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Nature Communications volume 17, Article number: 2897 (2026)
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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.
The emergence of metal halide perovskites has catalyzed a paradigm shift in photovoltaics1,2,3. 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 benchmark4. Despite the encouraging progress achieved, energy loss remains an intractable challenge within PSCs, restricting further efficiency climbing towards the Schockley-Queisser (SQ) limit5,6,7. 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 devices8,9.
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 loss10. 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 (VOC) deficit11. 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 recombination12. As such, research efforts to achieve efficiency breakthroughs are centered on addressing the losses in VOC and FF, both of which bridge the gap of at least 5% from the theoretical values13.
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 sources14. 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 residue15. 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 efficiency16. Additionally, the deep-level defects can also induce charge accumulation and yield the back electric field inside devices, thereby compromising device performance11. 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 PSCs17. 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 layer18. 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 extraction19. Theoretically, ferroelectric perovskites, which exhibit dipole characteristics under external fields, are expected to enhance the built-in potential of PSCs through self-polarization20. 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 phase21. This approach mitigated interfacial recombination in PSCs and achieved a significant VOC gain of 1.21 V. However, despite this VOC 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 VOC 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 VOC × 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%.
We engineered two distinct FBHJ architectures to validate the universality of our strategy: FAPbI3 integrated with DJ-phase ferroelectric perovskite (4-(aminomethyl)piperidinium)PbI4 (denoted as 4AMP-DJ), and FAPbI3 coupled with RP-phase ferroelectric (4,4-difluoropiperidinium)2PbI4 (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.
a The device structure of the inverted PSCs. b The packing structure of (4AMP)PbI4 and (DDFP)2PbI4 perovskites in their ferroelectric phases. c J–V curves and d statistic distribution of PCEs of the control and FBHJ PSCs. Error bars are the standard deviation of PCEs for 30 devices in each case. The lines in the middle of the box plots indicate the median values. e The stabilized power output curves of FBHJ PSCs. f EQE curves of the control and FBHJ PSCs. g The J–V curve of the champion FBHJ-based flexible PSCs. Illustrations provide specific device performance parameters. h The certified J–V curve of the champion FBHJ PSCs from the Fujian Metrology Institute (FIL) of the National Photovoltaic Industry Metrology Center (NPVM). The institutional logo @ 2025 Fujian Metrology Institute. All rights reserved. Reprinted with permission. i Plots of VOC × FF products for the state-of-the-art inverted PSCs with PCE over 26%, derived from Supplementary Table 7. j Operational stability of unencapsulated devices at the maximum power point (MPP) tracking with continuous one-sun illumination under N2 atmosphere at 45 °C.
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 VOC from 1.16 V to 1.21 V (ΔVOC = +50 mV) and FF from 83.57% to 86.79% (ΔFF = +3.22%), while maintaining a stable short-circuit current density (JSC ≈ 25.4 mA cm-2) (Supplementary Fig. 1). Notably, the VOC enhancement exhibited concentration-dependent saturation behaviors. Further increasing the 4AMP-DJ concentration preserved a high VOC (1.21 V), while it triggered reductions in FF and JSC (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 (VOC = 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 JSC 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 JSC variation (<0.4%) corroborates that the efficiency gains are primarily originated from improvements in VOC 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 VOC 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 cm2 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 cm2 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 VOC × FF product of 1.053 V approaches 90.3% of the SQ theoretical limit for the 1.56 eV bandgap perovskite22,23, 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)5,6,7,24,25,26,27,28,29,30,31,32,33,34,35,36,37,38,39,40,41. This near-ideal VOC × FF product stems from low VOC 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 atmosphere42. 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 T80 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)PbI4 belongs to a typical 2D DJ-phase perovskite, in which PbI64− 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)2PbI4, featuring an RP configuration with infinite corner-sharing [PbI4]n2− 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)PbI4 and (DDFP)2PbI4. For (4AMP)PbI4, 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)2PbI4 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 (ɛ′) of (DDFP)2PbI4 under the heating treatment exhibited an abnormal surge at 428 K, labeled as Curie temperature (TC), 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)PbI4 and (DDFP)2PbI4 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)43. 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)PbI4 and (DDFP)2PbI4, showing clear and well-defined domains (Fig. 2d–i). PFM analyses revealed that FBHJ films preserve ferroelectric domains within the 2D ferroelectric composites43. 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 FAPbI3 films.
Topography, amplitude, and phase images of a–c control, d–f 4AMP-DJ and g–i DDFP-RP films. Scalebar of 600 nm.
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 mode44,45. 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 interface46,47, 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 rate48. 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 retardance49. Notably, FBHJ also suppressed defect-induced PL quenching, presenting more stable emission intensity over time.
a in situ PL and b in situ UV-vis spectra of perovskite films during the annealing process. c Evolution of PL intensity at 775 nm for different perovskite films during the annealing process. d Evolution of absorption intensity at 550 nm for different perovskite films during the annealing process. e, f LaMer model illustrating the correlation between nucleation/growth rate and concentration evolution of precursor solution in (e) the bulk and f the top surface for the perovskite films.
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 dissolution46,50, 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 FAPbI348, thereby promoting structurally ordered recrystallization (Stage III). However, excessive dopant concentrations may push the precursor solution beyond its critical supersaturation limit (Cmax*), 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 FAPbI3 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 PbI2 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)51. Furthermore, space charge limited current (SCLC) measurements quantified the trap density within perovskite films, showing the average value reduced from 5.88 ± 0.94 × 1015 cm−3 (control) to 2.70 ± 0.34 × 1015 and 1.45 ± 0.12 × 1015 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.
a Top-view SEM images for different perovskite films. Scalebar of 1 μm. b ToF-SIMS depth profiles of DDFP-RP film. c, d PL intensity in PL mapping for control c and 4AMP-DJ d films. Scalebar of 50 μm. e–g steady-state PL, Time-resolved PL, and trap density statistics of perovskite films. Error bars are the standard deviation of trap density derived from 5 samples for each case.
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 Eg = 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 (EF) was downshifted from −4.40 (control) to −4.42 (DDFP-RP) and −4.44 eV (4AMP-DJ). The energetic gap between EF 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 films52. 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).
a KPFM image of perovskite films. Scalebar of 500 nm. b The line-scan CPD data for the perovskite films extracted from KPFM images. c The CPD distribution histograms of different perovskite films. d Valence band edges and cutoff regions of UPS spectra for the perovskite films. e Energy level alignment within complete devices. f The dependence of VOC on light intensity for PSCs. g Mott-Schottky plots of PSCs. h Schematic of ferroelectric perovskite doping on improving the built-in electric field (BEF) and defect passivation in FBHJ PSCs.
To further elucidate the charge recombination behavior in FBHJ cells, the dependence of VOC under varying light intensity (Plight) was studied. Based on the formula of53
the corresponding slopes were calculated to be 1.47kBT/q, 1.12kBT/q, and 1.21kBT/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 analysis54:
where C represents the capacitance of the depletion layer, N is the carrier density, A is the active area, Vbi is the built-in potential at equilibrium, and V is the applied voltage. The Vbi extracted from the intercept of 1/C2 = 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 VOC 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 FAPbI3 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 ionic 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 FAPbI3, 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 VOC × FF value of 1.05 V, surpassing 90% of the SQ thermal limit. Our study deepens 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.
FAPbI3, MAPbBr3, (4AMP) PbI4 and (DDFP)2PbI4 single crystal powders were derived from self-synthesis, patterned indium tin oxide (ITO) glass, cesium iodide (CsI), lead(II) iodide (PbI2), and methylammonium chloride (MACl) were obtained from Advanced Election Technology Co., Ltd. Phenylethylammonium iodide (PEAI), C60, 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.
Devices were fabricated with a planar p-i-n structure of fluorine-doped tin oxide (FTO)/SAM/Perovskite/PEAI/C60/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 N2-filled glovebox. After annealing at 100 °C on a hotplate 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 FAPbI3, 0.06 M MAPbBr3, 0.05 M CsI, and 10 mol% excess MACl in a solvent mixture of DMF and DMSO (volume ratio: 4:1). (4AMP) PbI4 and (DDFP)2PbI4 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 C60, 7 nm BCP, and 100 nm silver were thermally evaporated onto the perovskite films.
Differential scanning calorimetry (DSC) was performed with a differential scanning calorimeter (NETZSCH DSC 214 Polyma) under an N2 atmosphere. The temperature cycles during both the heating and cooling phases were set at a rate of 10 K min−1. Dielectric constants (ɛ’) were determined based on the variable-temperature dielectric permittivity, using a Tonghui TH2828 A 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.
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.
J–V characteristics of the solar cells were analyzed by solar simulator equipment (Enlitech, SS-F5) with an illumination intensity (AM 1.5 G, 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 cm2. The external quantum efficiency (EQE) was measured on a QE-R system (Enli Technology Co., Ltd.) using a 300-WXe 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 JSC of PSCs as the values measured under the standard solar simulator (AM 1.5 G, 100 mW cm−2). During MPP tracking (N2 atmosphere, 45 ± 5 °C, equal to cell temperature), J–V characteristic curves of the devices were periodically monitored every 2 h.
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This work was supported by the Scientific Research Project of China Three Gorges Corporation (Grant No. 202303014, K.Z.), Key Research and Development Program of Shaanxi (Program No. 2025CY-YBXM-170, K.Z.), National Natural Science Foundation of China (52402285, T.N.), National University Research Fund (GK202201005, K.Z.), 111 Project (B21005, S.L.), Fundamental Research Funds for the Central Universities (GK202503001, K.Z. and GK202304047, T.N.), Young Talent Fund of Xi’an Association for Science and Technology (959202413046, T.N.).
These authors contributed equally: Nan Wu, Haofei Ni.
Key Laboratory of Applied Surface and Colloid Chemistry, National Ministry of Education; Shaanxi Key Laboratory for Advanced Energy Devices, Shaanxi Engineering Lab for Advanced Energy Technology and School of Materials Science and Engineering, Shaanxi Normal University, Xi’an, China
Nan Wu, Tianqi Niu, Tinghuan Yang, Ru Qin, Lei Lang, Shuang Wang, Di Zhao, Chenqing Tian, Erxin Zhao, Chenxin Zhao & Kui Zhao
Institute for Science and Applications of Molecular Ferroelectrics, Key Laboratory of the Ministry of Education for Advanced Catalysis Materials, Zhejiang Normal University, Jinhua, China
Haofei Ni, Changfeng Wang & Yi Zhang
KAUST Solar Center (KSC), Physical and Engineering Division (PSE), King Abdullah; University of Science and Technology (KAUST), Thuwal, Kingdom of Saudi Arabia
Xiaoming Chang
Key Laboratory of Photoelectric Conversion and Utilization of Solar Energy, Dalian Institute of Chemical Physics, Chinese Academy of Sciences, Dalian, Liaoning, China
Shengzhong Frank Liu
Center of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing, China
Shengzhong Frank Liu
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K.Z. conceived and designed the study. K.Z. and T.N. supervised the project. N.W. conducted most of the experiments and performed the data analysis. H.N. and C.W. synthesized and characterized ferroelectric materials. T.Y. and X.C. optimized experimental recipes. R.Q. conducted in situ PL and in situ UV-vis measurements. L.L., D.Z., and C.Z. assisted with SEM testing. S.W. performed GIWAXS. C.T. and E.Z. carried out KPFM measurements. Y.Z. and C.W. helped with experimental design and manuscript preparation. All the authors discussed the results and commented on the paper.
Correspondence to Tianqi Niu, Yi Zhang or Kui Zhao.
The authors declare no competing interests.
Nature Communications thanks Jianhua Hao and the other anonymous reviewer(s) for their contribution to the peer review of this work. A peer review file is available.
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Wu, N., Ni, H., Niu, T. et al. High Open-Circuit Voltage–Fill factor product in perovskite solar cells enabled by ferroelectric heterojunction modulation. Nat Commun 17, 2897 (2026). https://doi.org/10.1038/s41467-026-69391-3
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