High Open-Circuit Voltage–Fill factor product in perovskite solar cells enabled by ferroelectric heterojunction modulation

Domain: chemistry
Published: 2026-02-18
Editorial outcome: major_revision

Abstract
Overcoming the inherent limitations of perovskite-perovskite heterojunctions in simultaneously boosting built-in potential and suppressing non-radiative recombination remains a critical challenge in perovskite solar cells. Here, we introduce a ferroelectric-based heterojunction architecture that addresses this dual challenge through synergistic mechanisms. Firstly, the spontaneous polarization inherent to the ferroelectric-based heterojunction significantly amplifies the built-in electric field, enhancing charge separation and transport, thereby increasing open-circuit voltage from 1.16 to 1.21 V. Secondly, ferroelectric nuclei effectively regulate perovskite crystallization kinetics via dissolution-recrystallization modulation, effectively suppressing trap states and elevating fill factor from 83.6% to 86.8%. The champion devices achieve a power conversion efficiency of 26.62% (certified 26.07%) with an open-circuit voltage-fill factor product of 1.05 V, reaching 90.3% of the Shockley-Queisser limit. Furthermore, the modified devices demonstrate enhanced operational stability with over 85% efficiency reservation after 500 h of maximum power point tracking, charting clear pathways towards high-performance photovoltaic cells.

Introduction
The emergence of metal halide perovskites has catalyzed a paradigm shift in photovoltaics. Leveraging their exceptional optoelectronic characteristics, perovskite solar cells (PSCs) have achieved rapid breakthroughs in power conversion efficiency (PCE), with the current record of 27.3% positioning them as a compelling contender to the monocrystalline silicon benchmark. Despite the encouraging progress achieved, energy loss remains an intractable challenge within PSCs, restricting further efficiency climbing towards the Schockley-Queisser (SQ) limit. The performance of PSCs with certain functional layers is closely related to the difference between quasi-Fermi levels for electrons and holes, which determines the established built-in potential and the achievable energy conversion in devices. The built-in electric field can serve as the driving force for charge separation and transport from the perovskite absorber to the corresponding electrodes, with its reinforcing assistance to enhance carrier extraction and mitigate energy loss. However, non-radiative recombination loss stemming from defect capture and interface mismatch poses significant obstacles, potentially threatening the splitting of quasi-Fermi levels and the open-circuit voltage (V_OC) deficit. Furthermore, the lower fill factor (FF) of PSCs, as compared to the commercial silicon and GaAs photovoltaics, is primarily attributed to carrier transport limitations and interfacial recombination. As such, research efforts to achieve efficiency breakthroughs are centered on addressing the losses in V_OC and FF, both of which bridge the gap of at least 5% from the theoretical values.

Among the non-radiative recombination pathways identified in PSCs, the high-density defect within the solution-processed perovskite films and mismatched band structure are the predominant sources. The rapid nucleation and growth process generally causes disordered crystallization issues on both spatial and time scales, resulting in films characterized by inhomogeneous compositional distributions and severe defect residue. The high-density charged defects and lattice dislocations serve as the charge recombination centers to capture the carrier by Coulombic interactions and restrain the charge transport efficiency. Additionally, the deep-level defects can also induce charge accumulation and yield the back electric field inside devices, thereby compromising device performance. Besides the defect influences, the considerable energy offset between the highest occupied molecular orbital (HOMO) level of the general carbazole-based self-assembly monolayers (SAMs) and the valence band edge of perovskite causes a weak built-in electric field for driving the carrier separation and extraction in inverted PSCs. This results in unbalanced charge transfer and significant carrier accumulation at the contact interfaces, further exacerbating the energy loss and compromising the device performance. Therefore, to overcome the theoretical limit of solar cells, optimizing the crystallization kinetics to minimize defect density and strengthening the built-in field to accelerate charge transport in PSCs tends to be imperative.

The perovskite-perovskite heterojunction emerges as an effective strategy to enhance the built-in electric field in PSCs. This approach involves blending two types of perovskite compositions with distinct band gaps and optoelectronic properties, achieved through in situ phase segregation within the absorber layer. Low-dimensional perovskite materials, including 2D perovskite crystals, 1D perovskitoids, and 0D quantum dots, have been incorporated into bulk 3D perovskites to introduce an additional electric field, thus promoting more efficient charge extraction. Theoretically, ferroelectric perovskites, which exhibit dipole characteristics under external fields, are expected to enhance the built-in potential of PSCs through self-polarization. Well-controlled intrinsic polarization directions and the precise spatial distribution of ferroelectrics are beneficial for aligning the local electric field, ultimately impacting charge collection efficiency. Quite recently published pioneering work has demonstrated this concept by introducing a 2D photo-ferroelectric perovskite to form a 2D/3D layered junction, which induced the spontaneous accumulation of the 2D ferroelectric phase. This approach mitigated interfacial recombination in PSCs and achieved a significant V_OC gain of 1.21 V. However, despite this V_OC enhancement, substantial energy loss persists in such 2D/3D layered perovskite junctions. These losses are primarily attributed to the high density of non-radiative recombination sites in the bulk 3D perovskite layer, leading to a maximum PCE below 25%. This underscores the challenge of simultaneously achieving a high built-in electric field and a reduction in non-radiative recombination sites in the pursuit of enhanced cell performance.

Herein, we demonstrate ferroelectric-based heterojunctions (FBHJ), constructed by incorporating either Dion–Jacobson (DJ)- or Ruddlesden–Popper (RP)-phase 2D perovskites, synergistically decrease V_OC and FF deficits in PSCs through coupled built-in field amplification and crystallization modulation. We experimentally verify the effects of FBHJ on optimizing nucleation rate, growth uniformity, dissolution-recrystallization dynamics, and defect suppression in films by in situ optical spectrometry techniques. These benefits demonstrate substantial enhancements in the built-in electric field and charge collection efficiency in devices. As a result, PSCs achieve an impressive PCE of 26.62% (certified 26.07%) with a V_OC × FF product of 1.05 V, surpassing 90% of the SQ limits and representing a low energy loss among reported PSCs with PCEs exceeding 26%.

Results
We engineered two distinct FBHJ architectures to validate the universality of our strategy: FAPbI_3 integrated with DJ-phase ferroelectric perovskite (4-(aminomethyl)piperidinium)PbI_4 (denoted as 4AMP-DJ), and FAPbI_3 coupled with RP-phase ferroelectric (4,4-difluoropiperidinium)_2PbI_4 (denoted as DDFP-RP). Figure 1a illustrates the device configuration of the FBHJ solar cell using ferroelectric decoration, and Fig. 1b presents the molecular architectures of the DJ- and RP-phase ferroelectric perovskites reported in current research.

Systematic device characterizations were conducted to validate the efficacy of FBHJ in overcoming the SQ efficiency limit. Incorporation of 4AMP-DJ at optimized concentrations induced remarkable performance enhancements, in which the PCE increased from 24.46% (control) to 26.62%, accompanied by substantial improvements in V_OC from 1.16 V to 1.21 V (ΔV_OC = +50 mV) and FF from 83.57% to 86.79% (ΔFF = +3.22%), while maintaining a stable short-circuit current density (J_SC ≈ 25.4 mA cm^-2) (Supplementary Fig. 1). Notably, the V_OC enhancement exhibited concentration-dependent saturation behaviors. Further increasing the 4AMP-DJ concentration preserved a high V_OC (1.21 V), while it triggered reductions in FF and J_SC (Supplementary Fig. 2, Supplementary Table 1), suggesting compromised charge transport at an excessive ferroelectric content. This nonlinear response highlights the requirement of balance modulation between charge transport and built-in electric field, which will be systematically discussed in the following sections. The universality of FBHJ engineering was further confirmed as the DDFP-RP also achieved comparable performance metrics (V_OC = 1.20 V, PCE = 26.18%) as depicted in Fig. 1c. The PCE statistics further validated the performance improvement in devices through ferroelectric doping (Fig. 1d, Supplementary Fig. 3 and Supplementary Table 2).

Steady-state power output measurements under maximum power point tracking (1.04 V bias) confirmed exceptional operational stability, with PCEs stabilizing at 26.31% (4AMP-DJ) and 26.01% (DDFP-RP) over 600 s of illumination (Fig. 1e). Spectral response analysis revealed near-unity external quantum efficiency (EQE) across 300–800 nm, yielding integrated J_SC values of 25.01, 25.10, and 25.06 mA cm^-2 for the control, 4AMP-DJ, and DDFP-RP devices, respectively (Fig. 1f). The minimal variation in J_SC variation (<0.4%) corroborates that the efficiency gains are primarily originated from improvements in V_OC and FF. To evaluate the compatibility of the FBHJ strategy in variations of device structures, we extended this approach to flexible devices, obtaining a champion PCE of 25.15% with a V_OC of 1.19 V (Fig. 1g, Supplementary Fig. 4 and Supplementary Table 3). In addition, the 4AMP-DJ PSCs at an active area of 1.04 cm^2 delivered a champion PCE of 25.45%, illustrating good scalability (Supplementary Fig. 5, Supplementary Table 4).

The champion device with an aperture area of 0.0718 cm^2 achieved a certified PCE of 26.07% (Fig. 1h, Supplementary Fig. 6) from the Fujian Metrology Institute (FIL) of the National Photovoltaic Industry Metrology Center (NPVM). Crucially, the V_OC × FF product of 1.053 V approaches 90.3% of the SQ theoretical limit for the 1.56 eV band gap perovskite, corresponding to a low energy loss of 0.35 eV with respect to the SQ limit of 0.29 eV, which represents a small reported value for inverted PSCs exceeding 26% efficiency (Fig. 1i, Supplementary Table 5). This near-ideal V_OC × FF product stems from low V_OC deficit (0.35 eV) and FF deficit (3.41%) simultaneously. Moreover, the hysteresis index of PSCs was significantly mitigated through the ferroelectric doping (Supplementary Fig. 7, Supplementary Table 6).

The operational stability of PSCs was investigated under constant 100 mW cm^-2 illumination in a nitrogen atmosphere. FBHJ devices retained 85% initial PCE after 500 h of maximum power point tracking (MPPT), outperforming that of the control with 50% retention (Fig. 1j). The T_80 lifetime extrapolation suggests over 700 h operational stability under 1-sun equivalent stress (Supplementary Fig. 8). This synergistic enhancement of efficiency and stability conclusively indicates a dual-function potential of the FBHJ strategy, simultaneously engineering built-in electric fields through ferroelectric polarization while passivating defect sites via improved perovskite crystallization.

To elucidate the optimization mechanism underlying the energy loss in FBHJ PSCs, we then explored the influences of ferroelectric doping on the ferroelectricity of perovskite films. (4AMP)PbI_4 belongs to a typical 2D DJ-phase perovskite, in which PbI_6^4- octahedral adopts a corner-shared model to form 2D inorganic perovskite frameworks, while 4AMP cations acting as interlayers are arranged in an alternating up-and-down orientation (Supplementary Fig. 9, Supplementary Note 1). The DDFP cation with asymmetrical structure and fluoridation substituent was chosen as the spacer ligand to construct the 2D perovskite ferroelectric (DFPD)_2PbI_4, featuring an RP configuration with infinite corner-sharing [PbI_4]_n^2- layers separated by two DFPD-based spacer layers (Supplementary Fig. 10, Supplementary Note 2).

We then conducted the differential scanning calorimetry (DSC) tests to determine the ferroelectric-paraelectric phase transition behaviors of (4AMP)PbI_4 and (DDFP)_2PbI_4. For (4AMP)PbI_4, an endothermic peak was observed at 351 K in the heating run, accompanied by a corresponding exothermic peak detected at 348 K in the cooling run (Supplementary Fig. 11). The DSC curves of (DDFP)_2PbI_4 revealed an endothermic anomaly peak at 428 K during the heating cycle with an exothermic anomaly peak at 405 K (Supplementary Fig. 12a), indicating a reversible phase transition. Notably, the relative permittivity (ε') of (DDFP)_2PbI_4 under the heating treatment exhibited an abnormal surge at 428 K, labeled as Curie temperature (T_C), further confirming the paraelectric-ferroelectric phase transition (Supplementary Fig. 12b).

Piezoresponse force microscopy (PFM) is a powerful tool for identifying ferroelectricity. As illustrated in Fig. 2a–c, the control film demonstrated a homogeneous phase without distinct ferroelectric domain walls, implying nonintrinsic ferroelectricity. By contrast, (4AMP)PbI_4 and (DDFP)_2PbI_4 presented obvious domain structures along the out-of-plane direction, where domain walls with weak piezoelectric response separate domains with different orientations (Supplementary Figs. 13, 14). Critically, the observed ferroelectric domain structure is unrelated to the topography, confirming that the signal comes from piezoelectric responses with different orientation domains rather than morphology crosstalk. 4AMP-DJ and DDFP-RP FBHJ films exhibited a ferroelectric domain structure similar to the (4AMP)PbI_4 and (DDFP)_2PbI_4, showing clear and well-defined domains (Fig. 2d–i). PFM analyses revealed that FBHJ films preserve ferroelectric domains within the 2D ferroelectric composites. In addition, the FBHJ films also presented butterfly-shaped amplitude loops in response to variations in the electric field (Supplementary Fig. 15), further confirming the manifestation of ferroelectric behaviors in FBHJ films. Macroscopic P–E hysteresis loops of FBHJ film have been supplemented as Supplementary Fig. 16, revealing their ferroelectricity at the enlarged scale. By contrast, the intrinsic ferroelectricity was absent in the pristine FAPbI_3 films.

To elucidate the crystallization behavior of FBHJ perovskite, we conducted in situ spectroscopy characterizations during film formation (Fig. 3a, b). In situ photoluminescence (PL) spectrometry was applied to monitor the crystallization evolution behaviors of perovskite films during the spin-coating and annealing stages, especially for the top surface under an optical reflection mode. The intensity evolution against time was plotted in Fig. 3c for the PL peak at 775 nm. During spin-coating, anti-solvent dripping triggers rapid perovskite nucleation, whereas no significant difference in PL intensity was observed among different films (Supplementary Fig. 17). However, distinct kinetics emerged during the annealing stage. Initially, residual solvent evaporation induced a rapid dissociation of nuclei at the liquid-air interface, leading to a sharp decay in PL intensity within the first ca. 7 s. Subsequently, the dissolved crystal at the top surface began to recrystallize, which took ca. 16 s, 28 s, and 49 s for the PL intensity to reach the peak for the control, DDFP-RP, and 4AMP-DJ films, respectively. The extended recrystallization window suggested that FBHJ involves the 2D ferroelectric crystal seeds to regulate the growth behaviors of the intrinsic 3D component, thereby retarding the crystallization rate. The prolonged growth duration in 4AMP-DJ FBHJ film was associated with the enhanced coordination affinity of diammonium spacer ligands and potential intermediate participation in the crystallization retardance. Notably, FBHJ also suppressed defect-induced PL quenching, presenting more stable emission intensity over time.

In situ UV-vis spectroscopy further revealed solvent-evaporation-driven bulk nucleation dynamics in a transmission mode. In the control film, the absorption intensity gradually increased and reached the edge of the bulk phase after an initial 15 s (Fig. 3d), coinciding with top-surface nuclei total dissolution, followed by secondary crystallization reaching the absorption peak at ca. 23 s. Ferroelectric doping accelerated nucleation, shortening the appearance of the first absorption peak to ca. 7 s (DDFP-RP) and ca. 9 s (4AMP-DJ), while the secondary crystallization window terminated at ca. 15 s and 19 s. This acceleration stems from a Pb-rich environment induced by partial dissolution of ferroelectric perovskite in precursor solution, which promoted the mobile ion diffusion and reduced the perovskite nucleation barrier. Critically, in the FBHJ case, the nucleation termination in the bulk aligned well with surface recrystallization onset at the liquid-air interface, alleviating strain-induced defects via homogeneous growth.

In situ spectroscopic analysis revealed the FBHJ-modulated perovskite crystallization kinetics during thermal annealing, as rationalized by the principles of the LaMer model. In the bulk film, the nucleation process is initiated with solvent evaporation driving rapid dissolution of transient nuclei at the film surface (Fig. 3e), concomitant with bulk solution concentration surpassing the supersaturation threshold (Stage I). Intriguingly, the ferroelectric doping supported a Pb-rich environment in the FBHJ bulk, markedly enhancing perovskite nucleation rates compared to the control (Stage II). The high-polarity organic ligands, featuring strong coordination affinity with Pb-I octahedra framework, would induce the rapid formation of 2D seeds to template crystal growth of FAPbI_3, thereby promoting structurally ordered recrystallization (Stage III). However, excessive dopant concentrations may push the precursor solution beyond its critical supersaturation limit (C_max), triggering uncontrolled nucleation that compromises size uniformity and film morphology.

At the liquid-air interface, solvent evaporation induces the dissolution of metastable nuclei while progressively elevating the solution concentration, with all film compositions exhibiting similar dissolution-recrystallization initiation thresholds (Fig. 3f). Upon depletion of these interfacial nuclei, a dynamic evolution was established between evaporative flux and bulk solution replenishment, which presented a down-up stage of PL intensity with the extended annealing time. Subsequent solvent depletion shifted the predominant phase transformation mechanism toward regular surface rearrangement. Critically, ferroelectric incorporation facilitates the precise regulation of growth kinetics throughout the FAPbI_3 perovskite. The ferroelectric perovskite effectively serves as a multifunctional crystallization modulator by simultaneously lowering the nucleation energy barrier, prolonging the interfacial recrystallization timeframe, and homogenizing crystal growth dynamics across films. This synergistic regulation affords heterojunction films with exceptional crystalline quality and denser crystalline stacking, as key metrics for the optoelectronic performance of PSCs.

The scanning electron microscopy (SEM) analysis demonstrated effective suppression of PbI_2 residues in films upon ferroelectric incorporation, accompanied by a marginal increase in average grain size and improved spatial homogeneity (Fig. 4a). Cross-sectional SEM revealed vertically continuous grain structures spanning the entire ca. 835 nm-thick perovskite layer (Supplementary Fig. 18), suggesting optimized crystallization pathways conducive to unimpeded charge carrier transport. For the spatial distribution of ferroelectric perovskite in final FBHJ films, time-of-flight secondary-ion mass spectroscopy (ToF-SIMS) and X-ray diffraction (XRD) characterizations revealed that the 2D ferroelectric perovskites retained their inherent structure within resulting films, which distributed throughout the bulk film and dominantly enriched at the bottom interface (Fig. 4b, Supplementary Figs. 19, 20). Moreover, XRD patterns exhibited characteristic perovskite phase amplification (at 2θ = 14.2°) with 4AMP-DJ and DDFP-RP incorporation, corroborating the enhanced crystallinity observed in morphological studies. Note that the ferroelectric doping distinctly enhanced PL intensity and signal uniformity of perovskite films, indicative of suppressed non-radiative recombination process (Fig. 4c–e). Time-resolved PL spectra illustrated that the average carrier lifetime increased from 9.4 μs (control) to 12.8 and 17.7 μs for the DDFP-RP and 4AMP-DJ samples, respectively, representing one of the state-of-the-art results for the polycrystalline perovskite films (Fig. 4f). Furthermore, space charge limited current (SCLC) measurements quantified the trap density within perovskite films, showing the average value reduced from 5.88±0.94×10^15 cm^-3 (control) to 2.70±0.34×10^15 and 1.45±0.12×10^15 cm^-3 of the DDFP-RP and 4AMP-DJ cases, respectively (Fig. 4g, Supplementary Fig. 21). Consequently, ferroelectric incorporation effectively modulates the perovskite crystallization kinetics, facilitating enhanced crystalline quality and suppressed non-radiative recombination in films.

Surface potential distributions across the perovskite films were characterized using Kelvin probe force microscopy (KPFM) (Fig. 5a). Compared with the control film (–55 mV), both the 4AMP-DJ and DDFP-RP FBHJ films showed lower contact potential difference (CPD) values of –212 mV and –142 mV, respectively (Fig. 5b). The statistics of CPD exhibited a more homogeneous distribution of surface potential in the FBHJ films (Fig. 5c), which is favored for interface contact and charge extraction. UV-vis absorption spectra maintained invariant band-edge characteristics (Tauc plot-derived E_g = 1.56 eV) across modified compositions (Supplementary Fig. 22). The energetic alignments extracted from UV-vis absorption and ultraviolet photoelectron spectroscopy (UPS) were plotted in Fig. 5d, e. Compared to the control film, the Femi level (E_F) was downshifted from –4.40 (control) to –4.42 (DDFP-RP) and –4.44 eV (4AMP-DJ). The energetic gap between E_F and valence band maximum (VBM) was shortened from 1.22 (control) to 1.15 (DDFP-RP) and 1.10 eV (4AMP-DJ), indicating the enhanced p-type character of FBHJ films. Notably, the VBMs of FBHJ perovskites were slightly upshifted and aligned better with the highest occupied molecular orbital (HOMO) level of MeO-4PACz, with the energy offset reduced from 0.30 (control) to 0.25 (DDFP-RP) and 0.22 eV (4AMP-DJ), endowing the more matched energetic alignment at the anode interface. Moreover, the band bending at the junction interface could yield an additional electric field to strengthen the built-in potential, ensuring suppressed energy loss in devices (Supplementary Fig. 23, Supplementary Note 3). TA spectra further illustrated the accelerated carrier transport at the hole extraction interface in FBHJ cases, which verified the synergy of ferroelectric doping on crystallization modulation and interfacial optimization (Supplementary Fig. 24).

To further elucidate the charge recombination behavior in FBHJ cells, the dependence of V_OC under varying light intensity (P_light) was studied. Based on the formula of V_OC = (nk_BT/q) ln(P_light), the corresponding slopes were calculated to be 1.47 k_BT/q, 1.12 k_BT/q, and 1.21 k_BT/q for the control, 4AMP-DJ, and DDFP-RP PSCs, respectively (Fig. 5f). By quantifying the FF loss in PSCs, the ferroelectric doping collectively reduced the charge recombination and non-radiative recombination losses in devices (Supplementary Fig. 25, Supplementary Note 4). In transient photocurrent decay (TPC) measurements (Supplementary Fig. 26), the FBHJ devices exhibited a faster decay rate compared to the control, implying the improved charge transport behaviors attributable to a reduction in the trap-mediated recombination process. From the C–V measurements in Fig. 5g, the effect of ferroelectric doping on the corresponding built-in potential at the contact interface was estimated by the Mott-Schottky analysis: 1/C^2 = 2(V_bi - V) / (A^2 ε_0 ε_r q N), where C represents the capacitance of the depletion layer, N is the carrier density, A is the active area, V_bi is the built-in potential at equilibrium, and V is the applied voltage. The V_bi extracted from the intercept of 1/C^2 = 0 was determined to be 1.12, 1.17, and 1.19 V for the control, DDFP-RP, and 4AMP-DJ devices, respectively. The higher built-in potential further confirmed the higher driving force for charge separation and transport, endowing the synergistic enhancements in V_OC and FF in the FBHJ PSCs.

Overall, our observations indicate that FBHJ architecture features coupled crystallization optimization and built-in field amplification, as shown in Fig. 5h. Ferroelectric doping contributes to the improved energetic alignment at the hole-extraction interface. The reinforced built-in potential can support a stronger driving force for charge separation and extraction. Moreover, the ferroelectric actively participates in the nucleation and growth of FAPbI_3 perovskite and subsequently aggregates at the buried interface in resulting films, achieving significant trap suppression, thereby contributing to the suppressed carrier capture by defects and the charge accumulation at the heterointerface. The synergistic effect of ferroelectric doping on crystallization modulation and internal electric field of perovskite film effectively decreases energy loss in PSCs. Furthermore, ferroelectric-based FBHJ films demonstrate the capacity to inhibit ion migration under constant illumination and enhance the long-term operational stability of devices (Supplementary Fig. 27). To fully harness the performance potential of devices, further ferroelectric design for improved ferroelectricity, including dipole regulation, intermolecular interactions, and polarization properties, is imperative. Delicate crystallization control to balance the domain distribution and polarization direction within the ferroelectrics would assist in effectively reinforcing the built-in field in perovskite-ferroelectric heterojunctions and minimizing energy loss in PSCs.

In summary, we demonstrate an effective ferroelectric-based heterojunction strategy to minimize the energy loss of PSCs and reveal the underlying mechanisms. The results indicate that ferroelectric heterojunction facilitates the broadening of the built-in potential under spontaneous polarization and supports an enhanced driving force for carrier transport. Furthermore, the ferroelectrics can assist the crystallization process of FAPbI_3, thus contributing to the shortened nucleation duration and prolonged recrystallization window for perovskite films. These benefits promote the growth balance across the films and optimize the crystalline quality, therefore endowing less non-radiative recombination loss in heterojunction films. Consequently, the FBHJ PSCs achieved a champion efficiency of 26.62% (certified 26.07%) with a V_OC × FF value of 1.05 V, surpassing 90% of the SQ thermal limit. Our study deepen the understanding of the optimization effects of ferroelectric perovskite on charge transport and crystallization modulation, thus shedding light on the further development of high-efficiency and stable perovskite photovoltaics.

Methods
Materials
FAPbI_3, MAPbBr_3, (4AMP)PbI_4 and (DDFP)_2PbI_4 single crystal powders were derived from self-synthesis, patterned indium tin oxide (ITO) glass, cesium iodide (CsI), lead(II) iodide (PbI_2), and methylammonium chloride (MACl) were obtained from Advanced Election Technology Co., Ltd. Phenylethylammonium iodide (PEAI), C_60, and [4-(3,6-Dimethoxy-9H-carbazol-9-yl)butyl]phosphonic Acid (MeO-4PACz) were obtained from Xi’an Solar Co., Ltd. Chlorobenzene (CB) (anhydrous, 99.8%), Ethanol (anhydrous, 99.8%) and isopropanol (IPA) (anhydrous, 99.5%) were purchased from Sino-pharm Chemical Reagent Co., Ltd. N,N-dimethylformamide (DMF) (anhydrous, 99.8%) and dimethyl sulfoxide (DMSO) (anhydrous, ≥99.9%) were obtained from Sigma-Aldrich.

Device fabrication
Devices were fabricated with a planar p-i-n structure of fluorine-doped tin oxide (FTO)/SAM/Perovskite/PEAI/C_60/BCP/Ag. The FTO glass substrates (2.5 cm × 2.5 cm) underwent a cleaning process involving ultrasonication in a specialized cleaning concentrate (Hellmanex III) mixed with ultrapure water (v:v=1~1.5:100) for 30 min, followed by three additional ultrasonication cycles in ultrapure water for 30 min each. Subsequently, the substrates were dried under a nitrogen flow and treated with UV-ozone for 15 min. The MeO-4PACz was dissolved in anhydrous ethanol at a concentration of 0.5 mg ml^-1. The as-prepared solution was spin-coated on the glass/FTO substrates at 3000 rpm for 30 s in the N_2-filled glovebox. After annealing at 100 °C on a hot plate for 10 min, the HTL-coated substrates were cooled to room temperature for subsequent fabrication.

The perovskite precursor (1.5 M) was prepared by mixing 1.5 M FAPbI_3, 0.06 M MAPbBr_3, 0.05 M CsI, and 10 mol% excess MACl in a solvent mixture of DMF and DMSO (volume ratio: 4:1). (4AMP)PbI_4 and (DDFP)_2PbI_4 were incorporated into the precursor solutions at different concentrations (ranging from 0 to 5.0 mg ml^-1). Subsequently, 100 μl of the perovskite precursor was spin-coated onto the SAM layers at 1000 rpm for 10 s and 4000 rpm for 30 s, with the addition of 200 μl chlorobenzene dripped onto the film at 10 s before the end of the procedure. The film was then transferred to the hot plate and annealed at 100 °C for 40 min. After cooling to room temperature, the perovskite film was spin-coated with PEAI dissolved in IPA (20 mM) at 4000 rpm for 30 s without further processing. Finally, 20 nm C_60, 7 nm BCP, and 100 nm silver were thermally evaporated onto the perovskite films.

Ferroelectricity characterization
Differential scanning calorimetry (DSC) was performed with a differential scanning calorimeter (NETZSCH DSC 214 Polyma) under an N_2 atmosphere. The temperature cycles during both the heating and cooling phases were set at a rate of 10 K min^-1. Dielectric constants (ε') were determined based on the variable-temperature dielectric permittivity, using a Tonghui TH2828A impedance analyzer. P−E hysteresis loops were recorded using a Sawyer−Tower circuit with Precision Premier II (Radiant Technologies, Inc.). Nanoscale polarization imaging and local switching spectroscopy were performed using resonant-enhanced piezoresponse force microscopy (MFP–3D, Asylum Research). Domain imaging and polarization switching studies were conducted using conductive Pt/Ir-coated silicon probes (EFM–50, Nanoworld). All PFM measurements were performed in the out-of-plane (vertical) mode (VPFM) to probe the vertical-aligned polarization component. To confirm the piezoresponse, a 2 V AC driving voltage was applied, measuring normal and shear responses at the second resonant peak of the cantilever-sample system to enhance sensitivity.

Material characterization
UV-visible absorption spectra were acquired on a PerkinElmer UV-Lambda 950 instrument. Steady-state photoluminescence (PL) was recorded with a PicoQuant FT-300 spectrometer. X-ray diffraction (XRD) studies were performed using a DX-2700BH diffractometer (Dandong Haoyuan Instrument Co., Ltd.). Scanning electron microscopy (SEM) images were obtained using a Hitachi SU-8020 field-emission scanning electron microscope (Japan). Atomic force microscopy (AFM) images and Kelvin Probe Force Microscopy (KPFM) images were obtained by a Bruker Dimension Icon instrument (USA). Ultraviolet photoelectron spectroscopy (UPS) was performed on a photoelectron spectrometer (ESCALAB Xi+, Thermo Fisher Scientific). In situ PL and UV-vis were tested by self-assembling equipment in our laboratory.

Device characterization
J–V characteristics of the solar cells were analyzed by solar simulator equipment (Enlitech, SS-F5) with an illumination intensity (AM1.5G, 100 mW·cm^-2) calibrated via a reference silicon cell with a KG5 filter. The scan range was set from 1.4 V to –0.1 V with a 0.02 V bias step and a 20 ms delay time. The active area of the non-refractive mask is 0.072 cm^2. The external quantum efficiency (EQE) was measured on a QE-R system (Enli Technology Co., Ltd.) using a 300-W Xe lamp as the light source. Capacitance–voltage (C–V) measurements were performed using an electrochemical workstation (Modulab XM, USA) with a frequency of 200 kHz and a scan voltage ranging from 0 to 1.6 V. Electrochemical impedance spectroscopy (EIS) was measured in the dark using a ModuLab XM CHAS08 with a frequency range from 0.1 Hz to 100 MHz. The operational stability of the unencapsulated PSCs was evaluated by a multichannel stability test system operating in the MPP tracking mode. The light intensity of the xenon lamp chamber was calibrated to achieve the same J_SC of PSCs as the values measured under the standard solar simulator (AM1.5G, 100 mW cm^-2). During MPP tracking (N_2 atmosphere, 45 ± 5 °C, equal to cell temperature), J–V characteristic curves of the devices were periodically monitored every 2 h.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In this manuscript, the author reported a ferroelectric bulk heterojunction (FBHJ) to enhance both open-circuit voltage (Voc) and fill factor (FF) in perovskite solar cells (PSCs). The spontaneous polarization in FBHJ simultaneously amplifies the built-in electric field to improve charge transport and modulates the crystallization kinetics to passivate trap states. Benefiting from this dual optimization, the fabricated PSCs delivered an exceptional power conversion efficiency (PCE) of 26.62%% with a certified value of 26.07%. Overall, the manuscript is well-organized, and the presented data support the main conclusions. The reported results might guide the development of high-performance PSCs towards theoretical limits. I would like to recommend it for publication in Nature Communications after considering the following comments.\n1. The introduction of FBHJs demonstrates a significant boost in device performance of small-area devices. To enhance the impact of this paper, the scalability of this optimization strategy could be validated. Therefore, it is recommended that device performance with a reasonably larger active area (e.g., 1.0 cm2) be provided.\n2. The authors demonstrated a synergistic improvement of Voc and FF for the FBHJ-based PSCs. While a clear discussion regarding the increase in Voc has been presented, the analysis related to the FF improvement lacks depth. To gain a better understanding of the underlying loss mechanisms of PSCs, it is essential to supplement the discussion with quantified contributions from non-radiative recombination and charge transport losses relative to the Shockley–Queisser (SQ) limit, based on the measured values.\n3. Clarify whether pure FAPbI3 exhibits ferroelectric property, and assess the necessity of ferroelectric perovskite doping to infer the ferroelectricity of the FBHJ structure?\n4. The characterization data concerning the carrier dynamics of the FBHJ structure can be strengthened. The data analysis about the time-resolved photoluminescence or transient-absorption spectra could be added, and the authors are encouraged to engage in some discussion surrounding the supplementary data.\n5. Ferroelectricity in metal halide perovskite typically correlates with the device hysteresis. [J. Phys. Chem. Lett. 2015, 6, 22, 4633., Energy Environ. Sci., 2015,8, 1602.] In this manuscript, the related device parameters are missing. The authors are recommended to provide some discussion on this concern.\n6. Figure formatting should be corrected to enhance the manuscript presentation. For example, the labels of “RP-4AMP/DJ-DDFP” and “FAPbI3” in Fig. 5g and 5h are not standardized. The units of scale bar within KPFM images in Fig. 5a are unclear.", "reviewer_2": "Wu et al. proposes a novel ferroelectric-based heterojunction (FBHJ) strategy for enhancing the built-in potential and simultaneously passivating non-radiative recombination sites of FAPbI3 perovskite solar cells (PSCs), achieving an impressive Voc× FF product approaching to the theoretical limit. The FBHJ structure facilitates spontaneous polarization and charge transport, elevating Voc from 1.16 V to 1.21 V. Furthermore, this optimization technique accelerates the dissolution-recrystallization process during film formation, leading to improved crystallization uniformity and the increase of FF from 83.6% to 86.8%. Overall, the characterizations presented in the manuscript are comprehensive, and the central conclusion is well-supported by the accompanying data. This manuscript is well-organized and the key doping effects on charge transport and perovskite crystallization are carefully investigated. The proposed FBHJ structure holds promise for facilitating further breakthroughs in the efficiency limits of PSCs. At this stage, there are still some details in the study that need to be improved. After addressing the following concerns, I recommend the publication of this manuscript.\n1. The manuscript demonstrates that the incorporation of ferroelectric perovskites enhances the built-in potential and carrier kinetics, thereby yielding significant improvements in Voc and FF values. However, it is noteworthy that Jsc exhibits no distinct variation between the control and FBHJ devices. It would be beneficial for the authors to provide a more in-depth analysis to elucidate the underlying insight of this observation? A comprehensive examination is recommended to strengthen the manuscript.\n2. The hysteresis behavior of FBHJ-based PSCs requires investigation. Ferroelectric polarization in perovskites is typically a primary contributor to hysteresis in PSCs. Clarification of the intrinsic ferroelectricity of pristine FAPbI3 and experimental evidence demonstrating ferroelectric doping on device hysteresis are recommended to support.\n3. Ferroelectric 2D perovskite was incorporated into the bulk FAPbI3 perovskite to construct the FBHJ structure. However, the final existing form of such ferroelectric materials in resulting films is less investigated. Further experimental characterization and discussion could be better supplemented.\n4. The doping amount of ferroelectric perovskite should be quantified to enhance the clarity of experimental details. Moreover, the impacts of doping concentrations on device performance should be supported.\n5. Distinct from the conventional bulk passivation agents, what are the specific advantages of ferroelectric materials for synergistic crystallization modulation and potential enhancement? Please comment on the further optimization routes for this FBHJ structure towards the efficiency limit of PSCs.\n6. Figure 3d should maintain a uniform time scale consistent with Figure 3c to clarify the crystallization kinetic evolution. Moreover, relocating Figure 3i before the mechanistic diagram in Figure 3h would improve logical flow.\n7. The Fermi level changes from -4.40 eV to -4.42 eV and 4.44 eV. How can the very small numerical change be excluded from the calculation error?\n8. The author mentioned ‘band bending’, and more evidence is needed to confirm it.\n9. Please list detailed comparisons for some key parameters, such as those mentioned in Matter (10.1016/j.matt.2022.05.040).", "reviewer_3": "The manuscript by Wu et al. reports a ferroelectric-based heterojunction strategy that yields perovskite solar cells with a power conversion efficiency of 26.62%. They demonstrate a record Voc×FF product of 1.05 V among cells with over 26% efficiency. The results are impressive. However, additional data is needed to clarify the effect of the FBHJ, specifically in the crystallization kinetics section. And the arrangement and readability of the manuscript should be much improved.\n1. The presented ferroelectricity characterization of perovskite materials is confusing and not properly discussed. According to the authors’ description, the polarization orientations are different for (4-AMP)PbI4 (parallel to the (010) plane, Supplementary Figure 7) and (DDFP)2PbI4 (along c axis, Supplementary Figure 8). However, the PFM measurement fails to specify whether vertical or lateralcomponents were probed. Also, the PFM hysteresisloops should be given against electrical fields, instead of voltage. Critically, although mentioned in the Methods section, macroscopic P-E hysteresis loops of FBHJ are not given in the text. The current PFM-based measurements are skeptical to support the presence of ferroelectricity.\n2. The authors claim that “Moreover, the ferroelectric actively participates in the nucleation and growth of FAPbI3 perovskite and subsequently accommodate within the grain boundaries in resulting heterojunction films”. The direct experimental evidence such as the EDX mapping to show the distribution of the ferroelectric perovskite is important to confirm the mechanism of the FBHJ.\n3. The author demonstrates impressive hysteresis suppression in Supplementary Fig.6. Please also include the J-V curve of the control device to demonstrate the effect of the FBHJ in suppressing hysteresis.\n4. Moreover, the C-V measurements for Mott-Schottky plot at various frequencies are recommended to add to confirm the suppression effect of the ion migration in perovskite and to validate the extracted built-in potential.\n5. In the crystallization kinetics part, the authors attribute the slowing down crystal growth to the strong coordination affinity of amine ligands with lead ions. This strong coordination affinity should anchor the lead ions, creating a kinetically \"Pb-poor\" environment. On the other hand, the authors proposed a \"Pb-rich environment\" to explain the accelerated bulk nucleation. Can the authors clarify the origin of this \"Pb-rich environment\" and how these two apparently contradictory effects can coexist and operate within one system?\n6. How about the reproducibility of the significant trap density reduction shown in Figure 5i? The statistical SCLC data from multiple devices for each condition is important to confirm the reproducibility.\n7. The overall organization and readability of the manuscript should be improved. Several key figures in the main text and Supplementary Information appear to be unreferenced in the discussion, such as Figure 3e and 3f, and Supplementary Figures 15-17 and the Supplementary Tables, while Supplementary Figure 18is missing Furthermore, there are some apparent citation errors that should be corrected. Lastly, the images in Figure 2 are of low resolution, appearing as raw instrument outputs with illegible text and scale bars. It is strongly recommended to remake this figure.", "reviewer_4": "In this manuscript, authors integrated two paradigmatic ferroelectric materials, (4-aminomethyl)piperidinium(PbI₃) (4AMP-DJ) and (4,4-difluoropiperidinium)₂PbI₄ (DDFP-RP), with FAPbI3 to form ferroelectric-based heterojunction (FBHJ) framework. As a result, over 90% of the Shockley-Queisser limit for Voc × FF and a front-line PCE of 26.62% (certified 26.07%) were obtained. The authors attributed this enhancement to two contributors: 1) the ferroelectric nuclei regulate the crystallization kinetics of perovskites, reducing defect density and enhancing film quality; 2) the spontaneous polarization of ferroelectric materials amplifies the built-in electric field, improving charge separation and transport. However, several gaps must be addressed to ensure the work’s robustness and credibility.\n1) Typing errors and problems with the article structure and (sub-)headings remain to be resolved. For example, a spelling error of ‘thr’ shown in Line 112; a subheading issue of ‘Discussion’ shown in Line 330. Authors should carefully check the correctness of other writings in the manuscript.\n2) It is noted that the manuscript inconsistently references its figures in the main text. For instance, Fig. 2a-c are explicitly mentioned and their intent is clear. Yet, other subfigures such as Fig. 2d-f and Fig. 2g-i are not referenced at all, leaving their purpose and relevance ambiguous. Similarly, Fig. 3e and Fig. 3f are missing their corresponding specific references in the text. The authors are encouraged to offer a clear explanation of each subfigure's role in supporting the study's arguments and results.\n3) The author performed PFM to identify the ferroelectricity of each control sample. Can authors explain why the topography of the DDFP-RP film appears so blurry and anomalous? Also, why does the phase contrast of its phase map differ by 60° from that of other samples? Compared to the SEM top views, shouldn't the morphology of DDFP-RP show distinct and well-defined grains?\n4) The author conducted an in-depth study on the crystallization kinetics of perovskite films with and without ferroelectric materials during the annealing process, utilizing in-situ photoluminescence (PL) and in-situ UV-vis spectroscopy. Firstly, could the author elaborate on why high-quality ferroelectric-based perovskite thin films demonstrate longer crystallization time domains in PL measurements but shorter nucleation time domains in absorption measurements? What accounts for the significant discrepancy in the total time required for solvent evaporation and film growth termination between these two characterization techniques? Secondly, this interpretation appears to rely on the assumption that the ferroelectric perovskite phase forms more rapidly than the conventional 3D perovskite phase, yet this claim lacks direct experimental evidence. Thirdly, the proposed mechanism by which the ferroelectric phase influences the nucleation process seems questionable. The authors suggest that local electric fields generated via spontaneous polarization increase the nucleation energy barrier. However, in ferroelectric materials, the internal electric field is confined within the material itself and is unlikely to extend externally to influence the nucleation of the 3D perovskite phase. This raises fundamental concerns about the scientific validity of the explanation. To substantiate the proposed mechanism, the authors should provide robust theoretical justification, such as finite element simulations or other computational modeling approaches.\n5) In Line 288-290, the author states that the enhanced p-type characteristics of FBHJ films are due to an increase in the energetic gap between Ef and VBM, which obvisouly contradicts semiconductor physics theory.\n6) When describing Figure 5h, the author believes that the spontaneous polarization of ferroelectricity contributes to the band bending at the heterojunction interface and provides an additional electric field. Then, how is the spatial distribution of the ferroelectric phase in the overall perovskite active layer? How is its polarization orientation, and how does it affect the transport of charge carriers? The author might use more advanced characterization methods and clearer illustrations to explain these confusions.\n7) The effectiveness of low-dimensional ferroelectric perovskites based on 4AMP and DDFP in enhancing the performance of three-dimensional perovskite solar cells has been investigated previously. For reference, see ‘Pica G., et al. Nat. Commun., 2024’ (doi:10.1038/s41467-024-53121-8); ‘Ma Y., et al. ACS Appl. Mater. Interfaces, 2023’ (doi:10.1021/acsami.3c05277), etc. Could the authors please explain the innovative aspects of this work and how it differs from previous studies in terms of research significance?\n8) A record-breaking Voc × FF stems from the dual improvement of these two performance metrics. Could you explain more about the reason why the product of Voc and FF should be compared between different works (Fig 1i)? The current manuscript seems to solely emphasize on discussing the underlying mechanisms behind the Voc enhancement. Then, why did FF improve significantly from 83.6% to 86.8%? What models/characterization approaches can be used to explain this improvement in FF?"}, "author_response": "List of point-to-point response of reviewers’ comments\n_____________________________________\nReviewer #1: In this manuscript, the author reported a ferroelectric bulk heterojunction (FBHJ) to enhance both open-circuit voltage (V ) and fill factor (FF) in perovskite solar cells (PSCs). The spontaneous polarization in FBHJ simultaneously amplifies the built-in electric field to improve charge transport and modulates the crystallization kinetics to passivate trap states. Benefiting from this dual optimization, the fabricated PSCs delivered an exceptional power conversion efficiency (PCE) of 26.62% with a certified value of 26.07%. Overall, the manuscript is well-organized, and the presented data support the main conclusions. The reported results might guide the development of high-performance PSCs towards theoretical limits. I would like to recommend it for publication in Nature Communications after considering the following comments.\nResponse: We are encouraged by the reviewer’s overall assessment of our manuscript and the specific comments to improve our paper. We have modified the manuscript based on the suggestions from the reviewer. The detailed responses to the reviewer’s comments and the related modifications in the revised manuscript are shown below.\nComment 1. The introduction of FBHJs demonstrates a significant boost in device performance of small-area devices. To enhance the impact of this paper, the scalability of this optimization strategy could be validated. Therefore, it is recommended that device performance with a reasonably larger active area (e.g., 1.0 cm2) be provided.\nResponse: We thank the reviewer for this important comment. We have prepared the 4AMP-DJ PSCs at an active area of 1.04 cm2, which delivered a champion PCE of 25.45% (Fig. R1a). The PCE loss between small- and large-area devices was only 1.17% (26.62% versus 25.45%). Furthermore, the 4AMP-DJ devices presented a narrow PCE distribution below 0.3%, with an average PCE of 25.35±0.05%, illustrating exceptional reproducibility(Fig. R1b).\n[... content truncated for brevity ...]"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have already addressed well the issues of the reviewers, and the paper could be accepted for publication.", "reviewer_2": "The author has made sufficient improvements, now I agree to publish.", "reviewer_3": "The authors have made attampts to revised this work. There is no further comment.\nFor Reviewer #4's comments,\nIn the item 3, the question of “why does the phase contrast of its phase map differ by 60° from that of other samples” is not directly justified. The revised image and data are not clearly explained and marked."}, "author_response": "Response to Reviewer #3’s further comment: In the item 3, the question of “why does the phase contrast of its phase map differ by 60° from that of other samples” is not directly justified. The revised image and data are not clearly explained and marked.\nResponse: We are glad to give further explanation. In PFM measurements, application of an alternating voltage to a conductive AFM probe enabled detection of a local electromechanical response on the film surface (Fig. R1). This response originates from the piezoelectric effect, wherein the sample undergoes periodic stretching deformation under an applied electric field. [J. Phys. D: Appl. Phys. 2011, 44, 464003] By detecting the probe deflection signal, both the amplitude and phase of the electromechanical response can be obtained. The phase ϕ reflects the orientation of the sample’s polarization direction relative to the external electric field. Ideally, in a c⁻ domain (polarization oriented downward, perpendicular to the surface), the applied positive bias would induce sample expansion, resulting in surface vibration in phase with the applied voltage (ϕ ≈ 0°). In contrast, c+ domains exhibit an out-of-phase response, labeled as ϕ = 180°. However, under practical PFM conditions, the phase signal frequently deviates from theses ideal values (from 0° 180°) to due to several influencing factors: (1) the low signal-to-noise ratio associated with minute piezoelectric displacements; (2) contributions from electrostatic forces and tip-sample capacitive coupling; (3) variations in local contact stiffness and damping indued by surface morphology; and (4) systematic phase offsets arising from lock-in amplifier reference settings and instrumental delays.[Nanotechnology,2006, 6, 0957./// Journal of Applied Physics, 1998, 11, 5973] Therefore, the measured phase images generally exhibit a constant offset from the ideal values due to these combined effects, which have been widely recognized in ferroelectric research.[Science 2019, 363, 1206./// Science 2014, 383, 1492./// Nat. Commun. 2024, 15, 10221./// J. Am. Chem. Soc. 2017, 139, 6369.//// Nat. Commun. 2024, 15, 6738./// Nat. Commun. 2023, 14, 5854./// Nat. Commun. 2025, 16, 3071./// J. Am. Chem. Soc. 2020, 142, 13989./// Nat. Commun. 2023, 14, 2863.]"}, {"round": 3, "reviewer_comments": {"reviewer_3": "More details are provided to address the reviewer's concern."}, "author_response": ""}]

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