A self-assembled molecule directs ordered α-FAPbI_3 for n-i-p perovskite solar cells

Domain: physics
Published: 2026-01-10
Editorial outcome: minor_revision

Abstract
Formamidinium lead iodide (FAPbI3) perovskite is a leading candidate for high-efficiency solar cells, but its application is hindered by the kinetic instability of its α-phase crystallization. Here, a highly effective strategy for directing the ordered crystallization of α-FAPbI3 in the n-i-p architecture is developed by incorporating a self-assembled molecule (SAM) into the anti-solvent, unlike conventional phosphate-based SAMs requiring polar protic solvents, the -SCN SAM are processed from non-polar, aprotic solvents compatible with perovskite surfaces. An in-situ formed, self-assembled layer acts as a dynamic template, guiding the top-down crystallization process to yield highly ordered α-phase films. Through a series of in-situ measurements, it is rigorously demonstrated that this SAM treatment suppresses undesirable intermediate phases, accelerates the δ-to-α transition, and strategically slows crystal growth, promoting highly ordered films. This meticulously designed strategy culminates in high-quality perovskite films, enabling single-junction devices with a champion power conversion efficiency (PCE) of 26.18% (certified at 25.67%) and mini-modules with a PCE of 21.70%. These results underscore the power of directed self-assembly in fabricating highly efficient and stable n-i-p perovskite solar cells.

Introduction
Formamidinium lead iodide (FAPbI3) is a leading candidate for high-efficiency solar cells due to its near-ideal bandgap and superior stability. However, its practical application is fundamentally limited by the challenge of its crystallization. The required transition from a non-photoactive δ-phase to the desired photoactive α-phase is kinetically unstable and often incomplete. This results in films plagued by residual δ-phase impurities and high defect densities, which ultimately compromises both device performance and operational stability, making the precise control of the δ-to-α phase transformation the critical challenge.

Numerous efforts have been dedicated to controlling the crystallization kinetics of FAPbI3, aiming to achieve the desired black α-phase and enhance device performance. These strategies include manipulating precursor adducts, optimizing solvent and anti-solvent systems, controlling intermediate phase formation, and employing various deposition techniques. However, a common limitation of these strategies, especially in the widely used one-step method coupled with anti-solvent treatment, is the difficulty in achieving precise control over the crystallization process. While anti-solvents accelerate nucleation, they can also lead to uncontrolled crystal growth. This is further exacerbated by the inherently top-down crystallization process of perovskites, where rapid solvent evaporation from the wet film surface leads to preferential nucleation at the surface, potentially resulting in non-uniformity and a higher density of defects. This top-down crystallization process, while posing challenges, also presents a unique opportunity for manipulating the perovskite film formation. Previous research has shown that top-surface modulation is more effective than traditional additive engineering in reconstructing the perovskite lattice and achieving precise control over crystal quality. By acting as a template, the modified perovskite film surface can induce epitaxial growth and promote the formation of well-crystallized, low-defect perovskite domains. Therefore, it is crucial to develop strategies that effectively leverage this top-down crystallization process to guide the formation of a highly ordered perovskite crystal structure while minimizing defects.

Self-assembled molecules (SAMs) have recently become a powerful tool in perovskite solar cell engineering, especially as hole selective layers (HSLs) in inverted (p-i-n) devices. As HSLs in p-i-n structures, SAMs effectively reduce interfacial recombination and enhance hole extraction by modifying work function, passivating defects, and improving interface contact. However, unlike their established success in p-i-n devices, the use of SAMs to direct perovskite crystallization in conventional (n-i-p) perovskite solar cells remains largely unexplored because inherent structural differences in the n-i-p architecture, particularly the position of the HSL above the perovskite layer, limit the realization of effective SAM anchoring and uniform coverage required for crystallization control.

To overcome this challenge, this work introduces a pioneering top-down strategy that leverages the incorporation of SAMs into the anti-solvent to direct the ordered crystallization of α-FAPbI3 in the n-i-p architecture. We demonstrate that a rationally designed SAM (Synthesis route in Supplementary Fig. 1), which consists of a carbazole core, a flexible butane linker, and a terminal anchoring group, spontaneously forms an ordered, self-assembled layer on the nascent perovskite surface during film formation. This in-situ formed self-assembled layer then acts as a dynamic template, guiding the subsequent top-down crystallization process to yield highly ordered α-phase films (Fig. 1a). Employing a series of in-situ measurements, we rigorously demonstrate that the incorporated SAMs effectively suppress undesirable intermediate phases, accelerate the efficient δ-to-α phase transition, and strategically slow down crystallization, ultimately promoting highly ordered perovskite film growth. A systematic investigation comparing the -SCN group to other anchoring groups (-NH2, -CN, -COOH) confirms that this self-assembly-driven crystallization control is unique to the -SCN variant, owing to its superior interaction with the perovskite lattice. This meticulously designed strategy culminates in high-quality, ordered perovskite films that exhibit significantly enhanced efficiency and stability, yielding a champion power conversion efficiency (PCE) of 26.18% (Certified PCE=25.67%) for single-junction devices and 21.28% for mini-modules, thus demonstrating the significant potential of SAMs in controlling perovskite crystallization within the n-i-p architecture.

Results
Self-assembled molecular templates for ordered perovskite growth
9-(4-thiocyanatobutyl)−9H-carbazole (Cz-SCN) (denoted as SAM in Figs. 1 and 2) was successfully synthesized and incorporated into the anti-solvent (optimal 10 mmol/L in Ethyl acetate (EA), Supplementary Fig. 2, Supplementary Note 1) to manipulate perovskite crystallization. Cross-sectional scanning electron microscopy (SEM) images (Fig. 1b) reveal a significant difference in perovskite film morphology upon SAM treatment (w/SAM). The irregular crystals observed in the control film are replaced by monolithic grains spanning the entire film thickness, suggesting enhanced crystal growth and a reduction in grain boundaries. Top-view SEM images (Fig. 1c) and grain size distribution histograms (Supplementary Fig. 3) further confirm this enhancement, showing an increase in average grain size from 0.5 to 1.3 μm and a complete suppression of residual PbI2 impurities in the w/SAM film.

Furthermore, atomic force microscopy (AFM) images reveal a significant decrease in surface roughness for the w/SAM perovskite films, from 23.9 to 18.8 nm (Supplementary Fig. 4). Collectively, this transition to larger, phase-pure grains with fewer boundaries and a smoother surface creates a highly ordered morphology. This is crucial for superior device performance as it minimizes defect-driven recombination losses while providing unobstructed pathways for efficient charge transport and extraction. Grazing-incidence wide-angle X-ray scattering (GIWAXS) measurements performed at different incidence angles offer further insights into how SAM affects perovskite crystallization (Fig. 1d and Supplementary Fig. 5). The w/SAM perovskite film exhibits a pure black α-phase with no detectable PbI2 signals, compared to the control counterpart. Notably, the w/SAM film displays a distinct signal at qz = 0.13 Å⁻¹ along the direction perpendicular to the substrate, corresponding to a lamellar spacing of approximately 4.5 nm (Fig. 1e and Supplementary Fig. 6). This observation suggests the presence of ordered lamellar structures within the film, potentially arising from the self-assembly of SAM molecules. Increasing the concentration of the SAM reveals additional lamellar stacking peaks at qz = 0.13, 0.35, 0.57, and 0.79 Å⁻¹ (Supplementary Fig. 7). This indicates a preferential edge-on orientation of SAMs relative to the substrate, facilitating efficient vertical charge transport across the perovskite film. To further investigate the spatial distribution of the SAM within the perovskite films, we employed time-of-flight secondary-ion mass spectrometry (ToF-SIMS) (Fig. 1f, g). The SIMS depth profiles, acquired in positive ion mode, revealed characteristic fingerprint ions, including NH4+ (m/z=18), FA+ (m/z=45), GA+ (m/z=69), Sn+ (m/z=120), Pb+ (m/z=208), and a fragment of SAM (Cz+, m/z=225, corresponding to the carbazole moiety after loss of SCN). These fingerprint ions provided direct evidence for the presence of SAM and perovskite components throughout the film. Furthermore, the cross-sectional ToF-SIMS images confirmed a uniform distribution of Cz-SCN throughout both the bulk and surface of the perovskite film. This homogeneous distribution suggests that SAM effectively permeates the perovskite layer during film formation, potentially influencing crystallization throughout the entire film. It should be noted, however, that the apparent uniformity in ToF-SIMS depth profiles mainly reflects fragment distribution generated during sputtering, together with depth resolution limits that can smear sharp interfacial features and broaden organic signals. Consistent with this interpretation, UPS measurements show a clear increase in work function at the Cz‑SCN‑modified top surface but negligible change at the buried perovskite interface, indicating minimal SAM accumulation there (Supplementary Fig. 8). In addition, SEM analysis of the buried interface reveals improved crystallinity, which is attributed to more complete perovskite crystallization induced by SAM (Supplementary Fig. 9). Ultraviolet–visible (UV–vis) absorption spectra indicate that SAM treatment does not alter the bandgap of the perovskite films (Supplementary Fig. 10).

In-situ probing of the SAM-directed crystallization pathway
Figure 2 provides a comprehensive analysis of the perovskite crystallization process using in-situ GIWAXS, in-situ photoluminescence (PL) spectroscopy, and in-situ UV–vis absorption spectroscopy, revealing the profound impact of the SAM on nucleation, crystal growth, and phase transformation. In-situ GIWAXS measurements (Fig. 2a, b and Supplementary Note 2) reveal distinct crystallization pathways for the control and w/SAM perovskite films. During spin-coating, after the anti-solvent dripping step, the control film exhibits signals corresponding to the hexagonal 2H (δ-phase, qz = 0.80 Å−1), hexagonal 4H polycrystal (qz = 0.84 Å−1), and 3C (cubic) phase (qz = 1.0 Å−1), indicating a heterogeneous mixture of intermediate phases. In contrast, the w/SAM perovskite film predominantly shows the 2H phase, suggesting that SAM influences the initial perovskite nucleation, promoting a more uniform intermediate phase. During annealing, the in-situ GIWAXS data reveals two distinct stages: Stage I, representing the δ-to-α phase transition as residual solvents evaporate, and Stage II, corresponding to the growth and stabilization of the α-phase perovskite. In the control film, the δ-to-α phase transition is incomplete, with residual hexagonal 6H polytype (qz = 0.87 Å⁻¹) persisting even after 260 s of annealing (Fig. 2a). The presence of MAPbI3·2DMSO during Stage I further complicates the crystallization process and hinders the complete phase transformation, as our previous report. In contrast, the w/SAM perovskite film exhibits complete conversion to the α-phase within 35 s (Fig. 2b). Intriguingly, a strong signal appears at qz = 0.13 Å⁻¹ during Stage II, indicative of ordered lamellar structures formed by SAMs. During the early stage of annealing, rapid solvent loss and preliminary crystallization produce a disordered grain structure. Continued thermal input drives the SAM toward an ordered arrangement that guides crystal growth. This spatiotemporal evolution, driven by thermodynamic stabilization and facilitated by molecular diffusion, explains why long-range lamellar stacking appears only after 20–30 s of annealing, as evidenced by Bragg diffraction spots in the GIWAXS pattern. These results indicate that SAM-assisted crystallization requires both thermal activation and structural relaxation before exerting its templating effect on perovskite grain growth. This observation confirms that SAM self-assembly occurs concurrently with the α-phase transition, suggesting a templating effect on crystal growth. This process likely disrupts the face-sharing octahedral structure (6H) of the δ-phase, facilitating the transition to the corner-sharing architecture (3C) of α-FAPbI3. The formed surface crystallization of α-FAPbI3 serves as a template, propagating the phase transition towards the bulk of the film. The time evolution of peak intensity ratios between the α-phase and δ-phase (α/δ) further highlights the impact of SAM. The steeper slope for the w/SAM film indicates a faster and higher transformation rate from δ to α-phase (Fig. 2c). Moreover, the near-linear evolution of the α/δ ratio for the w/SAM film suggests a more direct conversion pathway, whereas the staged evolution observed in the control film indicates competing transformation processes, involving both δ-to-α phase transition and the conversion of MAPbI3·2DMSO to the α-phase. Density functional theory (DFT) calculations provided further insights into the thermodynamics of perovskite crystallization and phase transition (Fig. 2d). The presence of SAM significantly reduced the surface energy of α-phase perovskite (from 3.32 to 1.45 eV), thermodynamically driving the preferential formation of the black phase (Supplementary Fig. 11 and Supplementary Note 3).

In-situ PL measurements reveal complex nucleation and growth dynamics of the control perovskite film (Fig. 2e and Supplementary Note 4). Upon antisolvent dripping, an intense, broad PL peak emerges at ~700 nm, attributed to Methylamine-rich (MA-rich) nanocrystal nuclei. This peak rapidly redshifts to ~735 nm within 8 s as FA-rich nuclei form and grow. The consistently large full-width at half-maximum of the emission throughout this initial stage indicates a highly polydisperse size distribution of these luminescent species. The subsequent annealing process is marked by three distinct and dramatic stages. As shown in the enlarged PL spectrum in Fig. 2g, during the initial 15 s, the PL intensity undergoes pronounced oscillations as it redshifts towards 760 nm. This heterogeneous behavior reflects a competition between PL enhancement from crystalline growth and severe quenching caused by defect generation during the disordered coalescence of nanocrystals. Second, the intensity becomes severely quenched for an extended period (~55–100 s). This prolonged PL quenching, consistent with the incomplete δ-to-α phase transformation observed in the in-situ GIWAXS data, is attributed to the presence of the photoinactive δ-phase, which quenches the α-FAPbI3 luminescence. Finally, after 60 s annealing, a gradual PL recovery occurs as grains coarsen. However, the PL intensity does not fully recover its initial peak value due to the persistent presence of defects and grain boundaries in the fully annealed film. The crystallization dynamic of the w/SAM film presents a stark contrast, indicative of a highly ordered process (Fig. 2f, h). Upon antisolvent treatment, no MA-rich emission is observed. Instead, a single, weaker PL peak appears at ~680 nm, which can be attributed to an intermediate phase involving the SAM and perovskite precursors, consistent with our in-situ UV–vis absorption data as discussed below. During annealing, the PL emission undergoes a smooth and continuous redshift while its intensity steadily intensifies, tracking the formation and growth of α-FAPbI3. This controlled evolution likely stems from a “slow-release” effect, where SAM molecules temporarily chelate with perovskite precursors, retarding their immediate reaction and facilitating a more orderly crystallization process. Crucially, the redshift in the w/SAM film initiates earlier (5 s) than in the control film (7 s), indicating that the SAM effectively lowers the formation energy barrier of the α-phase and accelerates the δ-to-α transformation (enlarged spectrum in Fig. 2h). Unlike the control, the w/SAM film exhibits only a single, shallow dip in PL intensity, which corresponds to the minimal non-radiative energy loss during the orderly phase transition. Subsequently, the PL intensity increases monotonically and surpasses its initial value, a direct consequence of the high-quality, defect-minimized crystal lattice formed under the guidance of the SAM monolayer.

In-situ UV-vis absorption spectroscopy (Fig. 2i, j) provides further insights into the crystallization process, particularly the role of the antisolvent in promoting nucleation. Immediately after antisolvent dripping during spin-coating, the control film exhibits several strong absorption peaks in the 410–440 nm range, attributable to various intermediate species present in the perovskite precursor solution. As the solvent evaporates and nucleation proceeds, multiple strong absorption peaks appear between 450 and 500 nm, due to the formation of multiple types of nuclei. In contrast, during the spin-coating process, the w/SAM perovskite film initially shows a single, strong absorption peak at 420 nm, corresponding to PbI4 2−• SAM intermediate phase. This suggests that SAM preferentially adsorbs onto the Pb-I framework, displacing iodide ions and promoting a more homogenous intermediate phase. As spin-coating continues, a prominent peak appears at 460 nm in the w/SAM perovskite film, assigned to the δ-phase perovskite, consistent with the formation of a single dominant type of nucleus, as also observed in the GIWAXS data. During the initial annealing stage (Stage I, Fig. 2k, l), the absorption onset of the w/SAM perovskite film redshifts earlier (5 s after annealing) compared to the control film (7 s). This faster redshift, indicating an accelerated δ-to-α phase transition, is consistent with the in-situ PL observations and is attributed to Cz-SCN reducing the energy barrier for α-phase formation. Moreover, as highlighted in Fig. 2k, l the w/SAM perovskite film exhibits a near-linear redshift from 420 to 750 nm, suggesting an ordered and gradual δ-to-α phase transformation guided by the SAM. In contrast, the control film undergoes a multi-stage redshift, indicative of a disordered phase transformation arising from the presence of multiple nuclei species. Furthermore, the delayed emergence of the long-wavelength absorption beyond 770 nm in the w/SAM perovskite film, compared to the control, suggests that SAM retards the crystallization of α-FAPbI3 during Stage II. This slower, more controlled crystallization is conducive to the formation of high-quality perovskite films. We speculate that Pb2+-SCN coordination is the dominant driving force for α-phase stabilization and templated crystallization, while π-π stacking among carbazole backbones cooperatively sustains ordered packing and enhances the robustness of the templating effect.

Mechanism of anchoring group-dependent crystallization control
To elucidate whether the successful crystallization control demonstrated by Cz-SCN is a general property of this molecular class or a specific function of this molecule, three analogous Cz-molecules (Cz-NH2, Cz-CN, and Cz-COOH) are investigated. A comparison of the resulting films reveals a stark contrast. GIWAXS and SEM analysis show that films treated with these alternative molecules not only suffer from residual PbI2 impurities but also completely lack the self-assembled lamellar structure unique to the Cz-SCN case (Fig. 1d, Supplementary Figs. 4 and 12–14). This pivotal finding confirms that the ability to simultaneously ensure phase purity and drive self-organization is a capability exclusive to the Cz-SCN molecule. Quantitative GIWAXS analysis further demonstrates this distinction, showing that the Cz-SCN film exhibits the strongest and sharpest (100) diffraction peak together with the highest (100)/PbI2 intensity ratio, confirming its superior ability to suppress PbI2 formation and promote highly ordered crystallization (Supplementary Fig. 15). To understand the fundamental origin of this unique ability, a synergistic investigation is undertaken to dissect the underlying molecule-perovskite interactions. DFT calculations provide the initial theoretical analysis. By calculating the adsorption energy (Eads) of each molecule onto a perovskite surface with an iodine vacancy, we found that Cz-SCN possesses a remarkably strong adsorption energy of −0.77 eV, substantially more favorable than its counterparts (Fig. 3a). This superior energetic stabilization, rooted in the large dipole moment and highly negative electrostatic potential of the molecule, provides the robust anchoring force necessary to form an ordered template and resist impurity formation (Fig. 3b). Fourier-transform infrared spectroscopy (FTIR) further tracks changes in the chemical bonds of the anchoring groups (Fig. 3c). The S-C≡N stretching vibration in Cz-SCN undergoes a substantial blue shift of 10 cm−1 upon interaction with the perovskite, indicative of a strong coordination bond. In contrast, the other groups exhibit weaker interactions, reflected by smaller blue shifts for the C≡N (3 cm⁻¹) and C=O (5 cm⁻¹) stretches, while the N-H stretch remain effectively unchanged. XPS analysis of the Pb 4f peak (Fig. 3d, Supplementary Fig. 16 and Supplementary Note 5) for the Cz-molecule-treated films reveal a downshift in binding energy, indicative of a Lewis acid-base interaction with undercoordinated Pb2+ ions. Cz-SCN induces the largest energy shift (0.41 eV), significantly greater than the shifts from Cz-COOH (0.30 eV), Cz-CN (0.19 eV), and Cz-NH2 (0.09 eV). To further substantiate this observation, XPS measurements of the relevant core levels (N 1s and O 1s) reveal clear binding energy shifts upon SAM introduction (Supplementary Fig. 17 and Supplementary Note 6). Among them, Cz-SCN shows the largest shift (N 1s: +0.46 eV), indicative of strong Pb²⁺-SCN coordination. The varying interaction strengths are corroborated by complementary characterizations. UV-Vis absorption spectroscopy (Supplementary Fig. 18) shows that Cz-SCN treatment most effectively suppresses the characteristic absorption of PbI2-related species, which is visually confirmed by the bleaching of the yellow PbI2 film to transparent (Supplementary Fig. 19). Consistently, XRD analysis (Supplementary Fig. 20) reveals that the PbI2 diffraction peak at 12.8° disappears in films prepared from Cz-SCN/PbI2 solutions, whereas control and other Cz-analogue samples still exhibit residual PbI2 signals. These results collectively suggest that PbI2 is not preserved as a separate phase but is rapidly consumed, most likely through the formation of a transient intermediate. Collectively, although the four Cz derivatives exhibit distinct aqueous pKa values (Supplementary Table 1, Supplementary Note 7), the direct influence of acidity/basicity on crystallization is limited in the non-aqueous DMF/DMSO precursor environment, where coordination strength with undercoordinated Pb2+ dominates.

The profound consequences of this superior, Cz-SCN-guided crystallization are evident in the optoelectronic properties and formation kinetics of the film. The Cz-SCN-treated film exhibit the lowest trap-state density (Nt) in both hole-only or electron-only devices, the values approximately 60% lower than the control and markedly superior to the other modified films (Supplementary Figs. 21 and 22, Supplementary Tables 2–3, and Supplementary Note 8). These findings on static trap states are further corroborated by an analysis of charge carrier dynamics using time-resolved photoluminescence (TRPL) spectra (Supplementary Fig. 23 and Supplementary Note 9). The Cz-SCN-based film display a substantially longer average PL lifetime (τave) of 1980.9 ns, a 2.9-fold increase compared to the control film (677.4 ns) (Supplementary Table 4). The reduced trap density and extended carrier lifetime demonstrate that the Cz-SCN treatment effectively suppresses non-radiative recombination. To quantitatively probe the crystallization kinetics during thermal annealing, the temporal evolution of PL intensity at 770 nm was monitored in-situ (Fig. 3e, f). All as-deposited films initially exhibit a transient PL peak within the first 25 s of annealing, where the brief emission from nascent crystallites is rapidly quenched by defects as non-radiative recombination pathways. The subsequent recovery and rise in PL intensity directly track the growth and perfection of the emissive α-phase crystalline domains. The control, Cz-NH2, and Cz-CN films display a slow and inefficient PL recovery to low final intensities, indicating a crystallization pathway plagued by defect formation. While Cz-COOH offers a modest improvement, the Cz-SCN film exhibits a remarkably superior kinetic evolution. It demonstrates a rapid and sustained PL recovery to a final intensity far surpassing all other samples. This steep rise and high final efficiency provide compelling dynamic evidence that the Cz-SCN-guided process most effectively suppresses non-radiative recombination during annealing, yielding a film of the highest electronic quality. Ultimately, this well-ordered, defect-minimized interface, formed by the strong Cz-SCN dipole, also establishes the most favorable energy level alignment, upshifting the perovskite valence band maximum to an optimal −5.18 eV for barrier-free hole extraction, as confirmed by Ultraviolet photoelectron spectroscopy (UPS) (Supplementary Fig. 24).

Photovoltaic performance and long-term stability
N-i-p type PSCs are fabricated, employing the structure of glass/FTO/CBD-SnO2/perovskite/Spiro-OMeTAD/Au. A comparative evaluation of devices incorporating the various Cz-molecules reveal the clear superiority of the Cz-SCN treatment (Supplementary Fig. 25 and Supplementary Table 5). While control devices yield an average PCE of 23.5 ± 0.7%, the Cz-SCN-based cells achieve a significantly higher average PCE of 25.8 ± 0.4% (Supplementary Figs. 26 and 27). The champion device featuring the Cz-SCN-modified perovskite layer deliver a remarkable PCE of 26.18%, driven by an enhanced open-circuit voltage (VOC) of 1.190 V and a high fill factor (FF) of 83.65% (Fig. 4a). This performance is confirmed by a stabilized power output (SPO) of 25.9% (Supplementary Fig. 28). A device sent for external certification achieved a PCE of 25.67%, which is among the top certified efficiencies reported for n-i-p architectures (Supplementary Fig. 29 and Supplementary Table 6). The integrated short-circuit current densities (JSC) from IPCE measurements (Fig. 4b) of PSCs are within 5% mismatch of the J-V results. The enhanced VOC and FF in Cz-SCN-based devices are attributed to a synergistic improvement in device physics. Mott-Schottky and Electrochemical impedance spectroscopy (EIS) measurements collectively confirm that the Cz-SCN treatment enlarges the built-in potential, which both facilitates more efficient charge extraction and suppresses carrier recombination (Supplementary Figs. 30 and 31, Supplementary Tables 7–8, and Supplementary Notes 10–11). To demonstrate the scalability of this approach, 5 cm × 5 cm mini-modules comprising six series-connected subcells are fabricated. The champion mini-module, with an active area of 14.21 cm2, achieved a high PCE of 21.7%, with a VOC of 7.024 V, a JSC of 4.01 mA cm−2, and an FF of 77.05% (Fig. 4c). Such efficiency is lower than that of small-area devices due to inevitable series resistance and inactive area losses in module fabrication.

The practical viability of the Cz-SCN treatment is validated through a series of rigorous long-term stability tests on unencapsulated devices. The Cz-SCN-treated PSC demonstrate exceptional environmental robustness, retaining 95% of its initial PCE after 1500 h in ambient air (23 ± 3°C, 50 ± 10% RH). This stands in stark contrast to the control device, which catastrophically degrade to just 60% of its initial performance under the same conditions (ISOS-D-1, Fig. 4d). This enhanced humidity stability stems from the superior hydrophobicity of the Cz-SCN-treated surface, where the dense, uniform monolayer effectively shields the perovskite from moisture ingress (Supplementary Fig. 32). Furthermore, the Cz-SCN treatment impart outstanding photostability. Under continuous 1-sun illumination in a nitrogen atmosphere, the device maintains 90% of its initial PCE after 2000 h, outperforming all others devices (ISOS-L-1, Fig. 4e). Even more impressively, during continuous maximum power point tracking (MPPT) for 1000 h, the Cz-SCN device retain a remarkable 92% of its initial efficiency; in comparison, the control device decay to 80% of its initial value after only 450 h (ISOS-LC-1, Fig. 4f). Post-aging SEM and XRD analyses further confirm that Cz-SCN treated devices retain more compact morphology and strong α-phase crystallinity with negligible PbI2 formation, in sharp contrast to the degraded control counterpart (Supplementary Figs. 33 and 34). These collective results unequivocally demonstrate that the self-assembly-directed, ordered crystallization induced by Cz-SCN is a highly effective strategy for producing n-i-p solar cells with the exceptional efficiency and long-term stability required for practical applications.

Methods
Materials
The reagents for Cz-SCN, Cz-COOH, Cz-CN, or Cz-NH2 synthesis were purchased from Aladdin and Energy Chemical (China). Formamidinium iodide (FAI) was sourced from Greatcell (Australia), while other reagents such as methylammonium chloride (MACl), cesium iodide (CsI), methylammonium bromide (MABr), lead bromide (PbBr2), Spiro-OMeTAD, lithium bis(trifluoromethanesulfonyl)imide (Li-TFSI), and FK209-Co(III)-TFSI were obtained from Xi’an Yuri Solar Co. (China). Lead iodide (PbI2) was supplied by TCI Shanghai (China). Solvents, including N,N-dimethylformamide (DMF), dimethyl sulfoxide (DMSO), chlorobenzene (CB), ethyl acetate (EA), isopropanol (IPA), acetonitrile (ACN), and 4-tert-butylpyridine (TB), were purchased from Thermo Fisher Scientific (USA). Stannous chloride (SnCl2·2H2O), thioglycolic acid (TGA), and urea were acquired from Sigma-Aldrich (USA). All reagents had a purity of at least 99%.

Device fabrication
The glass substrates with pre-patterned fluorine-doped tin oxide (FTO) were thoroughly cleaned. Ultrasonication was employed, with the substrates being successively immersed in detergent, deionized (DI) water, and ethanol, each immersion lasting for 20 min. After that, prior to utilization, the purified FTO-coated glass underwent a 15 min UV-ozone treatment. The electron-transporting layer consisted of a dense SnO2 film. This film was deposited on the FTO layer via a chemical bath deposition (CBD) method. The CBD solution was formulated as follows: per 100 mL of deionized (DI) water, 1.25 g of urea, 1.25 mL of HCl, 25 μL of TGA, and 275 mg of SnCl2·2H2O were mixed. The FTO-coated glass was then submerged in a diluted version of this mixed solution at 90°C for 3 h. Subsequently, it was cleaned with DI water and IPA through sonication for 5 min, and thereafter annealed at 170°C for 1 h.

A 1.5 M perovskite precursor, FA0.98MA0.02Pb(I0.98Br0.02)3, was prepared. To achieve this, 252.79 mg of FAI, 711.56 mg of PbI2, 33.76 mg of MACl, 3.36 mg of MABr, and 11.01 mg of PbBr2 salts were dissolved in 1 mL of a DMF/DMSO (8:1, v/v) mixed solvent. A volume of 50 μL of the perovskite precursor solution was then dropped onto the CBD-SnO2 film. Consecutive spin-coating was carried out, first at 1000 rpm for 10 s and then at 5000 rpm for 30 s. During the last 10 s of spin-coating, 120 μL of ethyl acetate (EA) was added as an anti-solvent to the center of the perovskite film. The Cz-SCN, Cz-COOH, Cz-CN, or Cz-NH2 solution (3 mg/mL in EA) was then applied. The formed FAPbI3 film was annealed at 100°C for 60 min. After this, a Spiro-OMeTAD layer was prepared. It had a concentration of 0.059 mmol mL⁻¹ in chlorobenzene (CB) and was doped with Li-TFSI (0.033 mmol), 4-tert-butylpyridine (0.205 mmol), and FK209-Co (III)-TFSI (0.0057 mmol). This layer was spin-coated onto the perovskite film at 3000 rpm for 30 s. Finally, an 80 nm gold electrode was thermally evaporated onto the film, followed by a 100 nm MgF₂ coating on the glass. A perovskite module composed of six series-connected sub-cells was constructed on a 5 cm × 5 cm glass/FTO substrate. The interconnection were carried out by laser scribing at a wavelength of 532 nm. The active area of the module is 14.21 cm2. The geometric fill factor (GFF) calculation result is GFF = (4.0 cm − 5 × 680 μm)/4.0 cm = 91.4%. The line widths of P1, P2, and P3 were 30 μm, 200 μm, and 150 μm, respectively, with inter-line spacings of 100 μm (P1-P2) and 200 μm (P2-P3).

Material characterization
1H-NMR and 13C-NMR were conducted in d6-DMSO using a Bruker 400 MHz instrument. High-resolution mass spectrometry (HRMS) tests are performed by Thermo Scientific Q Exactive Orbitrap instruments to provide accurate molecular mass measurements. Fourier Transform Infrared Spectroscopy (FTIR) measurement was performed by the INVENIO S Fourier Transform Infrared (FTIR). Measuring 64 scans per sample, from 4000 to 400 cm−1 with a resolution of 2 cm−1. Thermogravimetric analysis (TGA) was performed on Shimadzu DTG-60H at a heating rate of 10°C min−1 under a nitrogen atmosphere, respectively. The absorbance spectra were measured by a UV–vis spectrophotometer with an integrating sphere (PerkinElmer, Lambda 950). 1H-NMR, 13C-NMR spectra, HRMS tests, and TGA characterization of Cz-molecules were showed in Supplementary Figs. 35–53.

Film characterization
SEM measurement was conducted using the JEOL JSM-7800F. Grazing-incidence wide-angle X-ray scattering (GIWAXS) experiments were conducted at the BL14B1 beamline of the Shanghai Synchrotron Radiation Facility, with samples prepared following the same film preparation procedure in the SSRF chemistry lab. Time of flight secondary ion mass spectrometry (ToF-SIMS) measurement. The instrument used was the Focused Ion Beam-scanning electron microscope at Westlake University. The in-situ PL spectra were recorded using an ISPL-HI001 instrument equipped with a 445 nm laser source (Nanjing Ouyi Optoelectronics Technology). In situ UV–vis absorption spectra were recorded using a setup where Y-type optical fibers were connected to the glovebox through a flange interface. A Hamamatsu EQ-99-FC served as the white light excitation source positioned outside the glovebox, while an Ocean Optics QE-PRO spectrometer was employed to monitor the real-time absorption changes during both the spin-coating and annealing stages. The surface morphology of the perovskite film was analyzed using AFM with the Park XE7 instrument. XPS measurements using the Thermo-Fisher ESCALAB 250Xi system equipped with a monochromatized Al Kα X-ray source were utilized. X-ray diffraction (XRD) measurements were acquired using a Smartlab 3KW diffractometer (Rigaku Corporation, Japan). Space charge limited current (SCLC) measurements were performed using the Keithley 2450 SMU. Time-resolved photoluminescence (TRPL) measurements were obtained using the DeltaFlex TCSPC system from Horiba, with data collected at 820 nm under 530 nm excitation. UPS measurements were conducted using the PHI 5000 Versa Probe III instrument, equipped with a He I source (21.22 eV).

Device characterization
The current density-voltage (J-V) characteristics of the devices were measured by a Keithley 2400 source meter under a Class AAA sunlight simulator (XES-40S3, SAN-EI). The light intensity was calibrated to 100 mW cm−2 by a National Institute of Metrology (NIM) certified standard silicon reference cell (QEB1) at AM 1.5G. Incident photon to electron efficiency (IPCE) testing was accomplished using a QER-900AD system (Nanjing Ouyi Photonics Technology) with a NIM-calibrated Hamamatsu Model S1337-1010BQ silicon detector for response correction. EIS and Mott-Schottky analysis were accomplished with the aid of a CHI760E electrochemical workstation (Chenhua Instruments, China). SPO tests were evaluated by biasing the device at the maximum power point (MPP) determined by its J-V curve and recording the steady-state current; the steady-state photoelectric conversion efficiency was calculated from the product of the applied voltage and the stabilized current density. Device humidity stability testing was conducted at 23 ± 2°C and 40 ± 10% relative humidity, during which a Class AAA sunlight simulator was used to periodically monitor performance changes. In addition, MPPT tests were conducted under 1 sun irradiation in a nitrogen atmosphere at 20°C ± 3°C using a self-assembled multichannel potentiostat to further evaluate the device stability under operating conditions. For these measurements, we employed poly[bis(4‑phenyl)(2,4,6‑trimethylphenyl)amine] (PTAA) doped with 4‑isopropyl‑4′‑methyldiphenyliodonium tetrakis(pentafluorophenyl)borate (TPFB) instead of spiro‑OMeTAD, with a PTAA concentration of 30 mg mL−1 and a PTAA:TPFB weight ratio of 10:1.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In this manuscript, the authors report a strategy for controlling the crystallization of α-FAPbI₃ perovskite films in a\nconventional n-i-p solar cell architecture. A carbazole-based molecule (Cz-SCN) is introduced via the anti-solvent, where it\nforms a self-assembled layer on the perovskite surface. Through in-situ and ex-situ characterizations, the authors propose\nthat this template guides a top-down crystallization process, resulting in highly ordered, phase-pure films with low defect\ndensities. The devices achieve a certified efficiency of 25.67% and improved operational stability. The approach of using the\nanti-solvent to deliver a self-assembling template for in-situ crystallization control is novel. I recommend acceptance of the\nmanuscript after carefully addressing the following points to further strengthen the manuscript:\n1. The in-situ GIWAXS, UV-Vis, and PL data in Figure 2 provide insight into the crystallization mechanism. A more detailed\ndiscussion of how the self-assembled molecule influences crystallization kinetics and phase evolution would enhance the\nmechanistic analysis.\n2. Regarding the encapsulated mini-module in Figure 4c, please provide additional fabrication details (e.g., aperture area) in\nthe Experimental Section to improve reproducibility.\n3. The absence of PbI₂ in GIWAXS and SEM suggests its consumption during crystallization. The authors should clarify\nwhat PbI₂ is converted into and whether the detection conditions (e.g., shallow incidence angle) could limit interpretation to\nsurface regions.\n4. Please describe the distinguishing characteristics of the SAM used in this work compared to those commonly employed in\np-i-n architectures, particularly in the Abstract.\n5. Define all abbreviations upon first use (e.g., EA, Cz-SCN, EIS).\n6. Comment on the effect of excess SAM on film formation and properties.\n7. The term “Cz-SCN interlayer” (Line 262) may not be accurate if the molecule is distributed throughout the film. Please\nrevise accordingly.\n8. Specify the measurement conditions (e.g., reverse scan or stabilized power output) for both champion lab-cell and\ncertified efficiencies.\n9. The presentation of Figure 2 should be reorganized for clarity. Include units for the Y-axis in Figure 3b, and relabel the\nlateral axis in Figure 3c to “wavenumber”.\n10. For the MPP tracking stability test (1000 h), please specify the HTL used and how Li-TFSI migration in spiro-OMeTAD\nwas mitigated. Also, indicate the testing temperature.", "reviewer_2": "This manuscript presents a novel and effective strategy by introducing self-assembled molecules (SAMs) into the antisolvent\nto direct the ordered crystallization of α-FAPbI3.Through a series of in situ characterizations (such as in-situ GIWAXS and in-\nsitu PL), the authors systematically reveal the role of this strategy in regulating crystallization kinetics, confirming its\ncapability to suppress undesired intermediate phases, accelerate the δ-to-α phase transition, and modulate crystal growth to\npromote the formation of highly ordered films. This fundamental understanding is convincingly linked to the fabrication of\nexceptionally high-quality perovskite films. The resulting device performance, with a certified efficiency of 25.67% for single-\njunction cells and 21.70% for mini-modules, is state-of-the-art and provides unequivocal validation of the method's power.\nOverall, the study demonstrates outstanding innovation, rigorous scientific investigation, and significant practical impact. The\nmanuscript is well-written, logically structured, and the conclusions are strongly supported by the data. Overall,I recommend\nits publication in Nature communications after the following questions are well responsed.\n1. The manuscript focuses on the success of Cz-SCN. For completeness, could the authors briefly comment on the\nconcentration optimization process for the other analogous molecules (Cz-NH₂, Cz-CN, Cz-COOH)? Clarifying if their\nperformance was also optimized would strengthen the conclusion that the anchoring group, not suboptimal concentration, is\nthe differentiating factor.\n2. In the DFT results presented in Fig. 2d, the energy labels (E_control and E_Cz-SCN) are difficult to read. Please increase\nthe font size for improved clarity.\n3. The authors emphasize that the SAM serves as a dynamic template to guide crystallization. To elevate this mechanistic\ninsight, could the authors further speculate on the dominant intermolecular forces driving the self-assembly and templating\neffect? For instance, is the stabilization of the α-phase primarily governed by specific chemical coordination with the -SCN\ngroup, or do van der Waals forces and hydrogen bonding play a significant role in the ordered packing?\n4. Figure 3d presents the XPS spectra of Pb 4f. Could the authors provide the binding energy changes of the relevant peaks\nbefore and after SAM introduction, in order to more directly illustrate the effect of SAM on perovskite crystallization?\n5. The in-situ PL and UV-Vis measurements provide exceptional insight into the crystallization dynamics and are a highlight\nof this work. To ensure the full reproducibility of these crucial experiments, the authors are requested to provide more\nspecific details in the Methods section. This could include, for example, the laser excitation wavelength for PL, the type of\nspectrometer and integration time for UV-Vis, and the precise timing protocol relative to the anti-solvent deposition and start\nof annealing.\n6. Compared to the certified efficiency of 25.67% for single-junction cells, the PCE of 21.70% for mini-modules is a little bit\nlow. The authors are suggested to provide the discussion why there is so big gap between the PCEs of point PSC and mini-\nmodule."}, "author_response": "Reply to reviewer #3\nThis work reports a facile strategy to induce the ordered crystallization of FAPbI3 by\nincorporating a SAM as a dynamic template to guide the top-down crystallization\nprocess. This enables a champion PCE of 26.18% with excellent stability in an ambient\nenvironment. However, fundamental reinterpretation of the proposed crystallization\nmodel and comprehensive re-evaluation of all device metrics are essential prerequisites\nfor reconsideration. Concerns remain regarding the explanation of results and\nperformance certification. Therefore, this work could only be reconsidered after\naddressing the following issues:\nGeneral Response: We sincerely thank the reviewer for the careful evaluation of our\nmanuscript and for raising critical concerns regarding the crystallization model, device\nmetrics, and performance certification. We fully acknowledge the importance of these\npoints and revised the manuscript point by point according to the reviewer’s comments,\nensuring that all concerns have been carefully addressed and the overall clarity, rigor,\nand impact of the work have been improved. We are grateful for the reviewer’s\nconstructive feedback, which has helped us refine both the scientific narrative and the\npresentation of our results. We believe that we have thoroughly addressed the\nreviewer’s concerns and that the revised manuscript now meets the standards for\nreconsideration.\nComment #1: For the in-situ GIWAX results: (Q1) the authors should label the time\nwhen the anti-solvent was dropped; (Q2) the mechanism behind the lamella stacking of\nthe SAM appearing only after annealing for 20s-30s should be explained; (Q3) the\nclaim that “the formed surface crystallization of α-FAPbI serves as a template to\n3\npropagating the phase transition towards the bulk of the film” appears controversial\nwith the TOF-SIMs data indicating uniform SAM distribution across the film; (Q4)\nfurther analysis of GIWAX results is required to determine if the SAM induced a\npreferable orientation of the perovskite's (100) plane.\nResponse for Q1. We thank the reviewer for this helpful comment. In the spin-coating\nprocess, the total duration was 40 s, consisting of an initial step at 1000 rpm for 10 s\nfollowed by a second step at 5000 rpm for 30 s. During the final 10 s of spin-coating\n(30–40 s), 120 μL of ethyl acetate was dripped onto the film as an antisolvent.\nAccordingly, the in-situ absorption and GIWAXS data presented in Fig. 2 were selected\nfrom this 30-40 s interval. To clearly illustrate the subsequent changes, the annealing\nprocess is shown immediately after spin-coating, with its onset set at 40 s. We have\nnow labeled the antisolvent dripping time in the figure and clarified this point in the\ncaption.\nChanges in the manuscript:\n“Fig. 2. In-situ analysis of the SAM-directed perovskite crystallization process. a-b, In-situ\nGIWAXS patterns of control and SAM-treated perovskite films during spin-coating and annealing.\nSpin-coating lasted 40 s (10 s at 1000 rpm, then 30 s at 5000 rpm), with 120 μL ethyl acetate dripped\nas antisolvent during the final 10 s (30–40 s). Data from this interval were selected, and annealing\nis shown immediately afterward, starting at 40 s.”\nResponse for Q2. In the early stage of annealing, rapid volatilization of residual solvent\noccurs as the substrate temperature rises, while the perovskite precursor undergoes\npreliminary crystallization and reorganization into a relatively disordered grain\nstructure. With continued thermal input, the system gradually approaches the critical\ncondition for the self-assembled monolayer (SAM) to arrange in an ordered fashion and\nguide crystal growth. This ordering represents a spatiotemporal evolution process\ndriven by thermodynamic stabilization and facilitated by molecular diffusion. Thus,\nalthough initial crystallization occurs early, the long-range ordered lamellar stacking\ncharacteristic of the SAM emerges only after sufficient annealing time (20–30 s), as\nreflected by the appearance of Bragg diffraction spots in the GIWAXS pattern. We have\nadded this mechanistic explanation to the revised Discussion section.\nChanges in the manuscript:\n“During the early stage of annealing, rapid solvent loss and preliminary crystallization produce a\ndisordered grain structure. Continued thermal input drives the self-assembled monolayer (SAM)\ntoward an ordered arrangement that guides crystal growth. This spatiotemporal evolution, driven\nby thermodynamic stabilization and facilitated by molecular diffusion, explains why long-range\nlamellar stacking appears only after 20–30 s of annealing, as evidenced by Bragg diffraction spots\nin the GIWAXS pattern. These results highlight that SAM-assisted crystallization requires both\nthermal activation and structural relaxation before exerting its templating effect on perovskite grain\ngrowth.”\nResponse for Q3. Based on our experimental design—where the self-assembled\nmonolayer (SAM) is added to the surface of the wet film along with the anti-solvent—\nwe propose a dynamic template extension mechanism to explain the phenomenon of\n“surface crystallization guiding bulk phase transition.” In the early stage of annealing,\nrapid solvent volatilization at the film surface causes SAM molecules to preferentially\naccumulate and anchor on the surface layer, forming an ordered template that induces\nα-FAPbI₃ crystallization. As annealing proceeds, thermal driving promotes further\nsolvent loss and facilitates redistribution of SAM molecules along grain boundaries and\ninto the film depth, gradually extending the templating effect inward. This dynamic\nprocess ultimately results in the uniform distribution of SAM observed by TOF-SIMS,\nwhile also realizing the progressive transition of the α-phase from surface to bulk.\nImportantly, “surface crystallization acting as a template to advance toward the bulk”\nand “uniform SAM distribution” are not contradictory; they represent different stages\nof the same dynamic process, with TOF-SIMS reflecting the final post-annealed state.\nWe also note that TOF-SIMS depth profiling on hybrid perovskites is susceptible to\nsputter-induced mixing, fragmentation, and depth-resolution limits that can smear sharp\ninterfacial distributions and broaden organic signals, especially for species localized at\ngrain boundaries and surface-assembled layers. Such artifacts can yield an apparently\nuniform profile even when the initial ordering was surface-dominant. In our processing\nroute, SAM is expected to localize predominantly at the top surface and along grain\nboundaries during growth. We have clarified this point in the revised Discussion,\nemphasizing that TOF-SIMS reflects the post-annealed state and should be interpreted\nwith consideration of sputtering-related artifacts and resolution limits.\nChanges in the manuscript:\n“It should be noted, however, that the apparent uniformity in TOF-SIMS depth profiles mainly\nreflects fragment distribution generated during sputtering, together with depth resolution limits that\ncan smear sharp interfacial features and broaden organic signals.”\nResponse for Q4. By comparing the GIWAXS one-dimensional line profiles of the\ncontrol and Cz-SCN-treated films, we observed that the PbI peak at q = 0.86 Å-1\n2 z\ndecreased significantly after Cz-SCN treatment, consistent with strong coordination\nbetween the –SCN group and Pb²⁺ ions. The integrated intensity of the perovskite (100)\npeak was indeed enhanced, indicating improved crystallization quality. However, no\ndistinct diffraction spots were observed in the GIWAXS patterns, suggesting that\nCz-SCN treatment did not induce a preferential orientation of the (100) plane.\nComplementary characterizations, including SEM (larger and more uniform grains),\nXRD (sharper diffraction peaks), steady-state PL (enhanced emission intensity), and\nTCSPC/TRPL (longer carrier lifetimes), consistently confirm that Cz-SCN promotes\nhigh-quality perovskite crystallization. Taken together, these results demonstrate\nimproved film quality rather than a change in crystallographic orientation.\nChanges in the Supplementary information:\nSupplementary Fig. 6. Comparative GIWAXS one-dimensional line profiles of control and\nCz-SCN-treated perovskite films. The Cz-SCN modification suppresses the PbI scattering peak at\n2\nq = 0.86 Å-1 and enhances the integrated intensity of the perovskite (100) reflection.\nz\nComment #2: For the different SAMs, the authors claim that the absorption energy is\nthe main origin of different crystallization of films. However, the pKa of the four\nmolecules are significantly different; this should also be considered as a potential factor\ninfluencing the crystallization.\nResponse: We thank the reviewer for raising the pKa point; it is an important\nconsideration. We experimentally measured pKa by UV-vis titration following\nestablished literature protocols for organic molecules: reliable values were obtained for\nCz-NH2 (≈9.77) and Cz-COOH (≈5.18), while repeated attempts for Cz-CN and Cz-\nSCN produced irregular, non-reproducible signals and thus no reliable experimental\npKa could be determined for those two compounds. To ensure a complete comparison\nwe therefore computed pKa values using DFT (Cz-NH 14.26; Cz-COOH 4.15; Cz-CN\n2\n10.57; Cz-SCN 12.49) and have reported these calculations and the methodology in the\nrevised Supplementary Information. More importantly, in the DMF/DMSO perovskite\nprecursor environment protonation–deprotonation equilibria are strongly suppressed,\nso Brønsted acidity/basicity plays a much reduced role compared with Lewis-base\ncoordination to undercoordinated Pb2+. Experimentally, only Cz-SCN produced\nlamellar self-assembly and full suppression of residual PbI , while Cz-NH and Cz-\n2 2\nCOOH did not induce the same ordered crystallization despite their very different pKa\nvalues. Together with our DFT adsorption-energy results and spectroscopic trends (FT-\nIR, XPS), these observations indicate that adsorption/coordination strength to Pb2+ is\nthe dominant factor governing the differing crystallization behaviors in our\nnon-aqueous system. We have added a concise discussion of these points to the main\ntext and included the experimental details and full DFT data in the Supplementary\nInformation.\nFigure 1. Figures a-d show the UV-Vis absorption spectra of Cz-COOH, Cz-NH , Cz-\nof absorbance at the characteristic peak of 344 nm versus pH are summarized in the\nfigures, clearly reflecting the dependence of this absorption peak on pH. To accurately\ndetermine the dissociation constants, we adopted a more precise measurement method\n(Method II), and the resulting pKa values are presented in the subsequent figures.\nTable 1\nCz-COOH 0.1mol/L-HCL/AHL 0.1mol/L-NaOH/AL- pH=2.44 buffer/A pH=5.44 buffer/A\npH 1.19 13.33 3.06 7.29\nAmax 2.68 2.49 2.62 2.55\nλmax 344.49 346.01 344.99 346.01\n(A-AHL)/(AL--A) / / 0.46 2.17\npKa / / 2.29 8.06\naverage value 5.18\nTable 2\nCz-NH2 0.1mol/L-HCL/AHL 0.1mol/L-NaOH/AL- pH=2.44 buffer/A pH=5.44 buffer/A\npH 1.17 13.34 7.03 11.34\nAmax 1.58 1.77 1.69 1.72\nλmax 344.01 345.01 343.95 345.05\n(A-AHL)/(AL--A) / / 1.375 2.8\npKa / / 7.17 12.37\naverage value 9.77\nMethods for pKa Determination\nMethod 1: UV Spectrophotometric Method\nSample preparation: The compound was dissolved in DMSO to prepare a 10 mM\nstock solution. A series of buffer solutions with different pH values were prepared, and\ntheir ionic strength was adjusted to 0.1 M using KCl, followed by filtration through a\n0.22 µm membrane. For UV–Vis measurements, 40 µL of the stock solution was added\nto 2.0 mL of buffer in a quartz cuvette, yielding a final concentration of approximately\n0.2 mM and a DMSO volume fraction below 2%. After gentle mixing, the absorption\nspectra were recorded in the range of 200–500 nm using a UV–Vis spectrophotometer,\nwith the corresponding buffer used as the blank. The spectra were baseline-corrected,\nand the pH-dependent absorption peaks were selected to construct Abs–pH curves.\nMethod 2 Classical UV–Vis Spectrophotometric Titration (Benzoic Acid protocol)\nThis traditional approach measures absorbance changes of acidic/basic species at\nselected wavelengths under different pH conditions.\nPrinciple:\nIonizable compounds exist in equilibrium between protonated (HA) and deprotonated\n(A⁻) forms:\n𝐻𝐴 ⇌ 𝐴− +𝐻+\nThe two species exhibit distinct UV absorption spectra.\nExperimental procedure:\nSolutions prepared in buffers spanning pH values around the expected pKa (typically\npKa ± 1). The pKa range of the compound was preliminarily determined using Method\n1. The compound was then dissolved in anhydrous ethanol to prepare a 0.25 mg/mL\nstock solution. After selecting appropriate pH buffer solutions, the stock solution and\nbuffer were mixed in a 1:1 volume ratio and gently shaken. Acid and base solutions of\n0.1 M HCl and 0.1 M NaOH, respectively, were used to adjust pH. The entire procedure\nfollowed the benzoic acid method for pKa determination, in which the acid–base\nequilibrium of the compound is inferred from the absorbance changes upon mixing with\nbuffer or during titration with acid or base. Absorbance measured at wavelengths where\nthe difference between HA and A⁻ spectra is maximal.\nReference spectra obtained under strongly acidic (fully HA) and strongly basic (fully\nA⁻) conditions.\nData analysis:\nThe absorbance at intermediate pH values is expressed as a linear combination:\n𝐴 = 𝜀 [𝐻𝐴]+𝜀 [𝐴−]\n𝐻𝐴 𝐴−\nRearrangement yields:\n𝐴−𝐴\n𝐴−\n𝑝𝐾 = 𝑝𝐻+log⁡\na 𝐴 −𝐴\n𝐻𝐴\nChanges in the manuscript:\n“Importantly, although the four Cz derivatives exhibit distinct aqueous pKa values (Supplementary\nTable 1, Supplementary Note 7), the direct influence of acidity/basicity on crystallization is limited\nin the non-aqueous DMF/DMSO precursor environment, where coordination strength with\nundercoordinated Pb2+ dominates.”\nChanges in the Supplementary information:\nSupplementary Table 1. Theoretical determination of acid dissociation constants (pKa) for the four\ntarget compounds.\n∆G (A-) ∆G (AH) ∆G (A-) ∆G (AH) pK\ngas gas sol sol a\n[Hartree] [Hartree] [kcal/mol] [kcal/mol]\nCz-NH -729.19 -729.83 -116.00 -2.2 14.26\nCz-COOH -862.50 -863.06 -82.56 -0.52 4.15\nCz-SCN -1164.31 -1164.92 -98.97 -2.13 12.49\nSupplementary Note 7\n𝐻𝐴(𝑔𝑎𝑠)⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡ ∆ → 𝐺 𝑔 𝑎𝑠 𝐴− ⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡+𝐻+\ngas gas\n↑−∆𝐺 (𝐴𝐻)⁡⁡⁡⁡⁡⁡⁡↓∆𝐺 (𝐴−)⁡⁡⁡⁡⁡⁡↓∆𝐺 (𝐴−)\nsol sol sol\n𝐻𝐴(𝑎𝑞)⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡ ∆ → 𝐺 𝑎 𝑞 𝐴− ⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡⁡+𝐻+\naq aq\nThe definition of pKa is given by:\n𝑝𝐾 =−𝑙𝑜𝑔𝐾 ⁡ (1)\na a\nand since Gibbs Free Energy\n∆𝐺𝑜 =−𝑅𝑇𝑙𝑛𝐾 =−2.303𝑅𝑇𝑙𝑜𝑔𝐾 ⁡ (2)\na a\n∆𝐺𝑜\n𝑝𝐾 = (3)\na 2.303𝑅𝑇\nFor this cycle, ∆𝐺𝑜 =∆𝐺 an and the aqueous-phase free energy change (∆𝐺 ) equals the gas-phase\naq aq\nfree energy change (∆𝐺 ) plus the difference in solvation free energy (∆∆𝐺 ).\ngas sol\n∆𝐺 =∆𝐺 +∆∆𝐺 (4)\naq gas sol\n∆∆𝐺 =∆𝐺 (𝐻+)+∆𝐺 (𝐴−)−∆𝐺 (𝐴𝐻) (5)\nsol sol sol sol\nThe calculation of ∆𝐺 uses a reference state of 1 atm, while the calculations of ∆𝐺 use a reference\ngas sol\nstate of 1 M. Converting the ∆𝐺 reference state (24.46 L at 298.15 K) from 1 atm to 1 M is\ngas\naccomplished using:\n∆𝐺 (1𝑀)=∆𝐺 (1𝑎𝑡𝑚)−𝑅𝑇𝑙𝑛(1/24.46) (6)\ngas gas\nThe values for ∆𝐺 (𝐻+) and ∆𝐺 (𝐻+) are derived from experiment. We have used the values\ngas sol\n∆𝐺 (𝐻+)=−6.28 kal/mol and ∆𝐺 (𝐻+)=−264.61 kal/mol. Based on the above derivation, the\ngas sol\nfinal formula is obtained as:\n𝑝𝐾 =[∆𝐺 (𝐴−)−∆𝐺 (𝐴𝐻)+∆𝐺 (𝐴−)−∆𝐺 (𝐴𝐻)−269]/1.3644(7)\na gas gas sol sol\n7 8\nGeometry optimization and vibrational frequency analysis were performed at the B3LYP /6-31G(d,p)\nlevel with the SMD solvation model for water. Electronic energies were subsequently computed at the\n6. accurate-pka-calculations-for-carboxylic-acids-using-complete-basis-set-and-\ngaussian-n-models-combined-with-cpcm.\n7. P. C. Hariharan, J. A. Pople, Accuracy of AH equilibrium geometries by single\nn\ndeterminant molecular orbital theory. Mol. Phys. 27, 209–214 (1974).\n8. G. A. Petersson, A. Bennett, T. G. Tensfeldt, M. A. Al‐Laham, W. A. Shirley, J.\nMantzaris, A complete basis set model chemistry. I. The total energies of closed‐shell\natoms and hydrides of the first‐row elements. J. Chem. Phys. 89, 2193–2218 (1988).\n9. Y. Zhao, D. G. Truhlar, The M06 suite of density functionals for main group\nthermochemistry, thermochemical kinetics, noncovalent interactions, excited states,\nand transition elements: two new functionals and systematic testing of four M06-class\nfunctionals and 12 other functionals. Theor. Chem. Acc. 120, 215–241 (2008).\nComment #3: Quantitively analysis of GIWAX based on different SAMs should be\nperformed to provide further insight.\nResponse: We fully agree with the reviewer that quantitative analysis of GIWAXS data\nprovides deeper mechanistic insights into how different SAM molecules regulate\nperovskite crystallization. Following this suggestion, we compared the GIWAXS\none-dimensional line profiles of perovskite films modified with various SAMs. In the\nCz-SCN film, the PbI peak at q = 0.86 Å-1 is markedly suppressed, while the\n2 z\nperovskite (100) peak intensity is significantly enhanced. To provide a clear\ncomparison, we calculated the intensity ratio of the perovskite (100) peak to the PbI\n2\npeak. The Cz-SCN film shows the highest (100)/PbI ratio, confirming its ability to\n2\nsuppress PbI formation and promote highly ordered crystallization. In addition, a\n2\ndistinct Debye-Scherrer diffraction ring at q = 0.13 Å-1 is observed only in the Cz-SCN\nz\nfilm, indicative of lamellar structural ordering. In contrast, films modified with other\nSAMs (Cz-NH , Cz-CN, Cz-COOH) show reduced PbI signals but no clear lamellar\n2 2\nordering. These quantitative results, now incorporated into the revised discussion,\nindicates the unique role of the SCN anchoring group in modulating the crystallization\npathway and improving film quality.\nChanges in the manuscript:\n“Quantitative GIWAXS analysis further demonstrates this distinction, showing that the Cz-SCN\nfilm exhibits the strongest and sharpest (100) diffraction peak together with the highest (100)/PbI\n2\nintensity ratio, confirming its superior ability to suppress PbI formation and promote highly\n2\nordered crystallization (Supplementary Figs. 15)”\nChanges in the Supplementary information:\nSupplementary Fig. 15. Quantitative GIWAXS one-dimensional line profiles of perovskite films\nwith different SAM modifications. The Cz-SCN-treated film exhibits the strongest and sharpest\nperovskite (100) diffraction peak together with a markedly reduced PbI signal at q = 0.86 Å-1,\n2 z\nresulting in the highest (100)/PbI intensity ratio. In contrast, films modified with Cz-NH , Cz-CN,\n2 2\nand Cz-COOH show weaker (100) peaks and residual PbI signals, highlighting the unique ability\n2\nof Cz-SCN to suppress PbI formation and promote highly ordered crystallization.\nin Fig S13-14.\nResponse: We thank the reviewer for this helpful suggestion. In the previous version\nof the manuscript, the values of VTFL and trap density (Nt) for the control and\nCz-molecule treated PSCs (based on the hole-only device architecture) were already\nprovided in Supplementary Table 1. To make the presentation clearer, we have now\nalso included these values directly in the captions of Supplementary Fig. 21-22.\nChanges in the Supplementary information:\nSupplementary Fig. 21. J-V characteristics of hole-only devices for control and\nCz-molecule-treated perovskite films. The extracted VTFL values are summarized in\nSupplementary Table 1.\nSupplementary Fig. 22. J-V characteristics of electron-only devices for control and Cz-molecule-\ntreated perovskite films. The corresponding trap densities (Nt), calculated from VTFL, are provided\nin Supplementary Table 1.\nComment #5: As there is also SAM at the buried interface, its potential to passivate\ndefects of SnO and benefit the device performance should be evaluated.\nSIMS depth profiling suggests SAM signals throughout the film, this mainly reflects\nsputtering-induced fragments and limited depth resolution. Since SAM molecules are\nintroduced via the anti-solvent step, they predominantly self-assemble at the top surface\nand along grain boundaries during the surface-initiated crystallization process, with\nonly minimal presence at the buried perovskite/SnO interface. This conclusion is\n2\nsupported by UPS measurements, which show a clear increase in work function at the\nCz-SCN-modified top surface but almost no change at the buried interface, and by SEM\nanalysis, which reveals improved crystallinity at the buried interface. Taken together,\nthese results suggest that Cz-SCN enhances perovskite crystallization quality rather\nthan directly passivating SnO . We agree that improved crystallinity at the buried\n2\ninterface can indirectly facilitate charge extraction, and we have added a note in the\nrevised discussion to highlight this point. However, systematic passivation of SnO\n2\ndefects is beyond the scope of this study, which focuses on the role of SAM in directing\nsurface-initiated crystallization. We consider this an interesting direction for future\nwork.\nChanges in the manuscript:\n“Consistent with this interpretation, UPS measurements show a clear increase in work function at\nthe Cz‑SCN‑modified top surface but negligible change at the buried perovskite interface, indicating\nminimal SAM accumulation there. In addition, SEM analysis of the buried interface reveals\nimproved crystallinity, which is attributed to more complete perovskite crystallization induced by\nSAM (Supplementary Figs. 8).”\nSupplementary Fig. 8. a. and b. present the UPS spectra (showing both the secondary electron\ncutoff and valence band regions) for the top interface of the control and Cz-SCN-treated films,\nrespectively, while Figures c. and d. show the corresponding spectra for their bottom interfaces.\nSupplementary Fig. 9. SEM images comparing the morphology of the bottom interface: a. control\nfilm and b. Cz-SCN-modified perovskite film.\nComment #6: The authors claim that “The champion device featuring the Cz-SCN\ninterlayer deliver a remarkable PCE of 26.18%”; however, the Cz-SCN was used as\nan additive. They should clarify the exact role and location of the SAM.\nResponse: We thank the reviewer for this helpful comment. We agree that the term\n“Cz-SCN interlayer” may cause confusion, as our results indicate that Cz-SCN\nmolecules are not confined to a discrete interfacial layer but are partially incorporated\ninto the perovskite film during crystallization. To avoid ambiguity, we have revised the\nterminology throughout the manuscript. Specifically, we now refer to the additive as\n“Cz-SCN incorporated films” or “Cz-SCN modified perovskite films,” depending on\nthe context.\nChanges in the manuscript:\n“The champion device based on the Cz-SCN-modified perovskite layer achieved ...”\nComment #7: For the certification report, proof of the active area certification (0.0800\ncm2) should be provided to validate the certified performance.\nResponse: We thank the reviewer for this important comment. We would like to clarify\nthat in the third-party certification process, the device active area (0.0800 cm²) was first\nmeasured and certified prior to the efficiency certification. The certified efficiency\ntherefore directly corresponds to this certified active area. To make this clearer, we have\nincluded the certification report photograph in the Supplementary Information\n(Supplementary Fig. 29), which explicitly shows the certified active area together with\nthe efficiency result.\nComment #8: For the module performance, details of the area characterization should\nbe provided, as the JSC of 4 mA/cm2 appears much higher than the theoretical value.\nResponse: We sincerely thank the reviewer for raising this important point. We\nacknowledge that the module photograph in Fig. 4c was erroneously inserted during\nfigure assembly, and we have now replaced it with the correct image corresponding to\nthe presented performance data. The experimental results (J–V curves, efficiency)\nremain unchanged and accurate, and this correction does not affect any conclusions of\nthe manuscript.\nTo clarify the area characterization, the perovskite mini-module prepared in this study\nconsists of six sub-cells connected in series, with a total active area of 14.21 cm2. Under\nstandard test conditions (AM 1.5G, 100 mW/cm2), the module demonstrated promising\nphotovoltaic characteristics. We measured an aperture-area short-circuit current density\n(J , aperture) of 4.01 mA/cm2. This value was calculated directly based on the total\nsc\nexternal area of the module and reflects the overall output performance, including\ninterconnect electrodes, scribing dead zones, and other inactive regions. To accurately\nevaluate the intrinsic quality of the perovskite active material itself, we further\ncalculated the active-area short-circuit current density (J , active). According to the\nsc\nfundamental relationship for serially connected sub-cells, J , active = J , aperture × N\nsc sc\n(where N=6 represents the number of sub-cells), we obtain J , active = 4.01 mA/cm2 ×\nsc\n6 = 24.06 mA/cm2. The geometric fill factor (GFF) calculation result is GFF = (4.0cm-\n5×680um) ÷ 4.0cm = 91.4%. The line widths of P1, P2, and P3 were 30 μm, 200 μm,\nand 150 μm, respectively, with inter-line spacings of 100 μm (P1-P2) and 200 μm (P2-\nP3). We have now included a detailed description in the revised Experimental Section\nof the manuscript.\nChanges in the manuscript:\nTo demonstrate the scalability of this approach, 5 cm × 5 cm mini-modules comprising six series-\nconnected subcells are fabricated. The champion mini-module, with an active area of 14.21 cm2,\nachieved a high PCE of 21.7%, with a V of 7.024 V, a J of 4.01 mA cm-2, and an FF of 77.05%\nOC SC\n(Fig. 4c).\nChanges in the Device fabrication of manuscript:\n“A perovskite module composed of six series-connected sub-cells was constructed on a 5 cm × 5\ncm glass/FTO substrate. The interconnection were carried out by laser scribing at a wavelength of\n532 nm. The active area of the module is 14.21 cm2. The geometric fill factor (GFF) calculation\nresult is GFF = (4.0cm-5×680μm)÷4.0cm = 91.4%. The line widths of P1, P2, and P3 were 30 μm,\n200 μm, and 150 μm, respectively, with inter-line spacings of 100 μm (P1-P2) and 200 μm (P2-P3).”\nComment #9: To explain the excellent ambient stability, the authors should provide\nfilm characterization results (SEM, XRD, etc) of the perovskites after aging.\nResponse:\nWe thank the reviewer for this constructive suggestion. Following the reviewer’s advice,\nwe have performed additional film characterization (SEM and XRD) on perovskite\nfilms after aging under the same ambient conditions used in our stability tests. The\nresults are now provided in the Supplementary Information Figs 33-34.\nThe SEM images show that the control films develop severe surface degradation and\npinholes after aging, whereas the Cz-SCN treated films retain a compact and uniform\nmorphology with significantly fewer defects. Consistently, XRD analysis reveals that\nthe control films exhibit pronounced PbI peaks and reduced α-phase intensity after\n2\naging, while the Cz-SCN films maintain strong α-phase diffraction peaks with\nnegligible PbI signals. These results confirm that the Cz-SCN additive effectively\n2\nsuppresses degradation pathways and preserves the structural integrity of the perovskite\nfilms under ambient conditions. We have added a brief discussion of these findings in\nthe revised manuscript and Supplementary information.\nChanges in the manuscript:\n“Post-aging SEM and XRD analyses further confirm that Cz-SCN treated films retain more compact\nmorphology and strong α-phase crystallinity with negligible PbI formation, in sharp contrast to\n2\nthe degraded control films (Supplementary Figs. 33-34)”\nChanges in the supplementary information:\n\nSupplementary Fig. 33. SEM images and XRD patterns of perovskite films after aging under\nambient conditions (23 ± 3 °C, 50 ± 10% RH, 600h) (a, control; b, w/Cz-NH ; c, w/Cz-CN); d,\n2\nw/Cz-COOH; e, w/Cz-SCN). Devices were aged intact; before SEM and XRD characterization the\nmetal electrode and hole-transport layer were removed to expose the perovskite film for\nmeasurement.\nSupplementary Fig. 34. SEM images and XRD patterns of perovskite films after aging under\nambient conditions (continuous 1-sun illumination, 500h) (a, control; b, w/Cz-NH ; c, w/Cz-CN);\n2\nd, w/Cz-COOH; e, w/Cz-SCN). Devices were aged intact; before SEM and XRD characterization\nthe metal electrode and hole-transport layer were removed to expose the perovskite film for\nmeasurement."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The revised manuscript has addressed my concerns, and I recommend acceptance of the manuscript without further\nrevision.", "reviewer_2": "All the raised questions have been well addressed. I recommend the acception of the manuscript in current version.", "reviewer_3": "The revised manuscript has addressed my concerns"}, "author_response": ""}]

Decision letter