Misplaced-dipole engineered repairable fluoropolymer elastomer for flexible perovskite solar cell with excellent thermal-mechanical cycling resistance

Misplaced-dipole engineered repairable fluoropolymer elastomer for flexible perovskite solar cell with excellent thermal-mechanical cycling resistance | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Misplaced-dipole engineered repairable fluoropolymer elastomer for flexible perovskite solar cell with excellent thermal-mechanical cycling resistance Qi Chen, Feihu Liu, Jie Dou, Ying Li, Wei Zhang, Jiale Li, Yuchen Li, and 5 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-8773234/v1 This work is licensed under a CC BY 4.0 License Status: Under Review Version 1 posted You are reading this latest preprint version Abstract Flexible perovskite solar cells (F-PSCs) offer a compelling solution to the intrinic rigidity of silicon-based photovoltaics, enabling power generation on irregular surfaces. However, their practical application hinges on the perovskite layer’s ability to concurrent thermal cycling, so as to match the deformability of polymer substrates. Different from the conventional lattice solidification strategy for rigid devices, we incorporate a fluorinated misplaced-dipole engineered repairable elastomer into the perovskite film, which enhances perovskite intergranular toughness and mitigates thermal stress fatigue cracks. The resultant perovskite film exhibits suppressed lattice thermal fluctuation, thereby boosting enhanced environmental resilience. Consequently, the optimized F-PSCs deliver a champion efficiency of 25.54%, versus 26.83% for their rigid counterparts. More importantly, the F-PSC demonstrates exceptional durability under harsh operational stresses, retaining over 90% of the initial PCE after 11,000 bending cycles and maintaining a comparable retention rate following 500 thermal cycles, paving the way towards for the long-lasting flexible photovoltaic devices. Physical sciences/Energy science and technology/Renewable energy/Solar energy/Photovoltaics/Solar cells Physical sciences/Materials science/Materials for energy and catalysis/Solar cells Figures Figure 1 Figure 2 Figure 3 Figure 4 Introduction Organometallic halide perovskites, endowed with excellent low-temperature processibility, have emerged as promising candidates for assembling flexible perovskite solar cells (F-PSCs).1–4 In contrast to rigid counterparts fabricated on glass substrates, the intrinsic flexibility and lightweight nature of F-PSCs enables their conformal integration onto arbitrary surfaces, irrespective of morphological irregularties.5 , 6 To date, the power conversion efficiency (PCE) of F-PSCs has undergone rapid advancement, with record value reaching 25.74% for laboratory-scale single-junction devices and 21.6% for large-area modules.7–10 Nevertheless, substantial disparities persist between flexible and rigid devices in terms of both PCE and long-term perational stability. The state-of-the-art strategies for boosting efficiency and stability of F-PSCs predominantly center on buried interface engineering and additive modification to serve to strengthen interfacial adhesion, suppress defect formation, and even enhance the mechanical toughness of perovskite grain boundaries.11–14 Despite substantial advances, F-PSCs remain plagued by irreversible performance degradation triggered by deformation-induced exterinal stress. A primary reason stems from the utilization of flexible polymer substrates, such as polyethylene terephthalate (PET) and polyneopentamide glycolate (PEN), whose morphological characterstics exhibit strong dependence on external stress. This inherent substrate sensitivity readily induces buried perovskite interface delamination and bulk perovskite film cracking during both the annealing-cooling fabriation process and day-night operational cycles.11 , 12 Therefore, temperature fluctuation not only induces spatial distortion of the perovskite lattice but also drives the thermal expansion and contraction of flexible polymer substrates, thus making huge obstacles to the practical use of F-PSCs. Especially following dynamic operation under diurnal, seasonal, and climatic variations, F-PSCs are more prone to structural and chemical degradation of the perovskite active layer, which in turn impedes interfacial and intergranular charge transfer and fosters the formation of substantial defects, ultimately compromising device durability. In this context, strategies targeting the isolated reinforcement of the perovskite lattice such as molecular grain confinement or the incorporation of poly(methyl methacrylate)/graphene layers well-established for rigid scenarios are inadequate for F-PSCs.13 , 14 Therefore, engineering high-quality perovskite films featured with polymer-like mechanical behaviors including low Young’s modulus, high elasticity, and moderate rebound resilience, is equally critical to ensure compatibility with flexible substrates. Nevertheless, achieving this suite of mechanical requirements entails an inherent trade-off with the performance metrics and stability criteria aforementioned. To address this great challenge, herein we synthesized a misplaced-dipole engineered fluoropolymer elastomer by a hardness-softness synergistic coupling strategy, using a copolymer derived from the hard chain monomer 2,2,2-trifluoroethyl methacrylate (TFEMA) and the soft chain monomer 2,2,3,4,4,4-hexafluorobutyl acrylate (HFBA). Complementary molecular dynamics simulations and density functional theory (DFT) calculations confirm that this elastomer integrates high mechanical strength, robust hydrophobicity, and excellent self-healing capability. By incorporating this elastomer, we construct a hierarchically structured perovskite film with a distinctive grain-polymer-grain-polymer architecture, which endows the perovskite film with polymer-like mechanical traits including reduced Young’s modulus and efficient self-healing behavior to release the thermal and mechanical stresses. Consequently, the lattice deformation ratio of the perovskite film is significantly suppressed from 0.23% to mere 0.01% over a wide temperature range of -20 o C to 80 o C. Furthermore, the ion-dipole interactions between the fluoropolymer elastomer and perovskite lattice not only regulate the perovskite crystallization kinetics but also improve the film quality. Benefiting from these synergistic effects, the resultant F-PSCs deliver a champion PCE to 25.54% under 1-sun irradiation. More impressively, the devices show negligible degradation after 1100 h under 1-sun MPP tracking. They also exhibit exceptional environmental roubustness under harsh operational conditions, retaining 90% of their initial PCEs after 11,000 bending cycles and maintaining 90% following 500 thermal cycles (0°C ↔80°C). Results Molecular design and self-healing mechanisms The fluoropolymer elastomer was synthesized via the copolymerization of equimolar TFEMA and HFBA monomers, with 2,2-azobisisobutyronitrile (AIBN) serving as the initiator.15 Upon thermal curing at 60°C for 3 min, the homogeneous liquid reactive precursor solution gradually solidified into a solid elastomer (Fig. S1 ). As is well known, fluorinated polymers inherently exhibit exceptional chemical stability and superior hydrophobicity, stemming from the high bond dissociation energy of covalent C-F linkages.16 , 17 This hydrophobic characteristic is conducive to enhancing the moisture resistance of the resultant perovskite film (Fig. S2).18 Compared to the traditional polymer additives, the designed copolymer synergistically integrates the rigid segmental architecture of TFEMA chain with flexible chain conformation of HFBA moieties, thereby endowing the elastomer with balanced toughness, high stretchability, and efficient reparability. As shown in Fig. 1 a, two transparent polymer strips achieved a tensile elongation of 300% relative to their original length following a 3 min self-healing treatment at 60°C. After incorporated into the grain boundaries of perovskite film, the elastomer’s self-healing performance and stretchability effectively prolong the service lifetime of flexible perovskite by mitigating lattice strain under cyclic thermal and mechanical stress. The underlying mechanism governing this behavior will be discussed in the following sections. As documented in previous reports,19 , 20 the self-healing capability of the fluoropolymer elastomer is predominantly governed by robust dipole-dipole interactions among the highly polar groups namely -CF 3 , -CF 2 , and -C = O (Fig. 1 b).21 , 22 Specifically, the much larger electronegativity of F atom (χ = 4.0) than that of C atom (χ = 2.5) drives electron transfer from C atom to F atom, generating permanent dipoles with negatively charged center (δ⁻) and positively charged center (δ + ). An analogous charge polarization effect arises from the carbon-oxygen pair within C = O moiety. When the fluoropolymer elastomer suffers damage, the robust intermolecular interactions between fractured interfaces compress and shorten the interchain distance by reducing the entropy effect, which is governed by the interfacial interactions. More importantly, we discover that the interpenetrating network derived from the copolymerization of TFEMA and HFBA monomers enables the maximum intermolecular interaction to boost the self-healing kinetics.23 To in-depth understand this mechanism, we performed the molecular dynamics simulation to quantify intermolecular interactions by contrasting TFEMA homopolymer, HFBA homopolymer, and their 1:1 copolymer. As shown in Fig. S3, Fig. S4, and Table S1 , the TFEMA homopolymer exhibits the strongest Coulombic attraction yet the weakest van der Waals force, which accounts for its rigid solid-state behaviour. In contrast, the HFBA homopolymer shows repulsive Coulombic interactions (positive value), favouring a liquid-state characteristic.24 Neither of these homopolymers is thus suitable for buffering thermomechanical stresses in F-PSCs. In stark contrast, the TFEMA-co-HFBA copolymer achieves a balanced trade-off between Coulombic and van der Waals interactions, which fosters the formation of a fully amorphous elastic phase. Figure 1 c reveals that the TFEMA-co-HFBA copolymer possesses the superior self-healing capability among all samples. This is evidenced by the intimate interchain interpenetration within a 4 ns simulation timeframe, which demonstrates the pivotal role of misplaced-dipole engineering in enhancing self-healing process. The phenomenon can be rationalized by the formation of δ⁻-δ + ∙∙∙δ + -δ⁻ dipole-dipole crosslining interactions. By calculating the spatial charge distribution and quantifying the proportion of positive and negative domains (Fig. S5, Fig. S6, Fig. 1 d-f, and Table S2), we found that both TFEMA and HFBA homopolymers exhibit unbalanced surface charge distributions across polymerization degrees ranging from 1 to 4. In contrast, the TFEMA-co-HFBA copolymer achieves a 1 : 1 positive-to-negative charge ratio, thereby maximizing the intermolecular dipole-dipole interaction for self-healing performance (Fig. 1 g). Elastic self-healing perovskite film enabled by fluoropolymer copolymer The carbonyl (-C = O) and trifluoromethyl (-CF 3 ) groups in fluoropolymer offer binding sites to anchor perovskite species, such as [PbX 6 ] 4 - octahedron frameworks and formamidinium iodide (FAI) moieties via coordination and hydrogen-bonding interactions, which facilitates the formation of highly crystalline perovskite films. To this end, we characterized the resultant samples using X-ray photoelectron spectroscopy (XPS), Fourier transform infrared spectroscopy (FTIR), and nuclear magnetic resonance (NMR) spectroscopy. HFBA homopolymer is initially used as a model additive to explore the fundamental interaction mechanisms. As shown in Fig. S7, the Pb 4f peak exhibits a positive binding energy shift by incorporating HFBA relative to the control sample, meaning that F-Pb bond formation dominates the ion-dipole interactions between the fluoropolymer copolymer and perovskite matrix. This is in contrast to the negative binding energy shift induced by only -C = O ∙∙∙ Pb 2+ coordination.25 , 26 This is further evidenced by the chemical shift observed in the 1 H NMR spectra of HFBA molecules (Fig. S8). Meanwhile, as for the FAI, the characteristic peaks centered at 8.80 ppm (protonated ammonium groups) and 7.85 ppm (hydrogen-bonded carbon atoms in FA + cation) split into two or multiple new peaks upon the incorporation of HFBA (Fig. S9), confirming the formation of hydrogen bonds between -CF 3 groups and FA + cations.27 When extending this analysis to the copolymerized elastomer, analogous interaction signatures are detected in FTIR spectra (Fig. S10). The stretching vibration peaks of C-F bonds shift from 1236 cm - 1 to 1232 cm - 1 and 1189 cm - 1 , while those of C = O bonds shift from 1756 cm - 1 to 1749 cm - 1 and 1713 cm - 1 in both PbI 2 and perovskite film systems. Collectively, these spectral characterizations confirm that the copolymer’s -C = O and -CF 3 groups interact synergistically with Pb 2+ and FA + species to regulate perovskite crystallization. This is evidenced by the enlarged perovskite grain size (Fig. S11) and the elimination of undesirable yellow phase (Fig. S12), thereby yielding enhanced photoluminescence (PL) intensity and prolonged carrier lifetime (Fig. S13). Using the space charge limited current (SCLC) method, we quantify the reduced electron defect density of perovskite film from 4.90 × 10 15 cm -3 to 4.43 × 10 15 cm - 3 as well as the decreased hole defect density from 5.48 × 10 15 cm -3 to 4.10 × 10 15 cm - 3 (Fig. S14), providing direct evidence of efficient multi-type defect passivation. We further confirmed the spatial distribution of the fluoropolymer copolymers within perovskite film via atomic force microscopy-infrared spectroscopy (AFM-IR) by mapping the characteristic C-F vibrational signal at 1189 cm - 1 (Fig. 2 a, b). In contrast to the control film (Fig. S15), the target sample exhibits enhanced signal intensity at perovskite grain boundaries, demonstrating that the fluoropolymer copolymers preferentially locate at grain boundaries and thus construct a structured grain-polymer-grain-polymer block perovskite film. Transmission electron microscopy (TEM) characterization further reveals that the amorphous copolymers encapsulate individual perovskite lattice by the robust interfacial interactions aforementioned (Fig. 2 c). This encapsulation not only effectively impedes the infiltration of ambient moisture and oxygen but also reduces the lead leaching rate from 9.97 ppm min - 1 to 6.03 ppm min - 1 upon immersing film into water (Fig. S16).28 More importantly, the resulting block perovskite film displays excellent polymer-like mechanical properties including a reduced Young’s modulus and robust self-healing capability. As shown in Fig. 2 d-f, non-destructive peak force quantitative nanomechanical mapping test yields a surface averaged Derjaguin-Müller-Toporov (DMT) modulus that decreases from 26.69 GPa to 19.42 GPa for the copolymer modified perovskite film relative to the control sample. The modulus reduction provides a buffer to release the thermal and mechanical stresses. As evidenced by the finite element modelling (Fig. 2 g and Table S3) and grazing-incidence X-ray diffraction (GIXRD) measurement (Fig. S17), the fluoropolymer tailored perovskite film exhibits suppressed interfacial stress accumulation during cyclic bending, and the minimal downshift of the characteristic XRD diffraction peak confirms the copolymer’s efficacy in relieving residual lattice stress. This stress-alleviating effect can effectively suppress the formation of detrimental transgranular cracks after repeated bending cycles (Fig. S18). Similar to organic polymers, the intrinsic self-healing capacibility of perovskite film is critical for mitigating defect propagation under operational stresses, thereby enhancing the operational stability of PSCs.29 , 30 Given that the TFEMA-co-HFBA copolymer preferentially locates at perovskite grain boundaries, we constructed a perovskite-copolymer hybrid model to theoretically investigate its impact on crack evolution dynamics. (MD) simulations demonstrate that robust ion-dipole (perovskite-polymer) and dipole-dipole (polymer-polymer) interactions promote the gradual convergence of adjacent polymer chains. This drives the formation of a new polymer interpenetration network and ultimately achieves efficient crack closure and healing (Fig. 2 h). To experimentally validate this mechanism, we tracked the dynamic morphological evolution of an artificially introduced surface crack on the perovskite film using a three-dimensional (3D) optical microscopy. As shown in Fig. 2 i and Movie S1, the depth of the crack groove decreases significantly after thermal treatment at 60°C for 3 min, indicative of exceptional self-healing behavior. This structural recovery is further corroborated by atomic force microscopy (AFM) characterization. As depicted in Fig. 2 j, the mechanically fatigued perovskite film initially features numerous transverse cracks and a roughness of 20.5 nm, and subsequently becomes smooth after thermal treatment, with the roughness reduced to 10.5 nm due to the disappearance of microfractures. Collectively, these findings verify that TFEMA-co-HFBA copolymer modification not only regulates the perovskite film quality but also serves as an elastic buffer to mitigate residual stress accumulation. This dual functionality is thus anticipated to substantially enhance both the PCE and operational stability of PSCs, with particular efficacy for flexible device architectures. Solar cell performance and stability We fabricated inverted PSCs to systematically evaluate the impacts of the fluoropolymer copolymer on photovoltaic performances. For rigid device with an architecture of ITO/NiO x /4PABcz/Al 2 O 3 /perovskite/PiPl/PC 61 BM/BCP/Ag (Fig. 3 a), systematic optimization studies reveal that a copolymer concentration of 1 mg mL - 1 yields the best device performance (Fig. S19). Current density-voltage ( J - V ) characteristics and the corresponding photovoltaic data (Fig. 3 c and Table S4) show that the champion rigid device incorporating the target copolymer achieves a boosted PCE of 26.83% with a short-circuit current density ( J SC ) of 26.04 mA cm - 2 , an open-circuit voltage ( V OC ) of 1.195 V, and a fill factor (FF) of 86.29%. This performance outperforms the control rigid device, which exhibit a PCE of 25.70%, a J SC of 26.03 mA cm - 2 , a V OC of 1.188 V, and an FF of 83.10%. More importantly, the optimal copolymer dosage (Fig. S20 and Table S5) enables the F-PSCs with the structure of PEN/ITO/4PABcz/Al 2 O 3 /perovskite/PiPl/PC 61 BM/BCP/Ag to deliver a champion PCE of 25.54% (Fig. 3 e, g), which is substantially higher than of the control flexible device (24.72%). The increase of the average PCE from 25.38% to 26.45% for rigid devices (Fig. 3 b) and from 24.47% to 25.22% for flexible scenario (Fig. 3 f) demonstrates the exellent reproducibility of the fluoropolymer copolymer modification strategy.31–33 Furthermore, the integrated J SC derived from external quantum efficiency (EQE) measurements (Fig. 3 d, h) and the steady PCEs (Fig. S21) are in a good agreement with the J - V results, confirming the positive regulatory effect of the elastomeric fluoropolymer copolymer on device photovoltaic performances. We ascribe the enhanced photovoltaic performance to the improved perovskite film quality.34 A suite of electrochemical characterizations were conducted to reveal the charge transfer, extraction, and recombination dynamics in the PSCs. Benefiting from the reduced defect density, the solar cells incorporating TFEMA-co-HFBA copolymer exhibit a stronger built-in electric field (the intrinsic driving force for charge separation) and a larger recombination resistance (Fig. S22), indicative of suppressed charge carrier loss for maximizing power output. This conclusion is further corroborated by a reduced ideality factor closer to unity derived from illumination intensity dependent J - V measurements, and a marked decrease in leakage current (Fig. S23). To assess the practical viability of the devices, we monitored the performance evolution under steady-state external conditions, including persistent storage and operational stress. Under ambient air conditions (25°C, 20 ~ 40% relative humidity), the unencapsulated device modified with the moisture-resistant fluoropolymer maintains 90% of its initial efficiency after 2560 h (107 days), significantly outperforming the reference device with only 70% PCE retention after a mere 1200 h (50 days) (Fig. 3 i and Fig. S24). Under MPP tracking conditions with equivalent 1-sun illumination, the target device exhibits negligible PCE degradation after 1100 h of continuous operation. In stark contrast, the control device undergoes rapid performance decay, retaining only ~ 70% of its initial PCE after 850 h (Fig. 3 j). The enhanced device stability is further corroborated by the changes of PL properties of perovskite films, which is highly correlated with defect evolution and halogen segregation behaviour.35 Comparative analysis of PL data for the control and target films prior to aging demonstrates a blue-shifted PL emission wavelength (Fig. 3 j, k, m, n) and increased PL intensity (Fig. S25) in the copolymer modified films, as evidenced by PL mapping images and corresponding statistical distributions. These observations confirm effective defect passivation, consistent with the previously documented improvements in perovskite film crystallinity.36-38After accelerated aging tests under persistent light irradiation and high-humidity ambient exposure for over 90 days, the control sample exhibits the formation of heterogeneous impurities (e.g. PbI 2 and yellow perovskite phase), which are absent in the target film (Fig. S26). The emergence of these impurities in the control device induces PL intensity quenching and a red-shift PL emission peak. In contrast, the incorporation of TFEMA-co-HFBA copolymer mitigates the magnitude of PL peak shift, indicative of a more robust perovskite lattice that resists harsh environmental stresses and suppresses detrimental phase separation.39 , 40 Thermal-mechanical cycling stability of F-PSCs For F-PSCs, the tolerance to dynamic thermalmechanical stimuli is far more critical than that for their rigid counterparts.41 , 42 We conducted temperature-dependent XRD characterizations to understand the structural evolution of perovskite films and identify the impacts of fluoropolymer copolymer. As the temperature increases from − 20°C to 80°C, the control film exhibits a progressive shift of the characteristic (001) diffraction peak toward smaller angles (Fig. 4 a and Fig. S27), corresponding to lattice expansion of the (001) plane from 6.125 Å to 6.139 Å and a resultant perovskite lattice deformation ratio of 0.23%. In a PSC, such temperature-induced lattice expansion inevitably accelerates ion migration and promotes defect formation.43 In contrary, the XRD profiles for the fluoropolymer copolymer modified perovskite film are insensitive to temperature, with the (001) lattice deformation ratio of mere 0.01% (Fig. 4 b and Fig. S28). This exceptional thermal stability stems primarily from the preferential segregation of the copolymer at perovskite grain boundaries, which serves as an elastic buffer to suppress thermomechanical stress and thereby inhibits lattice expansion under elevated temperatures. As shown in Fig. 4 c, after 10 thermal cycles, the characteristic XRD peaks of the control perovskite film shift toward higher angles, indicating reduced interplanar spacing caused by out-of-plane compressive stress and in-plane tensile strain. In contrast, the target film exhibits negligible peak displacement, reflecting suppressed lattice strain evolution. This demonstrates that the carefully designed fluoropolymer copolymer effectively mitigates thermally induced structural deformation, thereby enhancing the mechanical resilience of the perovskite lattice under repeated thermal stress.44 SEM images indicate that an increasing number of bending cycles induces clearer mechanical cracks across the perovskite grains in Fig. 4 d. Further cycling causes these cracks to propagate and even trigger crystal exfoliation. In contrast to the control sample, the copolymer modified perovskite film features narrower and shallower cracks. Furthermore, the thermal treatment enables the target perovskite film to self-heal, which further enhances its mechanical stability. We further evaluated the device stability against mechanically and thermally induced stresses. Cyclic bending tests performed at a fixed radius of 4 mm (Fig. 4 e, Fig. S29) reveal that the F-PSC with copolymer modified perovskite has a higher PCE retention than that of control device after 3000 cycles, which is attributed to the superior polymer-like mechanical feature of the fluoropolymer modified perovskite film. Taking advantage of the self-healing ability of the modified perovskite film, we applied intermittent thermal treatment at 60°C for 3 min for effectively recovering the photovoltaic performances of F-PSCs. The control solar cell shows no obvious performance recovery under the same conditions. Notably, the target device retains over 90% of its initial PCE after 11,000 bending cycles, standing in sharp contrast to the control device that has mere 80.2% PCE retention after 5,000 cycles. This significant improvement in mechanical stability is attributed to the incorporation of elastic fluoropolymer copolymer, which not only reduces the Young’s modulus of the perovskite film but also enhances the interfacial interaction between the perovskite layer and adjacent functional layers (e.g. 4PABcz and PC 61 BM) (Fig. S30), thus suppressing crack formation under external stress. A comparison of the bending-induced efficiency loss of our devices with the data reported in previous studies demonstrates that our F-PSC exhibits superior mechanical stability (Fig. 4 f, Table S6). Thermal cycling stability was assessed by subjecting PSCs to cyclic temperature changes between 0°C (30 min) and 80°C (30 min). As shown in Fig. 4 g and Fig. S31, the control device retains only 85% of its initial PCE after 200 thermal cycles, whereas the TFEMA-co-HFBA copolymer modified PSC maintains over 90% PCE after 500 cycles. By comparing the thermal cycling-induced efficiency degradation of our devices with that documented in prior literatures (Fig. 4 h and Table S7), we find that the well-established F-PSC achieves the highest thermal cycling stability among similar devices reported to date. Collectively, these results confirm the exceptional operational stability of TFEMA-HFBA copolymer tailored F-PSCs with self-healing capacity under extreme temperature fluctuations, which lays a solid foundation for their practical application on complex curved surfaces. Discussion In summary, we synthesized a misplaced-dipole engineered fluoropolymer elastomer that enables the formation of high-quality perovskite films with polymer-like block characteristics, including a reduced Young’s modulus, elastic grain boundaries, and self-healing capability. These features significantly enhance the tolerance of perovskite lattice to mechanical deformation and temperature fluctuations. By leveraging this elastomer modification strategy, we achieve F-PSCs with a champion PCE of 25.54%, alongside 26.83% in rigid counterparts. Beyond improvement in static photostability and storage stability, the stress-buffering effect imparted by the elastomer at perovskite grain boundaries enables the F-PSC to retain 90% of its initial PCE after 11,000 bending cycles and 500 thermal cycles between 0°C and 80°C. Collectively, this work establishes a versatile and scalable pathway for the development of high-efficiency and ultra-stable F-PSCs, paving the way for their commercialization. Methods Materials and reagents All the commercial materials were used as received without any future purification, including lead iodide (PbI 2 , 99.99%, Xi’an Polymer Light Technology Corp.), lead bromide (PbBr 2 , 99.99%, Xi’an Polymer Light Technology Corp.), cesium Iodide (CsI, 99.9%, Sigma-Aldrich), formamidinium iodide (FAI, greatcell solar), methylammonium bromide (MABr, greatcell solar), methylamine hydrochloride (MACl, greatcell solar), (4-(9′-Phenyl-9H,9’H-[3,3′-bicarbazol]-9-yl)butyl)phosphonic Acid (4PABcz, Suzhou Liwei New Materials Technology Co., Ltd), [6,6]-phenyl C 61 butyric acid methyl ester (PC 61 BM, 99.5%, Xi’an Polymer Light Technology Corp.), 2,9-dimethyl-4,7-diphenyl-1,10-Phenanthroline (BCP, 99% Xi’an Polymer Light Technology Corp.), nickel oxide (NiO x , 99.99%, Aladdin), 2,2,2-Trifluoroethylmethacrylate (TFEMA, Aladdin, 98%), 2,2,3,4,4,4-Hexafluorobutyl acrylate (HFBA, Aladdin, > 95%), 2,2'-Azobis(2-methylpropionitrile) (AIBN, Aladdin, 99%), anhydrous N,N-dimethylformamide (DMF, 99.99%, Sigma-Aldrich), dimethylsulfoxide (DMSO, 99.99%, Sigma-Aldrich), methanol (MS, Aladdin) isopropyl alcohol (IPA 99.9%, Aladdin), chlorobenzene (CB 99.99%, Sigma-Aldrich), Flexible conductive substrate PEN/ITO (Liaoning Optimal New Energy Technology Co., Ltd). Preparation of perovskite precursors and perovskite films Preparation of rigid PSCs Preparation of hole/electron transport solutions NiOₓ dispersion was prepared by dissolving 7.5 mg of NiOₓ powder in 1 mL of deionized water. A 4PABcz solution was formulated by dissolving 0.54 mg of 4PABcz in 1 mL of anhydrous MS. For the Al 2 O 3 dispersion, 10 µL of nano-Al₂O₃ dispersion liquid was diluted with 990 µL of IPA to a total volume of 1 mL. The PiPl solution was prepared by dissolving 0.4 mg of PiPl in 1 mL of IPA. PC₆₁BM solution was obtained by dissolving 20 mg of PC₆₁BM in 1 mL of anhydrous CB. Finally, a BCP solution was prepared by dissolving 0.5 mg of BCP in 1 mL of IPA. Preparation of pristine and molecule-doped perovskite precursor solutions 1 : 71 uL PETMP, 85 µL PAE and 2 mg AIBN in 2 mL DMF and DMSO solution (volume ratio 4:1). Consider the above solution as a concentrated solution of 100 mg/mL, and continue to dilute it with a mixed solvent of DMF and DMSO (volume ratio 4:1) to 1 mg/mL for later use. Preparation of pristine and molecule-doped perovskite precursor solutions 2 : 6.7 mg MABr, 8.8 mg MACl, 18.2 mg CsI, 196.1 mg FAI, 22 mg PbBr 2 and 602 mg PbI 2 (Excess of PbI 2 by 5%) were dissolved in 865 µL DMF and DMSO solution (volume ratio 4:1), stirred overnight, and used to prepare an undoped precursor solution.For the doping solution, dissolve the perovskite precursor powder in 865 µL of 1 mg/mL cross-linked molecular mixture solution. Preparation of PSCs ITO substrates were sequentially cleaned with detergent, deionized water, and ethyl alcohol. The cleaned ITO substrates were dried with high-pressure N 2 gas flow, and then plasma modified for 1 min. The NiO x layer was formed by spin-coating the precursor solution at 3000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 150 ℃ for 20 min in the air. Then transfer it to the N 2 glove box for further preparation. The 4PABcz was formed by spin-coating the precursor solution at 3000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 100 ℃ for 10 min. The Al 2 O 3 was formed by spin-coating the precursor solution at 5000 rpm (acceleration 4000 rpm) for 30 s and then short-baking at 100 ℃ for 1 min. The as-prepared perovskite precursor solution was spin-coated onto the hole transport substrate with speed of 1000 rpm (acceleration 1000 rpm) for 5 s and 5000 rpm (acceleration 3000 rpm) for 30 s. During the last 15 s of the spinning process, the liquid film was treated by drop-casting chlorobenzene solvent (150 ul). The substrates were annealed on a hot plate at 100℃ for 30 min. After cooling the perovskite film, a precursor solution was spin coated at 5000rpm (acceleration 3000 rpm) for 30 seconds, and then baked briefly at 100 ℃ for 5 min to form a PiPl layer. The PC 61 BM layer was formed by spin-coating the precursor solution at 1500 rpm (acceleration 3000 rpm) for 35 s and then short-baking at 100 ℃ for 10 min. The BCP layer was formed by spin-coating the precursor solution at 4000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 100 ℃ for 3 min. Finally, 100 nm thick of Ag was deposited on top by thermal evaporation. The active area of PSCs is 0.03245 cm 2 as determined by the mask. Preparation of F-PSCs Preparation of hole/electron transport solutions the NiO x solution was prepared by dissolving 15 mg NiO x in 1 mL of deionized water. The solution was prepared by dissolving 1.08 mg 4PABcz in 1 mL of anhydrous MS (Add 10 µL of DMF). The Al 2 O 3 solution was prepared by dissolving 10 µL Nano Al 2 O 3 dispersion liquid in 990 µL of DMF. The remaining steps are consistent with the preparation steps of rigid perovskite devices. Preparation of flexible perovskite devices use thermosetting tape to bond a blank glass substrate onto a PEN/ITO substrate. Then wash with soapy water, deionized water, and ethanol, and sonicate in each solution for 15 minutes. Subsequently, before use, clean the substrate by blowing it dry with nitrogen and then treating it with ultraviolet ozone for 10 min. The NiOx layer was formed by spin-coating the precursor solution at 2000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 150 ℃ for 15 min in the air. Then transfer it to the N 2 glove box for further preparation. The 4PABcz was formed by spin-coating the precursor solution at 4000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 100 ℃ for 10 min. The Al 2 O 3 was formed by spin-coating the precursor solution at 5000 rpm (acceleration 4000 rpm) for 30 s and then short-baking at 100 ℃ for 1 min. The as-prepared perovskite precursor solution was spin-coated onto the hole transport substrate with speed of 1000 rpm (acceleration 1000 rpm) for 5 s and 5000 rpm (acceleration 3000 rpm) for 30 s. During the last 15 s of the spinning process, the liquid film was treated by drop-casting chlorobenzene solvent (150 ul). The substrates were annealed on a hot plate at 100℃ for 30 min. After cooling the perovskite film, a precursor solution was spin coated at 5000rpm (acceleration 3000 rpm) for 30 seconds, and then baked briefly at 100 ℃ for 5 min to form a PiPl layer. The PC 61 BM layer was formed by spin-coating the precursor solution at 1500 rpm (acceleration 3000 rpm) for 35 s and then short-baking at 100 ℃ for 10 min. The BCP layer was formed by spin-coating the precursor solution at 4000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 100 ℃ for 3 min. Finally, 120 nm thick of Ag was deposited on top by thermal evaporation. The active area of PSCs is 0.03245 cm 2 as determined by the mask. Characterizations: Fourier transform infrared (FTIR) spectra were performed by the infrared spectrometer (Nicolet iS50 FT-IR). The X-ray photoelectron spectroscopy (XPS) studies were performed using a Thermo-VG Scientific ESCALAB 250 photoelectron spectrometer equipped with a monochromated Al Kα (1486.6 eV) X-ray source. The composition was characterized by X-ray diffraction (Rigaku, with Cu-Kα radiation of λ = 0.15418 nm). Grazing incidence X-ray diffraction (GIXRD) measurements were carried out with a Bruker D8 grazing incidence X-ray diffractometer. Photoluminescence (PL) and time-resolved PL (TRPL) were measured by FLS980 (Edinburgh Instruments Ltd) with an excitation at 470 nm. The PL mapping was executed by a Laser Scanning Confocal Microscope (Enlitech, SPCM-1000) equipped with a 470 nm pulse laser and a galvo-based scanner. High resolution transmission electron microscopy (HRTEM) morphology was performed using a FEI Tecnai G2 F30 instrument manufactured by FEI Company, USA. Scanning electron microscopy was measured via a scanning electron microscope (Hitachi S8230). Atomic Force Microscopy (AFM) topography characterization was performed using a Bruker Dimension ICON Microscope. In-situ variable-temperature XRD measurements were performed using a SmartLab-9kW diffractometer. The instrument model for Nano IR testing is nano IR2-FS (Anasys Instruments), and this multifunctional nano-infrared spectrometer includes an atomic force microscope for detecting morphology. Optical microscope images were tested using a Leica polarizing microscope (DM4 P). The instrument model used for liquid nuclear magnetic resonance hydrogen spectroscopy is AVANCE NEO 400MHZ. Water contact angle measurements were performed using a Dataphysics OCA20 contact angle meter. The lead content in water for the lead leaching test was determined using an iCE™ 3500 atomic absorption spectrometer (AAS). The photovoltaic performance of the PSCs was collected with a source meter (Keithley 2400), with the light source (SS-F5-3A, Enlitech) via reverse scanning from 1.22 to -0.1 V or forward scanning from − 0.1 to 1.22 V at a scanning speed of 50 mV/s, and the quantum efficiency was tested using a quantum efficiency measurement system (QE-R, Enlitech). The illumination stability of devices was executed under one sun-equivalent white light. All devices were taken out from the chamber and tested at different time intervals under a separate solar simulator (AM 1.5G, 100 mW/cm²) for J-V characterization. The device MPP tracking was measured with a CHI 1000C potentiostat under a white LED lamp (100 mW/cm²) in an N₂ glove box. Then, the devices were measured with a maximum power point (MPP) tracking routine under continuous 1 sun illumination. The MPP was updated every 300 s by a standard perturb and observation method, and the average temperature was kept at ~ 50 ℃. Computational Methods: Molecular Dynamics (MD) Simulation Molecular dynamics (MD) simulations were performed using GROMACS 2021 software with the AMBER99SB force field. The simulation box was constructed via GROMACS tools. Prior to production simulations, energy minimization and a 100 ps pre-equilibration simulation were sequentially carried out to allow system relaxation. For production simulations, Newton's equations of motion were integrated using the leapfrog algorithm, and the V-rescale temperature coupling scheme was adopted with a simulation temperature of 298.15 K. Finite Element (FE) Simulation Finite element (FE) simulations were performed using ABAQUS software. The model was simulated based on Hooke's Law, and a rotation about the center was applied to both ends of the model, with each end rotated by approximately 40°. Data Availability The data that support the findings of this study are available within the Article and its Supplementary Information. Source data are provided with this paper. Declarations Competing interests The authors declare no competing interests. Author contributions L. and J. D. conceived the idea. F. L., Y. L., W. Z., Y. L. and J. L. performed the solar cell fabrication, characterization, and optimization. B. H. contributed to the PL measurement. Q. G. and Y. Z. measured the EQE. J. D. involved in the data analysis and discussion. Q. C. and Q. T.directed this work. J. D. and J. D. wrote the first draft of the manuscript. All the authors revised and approved the manuscript. Acknowledgements The authors gratefully acknowledged financial support provided by the National Natural Science Foundation of China (62304124, 62374105, 62204098, 22179051, 22309107, 52472259), Natural Science Foundation of Shandong Province, China (ZR2025MS892, ZR2024QE036, ZR2023QB281, ZR2024QB021, ZR2025MS1014), National Key Research and Development Program of China (2022YFE4200500), Qingdao Natural Science Foundation (25-3-1-13-zyyd-jch), Special Fund of Taishan Scholar Program of Shandong Province (tstp20250728, tsqnz20221141). References Kojima A, Teshima K, Shirai Y, Miyasaka T (2009) Organometal Halide Perovskites as Visible-Light Sensitizers for Photovoltaic Cells. J Am Chem Soc 131:6050–6051 Lee MM, Teuscher J, Miyasaka T, Murakami TN, Snaith HJ (2012) Efficient hybrid solar cells based on meso-superstructured organometal halide perovskites. Science 338:643–647 Jeon NJ et al (2015) Compositional engineering of perovskite materials for high-performance solar cells. 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Adv Mater 35:e2300302 Cai W et al (2025) In Situ Dual-Region Selective Anchoring of Zwitterionic Gel Enables Efficient and Mechanically Durable Flexible Perovskite Solar Cell. Adv Energy Mater Song F et al (2025) Buried Interface Modification for High Performance and Stable Inverted Perovskite Solar Cells. Angew Chem Int Ed Engl 64:e202516012 Yuan G et al (2023) Inhibited Crack Development by Compressive Strain in Perovskite Solar Cells with Improved Mechanical Stability. Adv Mater 35:e2211257 Additional Declarations There is NO Competing Interest. Supplementary Files VideoS1.mp4 Video S1 Supportinginformation.docx Supporting information Cite Share Download PDF Status: Under Review Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-8773234","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":585867130,"identity":"06399e03-a631-438c-b185-3e3aba2ad5bb","order_by":0,"name":"Qi 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Molecular dynamics simulation of the self-healing behaviour for TFEMA homopolymer, HFBA homopolymer, and HFBA-co-TFEMA copolymer (1:1 molar ratio); Surface charge distribution profiles (positive and negative proportions) of the polymers with varying degrees ofpolymerization: \u003cstrong\u003ed\u003c/strong\u003e TFEMA homopolymer; \u003cstrong\u003ee\u003c/strong\u003e HFBA homopolymer; \u003cstrong\u003ef\u003c/strong\u003e HFBA-co-TFEMA copolymer (1:1 molar ratio).\u003cstrong\u003eg\u003c/strong\u003e Correlationbetween surface charge proportion and intermolecular force magnitudes.\u003c/p\u003e","description":"","filename":"floatimage1.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-8773234/v1/0530360362d2e798368d1bc4.jpeg"},{"id":101943742,"identity":"74311d1a-ec7a-4d8a-9577-9a1b52d48c37","added_by":"auto","created_at":"2026-02-05 09:43:12","extension":"jpg","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":3308060,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eElastic self-healing perovskite film enabled by fluoropolymer copolymer:\u003c/strong\u003e \u003cstrong\u003ea\u003c/strong\u003eAtomic force microscopy (AFM) and \u003cstrong\u003eb\u003c/strong\u003e Nano-IR mapping images of the fluoropolymer copolymer modified perovskite film; \u003cstrong\u003ec\u003c/strong\u003e TEM image of fluoropolymer copolymer modified perovskite halide. \u003cstrong\u003ed\u003c/strong\u003e Schematic diagram of the fluoropolymer copolymer cross-linking strategy. \u003cstrong\u003ee\u003c/strong\u003e Elastic modulus mapping images of the perovskite films without and with copolymer modification and \u003cstrong\u003ef\u003c/strong\u003ecorresponding Young’s modulus distribution profiles. \u003cstrong\u003eg\u003c/strong\u003e Schematic illustration of finite element simulation for F-PSCs under mechanical stress. \u003cstrong\u003eh\u003c/strong\u003e Molecular dynamics simulation for the self-healing process in copolymer modified perovskite films. \u003cstrong\u003ei\u003c/strong\u003e 3D morphological microscopic images depictingthe self-healing process ofperovskite films. \u003cstrong\u003ej \u003c/strong\u003eAFM images tracking the self-healing process of perovskite films.\u003c/p\u003e","description":"","filename":"2.jpg","url":"https://assets-eu.researchsquare.com/files/rs-8773234/v1/1d36a63cf92960fcccb0e862.jpg"},{"id":101943634,"identity":"387d2ba2-380a-4038-8291-f74eda9c5512","added_by":"auto","created_at":"2026-02-05 09:42:36","extension":"jpeg","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":1008703,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eSolar cell performance and stability: a\u003c/strong\u003e Device architecture schematic of rigid PSCs. \u003cstrong\u003eb\u003c/strong\u003e Satistical distributions of PCEs for the control and target rigid perovskite devices. \u003cstrong\u003ec\u003c/strong\u003e Current density-voltage (\u003cem\u003eJ\u003c/em\u003e-\u003cem\u003eV\u003c/em\u003e) curves of the champion rigid devices.\u003cstrong\u003e d\u003c/strong\u003e EQE spectra and integrated \u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e of target and control rigid devices. \u003cstrong\u003ee\u003c/strong\u003e Device architecture schematic of F-PSCs.\u003cstrong\u003e f\u003c/strong\u003e Statistical distributions of PCEs for control and target flexible perovskite devices. \u003cstrong\u003eg \u003c/strong\u003e\u003cem\u003eJ\u003c/em\u003e-\u003cem\u003eV\u003c/em\u003e curves of the champion flexible devices. \u003cstrong\u003eh\u003c/strong\u003e EQE spectra and integrated \u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e of the target and control flexible devices. \u003cstrong\u003ei\u003c/strong\u003e Normalized PCE evolution of the control and copolymer modified F-PSCs under persistent 1 sun illumination. \u003cstrong\u003ej\u003c/strong\u003e Statistical distribution statistics of PL peak intensities. \u003cstrong\u003ek\u003c/strong\u003e PL mapping images of F-PSCs under MPP tracking conditions with equivalent 1-sun illumination. \u003cstrong\u003el\u003c/strong\u003e Normalized PCE evolution of F-PSCs under ambient conditions (20-40% relative humidity, room temperature). \u003cstrong\u003em\u003c/strong\u003e Statistical distribution of PL peak intensities. \u003cstrong\u003en\u003c/strong\u003e PL mapping images of F-PSCs under ambient conditions (20-40% relative humidity, room temperature).\u003c/p\u003e","description":"","filename":"floatimage3.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-8773234/v1/b10a1160527a25af7404c743.jpeg"},{"id":101937796,"identity":"6849c226-dee3-4502-a37b-9ada379f127a","added_by":"auto","created_at":"2026-02-05 08:52:03","extension":"jpeg","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":1081928,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eThermal-mechanical cycling stability of F-PSCs:\u003c/strong\u003e In-situ variable-temperature x-ray diffraction (XRD) patterns of the \u003cstrong\u003ea\u003c/strong\u003e control and \u003cstrong\u003eb\u003c/strong\u003etarget perovskite films. \u003cstrong\u003ec\u003c/strong\u003e XRD patterns of the control and target perovskite films before and after 10 thermal cycles over the temperature range of -20 °C to 80 °C. \u003cstrong\u003ed\u003c/strong\u003e SEM images of perovskite films. \u003cstrong\u003ee\u003c/strong\u003e Mechanical stability test of unencapsulated F-PSCs. \u003cstrong\u003ef\u003c/strong\u003e Mechanical stability of F-PSCs reported in 2025. \u003cstrong\u003eg\u003c/strong\u003e Thermal cycling stability test of unencapsulated F-PSCs. \u003cstrong\u003eh\u003c/strong\u003e Thermal cycling stability of F-PSCs reported in recent years.\u003c/p\u003e","description":"","filename":"floatimage4.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-8773234/v1/c056685ec67946fe04e346cc.jpeg"},{"id":102301211,"identity":"3d23ba99-e179-4589-acd1-27adcfe4ad0c","added_by":"auto","created_at":"2026-02-10 11:20:32","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":7291015,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-8773234/v1/a8bb8368-17b9-4670-a408-0b043c203f25.pdf"},{"id":102298659,"identity":"38813cff-bb38-43c9-81b5-9ff007e95244","added_by":"auto","created_at":"2026-02-10 10:56:49","extension":"mp4","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":9603205,"visible":true,"origin":"","legend":"Video S1","description":"","filename":"VideoS1.mp4","url":"https://assets-eu.researchsquare.com/files/rs-8773234/v1/2711e1c5a7dd398246341fba.mp4"},{"id":101943701,"identity":"05978527-70a7-4598-ae68-ac706b3a91ea","added_by":"auto","created_at":"2026-02-05 09:42:55","extension":"docx","order_by":2,"title":"","display":"","copyAsset":false,"role":"supplement","size":22147196,"visible":true,"origin":"","legend":"Supporting information","description":"","filename":"Supportinginformation.docx","url":"https://assets-eu.researchsquare.com/files/rs-8773234/v1/7f48f78618c846fd3eebf151.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"Misplaced-dipole engineered repairable fluoropolymer elastomer for flexible perovskite solar cell with excellent thermal-mechanical cycling resistance","fulltext":[{"header":"Introduction","content":"\u003cp\u003eOrganometallic halide perovskites, endowed with excellent low-temperature processibility, have emerged as promising candidates for assembling flexible perovskite solar cells (F-PSCs).1\u0026ndash;4 In contrast to rigid counterparts fabricated on glass substrates, the intrinsic flexibility and lightweight nature of F-PSCs enables their conformal integration onto arbitrary surfaces, irrespective of morphological irregularties.5\u003csup\u003e,\u003c/sup\u003e6 To date, the power conversion efficiency (PCE) of F-PSCs has undergone rapid advancement, with record value reaching 25.74% for laboratory-scale single-junction devices and 21.6% for large-area modules.7\u0026ndash;10 Nevertheless, substantial disparities persist between flexible and rigid devices in terms of both PCE and long-term perational stability. The state-of-the-art strategies for boosting efficiency and stability of F-PSCs predominantly center on buried interface engineering and additive modification to serve to strengthen interfacial adhesion, suppress defect formation, and even enhance the mechanical toughness of perovskite grain boundaries.11\u0026ndash;14 Despite substantial advances, F-PSCs remain plagued by irreversible performance degradation triggered by deformation-induced exterinal stress. A primary reason stems from the utilization of flexible polymer substrates, such as polyethylene terephthalate (PET) and polyneopentamide glycolate (PEN), whose morphological characterstics exhibit strong dependence on external stress. This inherent substrate sensitivity readily induces buried perovskite interface delamination and bulk perovskite film cracking during both the annealing-cooling fabriation process and day-night operational cycles.11\u003csup\u003e,\u003c/sup\u003e12 Therefore, temperature fluctuation not only induces spatial distortion of the perovskite lattice but also drives the thermal expansion and contraction of flexible polymer substrates, thus making huge obstacles to the practical use of F-PSCs. Especially following dynamic operation under diurnal, seasonal, and climatic variations, F-PSCs are more prone to structural and chemical degradation of the perovskite active layer, which in turn impedes interfacial and intergranular charge transfer and fosters the formation of substantial defects, ultimately compromising device durability. In this context, strategies targeting the isolated reinforcement of the perovskite lattice such as molecular grain confinement or the incorporation of poly(methyl methacrylate)/graphene layers well-established for rigid scenarios are inadequate for F-PSCs.13\u003csup\u003e,\u003c/sup\u003e14 Therefore, engineering high-quality perovskite films featured with polymer-like mechanical behaviors including low Young\u0026rsquo;s modulus, high elasticity, and moderate rebound resilience, is equally critical to ensure compatibility with flexible substrates. Nevertheless, achieving this suite of mechanical requirements entails an inherent trade-off with the performance metrics and stability criteria aforementioned.\u003c/p\u003e \u003cp\u003eTo address this great challenge, herein we synthesized a misplaced-dipole engineered fluoropolymer elastomer by a hardness-softness synergistic coupling strategy, using a copolymer derived from the hard chain monomer 2,2,2-trifluoroethyl methacrylate (TFEMA) and the soft chain monomer 2,2,3,4,4,4-hexafluorobutyl acrylate (HFBA). Complementary molecular dynamics simulations and density functional theory (DFT) calculations confirm that this elastomer integrates high mechanical strength, robust hydrophobicity, and excellent self-healing capability. By incorporating this elastomer, we construct a hierarchically structured perovskite film with a distinctive grain-polymer-grain-polymer architecture, which endows the perovskite film with polymer-like mechanical traits including reduced Young\u0026rsquo;s modulus and efficient self-healing behavior to release the thermal and mechanical stresses. Consequently, the lattice deformation ratio of the perovskite film is significantly suppressed from 0.23% to mere 0.01% over a wide temperature range of -20 \u003csup\u003eo\u003c/sup\u003eC to 80 \u003csup\u003eo\u003c/sup\u003eC. Furthermore, the ion-dipole interactions between the fluoropolymer elastomer and perovskite lattice not only regulate the perovskite crystallization kinetics but also improve the film quality. Benefiting from these synergistic effects, the resultant F-PSCs deliver a champion PCE to 25.54% under 1-sun irradiation. More impressively, the devices show negligible degradation after 1100 h under 1-sun MPP tracking. They also exhibit exceptional environmental roubustness under harsh operational conditions, retaining 90% of their initial PCEs after 11,000 bending cycles and maintaining 90% following 500 thermal cycles (0\u0026deg;C \u0026harr;80\u0026deg;C).\u003c/p\u003e"},{"header":"Results","content":"\u003cp\u003eMolecular design and self-healing mechanisms\u003c/p\u003e \u003cp\u003eThe fluoropolymer elastomer was synthesized via the copolymerization of equimolar TFEMA and HFBA monomers, with 2,2-azobisisobutyronitrile (AIBN) serving as the initiator.15 Upon thermal curing at 60\u0026deg;C for 3 min, the homogeneous liquid reactive precursor solution gradually solidified into a solid elastomer (Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e). As is well known, fluorinated polymers inherently exhibit exceptional chemical stability and superior hydrophobicity, stemming from the high bond dissociation energy of covalent C-F linkages.16\u003csup\u003e,\u003c/sup\u003e17 This hydrophobic characteristic is conducive to enhancing the moisture resistance of the resultant perovskite film (Fig. S2).18 Compared to the traditional polymer additives, the designed copolymer synergistically integrates the rigid segmental architecture of TFEMA chain with flexible chain conformation of HFBA moieties, thereby endowing the elastomer with balanced toughness, high stretchability, and efficient reparability. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea, two transparent polymer strips achieved a tensile elongation of 300% relative to their original length following a 3 min self-healing treatment at 60\u0026deg;C. After incorporated into the grain boundaries of perovskite film, the elastomer\u0026rsquo;s self-healing performance and stretchability effectively prolong the service lifetime of flexible perovskite by mitigating lattice strain under cyclic thermal and mechanical stress. The underlying mechanism governing this behavior will be discussed in the following sections.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eAs documented in previous reports,19\u003csup\u003e,\u003c/sup\u003e20 the self-healing capability of the fluoropolymer elastomer is predominantly governed by robust dipole-dipole interactions among the highly polar groups namely -CF\u003csub\u003e3\u003c/sub\u003e, -CF\u003csub\u003e2\u003c/sub\u003e, and -C\u0026thinsp;=\u0026thinsp;O (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb).21\u003csup\u003e,\u003c/sup\u003e22 Specifically, the much larger electronegativity of F atom (χ\u0026thinsp;=\u0026thinsp;4.0) than that of C atom (χ\u0026thinsp;=\u0026thinsp;2.5) drives electron transfer from C atom to F atom, generating permanent dipoles with negatively charged center (δ⁻) and positively charged center (δ\u003csup\u003e+\u003c/sup\u003e). An analogous charge polarization effect arises from the carbon-oxygen pair within C\u0026thinsp;=\u0026thinsp;O moiety.\u003c/p\u003e \u003cp\u003eWhen the fluoropolymer elastomer suffers damage, the robust intermolecular interactions between fractured interfaces compress and shorten the interchain distance by reducing the entropy effect, which is governed by the interfacial interactions. More importantly, we discover that the interpenetrating network derived from the copolymerization of TFEMA and HFBA monomers enables the maximum intermolecular interaction to boost the self-healing kinetics.23 To in-depth understand this mechanism, we performed the molecular dynamics simulation to quantify intermolecular interactions by contrasting TFEMA homopolymer, HFBA homopolymer, and their 1:1 copolymer. As shown in Fig. S3, Fig. S4, and Table \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e, the TFEMA homopolymer exhibits the strongest Coulombic attraction yet the weakest van der Waals force, which accounts for its rigid solid-state behaviour. In contrast, the HFBA homopolymer shows repulsive Coulombic interactions (positive value), favouring a liquid-state characteristic.24 Neither of these homopolymers is thus suitable for buffering thermomechanical stresses in F-PSCs. In stark contrast, the TFEMA-co-HFBA copolymer achieves a balanced trade-off between Coulombic and van der Waals interactions, which fosters the formation of a fully amorphous elastic phase. Figure\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec reveals that the TFEMA-co-HFBA copolymer possesses the superior self-healing capability among all samples. This is evidenced by the intimate interchain interpenetration within a 4 ns simulation timeframe, which demonstrates the pivotal role of misplaced-dipole engineering in enhancing self-healing process. The phenomenon can be rationalized by the formation of δ⁻-δ\u003csup\u003e+\u003c/sup\u003e∙∙∙δ\u003csup\u003e+\u003c/sup\u003e-δ⁻ dipole-dipole crosslining interactions. By calculating the spatial charge distribution and quantifying the proportion of positive and negative domains (Fig. S5, Fig. S6, Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ed-f, and Table S2), we found that both TFEMA and HFBA homopolymers exhibit unbalanced surface charge distributions across polymerization degrees ranging from 1 to 4. In contrast, the TFEMA-co-HFBA copolymer achieves a 1 : 1 positive-to-negative charge ratio, thereby maximizing the intermolecular dipole-dipole interaction for self-healing performance (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eg).\u003c/p\u003e \u003cp\u003eElastic self-healing perovskite film enabled by fluoropolymer copolymer\u003c/p\u003e \u003cp\u003eThe carbonyl (-C\u0026thinsp;=\u0026thinsp;O) and trifluoromethyl (-CF\u003csub\u003e3\u003c/sub\u003e) groups in fluoropolymer offer binding sites to anchor perovskite species, such as [PbX\u003csub\u003e6\u003c/sub\u003e]\u003csup\u003e\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e-\u003c/sup\u003e octahedron frameworks and formamidinium iodide (FAI) moieties via coordination and hydrogen-bonding interactions, which facilitates the formation of highly crystalline perovskite films. To this end, we characterized the resultant samples using X-ray photoelectron spectroscopy (XPS), Fourier transform infrared spectroscopy (FTIR), and nuclear magnetic resonance (NMR) spectroscopy. HFBA homopolymer is initially used as a model additive to explore the fundamental interaction mechanisms. As shown in Fig. S7, the Pb 4f peak exhibits a positive binding energy shift by incorporating HFBA relative to the control sample, meaning that F-Pb bond formation dominates the ion-dipole interactions between the fluoropolymer copolymer and perovskite matrix. This is in contrast to the negative binding energy shift induced by only -C\u0026thinsp;=\u0026thinsp;O ∙∙∙ Pb\u003csup\u003e2+\u003c/sup\u003e coordination.25\u003csup\u003e,\u003c/sup\u003e26 This is further evidenced by the chemical shift observed in the \u003csup\u003e\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003eH NMR spectra of HFBA molecules (Fig. S8). Meanwhile, as for the FAI, the characteristic peaks centered at 8.80 ppm (protonated ammonium groups) and 7.85 ppm (hydrogen-bonded carbon atoms in FA\u003csup\u003e+\u003c/sup\u003e cation) split into two or multiple new peaks upon the incorporation of HFBA (Fig. S9), confirming the formation of hydrogen bonds between -CF\u003csub\u003e3\u003c/sub\u003e groups and FA\u003csup\u003e+\u003c/sup\u003e cations.27 When extending this analysis to the copolymerized elastomer, analogous interaction signatures are detected in FTIR spectra (Fig. S10). The stretching vibration peaks of C-F bonds shift from 1236 cm\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e to 1232 cm\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e and 1189 cm\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e, while those of C\u0026thinsp;=\u0026thinsp;O bonds shift from 1756 cm\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e to 1749 cm\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e and 1713 cm\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e in both PbI\u003csub\u003e2\u003c/sub\u003e and perovskite film systems. Collectively, these spectral characterizations confirm that the copolymer\u0026rsquo;s -C\u0026thinsp;=\u0026thinsp;O and -CF\u003csub\u003e3\u003c/sub\u003e groups interact synergistically with Pb\u003csup\u003e2+\u003c/sup\u003e and FA\u003csup\u003e+\u003c/sup\u003e species to regulate perovskite crystallization. This is evidenced by the enlarged perovskite grain size (Fig. S11) and the elimination of undesirable yellow phase (Fig. S12), thereby yielding enhanced photoluminescence (PL) intensity and prolonged carrier lifetime (Fig. S13). Using the space charge limited current (SCLC) method, we quantify the reduced electron defect density of perovskite film from 4.90 \u0026times; 10\u003csup\u003e15\u003c/sup\u003e cm\u003csup\u003e-3\u003c/sup\u003e to 4.43 \u0026times; 10\u003csup\u003e15\u003c/sup\u003e cm\u003csup\u003e-\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e\u003c/sup\u003e as well as the decreased hole defect density from 5.48 \u0026times; 10\u003csup\u003e15\u003c/sup\u003e cm\u003csup\u003e-3\u003c/sup\u003e to 4.10 \u0026times; 10\u003csup\u003e15\u003c/sup\u003e cm\u003csup\u003e-\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e\u003c/sup\u003e (Fig. S14), providing direct evidence of efficient multi-type defect passivation.\u003c/p\u003e \u003cp\u003eWe further confirmed the spatial distribution of the fluoropolymer copolymers within perovskite film via atomic force microscopy-infrared spectroscopy (AFM-IR) by mapping the characteristic C-F vibrational signal at 1189 cm\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ea, b). In contrast to the control film (Fig. S15), the target sample exhibits enhanced signal intensity at perovskite grain boundaries, demonstrating that the fluoropolymer copolymers preferentially locate at grain boundaries and thus construct a structured grain-polymer-grain-polymer block perovskite film. Transmission electron microscopy (TEM) characterization further reveals that the amorphous copolymers encapsulate individual perovskite lattice by the robust interfacial interactions aforementioned (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec). This encapsulation not only effectively impedes the infiltration of ambient moisture and oxygen but also reduces the lead leaching rate from 9.97 ppm min\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e to 6.03 ppm min\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e upon immersing film into water (Fig. S16).28 More importantly, the resulting block perovskite film displays excellent polymer-like mechanical properties including a reduced Young\u0026rsquo;s modulus and robust self-healing capability. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ed-f, non-destructive peak force quantitative nanomechanical mapping test yields a surface averaged Derjaguin-M\u0026uuml;ller-Toporov (DMT) modulus that decreases from 26.69 GPa to 19.42 GPa for the copolymer modified perovskite film relative to the control sample. The modulus reduction provides a buffer to release the thermal and mechanical stresses. As evidenced by the finite element modelling (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eg and Table S3) and grazing-incidence X-ray diffraction (GIXRD) measurement (Fig. S17), the fluoropolymer tailored perovskite film exhibits suppressed interfacial stress accumulation during cyclic bending, and the minimal downshift of the characteristic XRD diffraction peak confirms the copolymer\u0026rsquo;s efficacy in relieving residual lattice stress. This stress-alleviating effect can effectively suppress the formation of detrimental transgranular cracks after repeated bending cycles (Fig. S18).\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eSimilar to organic polymers, the intrinsic self-healing capacibility of perovskite film is critical for mitigating defect propagation under operational stresses, thereby enhancing the operational stability of PSCs.29\u003csup\u003e,\u003c/sup\u003e30 Given that the TFEMA-co-HFBA copolymer preferentially locates at perovskite grain boundaries, we constructed a perovskite-copolymer hybrid model to theoretically investigate its impact on crack evolution dynamics. (MD) simulations demonstrate that robust ion-dipole (perovskite-polymer) and dipole-dipole (polymer-polymer) interactions promote the gradual convergence of adjacent polymer chains. This drives the formation of a new polymer interpenetration network and ultimately achieves efficient crack closure and healing (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eh). To experimentally validate this mechanism, we tracked the dynamic morphological evolution of an artificially introduced surface crack on the perovskite film using a three-dimensional (3D) optical microscopy. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ei and Movie S1, the depth of the crack groove decreases significantly after thermal treatment at 60\u0026deg;C for 3 min, indicative of exceptional self-healing behavior. This structural recovery is further corroborated by atomic force microscopy (AFM) characterization. As depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ej, the mechanically fatigued perovskite film initially features numerous transverse cracks and a roughness of 20.5 nm, and subsequently becomes smooth after thermal treatment, with the roughness reduced to 10.5 nm due to the disappearance of microfractures. Collectively, these findings verify that TFEMA-co-HFBA copolymer modification not only regulates the perovskite film quality but also serves as an elastic buffer to mitigate residual stress accumulation. This dual functionality is thus anticipated to substantially enhance both the PCE and operational stability of PSCs, with particular efficacy for flexible device architectures.\u003c/p\u003e \u003cp\u003eSolar cell performance and stability\u003c/p\u003e \u003cp\u003eWe fabricated inverted PSCs to systematically evaluate the impacts of the fluoropolymer copolymer on photovoltaic performances. For rigid device with an architecture of ITO/NiO\u003csub\u003ex\u003c/sub\u003e/4PABcz/Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e/perovskite/PiPl/PC\u003csub\u003e61\u003c/sub\u003eBM/BCP/Ag (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea), systematic optimization studies reveal that a copolymer concentration of 1 mg mL\u003csup\u003e-\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e yields the best device performance (Fig. S19). Current density-voltage (\u003cem\u003eJ\u003c/em\u003e-\u003cem\u003eV\u003c/em\u003e) characteristics and the corresponding photovoltaic data (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ec and Table S4) show that the champion rigid device incorporating the target copolymer achieves a boosted PCE of 26.83% with a short-circuit current density (\u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e) of 26.04 mA cm\u003csup\u003e-\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e\u003c/sup\u003e, an open-circuit voltage (\u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e) of 1.195 V, and a fill factor (FF) of 86.29%. This performance outperforms the control rigid device, which exhibit a PCE of 25.70%, a \u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e of 26.03 mA cm\u003csup\u003e-\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e\u003c/sup\u003e, a \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e of 1.188 V, and an FF of 83.10%. More importantly, the optimal copolymer dosage (Fig. S20 and Table S5) enables the F-PSCs with the structure of PEN/ITO/4PABcz/Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e/perovskite/PiPl/PC\u003csub\u003e61\u003c/sub\u003eBM/BCP/Ag to deliver a champion PCE of 25.54% (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ee, g), which is substantially higher than of the control flexible device (24.72%). The increase of the average PCE from 25.38% to 26.45% for rigid devices (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb) and from 24.47% to 25.22% for flexible scenario (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ef) demonstrates the exellent reproducibility of the fluoropolymer copolymer modification strategy.31\u0026ndash;33 Furthermore, the integrated \u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e derived from external quantum efficiency (EQE) measurements (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ed, h) and the steady PCEs (Fig. S21) are in a good agreement with the \u003cem\u003eJ\u003c/em\u003e-\u003cem\u003eV\u003c/em\u003e results, confirming the positive regulatory effect of the elastomeric fluoropolymer copolymer on device photovoltaic performances.\u003c/p\u003e \u003cp\u003eWe ascribe the enhanced photovoltaic performance to the improved perovskite film quality.34 A suite of electrochemical characterizations were conducted to reveal the charge transfer, extraction, and recombination dynamics in the PSCs. Benefiting from the reduced defect density, the solar cells incorporating TFEMA-co-HFBA copolymer exhibit a stronger built-in electric field (the intrinsic driving force for charge separation) and a larger recombination resistance (Fig. S22), indicative of suppressed charge carrier loss for maximizing power output. This conclusion is further corroborated by a reduced ideality factor closer to unity derived from illumination intensity dependent \u003cem\u003eJ\u003c/em\u003e-\u003cem\u003eV\u003c/em\u003e measurements, and a marked decrease in leakage current (Fig. S23).\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eTo assess the practical viability of the devices, we monitored the performance evolution under steady-state external conditions, including persistent storage and operational stress. Under ambient air conditions (25\u0026deg;C, 20\u0026thinsp;~\u0026thinsp;40% relative humidity), the unencapsulated device modified with the moisture-resistant fluoropolymer maintains 90% of its initial efficiency after 2560 h (107 days), significantly outperforming the reference device with only 70% PCE retention after a mere 1200 h (50 days) (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ei and Fig. S24). Under MPP tracking conditions with equivalent 1-sun illumination, the target device exhibits negligible PCE degradation after 1100 h of continuous operation. In stark contrast, the control device undergoes rapid performance decay, retaining only\u0026thinsp;~\u0026thinsp;70% of its initial PCE after 850 h (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ej). The enhanced device stability is further corroborated by the changes of PL properties of perovskite films, which is highly correlated with defect evolution and halogen segregation behaviour.35 Comparative analysis of PL data for the control and target films prior to aging demonstrates a blue-shifted PL emission wavelength (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ej, k, m, n) and increased PL intensity (Fig. S25) in the copolymer modified films, as evidenced by PL mapping images and corresponding statistical distributions. These observations confirm effective defect passivation, consistent with the previously documented improvements in perovskite film crystallinity.36-38After accelerated aging tests under persistent light irradiation and high-humidity ambient exposure for over 90 days, the control sample exhibits the formation of heterogeneous impurities (e.g. PbI\u003csub\u003e2\u003c/sub\u003e and yellow perovskite phase), which are absent in the target film (Fig. S26). The emergence of these impurities in the control device induces PL intensity quenching and a red-shift PL emission peak. In contrast, the incorporation of TFEMA-co-HFBA copolymer mitigates the magnitude of PL peak shift, indicative of a more robust perovskite lattice that resists harsh environmental stresses and suppresses detrimental phase separation.39\u003csup\u003e,\u003c/sup\u003e40\u003c/p\u003e \u003cp\u003eThermal-mechanical cycling stability of F-PSCs\u003c/p\u003e \u003cp\u003eFor F-PSCs, the tolerance to dynamic thermalmechanical stimuli is far more critical than that for their rigid counterparts.41\u003csup\u003e,\u003c/sup\u003e42 We conducted temperature-dependent XRD characterizations to understand the structural evolution of perovskite films and identify the impacts of fluoropolymer copolymer. As the temperature increases from \u0026minus;\u0026thinsp;20\u0026deg;C to 80\u0026deg;C, the control film exhibits a progressive shift of the characteristic (001) diffraction peak toward smaller angles (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea and Fig. S27), corresponding to lattice expansion of the (001) plane from 6.125 \u0026Aring; to 6.139 \u0026Aring; and a resultant perovskite lattice deformation ratio of 0.23%. In a PSC, such temperature-induced lattice expansion inevitably accelerates ion migration and promotes defect formation.43 In contrary, the XRD profiles for the fluoropolymer copolymer modified perovskite film are insensitive to temperature, with the (001) lattice deformation ratio of mere 0.01% (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb and Fig. S28). This exceptional thermal stability stems primarily from the preferential segregation of the copolymer at perovskite grain boundaries, which serves as an elastic buffer to suppress thermomechanical stress and thereby inhibits lattice expansion under elevated temperatures. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec, after 10 thermal cycles, the characteristic XRD peaks of the control perovskite film shift toward higher angles, indicating reduced interplanar spacing caused by out-of-plane compressive stress and in-plane tensile strain. In contrast, the target film exhibits negligible peak displacement, reflecting suppressed lattice strain evolution. This demonstrates that the carefully designed fluoropolymer copolymer effectively mitigates thermally induced structural deformation, thereby enhancing the mechanical resilience of the perovskite lattice under repeated thermal stress.44 SEM images indicate that an increasing number of bending cycles induces clearer mechanical cracks across the perovskite grains in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ed. Further cycling causes these cracks to propagate and even trigger crystal exfoliation. In contrast to the control sample, the copolymer modified perovskite film features narrower and shallower cracks. Furthermore, the thermal treatment enables the target perovskite film to self-heal, which further enhances its mechanical stability.\u003c/p\u003e \u003cp\u003eWe further evaluated the device stability against mechanically and thermally induced stresses. Cyclic bending tests performed at a fixed radius of 4 mm (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ee, Fig. S29) reveal that the F-PSC with copolymer modified perovskite has a higher PCE retention than that of control device after 3000 cycles, which is attributed to the superior polymer-like mechanical feature of the fluoropolymer modified perovskite film. Taking advantage of the self-healing ability of the modified perovskite film, we applied intermittent thermal treatment at 60\u0026deg;C for 3 min for effectively recovering the photovoltaic performances of F-PSCs. The control solar cell shows no obvious performance recovery under the same conditions. Notably, the target device retains over 90% of its initial PCE after 11,000 bending cycles, standing in sharp contrast to the control device that has mere 80.2% PCE retention after 5,000 cycles. This significant improvement in mechanical stability is attributed to the incorporation of elastic fluoropolymer copolymer, which not only reduces the Young\u0026rsquo;s modulus of the perovskite film but also enhances the interfacial interaction between the perovskite layer and adjacent functional layers (e.g. 4PABcz and PC\u003csub\u003e61\u003c/sub\u003eBM) (Fig. S30), thus suppressing crack formation under external stress. A comparison of the bending-induced efficiency loss of our devices with the data reported in previous studies demonstrates that our F-PSC exhibits superior mechanical stability (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ef, Table S6). Thermal cycling stability was assessed by subjecting PSCs to cyclic temperature changes between 0\u0026deg;C (30 min) and 80\u0026deg;C (30 min). As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eg and Fig. S31, the control device retains only 85% of its initial PCE after 200 thermal cycles, whereas the TFEMA-co-HFBA copolymer modified PSC maintains over 90% PCE after 500 cycles. By comparing the thermal cycling-induced efficiency degradation of our devices with that documented in prior literatures (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eh and Table S7), we find that the well-established F-PSC achieves the highest thermal cycling stability among similar devices reported to date. Collectively, these results confirm the exceptional operational stability of TFEMA-HFBA copolymer tailored F-PSCs with self-healing capacity under extreme temperature fluctuations, which lays a solid foundation for their practical application on complex curved surfaces.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e"},{"header":"Discussion","content":"\u003cp\u003eIn summary, we synthesized a misplaced-dipole engineered fluoropolymer elastomer that enables the formation of high-quality perovskite films with polymer-like block characteristics, including a reduced Young\u0026rsquo;s modulus, elastic grain boundaries, and self-healing capability. These features significantly enhance the tolerance of perovskite lattice to mechanical deformation and temperature fluctuations. By leveraging this elastomer modification strategy, we achieve F-PSCs with a champion PCE of 25.54%, alongside 26.83% in rigid counterparts. Beyond improvement in static photostability and storage stability, the stress-buffering effect imparted by the elastomer at perovskite grain boundaries enables the F-PSC to retain 90% of its initial PCE after 11,000 bending cycles and 500 thermal cycles between 0\u0026deg;C and 80\u0026deg;C.\u003c/p\u003e \u003cp\u003eCollectively, this work establishes a versatile and scalable pathway for the development of high-efficiency and ultra-stable F-PSCs, paving the way for their commercialization.\u003c/p\u003e"},{"header":"Methods","content":"\u003cp\u003eMaterials and reagents\u003c/p\u003e \u003cp\u003eAll the commercial materials were used as received without any future purification, including lead iodide (PbI\u003csub\u003e2\u003c/sub\u003e, 99.99%, Xi\u0026rsquo;an Polymer Light Technology Corp.), lead bromide (PbBr\u003csub\u003e2\u003c/sub\u003e, 99.99%, Xi\u0026rsquo;an Polymer Light Technology Corp.), cesium Iodide (CsI, 99.9%, Sigma-Aldrich), formamidinium iodide (FAI, greatcell solar), methylammonium bromide (MABr, greatcell solar), methylamine hydrochloride (MACl, greatcell solar), (4-(9\u0026prime;-Phenyl-9H,9\u0026rsquo;H-[3,3\u0026prime;-bicarbazol]-9-yl)butyl)phosphonic Acid (4PABcz, Suzhou Liwei New Materials Technology Co., Ltd), [6,6]-phenyl C\u003csub\u003e61\u003c/sub\u003e butyric acid methyl ester (PC\u003csub\u003e61\u003c/sub\u003eBM, 99.5%, Xi\u0026rsquo;an Polymer Light Technology Corp.), 2,9-dimethyl-4,7-diphenyl-1,10-Phenanthroline (BCP, 99% Xi\u0026rsquo;an Polymer Light Technology Corp.), nickel oxide (NiO\u003csub\u003ex\u003c/sub\u003e, 99.99%, Aladdin), 2,2,2-Trifluoroethylmethacrylate (TFEMA, Aladdin, 98%), 2,2,3,4,4,4-Hexafluorobutyl acrylate (HFBA, Aladdin, \u0026gt;\u0026thinsp;95%), 2,2'-Azobis(2-methylpropionitrile) (AIBN, Aladdin, 99%), anhydrous N,N-dimethylformamide (DMF, 99.99%, Sigma-Aldrich), dimethylsulfoxide (DMSO, 99.99%, Sigma-Aldrich), methanol (MS, Aladdin) isopropyl alcohol (IPA 99.9%, Aladdin), chlorobenzene (CB 99.99%, Sigma-Aldrich), Flexible conductive substrate PEN/ITO (Liaoning Optimal New Energy Technology Co., Ltd).\u003c/p\u003e \u003cp\u003ePreparation of perovskite precursors and perovskite films\u003c/p\u003e\n\u003ch3\u003ePreparation of rigid PSCs\u003c/h3\u003e\n\u003cp\u003e \u003cstrong\u003ePreparation of hole/electron transport solutions\u003c/strong\u003e \u003cp\u003eNiOₓ dispersion was prepared by dissolving 7.5 mg of NiOₓ powder in 1 mL of deionized water. A 4PABcz solution was formulated by dissolving 0.54 mg of 4PABcz in 1 mL of anhydrous MS. For the Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e dispersion, 10 \u0026micro;L of nano-Al₂O₃ dispersion liquid was diluted with 990 \u0026micro;L of IPA to a total volume of 1 mL. The PiPl solution was prepared by dissolving 0.4 mg of PiPl in 1 mL of IPA. PC₆₁BM solution was obtained by dissolving 20 mg of PC₆₁BM in 1 mL of anhydrous CB. Finally, a BCP solution was prepared by dissolving 0.5 mg of BCP in 1 mL of IPA.\u003c/p\u003e \u003c/p\u003e \u003cp\u003e \u003cb\u003ePreparation of pristine and molecule-doped perovskite precursor solutions 1\u003c/b\u003e: 71 uL PETMP, 85 \u0026micro;L PAE and 2 mg AIBN in 2 mL DMF and DMSO solution (volume ratio 4:1). Consider the above solution as a concentrated solution of 100 mg/mL, and continue to dilute it with a mixed solvent of DMF and DMSO (volume ratio 4:1) to 1 mg/mL for later use.\u003c/p\u003e \u003cp\u003e \u003cb\u003ePreparation of pristine and molecule-doped perovskite precursor solutions 2\u003c/b\u003e: 6.7 mg MABr, 8.8 mg MACl, 18.2 mg CsI, 196.1 mg FAI, 22 mg PbBr\u003csub\u003e2\u003c/sub\u003e and 602 mg PbI\u003csub\u003e2\u003c/sub\u003e (Excess of PbI\u003csub\u003e2\u003c/sub\u003e by 5%) were dissolved in 865 \u0026micro;L DMF and DMSO solution (volume ratio 4:1), stirred overnight, and used to prepare an undoped precursor solution.For the doping solution, dissolve the perovskite precursor powder in 865 \u0026micro;L of 1 mg/mL cross-linked molecular mixture solution.\u003c/p\u003e \u003cp\u003e \u003cstrong\u003ePreparation of PSCs\u003c/strong\u003e \u003cp\u003eITO substrates were sequentially cleaned with detergent, deionized water, and ethyl alcohol. The cleaned ITO substrates were dried with high-pressure N\u003csub\u003e2\u003c/sub\u003e gas flow, and then plasma modified for 1 min. The NiO\u003csub\u003ex\u003c/sub\u003e layer was formed by spin-coating the precursor solution at 3000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 150 ℃ for 20 min in the air. Then transfer it to the N\u003csub\u003e2\u003c/sub\u003e glove box for further preparation. The 4PABcz was formed by spin-coating the precursor solution at 3000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 100 ℃ for 10 min. The Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e was formed by spin-coating the precursor solution at 5000 rpm (acceleration 4000 rpm) for 30 s and then short-baking at 100 ℃ for 1 min. The as-prepared perovskite precursor solution was spin-coated onto the hole transport substrate with speed of 1000 rpm (acceleration 1000 rpm) for 5 s and 5000 rpm (acceleration 3000 rpm) for 30 s. During the last 15 s of the spinning process, the liquid film was treated by drop-casting chlorobenzene solvent (150 ul). The substrates were annealed on a hot plate at 100℃ for 30 min. After cooling the perovskite film, a precursor solution was spin coated at 5000rpm (acceleration 3000 rpm) for 30 seconds, and then baked briefly at 100 ℃ for 5 min to form a PiPl layer. The PC\u003csub\u003e61\u003c/sub\u003eBM layer was formed by spin-coating the precursor solution at 1500 rpm (acceleration 3000 rpm) for 35 s and then short-baking at 100 ℃ for 10 min. The BCP layer was formed by spin-coating the precursor solution at 4000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 100 ℃ for 3 min. Finally, 100 nm thick of Ag was deposited on top by thermal evaporation. The active area of PSCs is 0.03245 cm\u003csup\u003e2\u003c/sup\u003e as determined by the mask.\u003c/p\u003e \u003c/p\u003e\n\u003ch3\u003ePreparation of F-PSCs\u003c/h3\u003e\n\u003cp\u003e \u003cstrong\u003ePreparation of hole/electron transport solutions\u003c/strong\u003e \u003cp\u003ethe NiO\u003csub\u003ex\u003c/sub\u003e solution was prepared by dissolving 15 mg NiO\u003csub\u003ex\u003c/sub\u003e in 1 mL of deionized water. The solution was prepared by dissolving 1.08 mg 4PABcz in 1 mL of anhydrous MS (Add 10 \u0026micro;L of DMF). The Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e solution was prepared by dissolving 10 \u0026micro;L Nano Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e dispersion liquid in 990 \u0026micro;L of DMF. The remaining steps are consistent with the preparation steps of rigid perovskite devices.\u003c/p\u003e \u003c/p\u003e \u003cp\u003e \u003cstrong\u003ePreparation of flexible perovskite devices\u003c/strong\u003e \u003cp\u003euse thermosetting tape to bond a blank glass substrate onto a PEN/ITO substrate. Then wash with soapy water, deionized water, and ethanol, and sonicate in each solution for 15 minutes. Subsequently, before use, clean the substrate by blowing it dry with nitrogen and then treating it with ultraviolet ozone for 10 min. The NiOx layer was formed by spin-coating the precursor solution at 2000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 150 ℃ for 15 min in the air. Then transfer it to the N\u003csub\u003e2\u003c/sub\u003e glove box for further preparation. The 4PABcz was formed by spin-coating the precursor solution at 4000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 100 ℃ for 10 min. The Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e was formed by spin-coating the precursor solution at 5000 rpm (acceleration 4000 rpm) for 30 s and then short-baking at 100 ℃ for 1 min. The as-prepared perovskite precursor solution was spin-coated onto the hole transport substrate with speed of 1000 rpm (acceleration 1000 rpm) for 5 s and 5000 rpm (acceleration 3000 rpm) for 30 s. During the last 15 s of the spinning process, the liquid film was treated by drop-casting chlorobenzene solvent (150 ul). The substrates were annealed on a hot plate at 100℃ for 30 min. After cooling the perovskite film, a precursor solution was spin coated at 5000rpm (acceleration 3000 rpm) for 30 seconds, and then baked briefly at 100 ℃ for 5 min to form a PiPl layer. The PC\u003csub\u003e61\u003c/sub\u003eBM layer was formed by spin-coating the precursor solution at 1500 rpm (acceleration 3000 rpm) for 35 s and then short-baking at 100 ℃ for 10 min. The BCP layer was formed by spin-coating the precursor solution at 4000 rpm (acceleration 3000 rpm) for 30 s and then short-baking at 100 ℃ for 3 min. Finally, 120 nm thick of Ag was deposited on top by thermal evaporation. The active area of PSCs is 0.03245 cm\u003csup\u003e2\u003c/sup\u003e as determined by the mask.\u003c/p\u003e \u003c/p\u003e \u003cp\u003eCharacterizations:\u003c/p\u003e \u003cp\u003eFourier transform infrared (FTIR) spectra were performed by the infrared spectrometer (Nicolet iS50 FT-IR). The X-ray photoelectron spectroscopy (XPS) studies were performed using a Thermo-VG Scientific ESCALAB 250 photoelectron spectrometer equipped with a monochromated Al Kα (1486.6 eV) X-ray source. The composition was characterized by X-ray diffraction (Rigaku, with Cu-Kα radiation of λ\u0026thinsp;=\u0026thinsp;0.15418 nm). Grazing incidence X-ray diffraction (GIXRD) measurements were carried out with a Bruker D8 grazing incidence X-ray diffractometer. Photoluminescence (PL) and time-resolved PL (TRPL) were measured by FLS980 (Edinburgh Instruments Ltd) with an excitation at 470 nm. The PL mapping was executed by a Laser Scanning Confocal Microscope (Enlitech, SPCM-1000) equipped with a 470 nm pulse laser and a galvo-based scanner. High resolution transmission electron microscopy (HRTEM) morphology was performed using a FEI Tecnai G2 F30 instrument manufactured by FEI Company, USA. Scanning electron microscopy was measured via a scanning electron microscope (Hitachi S8230). Atomic Force Microscopy (AFM) topography characterization was performed using a Bruker Dimension ICON Microscope. In-situ variable-temperature XRD measurements were performed using a SmartLab-9kW diffractometer. The instrument model for Nano IR testing is nano IR2-FS (Anasys Instruments), and this multifunctional nano-infrared spectrometer includes an atomic force microscope for detecting morphology. Optical microscope images were tested using a Leica polarizing microscope (DM4 P). The instrument model used for liquid nuclear magnetic resonance hydrogen spectroscopy is AVANCE NEO 400MHZ. Water contact angle measurements were performed using a Dataphysics OCA20 contact angle meter. The lead content in water for the lead leaching test was determined using an iCE\u0026trade; 3500 atomic absorption spectrometer (AAS). The photovoltaic performance of the PSCs was collected with a source meter (Keithley 2400), with the light source (SS-F5-3A, Enlitech) via reverse scanning from 1.22 to -0.1 V or forward scanning from \u0026minus;\u0026thinsp;0.1 to 1.22 V at a scanning speed of 50 mV/s, and the quantum efficiency was tested using a quantum efficiency measurement system (QE-R, Enlitech). The illumination stability of devices was executed under one sun-equivalent white light. All devices were taken out from the chamber and tested at different time intervals under a separate solar simulator (AM 1.5G, 100 mW/cm\u0026sup2;) for J-V characterization. The device MPP tracking was measured with a CHI 1000C potentiostat under a white LED lamp (100 mW/cm\u0026sup2;) in an N₂ glove box. Then, the devices were measured with a maximum power point (MPP) tracking routine under continuous 1 sun illumination. The MPP was updated every 300 s by a standard perturb and observation method, and the average temperature was kept at ~\u0026thinsp;50 ℃.\u003c/p\u003e \u003cp\u003eComputational Methods:\u003c/p\u003e \u003cp\u003e \u003cstrong\u003eMolecular Dynamics (MD) Simulation\u003c/strong\u003e \u003cp\u003eMolecular dynamics (MD) simulations were performed using GROMACS 2021 software with the AMBER99SB force field. The simulation box was constructed via GROMACS tools. Prior to production simulations, energy minimization and a 100 ps pre-equilibration simulation were sequentially carried out to allow system relaxation. For production simulations, Newton's equations of motion were integrated using the leapfrog algorithm, and the V-rescale temperature coupling scheme was adopted with a simulation temperature of 298.15 K.\u003c/p\u003e \u003c/p\u003e \u003cp\u003e \u003cstrong\u003eFinite Element (FE) Simulation\u003c/strong\u003e \u003cp\u003eFinite element (FE) simulations were performed using ABAQUS software. The model was simulated based on Hooke's Law, and a rotation about the center was applied to both ends of the model, with each end rotated by approximately 40\u0026deg;.\u003c/p\u003e \u003c/p\u003e\n\u003ch3\u003eData Availability\u003c/h3\u003e\n\u003cp\u003eThe data that support the findings of this study are available within the Article and its Supplementary Information. Source data are provided with this paper.\u003c/p\u003e"},{"header":"Declarations","content":" \u003cp\u003e \u003cstrong\u003eCompeting interests\u003c/strong\u003e \u003cp\u003eThe authors declare no competing interests.\u003c/p\u003e \u003c/p\u003e\u003ch2\u003eAuthor contributions\u003c/h2\u003e \u003cp\u003eL. and J. D. conceived the idea. F. L., Y. L., W. Z., Y. L. and J. L. performed the solar cell fabrication, characterization, and optimization. B. H. contributed to the PL measurement. Q. G. and Y. Z. measured the EQE. J. D. involved in the data analysis and discussion. Q. C. and Q. T.directed this work. J. D. and J. D. wrote the first draft of the manuscript. All the authors revised and approved the manuscript.\u003c/p\u003e\u003ch2\u003eAcknowledgements\u003c/h2\u003e \u003cp\u003eThe authors gratefully acknowledged financial support provided by the National Natural Science Foundation of China (62304124, 62374105, 62204098, 22179051, 22309107, 52472259), Natural Science Foundation of Shandong Province, China (ZR2025MS892, ZR2024QE036, ZR2023QB281, ZR2024QB021, ZR2025MS1014), National Key Research and Development Program of China (2022YFE4200500), Qingdao Natural Science Foundation (25-3-1-13-zyyd-jch), Special Fund of Taishan Scholar Program of Shandong Province (tstp20250728, tsqnz20221141).\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003eKojima A, Teshima K, Shirai Y, Miyasaka T (2009) Organometal Halide Perovskites as Visible-Light Sensitizers for Photovoltaic Cells. 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Adv Mater 35:e2211257\u003c/span\u003e\u003c/li\u003e\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"nature-portfolio","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"","title":"Nature Portfolio","twitterHandle":"","acdcEnabled":false,"dfaEnabled":false,"editorialSystem":"ejp","reportingPortfolio":"","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"","lastPublishedDoi":"10.21203/rs.3.rs-8773234/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-8773234/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eFlexible perovskite solar cells (F-PSCs) offer a compelling solution to the intrinic rigidity of silicon-based photovoltaics, enabling power generation on irregular surfaces. However, their practical application hinges on the perovskite layer\u0026rsquo;s ability to concurrent thermal cycling, so as to match the deformability of polymer substrates. Different from the conventional lattice solidification strategy for rigid devices, we incorporate a fluorinated misplaced-dipole engineered repairable elastomer into the perovskite film, which enhances perovskite intergranular toughness and mitigates thermal stress fatigue cracks. The resultant perovskite film exhibits suppressed lattice thermal fluctuation, thereby boosting enhanced environmental resilience. Consequently, the optimized F-PSCs deliver a champion efficiency of 25.54%, versus 26.83% for their rigid counterparts. More importantly, the F-PSC demonstrates exceptional durability under harsh operational stresses, retaining over 90% of the initial PCE after 11,000 bending cycles and maintaining a comparable retention rate following 500 thermal cycles, paving the way towards for the long-lasting flexible photovoltaic devices.\u003c/p\u003e","manuscriptTitle":"Misplaced-dipole engineered repairable fluoropolymer elastomer for flexible perovskite solar cell with excellent thermal-mechanical cycling resistance","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2026-02-05 08:51:20","doi":"10.21203/rs.3.rs-8773234/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"nature-communications","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"NCOMMS","sideBox":"Learn more about [Nature Communications](http://www.nature.com/ncomms/)","snPcode":"","submissionUrl":"https://mts-ncomms.nature.com/","title":"Nature Communications","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Nature Communications","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"4ab55276-ed54-4f55-b302-e3b9d14c8bf0","owner":[],"postedDate":"February 5th, 2026","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"under-review","subjectAreas":[{"id":62332933,"name":"Physical sciences/Energy science and technology/Renewable energy/Solar energy/Photovoltaics/Solar cells"},{"id":62332934,"name":"Physical sciences/Materials science/Materials for energy and catalysis/Solar cells"}],"tags":[],"updatedAt":"2026-05-05T13:36:03+00:00","versionOfRecord":[],"versionCreatedAt":"2026-02-05 08:51:20","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-8773234","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-8773234","identity":"rs-8773234","version":["v1"]},"buildId":"XKTyCvWXoU3ODBz1xrDgd","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}