Indium-Based Octahedra Coordinating Pb–I Termini for Stable Perovskites

Abstract

Halide perovskite degradation initiates at surfaces/interfaces, where defect accumulation disrupts octahedral connectivity and triggers structural collapse. Herein, we design an indium-based octahedral motifs, i.e., (PMA)₄InCl₇ (PIC) (PMA=phenylmethylammonium) featuring disordered octahedra and (PMA)₄NaInCl₈ (PNIC) with ordered octahedral frameworks, as robust coordination agent to bind Pb-I termini at the surface/interface. Distinct from the isolated cluster of PIC, PNIC forms a coherent inorganic framework at Pb-I termini, enabling coherent interfacial coupling with the 3D perovskite via octahedral-level coupling between In(Na)-Cl and Pb-I units. This rigid octahedral coupling immobilizes formamidinium cations and Pb–I octahedra at the surface/interface, suppressing structural freedom. Moreover, this coupling facilitates the relaxation of interfacial structural imperfections while simultaneously establishing favorable band alignment, together governing enhanced carrier transport dynamics. As a result, the resulting perovskite solar cells deliver an efficiency of 27.29% (certified 26.84%), with the minimodules reaching 24% (certified 23.54%). Moreover, these devices demonstrate significant operational stability, retaining over 98% of their initial efficiency after 3,000 h of maximum power point tracking.

Halide perovskites degrade at surfaces where defects break octahedral connectivity and destabilize the structure. Wen et al. design indium-based octahedral compounds that bind lead-iodide, forming rigid interfaces that enhance stability and carrier transport for efficient, durable solar cells.

Introduction

Perovskite solar cells (PSCs) have achieved unprecedented efficiency gains, enabling them as a potential leading photovoltaic (PV) technology^1–4. However, their widespread commercialization is critically hindered by unstable perovskite materials under operational conditions, such as continuous light illumination and thermal stress^5,6. Owing to the soft lattice, polycrystalline perovskites, particularly at surfaces and grain boundaries (GBs), are highly prone to structural degradation^7–9. For example, under photothermal stress, organic component volatilization and halide-vacancy accumulation cooperatively drive irreversible lattice collapse and further device performance degradation^10,11.

Low-dimensional (LD) perovskite capping layers have been widely terminated on three-dimensional (3D) absorbers to temporarily regulate unstable surface and interface lattice^12–15. Despite their effectiveness, conventional Pb-based LD perovskites suffer from inherent material limitations^16,17. The formation of Pb-based LD perovskites typically utilizes residual PbI₂ at the surface/interface as a reactive template for dimensional conversion. Ideally, this PbI₂ is fully consumed to form a stable layered structure. However, any unreacted PbI₂ residue can still undergo photodegradation into metallic Pb⁰ and other deep-level defects^18–20. Moreover, the organic ammonium components in these LD perovskites lack sufficient structural rigidity, rendering the resulting regulation layers highly dynamic and prone to desorption or diffusion into the bulk under continuous operation, finally compromising interface stabilization^21–25.

Herein, we introduce an indium-based octahedral motifs, i.e., (PMA)₄InCl₇ (PIC) (PMA=phenylmethylammonium, C₇H₁₀N⁺) featuring disordered octahedra and (PMA)₄NaInCl₈ (PNIC) with ordered octahedral frameworks, as lattice-anchoring motifs to reconstruct the coordination environment of Pb-I termini at the surface/GBs. Comprehensive results reveal that the dimensionality of the inorganic framework is decisive for surface/interfacial stabilization. PIC consists of isolated [InCl₆] octahedra encapsulated by organic ligands, offering only limited structural constraint. In contrast, the 2D PNIC forms continuous, corner-sharing octahedral sheets that establish coherent octahedral-level coupling between In(Na)-Cl and Pb-I units. This rigid octahedral coupling immobilizes both formamidinium cations and Pb-I octahedra at the surface/interface, effectively suppressing structural freedom. Such coupling not only enables the relaxation of interfacial structural imperfections but also induces favorable band alignment, collectively improving carrier dynamics. Consequently, PSCs with Pb-I termini coordinated by ordered In(Na)-Cl frameworks reach a remarkable PCE of 27.29% (certified 26.84%) with negligible degradation over 3000 h of maximum power point (MPP) tracking, while a 16.1 cm² modules delivers 24% efficiency (certified 23.54%).

Results

Characterization of indium-based octahedral motifs

Figure 1a shows the single crystal structure of PIC (left) and PNIC (right), as determined by single-crystal X-ray diffraction (SCXRD). Structural analysis reveals that zero-dimensional (0D) PIC crystallizes in the monoclinic crystal system with the space group of C2/m space group²⁶. By contrast, two-dimensional (2D) PNIC also adopts a monoclinic lattice but belongs to P2₁/c space group. Detailed crystallographic parameters are summarized in Supplementary Table 1. As shown in the structural diagrams, PIC consists of isolated [InCl₆] octahedra that are fully separated in three-dimensional space by PMA⁺ cations. In contrast, PNIC exhibits a classic 2D bilayer perovskite architecture, in which [NaCl₆] and [InCl₆] octahedra are interconnected via corner sharing to form continuous inorganic sheets, with PMA⁺ cations residing between the layers^27,28.

Fig. 1. Structural characterization of PIC and PNIC perovskite.

a Schematic illustrations of the single-crystal structures of PIC (left) and PNIC (right). (PMA=phenylmethylammonium, C₇H₁₀N⁺). b Experimental PXRD patterns of freshly prepared (PMA)₄InCl₇ (left panel) and (PMA)₄NaInCl₈ (right panel), and after one-year storage under ambient conditions, together with the simulated patterns derived from SCXRD data. a. u., arbitrary units. c Top-view SEM images of FACs, FACs@PIC, FACs@PNIC perovskite films. Scale bar represents 600 nm. (PMA)₄InCl₇ (PIC) and (PMA)₄NaInCl₈ (PNIC).

Powder X-ray diffraction (PXRD) confirms the phase purity of both PIC and PNIC, as their experimental patterns closely match those simulated patterns from the corresponding single-crystal structures (Fig. 1b). Notably, both materials exhibit robust environmental stability: no structural degradation or secondary phases are detected after one year of ambient storage at 25 °C and 40 ± 5% relative humidity (RH), as evidenced by unchanged PXRD patterns. Under combined light and thermal stress (one-sun illumination and 85 °C), the ¹H nuclear magnetic resonance (NMR) results show that the characteristic signals of PMA⁺ diminish (the benzene ring in PMA⁺ presents a chemical shift between 7.3 and 7.6 ppm), while the PIC and PNIC remain unchanged even after 80 h (Supplementary Figs. 1–3), highlighting the stabilizing role of their robust inorganic frameworks. Thermal gravimetric analysis (TGA) further reveals that both compounds exhibit high thermal stability, with onset decomposition temperature exceeding 200 °C (Supplementary Fig. 4). This intrinsic stability, arising from the rigid inorganic frameworks and effective protection by organic cations, highlights the strong potential of these lead-free LD (LFLD) perovskites as robust interfacial functional layers in perovskite optoelectronic devices.

To investigate the surface morphology of 3D perovskite with PIC/PNIC coordination, the isopropanol solution of phase-pure PIC and PNIC was applied to the perovskite films, yielding samples denoted as FACs, FACs@PIC, and FACs@PNIC, respectively. Compared to FACs, both FACs@PIC or FACs@PNIC films exhibit notable changes (Fig. 1c). The addition of PIC and PNIC leads to the formation of a thin interfacial layer with distinct spatial distributions arising from their different octahedral configurations. FACs@PIC films exhibit sparsely distributed nanocrystals, indicating discontinuous coverage due to the 0D nature of PIC and steric hindrance from bulky organic cations surrounding isolated In-Cl octahedra. Moreover, FACs@PNIC films feature a preferential distribution concentrated at the grain boundaries, where the exposed edges of the inorganic layer are tightly bound to the 3D perovskite surface.

Octahedral-level coupling enhanced stability

To resolve the nanoscale interfacial assembly and structural coupling between PIC/PNIC and the perovskite, scanning transmission electron microscopy (STEM) is performed. High-angle annular dark-field (HAADF-STEM) images of the FACs@PNIC film reveal a distinct crystalline capping layer (Fig. 2a, b): the region of 3D FACs shows a ~ 3.2 Å spacing indexed to the (200) plane²⁹, while PIC and PNIC shows well-defined lattice fringes with an indexed ~7.3 Å and ~ 7.4 Å spacing, respectively, consistent with their (001) and (010) planes from SCXRD data³⁰. Although the FACs@PIC film reveals a surface coverage layer, the coherent lattice observed for PNIC evidences a structurally compatible hetero-interface that potentially minimizes lattice distortion and suppresses residual strain in FA-based perovskite.

Fig. 2. Octahedral-level coupling between In(Na)-Cl and Pb-I units.

a HAADF-STEM image of the FACs@PIC perovskite film. Scale bar represents 100 nm. Corresponding STEM image of the enlarged area marked by blue frame (region 1). Scale bar represents 4 nm. b HAADF-STEM image of the FACs@PNIC perovskite film. Scale bar represents 100 nm. Corresponding STEM image of the enlarged area marked by a red frame (region 2). Scale bar represents 4 nm. c ¹H NMR spectra for FAI, PIC, PNIC, FACs@PIC, and FACs@PNIC films, in which FAI, PIC and PNIC powders, as well as FACs@PIC and FACs@PNIC films, were dissolved in d₆-DMSO solution. d Cl 2p core-level spectra for PIC SCs, PNIC SCs, FACs, FACs@PIC and FACs@PNIC. e Pb 4 f core level spectra for FACs, FACs@PIC and FACs@PNIC films. f Schematic illustration of the interfacial stabilization and strain relaxation mechanisms for 3D perovskites capped with 0D PIC and 2D PNIC. g Residual stress evaluated in initial and aged perovskite films for FACs, FACs@PIC, and FACs@PNIC.

Molecular-level interactions between PIC/PNIC and perovskite are further investigated by ¹H Nuclear Magnetic Resonance (NMR) spectra. Both FACs@PIC and FACs@PNIC show high-field shifts of the -NH of PMA⁺ resonance (Fig. 2c), with a larger shift for PNIC, indicating stronger N-H···I hydrogen bonding enabled by the perpendicular PMA⁺ orientation in the 2D framework, which disrupts π-π stacking between adjacent molecules and optimizes the spatial orientation for interaction. Notably, a pronounced -NH shift of FA⁺ is observed for FACs@PNIC, providing evidence that the In(Na)-Cl octahedral inorganic framework immobilizes interfacial FA cations within the interstitial voids of the inorganic framework, creating a high energy barrier against volatilization. In contrast, FACs@PIC shows a much weaker effect. This unique FA-locking capability of PNIC, synergizing with enhanced hydrogen bonding, is pivotal in suppressing ion migration and preserving lattice integrity.

X-ray Photoelectron Spectroscopy (XPS) results are well consistent with the above result (Fig. 2d, e and Supplementary Figs. 5, 6). For the Cl 2p core level, FACs@PNIC exhibits a strong signal with its binding energy shifted to lower values relative to the pristine PNIC single crystal. This negative shift indicates that Cl⁻ ions at the interface undergo halide exchange of fill halide vacancies, thereby altering the local chemical environment of Cl⁻ ions. In contrast, FACs@PIC shows only a weak Cl⁻ signal due to the shielding effect of bulky organic cations encapsulating the isolated [InCl₆] octahedra, which limits the exposure of Cl⁻ at the interface. The In 3d signal also shifts to lower binding energy in FACs@PNIC relative to pristine PNIC single crystal, further confirming that the In(Na)-Cl octahedral framework engages in coordination interaction with the Pb-I terminated surface. Meanwhile, the Pb 4f, I 3d, and N 1s peaks all shift to higher binding energy for both FACs@PIC and FACs@PNIC films, with the largest shifts observed for FACs@PNIC. The positive shifts of Pb²⁺ and I⁻ can be attributed to the formation of coordinate bonds between undercoordinated Pb²⁺ and electron-withdrawing Cl⁻ from the In(Na)-Cl octahedra, which reduces electron density around Pb and I atoms. The N 1s shift to higher binding energy arises from enhanced hydrogen bonding (N-H···Cl and N-H···I) at the interface, where electron density is withdrawn from the nitrogen atoms.

According to the above analyses, strong interfacial coupling between PIC/PNIC and the perovskite is confirmed (Fig. 2f). Thus, the octahedral-level coupling between In(Na)-Cl and Pb-I units reconstructs the strain-sensitive surface/grain-boundary regions and promotes redistribution of the pre-stored residual lattice strain within the perovskite film. Pristine FACs films exhibit substantial residual strain, as evidenced by the pronounced positive slope in the 2θ-sin²(Ψ) plot, which is partially relieved by PIC and nearly eliminated by PNIC (Supplementary Figs. 7–9). To exclude possible solvent-induced effects, we further examined a control film treated with pure IPA (FACs@IPA). The residual strain remains nearly identical to that of the pristine FACs film, confirming that the observed strain relaxation originates from octahedral coupling rather than the solvent treatment (Supplementary Fig. 10). The superior strain suppression achieved with PNIC originates from its continuous inorganic framework, which forms a coherent interface with the 3D perovskite and enables efficient redistributes of lattice stress. In contrast, PIC consists of isolated [InCl₆] octahedra encapsulated by organic cations; steric hindrance limits effective lattice coupling, resulting in weaker strain modulation. Stress evolution during aging further confirms this mechanism: while residual stress in FACs increases markedly, FACs@PNIC remains essentially unchanged (Fig. 2g). These results demonstrate that the continuous octahedral coupling layer affords both mechanical strain relaxation and interfacial chemical stabilization, thereby potentially enhancing photothermal stability.

The octahedral coupling significantly enhances the stability of FACs perovskite films. Photothermal degradation of unencapsulated FA-based perovskite is typically initiated by the volatilization of FA cations, triggering PbI₂ formation²⁹. Under combined thermal and light stress, FACs film shows severe degradation, evidenced by a strong PbI₂ diffraction peak (Supplementary Fig. 11). This degradation is substantially suppressed by the addition of PIC and nearly eliminated by PNIC, suggesting that octahedral coupling between In(Na)-Cl and Pb-I termini effectively stabilize FA cations along the decomposition pathway. Notably, conventional lead-based LD perovskites still undergo significant degradation under identical photothermal stress conditions (Supplementary Fig. 12), highlighting the superior stabilizing effect of the PNIC-induced octahedral-coupled interface.

Enhanced optoelectrical properties by octahedral coupling

Beyond photothermal stabilization, octahedral coupling provides a unified mechanism for healing structural imperfections across the FACs perovskite surface/interface. Steady-state and time-resolved photoluminescence (TRPL) are examined to quantitatively evaluate the suppression of non-radiative recombination of the PIC or PNIC layer. PL intensity increases sequentially from FACs to FACs@PIC and FACs@PNIC films (Supplementary Fig. 13). Consistently, TRPL measurements show a pronounced extension of carrier lifetime from 400 ns (FACs) to 880 ns (FACs@PIC) and further to 1120 ns (FACs@PNIC) (Fig. 3a). PL mapping further reveals that FACs@PNIC exhibits not only the highest emission intensity but also exceptional spatial homogeneity across the scanned area, in contrast to the localized enhancement observed for FACs@PIC (Fig. 3b). The superior surface electronic quality afforded by PNIC is attributed to its 2D nanostructure, which provides abundant inorganic anchoring site that preferentially bind and effectively coordinate with vacancy-rich regions. By contrast, 0D PIC interacts through discrete isolated octahedral units, while steric shielding by the organic ligands limits the efficiency and uniformity of interfacial energetic landscape flattening and trap-state eradication.

Fig. 3. Characterization of optoelectrical properties.

a Time-resolved PL spectra for FACs, FACs@PIC and FACs@PNIC films. b PL mapping results for FACs, FACs@PIC and FACs@PNIC films. Scale bar represents 1 μm. c DLCP results for FACs, FACs@PIC and FACs@PNIC devices before aging. C-AFM results for d FACs, e FACs@PIC and f FACs@PNIC films. Scale bar represents 1μm. g KPFM results for FACs, FACs@PIC and FACs@PNIC. Scale bar represents 500 nm. h Energy level diagram of perovskite with charge transport material.

Defect physics is further elucidated by thermal admittance spectroscopy (TAS). As shown in Supplementary Fig. 14, trap density of state (tDOS) measurements reveal a substantial reduction in deep-level defects (0.2–0.3 eV) for both FACs@PIC and FACs@PNIC, which is critical for minimizing non-radiative recombination losses and enhancing overall device performance (Supplementary Figs. 15–17)³¹. Drive level capacitance profiling (DLCP) confirms that PNIC layer yields the lowest interfacial defect density (1.40 × 10¹⁶ cm⁻³), compared to FACs (3.60 × 10¹⁶ cm⁻³) and FACs@PIC (2.19 × 10¹⁶ cm⁻³) (Fig. 3c). Consistently, the trap-filled limit voltage (V_TFL) derived from space-charge-limited current (SCLC) measurements decreases from 0.13 V (FACs) to 0.11 V (FACs@PIC) and to the lowest value of 0.09 V (FACs@PNIC), consistently indicating a reduction in bulk trap state density (Supplementary Fig. 18). These results collectively confirm that efficient octahedral coupling at the surface/interface establishes a high-quality, low-defect interface, which underpins the subsequent improvements in electronic properties and device performance.

Conductive atomic force microscopy (C-AFM) is used to probe localized electronic properties of the films (Fig. 3d–f). The pristine FACs film exhibits the highest current response, which can be attributed to the high density of surface defects acting as leakage pathways in the dark. In contrast, octahedral coupling markedly suppresses the current, with FACs@PNIC showing the lowest response, consistent with more effective elimination of conductive trap channels and reduced parasitic leakage.

Ultraviolet photoelectron spectroscopy (UPS) and Kelvin probe force microscopy (KPFM) are further employed to examine the interfacial electronic structure. KPFM mapping reveals a monotonic increase in surface contact potential from FACs to FACs@PIC and FACs@PNIC (Fig. 3g). UPS measurements (Supplementary Fig. 19) further show that PNIC shifts both the conduction band minimum (CBM) and the Fermi level (E_F) upward, from −4.49 eV to −4.29 eV and from −4.77 eV to −4.56 eV, respectively. As shown in Fig. 3h, this band realignment enables barrier-free electron extraction by positioning the CBM above the LUMO of (−4.30 eV). The agreement between the highest surface potential and the highest Fermi level for FACs@PNIC confirms the formation of energetically optimized interface for electron extraction. Together, these results demonstrate that the 2D PNIC layer establishes a low-defect, energetically optimized interface that underpins the enhanced photovoltaic performance.

PV performance and stability

To comprehensively assess the PV performance, inverted PSCs with an architecture of ITO/4PADCB/ Perovskite/C₆₀/SnO₂/Au are fabricated. The current density-voltage (J-V) curves demonstrate a significant enhanced device performance upon octahedra coupling effecting (Fig. 4a). FACs device delivers a PCE of 24.67%, which increases to 26.02% for FACs@PIC device, while the FACs@PNIC device achieves a champion PCE of 27.29% (certified 26.84%, Supplementary Fig. 20), representing the highest efficiency reported for LD/3D heterojunction-based PSCs (Supplementary Fig. 21). External quantum efficiency (EQE) spectra for all three device configurations are shown in Supplementary Fig. 22. Consistent with our J-V measurements, the integrated J_sc values from EQE show excellent agreement with the J_sc values obtained from J-V measurements. Notably, FACs@PNIC PSCs exhibit a smaller hysteresis (Supplementary Fig. 23 and Table 2) and a stabilized efficiency of 27.03% under steady-state operation (Fig. 4c and Supplementary Fig. 24). The efficiency improvement is primarily driven by concurrent increases in open-circuit voltage (V_oc) and fill factor (FF) (Fig. 4e), consistent with suppressed non-radiative recombination and improved interfacial charge extraction.

Fig. 4. Photovoltaic performances characterization.

a J-V curves of the FACs, FACs@PIC and FACs@PNIC PSCs. b EQE spectra and integrated J_sc of FACs@PNIC PSC. c Steady-state efficiency of FACs@PNIC PSC. d J-V curve of the champion perovskite module with an active area of 16.1 cm². e PV parameters statistic distribution of FACs, FACs@PIC and FACs@PNIC PSCs. Error bars are the standard deviation of PV parameters for 25 devices in each case. The lines in the middle of the box plots indicate the median values. f Operational stability of FACs, FACs@PIC and FACs@PNIC PSCs under steady-state light illumination and bias voltage for 3000 h in a nitrogen atmosphere at 55 °C. T₉₈ denotes the time required for the device to reach 98% of its initial PCE. g DLCP results for FACs, FACs@PIC and FACs@PNIC before and after aging.

To confirm the charge recombination dynamics, further photophysical characterization of the device is performed. Light intensity dependence of V_oc reveals that FACs@PNIC device exhibits the lowest ideality factor, approaching 1.3 k_BT/q, indicating suppressed trap-assisted recombination (Supplementary Fig. 25). Consistent with this result, dark J-V measurements show the smallest reverse saturation current density, confirming minimized leakage pathways (Supplementary Fig. 26). Electrochemical impedance spectroscopy (EIS) analysis reveals a substantially larger recombination resistance for the FACs@PNIC device (Supplementary Fig. 27), collectively demonstrating that the PNIC-engineered interface is most effective mitigates non-radiative losses while facilitating efficient charge extraction.

The scalability of the octahedral coupling strategy is further evaluated in large-area devices comprising six series-connected sub-cells. The best-performing module achieves a PCE of 24% (certified 23.54%, Supplementary Fig. 28), with a V_oc of 7.21 V, an I_sc of 65.5 mA, and an FF of 82% (Fig. 4d). The high efficiency retention relative to a small-area device highlights the exceptional uniformity and robustness of the octahedral coupling strategy, highlighting its strong potential for scalable and commercially relevant perovskite photovoltaics.

The long-term operational stability of the devices is systematically evaluated under rigorous aging conditions^32,33. Maximum power point (MPP) tracking was conducted on encapsulated devices under continuous 1-sun illumination at ~55 °C (Fig. 4f). The FACs@PNIC device exhibits outstanding operation stability, retaining 98% of its initial efficiency after 3000 h of continuous operation. By comparison, FACs@PIC device maintains 90% of its initial efficiency after 2200 h, whereas the FACs device rapidly degrades to 80% of its initial efficiency within 800 h. Furthermore, the damp-heat stability of encapsulated devices is assessed in accordance with the ISOS D-3 protocol (85 °C and 85% relative humidity). After 1500 h of damp-heat testing, FACs@PNIC device retains 95% of its initial PCE, significantly outperforming the FACs@PIC (85%) and FACs device, which degrades to 60% of initial efficiency after 800 h test (Supplementary Fig. 29).

To directly link macroscopic stability with interfacial integrity, DLCP measurements are performed on devices aged under identical MPP tracking conditions (Supplementary Fig. 30). The evolution of defect density provides a direct microscopic explanation for the observed stability trend (Fig. 4g). After prolonged aging, the defect density in the FACs device increases sharply to 4.53 × 10¹⁶ cm⁻³, indicating severe defect generation under operational stress. In contrast, FACs@PIC device shows a moderate increase to 2.74 × 10¹⁶ cm⁻³. In particular, the FACs@PNIC device preserves its low defect density of 1.42 × 10¹⁶cm⁻³, even after extended aging. Combined material- and device-level analyses reveal that the coherent 2D In(Na)-Cl framework of PNIC establishes continuous octahedral coupling at the interface, enabling synergistic strain relaxation, kinetic immobilization of organic cations, and chemical passivation of reactive sites, thereby suppressing defect generation and accumulation under illumination and thermal stress, eliminating deep-level recombination centers, and ultimately enabling durable, high-efficiency perovskite photovoltaics.

In summary, we present octahedral-level interfacial coordination as a decisive strategy for stabilizing halide perovskites and simultaneously enhancing device performance. By replacing conventional Pb-based LD interlayers with indium-based octahedral motifs, we demonstrate that the dimensionality and connectivity of the inorganic framework govern interfacial strain, trap density regulation, and ion immobilization. In contrast to disordered 0D PIC, the ordered 2D PNIC forms a coherent In(Na)–Cl framework that directly couples with Pb–I termini, locks volatile formamidinium cations, and enables favorable band alignment and charge transport through vertically oriented organic cations. This integrated structural and electronic stabilization translates into significantly enhanced device performance and operational durability. In all, our findings define octahedral modulation as a general design principle for lead-free interfacial materials, offering a pathway toward intrinsically stable, high-efficiency perovskite optoelectronics.

Methods

Materials

Materials

Formamidinium iodide (FAI), Cesium iodide (CsI), lead (II) iodide (PbI₂), methylamine hydrochloride (MACl), (4-(7H-dibenzo [c, g] carbazol-7-yl) butyl) phosphonic acid (4PADCB), C₆₀ and BCP ( > 99% sublimed) were purchased from Xi’an Polymer Light Technology Corp. Benzylamine (99%) was purchased from Adamas. Indium chloride (InCl3, 99.9%) and sodium chloride (NaCl, 99.99%) were purchased from Bide Pharmatech.

N, N-dimethylformamide (DMF, 99.8%), dimethyl sulfoxide (DMSO, 99.7%), chlorobenzene (CB, 99.8%), and isopropanol (IPA, 99.8% anhydrous) are purchased from J&K Scientific Ltd.

Synthesis of LFLD perovskite

(PMA) ₄InCl₇ was prepared by the hydrothermal method. Firstly, the precursor solvent was prepared. At room temperature and normal pressure, benzylamine (214 mg, 2 mmol) and indium chloride (InCl₃, 110.59 mg, 0.5 mmol) were added to a beaker in a 4:1 molar ratio. Then, the precursor powder was added to hydrochloric acid (HCl, 4 ml). Stirring was carried out at a speed of 400 revolutions per minute, and heating was conducted at 100 °C until the solute was uniformly dispersed in the solution. Subsequently, the prepared precursor solution was placed in a quartz-fiber-lined crucible and tightly covered with a quartz-fiber-lined cover. The high-temperature reactor was sealed and placed in an oven set at 120 °C. It cooled at a rate of 5 °C per hour. After the oven naturally cooled to room temperature, the reactor was removed. The excess hydrochloric acid solution in the crucible was poured into a waste container. The product was washed with anhydrous diethyl ether. Finally, the product was dried for 12 h at 60 °C in a vacuum oven to obtain (PMA)₄InCl₇.

The experimental procedure for synthesizing perovskite material (PMA)₄NaInCl₈ through the antisolvent method comprises two sequential steps: First, under ambient temperature and pressure, p-methylaniline (214 mg, 2 mmol), sodium chloride (NaCl, 29.22 mg, 0.5 mmol), and indium chloride (InCl₃, 110.59 mg, 0.5 mmol) are precisely weighed at a molar ratio of 4:1:1 and placed in a beaker. The precursor powders are then dissolved in 4 mL of hydrochloric acid solution under constant stirring at 400 rpm, while heating to 100 °C until complete dissolution of the halide salts yields a transparent and clear solution. Subsequently, crystallization is induced by transferring the solution into a wide-mouth bottle containing anhydrous ether, stored under light-protected conditions. The addition of this antisolvent reduces solute solubility to promote crystal formation. Upon completion of crystallization, the resulting crystals are washed with anhydrous ethanol and dried under vacuum at 60 °C for 3 h, ultimately yielding (PMA)₄NaInCl₈ crystals.

Perovskite precursor solution

A total of 1.63 M perovskite precursor solution was prepared by dissolving 21.17 mg CsI, 265.44 mg FAI, 20 mg MACl, 750 mg PbI₂ in 1 mL mixed DMF and DMSO solvents (volume ratio: 9:1) and stirred in a N₂ glovebox at room temperature overnight before film fabrication.

Perovskite solar cell fabrication

The ITO glass was washed by ultrasonication with water, acetone, and IPA sequentially and then treated with UV-ozone for 15 min before use. The 4PADCB solution with a concentration of 0.3 mg/mL in IPA was spin-coated on the ITO glass at 5000 rpm. for 30 s, followed by annealing at 100 °C for 10 min. 65 μL perovskite precursor was then spin-coated by a two-consecutive step program at 1000 rpm for 5 s and 5000 rpm for 30 s. During the second step, 200 μL CB was dripped on the film. Subsequently, the samples were annealed on a hotplate at 120 °C for 30 min. After that, PIC, PNIC in isopropyl alcohol solution with concentrations (1 mg/mL) was first spin-coated on the perovskite films at 5,000 rpm for 30 s and annealed at 100 °C for 10 min. All steps were performed in a N₂ glove box at room temperature. Finally, C₆₀ (40 nm), SnO₂ (15 nm), and Au (60 nm) were sequentially deposited on top of the perovskite by thermal evaporation.

Perovskite module fabrication

ITO substrates were patterned by a 1064 nm Yb fibre laser with six sub-cells connected in series. The ITO glass was then transferred to an inline cleaning machine to be washed with surfactant-containing deionized water and pure deionized water. The substrate is blown dry with compressed dry air, removed of surface static electricity with an ionizer bar and processed with an inline UV-ozone unit. All layers are identical to those in the small-area PSCs. Finally, C₆₀ (40 nm), BCP (7 nm) and Cu (100 nm) were sequentially deposited on top of the above perovskites by thermal evaporation.

Operation parameters of laser scribing

For P1 lines, the scribing rate was 300 cm s⁻¹, the frequency was 50 kHz, and the power was 3.6 W. As for P2 lines, the scribing rate was 500 cm s⁻¹, the frequency was 50 kHz, and the power was 2.40 W. P3 lines were operated under a scribing rate of 500 cm s⁻¹, a frequency of 50 kHz and a power of 2.28 W.

Material characterizations

The PXRD measurement method involved grinding the single crystals into powders in a mortar, XRD patterns of the products were recorded with an X-ray diffractometer (Empyrean) using Cu Kα radiation and a PIVcel3D 2 × 2 area detector, operating at 40 KV and 40 mA. The diffraction patterns were collected under ambient conditions using a parallel beam geometry and the symmetric reflection mode. E, Bruker) using a Mo Kα rotating anode. Data were collected at 221K, and the collected data were integrated and applied with multi-scan absorption correction using the APEX3 software. The TGA was performed using a TA differential scanning calorimetry and thermogravimetric analyzer (SDTQ600) at a heating rate of 10 °C min⁻¹ from room temperature to 800 °C in N₂ flow using an alumina crucible.

Film characterizations

The top-view SEM images were performed using a S-4800 (Hitachi) field-emission scanning electron microscope in high vacuum mode. TEM was performed via Talos F200X S/TEM (Thermo Fisher Scientific Inc). High-resolution ¹H NMR data were recorded on a Bruker NEO 600 MHz spectrometer at Larmor frequencies V(1H) = 600.35 MHz. Dried d₆-DMSO was used as the solvent and was also used to calibrate isotropic chemical shifts. Grazing Incidence X-ray Diffraction (GIXRD) analysis were carried out with the Rigaku SmartLab diffractometer using Cu Kα radiation. XPS and UPS, respectively, PHI5000 VersaProbe III with an X-ray spot size of 500 μm. Time-resolved photoluminescence (TRPL) was measured by FLS1000 Photoluminescence Spectrometer (Edinburgh Instruments Ltd.) with a 450 nm excitation laser. The KPFM and C-AFM tests were performed via Oxford Cypher ES (Oxford Instruments Inc.) The PL mapping was characterized by a confocal Raman microscope (WITec) with a 532 nm excitation laser and an emission signal was collected in the wavelength range from 765 to 815 nm with a center wavelength of 800 nm.

Device characterizations

J-V characteristics of the devices were measured with a Keithley 2400 source meter under the simulated AM 1.5G illumination (100 mW·cm⁻²) using an Enlitech 3 A light source. For small-area PSCs, the voltage scan rate and scan steps were 50 mV s⁻¹ and 20 mV, and the active area was 0.085 cm². For perovskite solar modules, the voltage scan rate was 200 mV s⁻¹ and the scan step was 50 mV, and the aperture areas were 16.1 cm². The light intensity was calibrated by a SRC2020-KG1 stand Si cell before the test. The dark J–V characteristics for SCLC measurements were recorded for devices with structures of FTO/TiO₂/SnO₂/Perovskite/C₆₀/SnO₂/Au using a Keithley 2400 source meter.

Stability test

For the maximum power point tracking measurements, the initial PCE values of the three devices under standard test conditions were 24.11% for the FACs device, 25.35% for the FACs@PIC device, and 26.60% for the FACs@PNIC device. The maximum power point tracking measurements are conducted on encapsulated devices under light-emitting diode illumination (100 mW cm⁻²) at 55 °C. The thermal stability test of the device, following the ISOS D-3 protocol, was conducted under 85 °C/85% RH dual stress. The stability tests of PIC and PNIC single crystals were conducted by storing the samples in ambient air under controlled conditions: 40 ± 5 % RH at a temperature of 25 °C. Throughout the study, thermal stress refers to a constant temperature of 85 °C, and light stress refers to continuous illumination under one-sun intensity.

Capacitance measurements

All the capacitance-related measurements (TAS and DLCP measurements) were performed 20 by a LCR meter (Keysight E4980A). For TAS measurements, the AC frequency scanning is from 20 Hz to 100 KHz with a δV of 20 mV. Temperature-dependent capacitance measurements were carried out via a customized temperature controlling stage (WUHAN CONGTICAL TECHNOLOGY Co., LTD). The calculation of demarcation energy and trap density of states is detailed in the supplementary text. For DLCP tests, the V is scanning from 0 V to the V_oc for PSCs. A series of variable δV from 20 to 120 mV were used to measure the junction capacitance and acquire the capacitance contribution from the trap states.

TAS measurements

According to the theory of TAS, the frequency at the peak (ω₀) is related to emission rates of a trap state as described in

$$\omega =2v_0\exp \left(-\frac{E_d}{k_{B}T}\right)$$ 1

where v₀ is the attempt-to-escape frequency, T is temperature, k_B is Boltzmann constant. E_d is the energy depth of a trap state. We note that v₀ can be described by ξ₀T², where ξ₀ is a temperature-independent constant. Based on Eq. (1), an Arrhenius equation can be derived as described in

$$\ln \left(\frac{{\omega }_0}{T^2}\right)=\ln \left(2{\xi }_0\right)-\frac{E_d}{k_{B}T}$$ 2

Therefore, we can calculate v₀ based on the intercept of the linear fitting line of the ln(ω₀/T²) – 1/T plot. Based on Eq. (1), a demarcation energy (E_ω) can be defined

$$E_{\omega }=k_{B}T\ln \left(\frac{{2v}_0}{\omega }\right)$$ 3

In addition, the density states of the trap state (tDOS) are described by

$$\text{tDOS}=-\frac{V_{\text{bi}}}{\text{A}q{k}_{B}Td}\frac{\text{d}C}{\text{dln}\left(\omega \right)}$$ 4

where A is the device area, d is the thickness of perovskites, which is ~600 nm for the fabricated films, V_bi is the built-in potential of the device. Then, the peak in the tDOS plot is integrated to get the density of defects (N_t).

DLCP tests

The calculation of the carrier density, including deep trap states, is based on the nonlinear relationship between the change of charges (δQ) and the perturbation AC bias (δV):

$$\frac{\mathrm{\delta }Q}{\mathrm{\delta }V}=C_0+C_1\delta V+C_2{\left(\delta V\right)}^2+\cdot \cdot \cdot$$ 5

With the determination of C₀ and C₁, the carrier density (N) that includes both free carrier density (N₀) and trap density (N_T) at a certain position X from the junction barrier is calculated by:

$$N=-\frac{C_0^3}{2q{\epsilon }_0{A}^2{C}_1}$$ 6

The profiling distance from the junction barrier was calculated by ε₀ε_rA/C₀, which was changed by tuning the V. For each AC bias, an additional offset DC voltage was applied to keep the maximum forward bias constant. The trap density within a certain trap depth range was calculated by: subtracting the total carrier density measured at a larger E_ω (lower AC frequency) with that measured at a smaller E_ω (higher AC frequency).

Reporting summary

Further information on research design is available in the Nature Portfolio Reporting Summary linked to this article.

Author contributions

Y.W conceived the idea. D.Y., X.Y., Y.W., and W.N. supported this work. Y. Wen conducted material characterization, performed stability test, and fabricated as well as characterized the devices. Y.S. synthesized (PMA)₄InCl₇ and (PMA)₄NaInCl₈ single crystal and carried out PXRD and SCXRD experiments. X.W. and D.Z. contributed to PL mapping analysis. B. Li performed the DLCP test and analysis. T.Z. fabricated the small-area device and minimodules. W.N. was actively involved in discussions and collaboratively edited the manuscript. Y. Wen and Y.W. wrote the original draft, and D.Y., X.Y., Y.W., and Y. Wen reviewed and edited the draft.

Peer review

Peer review information

Nature Communications thanks Rong-Jun Xie, Guichuan Xing, and the other, anonymous, reviewer(s) for their contribution to the peer review of this work. A peer review file is available.

Funding

Y.W. thanks National Natural Science Foundation of China (No. 62474157, and 52302315), National Science Fund for Excellent Young Scholars (Overseas), the Zhejiang Provincial Natural Science Foundation of China (LR26F040001), Central Guidance Funds for Local Science and Technology Development Projects (2025ZY01012), Top Talent Project of West Lake Pearl Project, and the talent project of ZJU-Hangzhou Global Scientific and Technological Innovation Center (No. 02170000-K02013017). X.G. Yu thanks the project of the National Natural Science Foundation of China (No. 62025403, U23A20354). X. Wang thanks the project of Huaneng Clean Energy Research Institute Foundation (QNYYJJ-24-07). W. Ning thanks the support from the Gusu Innovation and Entrepreneurship Leading Talent Program (ZXL2023188), the Jiangsu Key Laboratory for Carbon-Based Functional Materials and Devices (Z221311) and the Suzhou Key Laboratory of Functional Nano & Soft Materials, Collaborative Innovation Center of Suzhou Nano Science& Technology, the 111 Project.

Data availability

Source data are provided with this paper. All data generated or analyzed in this study are included in the Article and its Supplementary Information. Additional data are available from the corresponding author upon request. Source data are provided with this paper.

Competing interests

The authors declare no competing interests.

Supplementary information

The online version contains supplementary material available at 10.1038/s41467-026-74107-8.