAPbI ) found in our freshly prepared samples (Figure R10). So, we confirm that the A₂BX₄\n3\northorhombic non-perovskite phase is the dominant phase in the fresh fabricated FA-Cs based\nperovskite films and one of the most important factors affecting the stability of perovskite\ndevices.\nSimulated results show that the main XRD diffraction peaks of the orthorhombic A₂BX₄ phase\nare overlapped with those of the PbI , as shown in Figure R11.The formation of (FA,Cs)₂PbI₄\n2\nphase is also accompanied by the appearance of PbI₂ according to equation 2(FA,Cs)PbI₃ →\nPbI₂ + (FA,Cs)₂PbI₄. Therefore, the XRD method alone cannot effectively identify the existence\nof the A BX phase. But, the highlight of the peak at the 2θ = ~12° (corresponding to (002)\n2 4\nplane of PbI or (110) plane of (FA,Cs)₂PbI₄) in the normalized pattern according the (100)\n2\npeak of perovskite (around 14°) can reflect the change of the impurity content ((FA,Cs)₂PbI₄\nand/or PbI ). XRD measurements (Supplementary Fig. 25 in revised manuscript) confirmed\n2\nthat the impurity content decreases with increase of the NDPS amount. However, given the\nlimitation of the XRD, atomic resolution imaging is a necessary and effective means to identify\nthe types and distribution of impurities.\nWe fully understand the reviewers' concerns about the statistical validity of the results due to\nthe limited scope of STEM observations.Therefore, we observed several different batches of\nFA Cs PbI thin film samples, which are most prone to forming the A BX phase. Some of\n0.5 0.5 3 2 4\nthe results are shown in Figure R12. All results show that the addition of NDPS effectively\neliminates the A BX phase caused by Cs segregation within the grains. This result is consistent\n2 4\nwith the gradual decrease in the XRD peak intensity of impurity with increasing additive\ncontent in precursor solution.\nFigure R8 a-c Top-view SEM images of pristine FA Cs PbI , FA Cs PbI and\n0.9 0.1 3 0.8 0.2 3\n(FA Cs PbI ) (MAPbBr ) samples.\n0.95 0.05 3 0.975 3 0.025\nFigure R9a Low magnification HAADF-STEM images obtained from pristine FA Cs PbI\n0.9 0.1 3\nsample. b Atomic-resolution image of the PbI around GBs.\n0.5 0.5 3\nsamples. (FA Cs ) PbI is the main intragrain impurity, with only a small amount of\n0.6 0.4 2 4\nadditional PbI present.\n0.6 0.4 2 4 2\nFigure R12 a-d Atomic-resolution HAADF-STEM images obtained from FA Cs PbI\n0.5 0.5 3\nsamples with 1 mg/mL NDPS. No intragrain impurities were observed in the cation-\nhomogenized perovskite devices.\n6. In lines 219–221, the authors state that Cs PbI has a larger bandgap (2.84 eV) compared to\n2 4\nCsPbI (1.83 eV), implying that IGINs introduce charge-blocking regions. However, if Cs PbI\n3 2 4\nis embedded within the grain, further clarification is needed on how these impurity phases\nspatially impede charge transport.\nThank you for your suggestions. Our DFT calculations show that Cs PbI exhibits a\n2 4\nconsiderably larger bandgap (2.84 eV) than CsPbI (1.83 eV). After aligning the energy bands\n3\naccording to the vacuum level, a type I heterojunction distribution is formed as illustrated in\nFigure R13. The conduction band minimum of the perovskite phase is 0.24 eV lower than the\nIGIN phase, while the valence band maximum is 0.77 eV higher. Therefore, the charge carriers\ngenerated in the perovskite phase can hardly be transferred to the IGIN phase. Instead, they\nwill bypass the IGIN phase along the perovskite domain during their transfer. This causes the\nelectron transfer pathway to become longer, resulting in more consumption and electron-hole\nrecombination. This is the underlying mechanism of charge-blocking by IGIN phase, which\nwill obviously reduce the photoelectric performance of perovskite films.\nFigure R13 Schematic illustration of band alignment of the CsPbI and Cs PbI .\n3 2 4\n7. In lines 196-198, the authors describe that after one week of vacuum storage, FA Cs PbI\n0.5 0.5 3\ndevelops a high density of PbI and FA Cs PbI IGINs. However, reference 30 (J. Am. Chem.\n2 0.4 0.6 4\nSoc. 2022, 144, 4, 1910-1920) reports excellent shelf-life stability for FA Cs PbI (Figure\n0.5 0.5 3\nS5). The discrepancy between these findings requires a thorough explanation.\nThank you for your carefully reading and raising this question. Careful inspection confirmed\nthat the storage conditions of the two researches were different. In our previous work, storage\nstability experiments were conducted under strictly light-shielded conditions in a glove box\nwith an ultra-high purity nitrogen atmosphere (H O < 0.1ppm, O < 0.1ppm). Specifically, the\n2 2\ndevices were enclosed in multi-layer aluminum light-shielding bags and the temperature was\nstrictly controlled at 25°C.The efficiency measurement is also completed in a very short time,\nand the light shielding is performed immediately after the test is completed. In the ex situ\nexperimental study in present work, we stored the device in a transparent vacuum bag (vacuum\ndegree is about -60kPa) under laboratory conditions without temperature control. Negative\npressure conditions, residual water and oxygen, and laboratory lighting will trigger perovskite\ndegradation.In addition, the degradation rate of perovskite in vacuum package is lower than in\nthe atmosphere, but significantly faster than in a glove box. This relatively slow degradation\nrate helps us compare the stability of the A BX phase with the perovskite phase and reveal the\n2 4\nimpact of the presence of IGINs on device stability. Therefore, these two completely different\nexperimental conditions are the main reason for the significant difference in device storage\nstability.\n8. The rationale for selecting NDPS as an additive is not sufficiently justified. The authors\nshould provide a mechanistic explanation of how NDPS enhances miscibility, forms kinetically\nstabilized cationic clusters, and suppresses impurity formation. Moreover, the mechanism by\nwhich NDPS slows crystallization should be elaborated upon with supporting evidence.\nThank you for your valuable suggestions. It is reported that the O atom in tetrabutylammonium\nanion ([TFSI]) is able to act as monodentate and bidentate ligands to form dative bonds with\nalkali metal ions, though the chelate effect (Chem. Commun. 12, 1368-1369 (2004); Energy\nEnviron. Sci. 17, 8209-8218 (2024)). We first performed DFT calculations to reveal the\ninteraction between NDPS and Cs in the precursor solution, as shown in Table R1. As the\nGibbs free energy presented in this table indicates this coordination is an exothermic process,\nwhich is thermodynamically favorable and contributes significantly to the overall energy\nstabilization of the system. These results clearly indicate that the chelate effect occurs between\nDNPS and Cs+ in precursor solution and can effectively disperse Cs.\nTable R1 TheGibbs free energy for the coordination between Cs+ and NDPS/[N₄₄₄₄][TFSI].\nΔG\nSpecies bidentate and polydentate complexes 298K\n(kcal/mol)\nDNPS + Cs+ -12.8\n2*DNPS + Cs+ -23.8\n[N₄₄₄₄][TFSI] +\n-12.1\nCs+\n2*[N₄₄₄₄][TFSI]\n-28.6\n+ Cs+\nAs a surfactant additive, the NDPS can affect both the nucleation and growth of perovskite\ngrains. As is well known, surfactant molecules contain both hydrophilic and hydrophobic\nfunctional groups. They can self-assemble on the surface of the precursor solution with the\nhydrophobic group pointing to air, thus reducing the surface tension of the solution (Nat. Chem.\n7, 703-711 (2015); Matter 6, 2987-3005 (2023); Nat. Commun. 14, 2166 (2023)). The\nreduction in surface tension was evidenced by the reduced contact angle of perovskite solution\ndroplets on substrate on addition of surfactant Figure R14). Assuming the nucleus are spherical,\nthe nucleation barrier can be expressed as follows:\n= + + (1)\n=\n△𝐺 △𝐺 𝑣𝑜𝑙𝑢𝑚𝑒 △𝐺𝑠𝑢𝑟𝑓𝑎𝑐𝑒 − 3𝑉𝑚 △µ 4𝜋𝑟 𝛾\nenergy, r is the nucleus radius, V is the molecular volume, the chemical potential\nm\ndifference p△er𝐺 so 𝑣 l 𝑜 u 𝑙𝑢 te 𝑚 m 𝑒 olec △ule𝐺 w 𝑠𝑢 h 𝑟 e 𝑓 n 𝑎𝑐 i 𝑒 t is added into nucleus, and is the surface tension.\n△µ\nThe critical nucleus radius r* can be calculated though setting the = 0, as shown follows:\n𝛾\n𝜕△𝐺\nr* = (2)\n𝜕𝑟\n2𝛾𝑉𝑚\nAccording to equation (2), The reduction of surface tension brought about by the addition of\n△µ\nsurfactants can decrease the critical radius of nucleus, namely the quantity of solute in one\nnucleus.\nFurthermore, the stability of nucleus is dominated though Ostwald ripening process, as\ndescribed by\n(2)\n2𝛾𝑉𝑚\nwhere the is the solubility of a small crystal grain with radius r, and is the\n𝐶𝑟 = 𝐶𝑏exp (𝑟𝑘𝐵𝑇)\nsolubility of the bulk crystal. Thus, reduction of surface tension can improve the\nstability of 𝐶s 𝑟 mall clusters. 𝐶𝑏\nIn addition, the nucleation process of perovskite at interface is mainly heterogeneous rather\nthan homogenous, and the relation between these two can be described as:\n(3)\nwhere\n△𝐺ℎ𝑒𝑡𝑒𝑟𝑜𝑔𝑒𝑛𝑒𝑜𝑢𝑠 =△𝐺ℎ𝑜𝑚𝑜𝑔𝑒𝑛𝑒𝑜𝑢𝑠 ·𝑓(𝜃)\n(4)\n3\nand the θ is the contact angle. That me2a−n3s𝑐 𝑜t𝑠h𝜃e+ n𝑐𝑜u𝑠c𝜃leation barrier decrease gradually upon\n𝑓(𝜃) = 4\nincorporating surfactants as the contact angle decreases. According to the Arrhenius equation,\nJ = (5)\n△𝐺\nThus, the addition of surfactant can increase the nucleation rate of perovskite.\n𝐴 exp (−𝑘𝐵𝑇)\nIn the following growth processes, the NDPS is automatically expelled to the perovskite\nsurface and GBs, for its large size, as shown in Fig. 3d in the revised manuscript. The polar\nhydrophilic groups will anchor to the charge ions and crystal surface and the hydrophobic alkyl\nchain tails gathering around the perovskite grains. The long-alkyl-chain will spontaneously\nself-assemble on the perovskite grain and solution surfaces, acting as a “reservoir” to tetrad the\nintercalation of solute ions and “solvent molecular sieve” in alleviating the evaporation of polar\nsolvent. In addition, the high polar S=O can foster the strong H-bonded network in polar solvent\nmolecules, slowing down the evaporation, which further decrease the crystallization process\nduring film annealing (Figure R15). We have incorporated the new experimental results and\ndiscussion as necessary in the revised manuscript and supplementary information.\nFigure R14 Contact angles of control a and target b samples.\nFigure R15In situ absorption tracking under different annealing times of control a and target\nb samples.\n9. In Supplementary Figures 24–25, the authors show that perovskite films without IGINs are\nmore stable under ambient conditions. However, given that IGINs are embedded within the\ngrains and that moisture-related degradation typically initiates at the surface, it remains unclear\nhow IGINs accelerate the degradation process. Further discussion on the degradation\nmechanism is necessary.\nThank you for your valuable suggestion. Our quasi-in situ/ex situ TEM experiments and DFT\ncalculations demonstrate that IGINs are intrinsically unstable relative to perovskites, prefer to\ndegrading into PbI and triggering the decomposition of the surrounding perovskite phase\n2\nunder experimental stresses even with very low water and oxygen content, such as in a vacuum\npackage. It indicates that IGINs may automatically undergo spontaneous thermodynamic\ndegradation reactions even without contacting with water in the air, leading to the\ndecomposition of the perovskite film. In addition, we also examined the distribution of IGINs\nat the upper and lower interfaces of the perovskite film and around the grain boundaries, which\nare easily corroded by water and oxygen. As shown in Figure R16, we confirmed that IGINs\nexist in these locations. This means that water and oxygen in the air only need to diffuse a short\ndistance (several nanometers) to trigger the decomposition of IGINs and surrounding\nperovskite to form layered PbI , making it easier for water and oxygen to diffuse into the\n2\nperovskite and accelerating its degradation. We have incorporated the new experimental results\nand discussion as necessary in the revised manuscript and supplementary information.\nFigure R16 a-d Atomic-resolution HAADF-STEM images of interfaces and GBs obtained\nfrom pristine FA Cs PbI samples.\n0.5 0.5 3\nResponse to Reviewer #4:\nThis manuscript investigates the nature of intragrain impurity nanoclusters (IGINs) in metal\nhalide perovskite films and proposes a chemical additive strategy for their mitigation. The\nauthors employ a combination of advanced characterization techniques, notably low-dose\nscanning transmission electron microscopy (STEM), and density functional theory (DFT)\ncalculations to identify a metastable A₂BX₄-type impurity phase. While the quality of the\nexperimental data, particularly the atomic-resolution STEM imaging, is commendable, I have\nsignificant concerns regarding the manuscript's novelty and conceptual advance, especially in\nlight of the authors' own recent publication on a highly similar topic (Intragrain impurity\nannihilation for highly efficient and stable perovskite solar cells, Nature Communications,\n2024, 15, 2329).\nMy primary concern is the substantial overlap with their previous work in terms of the research\nsubject, core methodology, and ultimate scientific goal. Both studies focus on IGINs, rely\nheavily on STEM as the central characterization tool, and aim to improve device performance\nby addressing these impurities. While the current work identifies a specific A₂BX₄ phase, this\nfeels more like an incremental progression rather than the fundamentally new discovery that\nwould warrant publication in a top-tier journal. The conceptual framework remains largely the\nsame: identifying and eliminating intragrain impurities.\nFurthermore, several of the key elements presented as innovative in this manuscript are, in fact,\nwell-established in the field. The use of low-dose STEM for characterizing beam-sensitive\nperovskites has been reported in numerous prior studies and can no longer be considered a\nnovel technique in itself. Similarly, the use of inner salt additives to passivate defects and\ncontrol crystallization in perovskites is a widely explored strategy. While the authors apply it\nto mitigate IGINs, the underlying chemical approach lacks fundamental novelty.\nGiven these significant concerns about the manuscript's independent conceptual advance, I\ncannot recommend publication in its current form. The authors must convincingly differentiate\nthis work from their prior publication and from the broader literature.\nThank you for your insightful comments. We appreciate your emphasis on the importance of\ndifferentiating this work from our prior publication and the broader literature. In response, we\nhave significantly revised our manuscript to clearly present the novelty, significance, and rigor\nof our findings. The core novelty of this study lies in the discovery and atomic-level\nunderstanding of a new metastable impurity phase of (FA,Cs)₂PbI₄, identified as a type of\ndetrimental intragrain defect commonly existed in all Cs-incorporated perovskites that has not\nbeen previously reported in the field.\nThis finding is significant on two levels. First, we determined the precise atomic structure of\nthe new matastable phase through multi-view atomic-scale observations, which is not covered\nin other low-dose electron microscopy studies on perovskites. It addresses a critical gap in the\nmicrostructural knowledge of perovskites, encouraging further exploration of unusual defects\nwithin crystalline perovskite grains. Second, we acknowledge the connection to our previously\npublished paper in Nature Communications, in which we examined the intragrain PbI₂ impurity\nnanophase and the intriguing perovskite self-healing behavior. While the mechanisms behind\nthe formation of these intragrain impurity nanophases were unclear at that time, our current\nwork clarifies the role of the (FA,Cs)₂PbI₄ impurity phase by reconstructing the exact atomic\nstructures and elucidating the chemical and physical properties, allowing us to propose a\nplausible chemical mechanism for describing their formation and guiding targeted annihilation.\nSpecifically, we found that the formation of PbI₂ and (FA,Cs)₂PbI₄ impurity pairs results from\nintragrain cation segregation, represented by the reaction: 2(FA,Cs)PbI₃ → PbI₂ + (FA,Cs)₂PbI₄.\nThe healing process is the reverse reaction, driven by intragrain cation homogenization, which\ndeepens our understanding of the importance of cation segregation in hybrid perovskites.\nOur insights into the structure-activity relationship from the perspective of the internal grain\nmicrostructure are instructive, demonstrating that attention to the microstructural information\nwithin the grains is essential for further improving the efficiency of current perovskite devices.\nThus, these two studies complement each other and contribute to a mechanistic understanding\nof the rarely studied intragrain defect phenomena in the field. Additionally, our processing\nmethod, inspired by insights into intragrain cation heterogeneity, is designed to promote cation\nhomogeneity during the solution state. While we acknowledge that the use of inner salt\nmaterials is documented for different purposes, primarily grain boundary passivation, there is\nno prior research indicating their effect on the intragrain microstructure. Notably, the effective\ninner salt utilized in our study has not been reported elsewhere.\nWe appreciate your critical comments regarding the novelty of our work, and we hope this\njustification clarifies the distinct contributions of our study. We have also revised the\nintroduction to ensure that the novelty, significance, and rigor of our work are clearly\narticulated, effectively differentiating it from our previous study and existing literature.\n1. The authors' previous Nature Communications paper already established the existence of\nIGINs and a method for their “annihilation”. The core claim of this new manuscript is the\nidentification of a specific A₂BX₄ phase. Could the authors please clarify why the identification\nof this specific stoichiometry constitutes a fundamental conceptual leap, rather than a more\ndetailed characterization of the previously reported phenomenon? How does knowing the\nphase is A₂BX₄, rather than simply PbI₂ or another impurity, fundamentally change our\nunderstanding or the mitigation strategies in a way that is not incremental?\nThank you for your constructive comments. Two types of nano-scale intragrain impurities, PbI\n2\nand non-PbI and their intriguing healing behavior were revealed in our prior publication.\n2\nchemical mechanism for their formation and annihilation. In this study, we successfully\nreconstructed the atomic structure of the non-PbI phase by combining electron microscopy\n2\nwith image simulation.\nTypically, each atomic-resolution two-dimensional (2D) STEM image can be considered the\nresult of projecting a tree-dimensional (3D) object of a certain thickness along a specific\ndirection. During this process, the symmetry elements of the 3D crystal are also projected into\nthe 2D image. Therefore, by analyzing the symmetry elements in atomic-resolution images of\nthe same phase obtained in different directions, the crystallographic space group of that phase\ncan be determined. Subsequently, the positions of the atomic columns in the HAADF-STEM\nimage can be used to determine the atomic coordinates within the unit cell.\nAccording to the position and intensity distribution of image spots in the minimum translation\nunit of the projection images of non-lead iodide phases in the [001] , [100] , and [110]\nO O O\ndirections, the projection symmetry of the non-lead iodide phase in these directions is c2mm,\np2mm, and p2mm, respectively, as shown in Figure R17. Therefore, its three-dimensional\nspace group is determined to be Cmmm. The impurity phases in perovskites that have been\nreported so far, including lead iodide and the intermediate phases in the preparation process of\nperovskite films, do not have this symmetry. Based on the positions of the atomic columns in\nthe STEM images, we finally determined that the non-lead iodide phase has the atomic\nstructure as shown in Supplementary Fig. 4 in the revised manuscript, and the composition\nwas determined through image simulation method (Supplementary Fig. 5). Based on the\nstructural and compositional analysis of A BX phase, we obtained the information on its\n2 4\nchemical properties and electronic structures, which helps to fully understand the impact of\nA BX on the performance and stability of perovskites. In addition, we also discovered that the\n2 4\nA BX phase is ubiquitous in Cs-incorporated perovskites and confirmed that the A BX phase\n2 4 2 4\nis the main intragranular impurity phase.\nOur quasi-in situ and ex situ results revealed that the matastable A BX phase acts the core of\n2 4\nthe generation of PbI and the induction of further degradation of perovskites. Therefore, this\n2\nwork is not only a more detailed microstructural characterization, but also a leap in the\nunderstanding of the mechanism of intragranular impurity-induced perovskites. Since Cs\ndoping is currently common in high-performance perovskites, it is even more important to\nreveal the A BX phase structure and its impact on perovskites. Based on the above research\n2 4\nresults, we developed a method to regulate Cs homogenization, which illustrates the\nimportance of analyzing intragranular impurities and defects from an atomic structure\nperspective.\nFigure R17 a-c Atomic-resolution HAADF-STEM images of (FA Cs ) PbI phase and d-f\n0.6 0.4 2 4\ncorresponding projection symmetry.\n2. The manuscript repeatedly highlights the use of low-dose STEM. However, this is a well-\nestablished technique for perovskite characterization (e.g., Science 370, 548 (2020), Nature\nCommunications, 2023 14:1940, ACS Energy Lett. 2023, 8, 2620−2629, and a review paper\non low dose TEM: Energychem, 2023,5,100105). Could the authors please specify what is\nuniquely innovative about their application of this technique in the current study that has not\nbeen demonstrated before?\nThank you for your constructive comments. We agree it is necessary to specify the innovative\nuse of low-dose atomic resolution STEM imaging in present work. TEM technology is a\nrecognized and widely used method that can effectively reveal the structure and chemical\ndistribution of materials, especially in materials with perovskite structure and semiconductor\nmaterials. However, compared with traditional oxide perovskite structured materials, such as\nLi Sr Ta Zr O (Nat. Commun. 14,194 (2023)), TEM characterization of halide\n0.375 0.4375 0.75 0.25 3\nperovskites is very challenging for their highly sensitive to electron dose. Therefore, low-\nmagnification STEM imaging and selected area electron diffraction techniques are very\ncommonly used methods in the study of halide perovskite microstructures (Nature Energy 1\n(2016); Nat. Commun. 8, 14547 (2017);Nat. Commun. 10, 2196 (2019);Nature 607, 294-300\n(2022); ACS Energy Lett. 8, 2620-2629 (2023)). However, these methods can only provide\nunderstanding of materials at the nanoscale. In-depth research at the atomic scale requires more\nprecise control of electron beam dose and other imaging parameters.\nAlthough HRTEM has achieved atomic-resolution imaging of organic halide perovskites,\ndirectly interpretable atomic-resolution low-dose STEM imaging of organic halide perovskite\nmaterials was not achieved until the work of Rothmann et al. (Nat. Commun. 12, 5516 (2021);\nScience 370, 548 (2020)). Perovskite films with a [001] orientation distribution prepared by\nthe dual-source evaporation method provide an excellent platform for low-dose atomic\nresolution imaging because the films prepared by this method have fewer defects and better\ncrystallinity than the films prepared by the spin coating method used in practice, and can\neffectively avoid electron irradiation caused by the tilting process of the sample.\nIn present study, we achieved atomic-scale characterization of the perovskite microstructure of\nthe device samples though a method of combining FIB sample preparation with carbon plating\nby optimizing parameters. This method can more realistically reflect the microstructure of\nperovskite grains in actual devices. In addition, we successfully used a series focus imaging\nand three-dimensional reconstruction techniques to accurately locate the position of A₂BX₄\nphase in the sample and its atomic configuration. Furthermore, this study achieved the first\nquasi-in situ atomic-resolution tracking of the microstructural changes of mixed-cation halide\nperovskites in ambient conditions. Unlike previous in situ TEM experiments that used electron\nbeam irradiation to simulate light, we simulated the environmental degradation of devices by\nprecisely controlling the electron dose and tracking the structural changes of the sample in an\natmospheric environment, determining the detrimental effect of the metastable A₂BX₄ phase\non perovskite stability. Therefore, in this work we provide a more comprehensive (three-\ndimensional structure) and more realistic (degradation process) low-dose atomic resolution\ncharacterization method. Our approach to three-dimensional nanostructure reconstruction and\nquasi-in situ atomic resolution tracking of sample microstructure changes is believed to provide\nvaluable insights and ideas for further research, including halide perovskite research.\n3. The use of zwitterionic additives, such as the NDPS reported here, for defect passivation and\ncrystallization control is a common strategy in the perovskite field. The proposed mechanism\nthat interaction with precursor clusters to homogenize the solution and slow crystallization is\nalso a well-understood concept for many additives. What is the specific novelty of using NDPS\nfor IGINs that distinguishes it from the vast body of work on additives for grain boundary\npassivation and bulk defect reduction? Is there a chemical interaction that is uniquely effective\nagainst the formation of the A₂BX₄ phase specifically?\nWe thank the reviewer for raising this important question. Our results demonstrate that the\naddition of NDPS can effectively inhibit the segregation of cations, removing the A₂BX₄ phase\nin the final films. It is reported that the O atom in tetrabutylammonium anion ([TFSI]) is able\nto act as monodentate and bidentate ligands to form dative bonds with alkali metal ions,\nthough the chelate effect (Chem. Commun. 12, 1368-1369 (2004); Energy Environ. Sci. 17,\n8209-8218 (2024)). This prompted us to compare the ability of NDPS and [TFSI] to interact\nwith Cs to reveal the intrinsic mechanism of NDPS in inhibiting cationic segregation. DFT\ncalculations show that the cations (Cs⁺) form bidentate or polydentate complexes by coordinating\nwith the oxygen atoms in both the NDPS and the [TFSI]. Solvent effects of DMF were incorporated\nusing the SMD polarizable continuum model.\nAs the Gibbs free energy presented in the Table R2 indicates, this coordination is an\nexothermic process, which is thermodynamically favorable and contributes significantly to the\noverall energy stabilization of the system. Furthermore, our analysis shows that the DNPS\nmolecule exhibits stronger coordination ability compared to the [TFSI], aligning with our\nmicrostructural analysis. These results clearly indicate that the chelate effect occurs between\nDNPS and Cs+ in precursor solution and can effectively disperse Cs. In the subsequent spin\ncoating process, NDPS as a surfactant can reduce the size of the crystal nucleus and control the\ngrain growth, further benefiting the cation homogeneity.\nWe emphasize that the novelty of this work lies not in the specific use of NDPS, but in\nestablishing precursor coordination as an effective strategy to suppress intragrain impurities.\nBased on this principle, other additives beyond NDPS may also be effective, provided they\nexhibit sufficient coordination affinity toward Cs+ and similar capabilities in modulating\ncrystallization dynamics.\nTable R2 (same as Table R1) The Gibbs free energy for the coordination between Cs+ and\nNDPS/[N₄₄₄₄][TFSI].\nΔG\nSpecies bidentate and polydentate complexes 298K\n(kcal/mol)\nDNPS + Cs+ -12.8\n2*DNPS + Cs+ -23.8\n[N₄₄₄₄][TFSI] +\n-12.1\nCs+\n2*[N₄₄₄₄][TFSI]\n-28.6\n+ Cs+\n4. The manuscript needs to more explicitly and convincingly articulate its unique contribution\nrelative to the authors' own Nature Communications 2024 paper. Beyond identifying a different\nimpurity composition (A₂BX₄ vs. PbI₂) and proposing a different solution (chemical\nsuppression vs. physical annihilation), what is the overarching, new scientific message?\nWithout a clear distinction, this work appears to be a follow-up study that is too closely related\nto the previous one for a separate high-impact publication.\nThank you for your constructive comments. In our previous article, we identified two types of\nintracrystalline impurity phases and studied their healing behavior under external energy input.\nHowever, due to the lack of a clear understanding of the atomic structure of the A BX phase,\n2 4\nwe have an inadequate understanding of the formation and healing mechanism of\nintracrystalline precipitate phases. In addition, our complementary experiments demonstrate\nthat the A BX phase is ubiquitous in Cs-incorporated perovskites and is the primary intragrain\n2 4\nimpurity. Our DFT calculations and quasi-in situ experiments fully demonstrate that the A BX\n2 4\nphase plays a crucial role in the perovskite degradation process.\nThe successful reconstruction of the atomic structure of the A BX phase not only reveals the\n2 4\nself-healing process of the intragrain impurities, but also enables us to understand the\nmechanism by which cation segregation affects the performance of perovskite cells at the\natomic scale. We then proposed, based on first-principles calculations, to add surfactants to\nevenly distribute the cations in the precursor, thereby obtaining a stable perovskite film with\nuniform composition. We also demonstrated for the first time that surfactants can not only\npassivate grain boundaries but also effectively regulate the microstructure inside the grains. We\nhave incorporated these results and related discussion in the revised manuscript as necessary.\nOur current work provides a deeper understanding of the microstructural disorder induced by\ncationic segregation in perovskite crystals, which has not been mentioned in previous works.\nThis is more instructive for the development of stable and high-efficiency perovskite solar cells.\n5. The study focuses on FA-Cs based perovskites. How prevalent is this newly identified A₂BX₄\nphase in other important perovskite systems (e.g., FAPbI , MAPbI , triple-cation,\n3 3\nall-inorganic)? Is this a universal, previously overlooked impurity, or a phenomenon specific to the\ncomposition studied? The claimed broad impact hinges on this point.\nThank you for your valuable suggestions. We agree that the existence of A₂BX₄ phase in\ndifferent composition systems is crucial for clarifying its formation mechanism and the\ncorresponding inhibition strategy. We then conducted a systematic atomic-resolution studies of\nperovskite films of representative compositions including FAPbI , MAPbI , CsPbI and a\n3 3 3\ncomplex composition of ((FA Cs )PbI ) (MAPbBr ) that we used for the fabrication\n0.95 0.05 3 0.975 3 0.025\nof devices with high efficiency. In Figures R18-21, we show atomic-resolution HAADF-\nSTEM images obtained from different samples of these compositions. The presence of the\nA BX phase was observed in ((FA Cs )PbI ) (MAPbBr ) samples and we\n2 4 0.95 0.05 3 0.975 3 0.025\nconfirmed that no A BX was formed in perovskites with a single cation component, including\n2 4\nFAPbI , MAPbI , CsPbI , after observing more than 20 grains from multiple samples, This\n3 3 3\nresult suggests that the tolerance factor defined by , which is often used to\n𝑟𝐴+𝑟𝑋\nevaluate the stability of perovskites, may be one of the important parameters that determine the\n𝜏 = √2(𝑟𝐵+𝑟𝑋)\nformation of A BX phase. Therefore, A BX is a unique intragrain impurity phase in Cs-\n2 4 2 4\nincorporated perovskite systems FA-Cs system perovskite films. Considering that FA-Cs\nsystems is still the main options of efficient and stable perovskite device absorption layer, the\ndetermination of A BX phase structure and its elimination mechanism are of great significance.\n2 4\nFigure R18a-d Atomic-resolution HAADF-STEM images obtained from CsPbI samples.\nFA Cs PbI , FA Cs PbI , FA Cs PbI and (FA Cs PbI ) (MAPbBr ) samples.\n0.5 0.5 3 0.8 0.2 3 0.9 0.1 3 0.95 0.05 3 0.975 3 0.025"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The Authors addressed the points we raised in a satisfactory manner, including providing better context for the work and fixing some other issues.", "reviewer_2": "The authors have added substantial experimental and computational results and have addressed most of my concerns satisfactorily. I have no further questions.", "reviewer_3": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts.", "reviewer_4": "The author has basically responded to the review comments and can be accepted"}, "author_response": ""}]

Decision letter