A synergistic strategy of crosslinking and filler toughening enabling stretchable organic photovoltaics for wearable applications

Domain: neuroscience
Published: 2026-01-29
Editorial outcome: major_revision

Abstract
Improving the elongation of intrinsically stretchable organic electronics typically prioritizes flexibility, which may increase the crack-onset strain at the expense of ductility. Here, we present a synergistic design that combines covalent crosslinking and silica filler reinforcement to construct a photoactive layer of organic photovoltaics (OPVs) with both elevated fracture strain and modulus. This interpenetrating network boosts the crack-onset strain to over 40% and raises the modulus by 5-fold to 1090 MPa. The silica filler promotes enhanced aggregation and molecular ordering in both donor and acceptor materials, enabling a power conversion efficiency exceeding 16% for intrinsically stretchable devices, with 80% of the initial efficiency retained under nearly 40% strain, which is one of the highest values reported to date for stretchable OPVs. These findings provide insights for developing stretchable and mechanically robust OPVs toward practical wearable applications.

Introduction
Intrinsically stretchable organic electronics hold great promise for applications in wearable electronics1–5, volumetric displays6–8 and energy power supplies9–14. Yet, achieving high performance in these devices demands a delicate reconciliation of mechanical deformability with optoelectronic efficiency15–17. In the case of intrinsically stretchable organic photovoltaics (IS-OPVs), mechanical strain perturbs the morphology of the donor-acceptor (D/A) active layer at multiple length scales18, including molecular conformation19, crystalline packing20, orientation21, phase separation22,23, and interfacial alignment24. These structural changes have a strong influence on exciton diffusion, exciton splitting, and charge carrier transport, thereby limiting the power conversion efficiency (PCE) of OPV devices under strain25.

To address this challenge, recent efforts have focused on the flexibility of D/A blend films through molecular design, such as introducing flexible segments26–28, or blending with elastomers and plasticizers29,30. These approaches enhance the stretchability by increasing intra-chain disorder31 and inter-chain interactions32 via non-covalent or elastomeric networks33. While these strategies have led to crack-onset strains (COS) exceeding 30%, they were often obtained at the cost of reduced electronic performance. The initial PCEs typically remained below 15%, and the PCE under 30% strain tended to drop below 80% of its original value34–36. Parallel advances in deformable interlayers, such as stretchable electrodes and charge transport layers, have further aided mechanical resilience by redistributing applied stress37,38. Despite these advances, enhancing the intrinsic toughness of the active layer itself, which governs crack resistance and structural integrity under strain, remains underexplored. Improving this parameter could offer a more direct route to sustaining high efficiency and practical applications for stretchable devices.

Here, we report a molecular crosslinking and filler toughening (MCFT) strategy that synergistically enhances both the stretchability and toughness of OPV blend films. A covalent network formed between crosslinkable polymer donors and non-fullerene acceptors (NFAs) efficiently transmits stress across molecular domains, while silica fillers serve as crack stoppers and promote ordered aggregation. Together, these dual reinforcements yield IS-OPVs with improved mechanical robustness and high performance retention under strain, setting a benchmark for intrinsically stretchable solar cells.

Results
MCFT Structure design
To construct a robust crosslinked network, we first engineered both the donor and acceptor materials to be chemically crosslinkable with vinyl groups (Fig. 1a, Supplementary Fig. 1). The donor polymer, PzA5, was synthesized by incorporating 1,2-di(thiophen-2-yl)ethene (TVT) and the flexible 1,8-di(thiophen-2-yl)octane (TOT) segments31 into the conjugated backbone of the imide-functionalized benzotriazole (TzBI) polymer PTzBI-dF polymer39. This backbone engineering enables elastic chain dynamics and facilitates crosslink formation. Systematic variation of TVT and TOT ratios identified 2% TVT:3% TOT as the optimal composition, based on second-derivative analysis of performance trends (Supplementary Fig. 3 a–b).

To complement the crosslinkable donor, we developed a small-molecule acceptor, Y5TV, by functionalizing the widely studied L8BO series acceptors with 2-(pent-4-en-1-yl)thiophene units at both indene carbonitrile termini. These alkene moieties enable efficient thiol-ene “click” crosslinking. We employed pentaerythritol tetra(3-mercaptopropionate) (PETMP) as the tetra-thiol crosslinker (Supplementary Fig. 2). Upon mild thermal (100 °C) and UV (365 nm) activation for 5 min, a stable and covalently bonded donor-acceptor network of PzA5:L8BO:Y5TV, termed UVC, was formed, as verified by the disappearance of S–H vibrational bands in Fourier transform infrared (FTIR) spectra (Fig. 1b).

To reinforce the mechanical integrity of the UVC network, we introduced an aerogel silica nanoparticle filler with an average size of ~100 nm, modified with 1,1,1,3,3,3-hexamethyldisilazane (ASH) (Fig. 1c–e). The addition of ASH yielded a sharp trade-off in performance with a sudden increase in modulus and a corresponding drop in PCE at filler loadings above 3 wt%, as revealed by second-derivative analysis (Supplementary Fig. 3 c–d). This critical loading, 3 wt%, was selected as the optimal concentration for mechanical enhancement without compromising device performance. Surface energy analysis revealed that ASH incorporation markedly increased film hydrophobicity, as evidenced by contact angle measurements (Supplementary Fig. 4). This surface modification enhances compatibility with the organic semiconductor matrix, facilitating homogeneous dispersion in chloroform without sedimentation.

To probe intermolecular interactions, density functional theory (DFT) calculations (B3LYP-D3/6-31+G**) were performed. The ASH filler exhibited strong interaction region indicator (IRI) scattering plots and isosurfaces binding with both donor (−10.1 kcal mol⁻1 with PzA5) and acceptor (−28.0 kcal mol⁻1 with Y5TV) molecules (Fig. 1f and Supplementary Fig. 5). Consistently, X-ray photoelectron spectroscopy (XPS) results revealed downshifts in the Si 2p peak of ASH upon mixing with PzA5 (100.5 eV) and Y5TV (100.3 eV), relative to pristine ASH (100.8 eV) (Supplementary Fig. 6), confirming strong electrostatic interactions at the organic–inorganic interface. The Flory–Huggins interaction parameters (χ) derived from contact angle data were below 0.04 for both donor and acceptor with ASH, further supporting their favourable miscibility and stable dispersion (Supplementary Fig. 6). These findings underscore the feasibility of constructing a mechanically robust and molecularly cohesive network via the MCFT approach.

To elucidate the structural impact of the crosslinked D/A network and ASH filler, we conducted molecular dynamics (MD) simulations (Supplementary Fig. 7–8). The radius of gyration distributions revealed that replacing L8BO with Y5TV increased structural disorder (higher standard deviation, SD), which was further amplified by crosslinking and ASH addition. In blend systems, both donor and acceptor chains showed elevated SD values upon MCFT treatment, suggesting increased configurational heterogeneity and domain dispersion across length scales. This structural disorder likely contributes to enhanced entropic elasticity, which is critical for accommodating mechanical strain without catastrophic failure.

Enhanced stretchability while toughening
The mechanical enhancement conferred by the MCFT strategy was first assessed through uniaxial tensile testing of free-standing thin films (Fig. 2a), showing consistent results across three independent specimens for each sample, with the corresponding tensile parameters and standard deviations summarized in Supplementary Table 1. The pristine PzA5:L8BO blend exhibited brittle failure at low strain below 5%, while the UVC network significantly improved ductility, yielding a modulus of approaching 505 MPa and a crack-onset-strain (COS) of 44%. The further introduction of ASH resulted in a slightly reduced COS of 43%, likely due to local stiffness contrast from the rigid filler that promotes early strain localization. Nevertheless, interfacial debonding and viscoelastic dissipation around the filler enable efficient energy absorption, resulting in markedly enhanced toughness and modulus (1090 MPa for UVC+ASH). This combination of nearly unchanged COS with markedly reinforced modulus highlights the role of ASH in providing a toughening reinforcement effect.

To probe the filler-induced network formation, dynamic mechanical analysis was conducted (Fig. 2b). The change in dynamic modulus between 0.2% and 2% strain (ΔE′0.2-2%), an indicator of the Payne effect40, was 9.5 kPa for the UVC+ASH film, significantly higher than that of the control samples (Supplementary Fig. 9a). Temperature-sweep DMA further revealed a pronounced rise in both storage and loss modulus near 100°C, accompanied by a modest decrease in Tg (Supplementary Fig. 9b). The increase of E′ and E″ at elevated temperatures indicates the formation of a filler-induced reinforcing network that becomes active under strain, while the shift in Tg suggests altered chain dynamics associated with stronger filler–matrix interactions. These features consistently support the presence of a dynamic, strain-sensitive filler network that contributes to energy dissipation and mechanical reinforcement.

Finite element analysis (FEA) further clarified the role of ASH in mechanical stress distribution (Fig. 2c, Supplementary Fig. 10 a–c). Under tensile load, stress is concentrated at the filler–matrix interfaces, with ASH particles absorbing and redistributing stress while undergoing minimal deformation. This leads to a more uniform stress profile and the establishment of a rigid yet deformable mechanical skeleton within the composite. Additional FEA using a bilayer model revealed that the UVC+ASH blend also exhibits improved out-of-plane (OOP) toughness, with reduced bending-induced displacement compared to control films (Supplementary Fig. 10 d–f).

To investigate fracture dynamics at the molecular level, we performed molecular dynamics (MD) simulations under periodic boundary conditions, which suppressed edge effects and emphasize intrinsic tensile responses, with snapshots taken near the failure point (Fig. 2d, Supplementary Figs. 11 and 12). The PzA5:L8BO film showed abrupt displacement within L8BO-rich domains, indicative of a brittle fracture. By contrast, the UVC network delayed crack formation by enhancing covalent bonding and chain entanglement. When ASH was incorporated, cracks were arrested at filler interfaces while the surrounding network remained structurally intact, revealing a synergistic toughening mechanism where ASH particles act as crack stoppers and stress delocalizers. This is consistent with the optical imaging under strain (Supplementary Fig. 13), where ASH domains visibly interrupt fracture propagation.

Cyclic tensile testing further confirmed the improved mechanical performance of UVC+ASH films (Fig. 2e), which dissipated energy up to 144 MJ m⁻2, far exceeding that of PzA5:L8BO and UVC blends. A summary of the key mechanical parameters underscores the effectiveness of the MCFT strategy in simultaneously enhancing ductility, stiffness, and fracture toughness (Fig. 2f).

MCFT-based efficient IS-OPVs
The photovoltaic performance of the MCFT-based blends was first evaluated using a conventional device architecture: ITO/PEDOT:PSS/active layer/PNDIT-F3N/Ag. As shown in Fig. 3a, the reference PzA5:L8BO device achieved a moderate performance, with an open-circuit voltage (VOC) of 0.884 V, a short-circuit current density (JSC) of 24.4 mA cm⁻2, and a fill factor (FF) of 76.3%, yielding a power conversion efficiency (PCE) of 16.4%. The crosslinked UVC device exhibited a similar PCE of 16.7% but with an increased JSC and a decreased VOC. The lower VOC is consistent with the presence of traps introduced by UV crosslinking, resulting in higher non-radiative recombination and lower SCLC mobilities (Supplementary Fig. 14 a–b and Supplementary Table 2). This is reflected by the light-intensity dependences of VOC and JSC (Supplementary Fig. 14 c-d): n increases from 1.33 to 1.36, indicating higher trap-assisted recombination in UVC, whereas α increases from 0.886 to 0.956, indicating improved charge collection with reduced bimolecular recombination.

In contrast, the UVC+ASH devices delivered a higher PCE of 18.7% (VOC = 0.888 V, JSC = 27.2 mA cm⁻2, and FF = 77.5%). Upon adding ASH, the transport is restored to over 4.5×10−4 cm2 V-1 s-1 with near-unity hole/electron mobility balance (μh/μe = 0.93). Moreover, α further approaches unity (0.989) and n decreases to 1.17, indicating more efficient carrier collection and reduced trap-assisted non-radiative recombination.

To assess the effect of ASH on interfacial charge transfer, we analyzed the donor ground-state bleach (GSB) growth using a biexponential model, referenced at 0.2 ps. The extracted rise lifetimes (τ1, τ2) are 1.18 ps, 8.88 ps (PzA5:L8BO), 2.82 ps, 27.77 ps (UVC), and 0.63 ps, 12.89 ps (UVC+ASH). Crosslinking slows both components, consistent with UV-induced traps, whereas adding ASH mitigates this degradation, restoring and accelerating the dynamics: τ1 shortens to 0.63 ps (faster hole-transfer build-up) and τ2 decreases from 27.77 ps (UVC) to 12.89 ps, indicating a quicker establishment of the charge-separated population. These kinetic improvements align with the enhanced carrier collection and device metrics of the UVC+ASH films (Supplementary Figs. 15–17). Together with the red-shifted, intensified EQE (Fig. 3b and Supplementary Fig. 18), these kinetic and transport improvements account for the enhanced charge generation and extraction in UVC+ASH.

The balance between mechanical properties and photovoltaic performance is translated well into intrinsically stretchable OPVs (IS-OPV), which were fabricated using the architecture of TPU/PH1000/PEDOT:PSS (4083)/active layer/EGaIn (Fig. 3d and Supplementary Fig. 19). The UVC+ASH-based stretchable device delivered a PCE of 16.5%, with a VOC of 0.88 V, a JSC of 26.2 mA cm⁻2, and an FF of 71.53%, outperforming both PzA5:L8BO and UVC analogues (Supplementary Table 3). Crucially, these devices maintained 80% of their initial efficiency under a stretching loading of 1 kg, resulting in tensile strain of approximately 40% (Fig. 3e, Supplementary Fig. 19–20), representing one of the highest reported levels of strain endurance among IS-OPVs. In addition, cyclic stretching tests at defined strains (5%, 10%, or 20%) over 100–200 cycles further underline the advantage of UVC+ASH, which exhibits significantly higher performance retention than PzA5:L8BO throughout repeated loading–unloading processes. This balance of stretchability, toughness, and performance underscore the potential of the MCFT strategy for enabling deformable solar energy harvesting platforms (Fig. 3f and Supplementary Table 4). Depending on the target scenario, either higher strain tolerance or maximized efficiency at moderate strain may be more relevant, suggesting the need for complementary design strategies across different strain regimes.

To understand the underlying morphological evolution under strain, in-situ absorption spectroscopy was performed. The PzA5:L8BO blend showed rapid spectral disordering beyond 0.5–1% strain, indicating a fragile morphology (Supplementary Fig. 21, Fig. 3g). In contrast, UVC+ASH exhibited a strain-induced alignment before yield, followed by partial structural relaxation and recovery during strain hardening, signatures of reversible network reorganization. These adaptive behaviours were further validated by in situ photo-CELIV measurements. UVC+ASH devices showed minimal mobility decay and even a slight increase in peak current at low strain levels, highlighting the role of MCFT-enhanced morphology in preserving carrier transport under mechanical deformation (Fig. 3 h–i, Supplementary Fig. 22).

MCFT morphology
The morphological characteristics of the blend films were investigated using a combination of nanoscale mapping and scattering techniques.

To probe molecular ordering, grazing incidence wide-angle X-ray scattering (GIWAXS) was performed (Fig. 4a, Supplementary Figs. 26 and 27). The UVC blend, formed via UV-induced crosslinking of PzA5:L8BO:Y5TV, displayed a broadened and diminished (1 1 11) reflection of L8BO in the in-plane (IP) direction, indicative of reduced crystallinity and a shorter coherence length. By contrast, the UVC+ASH blend showed significantly enhanced (1 1 11) and (021) diffraction peaks, signifying improved L8BO crystallization. The strengthened π–π stacking would increase carrier mobility, benefiting both JSC and FF, while the improved crystalline order lowers trap-assisted recombination, providing a modest VOC gain. This enhancement likely arises from the gel-like rheological behaviour of ASH (Supplementary Fig. 28), which modulates phase separation kinetics during solvent evaporation and thus reshapes the thin film morphology.

Atomic force microscopy coupled with infrared spectroscopy (AFM-IR) revealed increased-scale donor-acceptor (D/A) networks in the crosslinked UVC blend (Fig. 4b, Supplementary Fig. 23), consistent with enhanced elastic behaviour. Upon ASH incorporation, the UVC+ASH film exhibited stronger molecular aggregation, finer phase separation, and a denser percolating network. Power spectral density (PSD) analysis of the AFM topographies shows a pronounced peak emerging at k≈0.0025 Å−1 (d=251 nm) upon crosslinking and furtherly shifting to ∼0.0015 Å−1 for UVC+ASH, accompanied by a marked suppression of spectral power at low (<0.001 Å−1) and high (>0.01 Å−1) frequencies (Supplementary Figs. 24 and 25). These features indicate a growth of the characteristic domain size together with the damping of large-scale inhomogeneity and nanoscale roughness, aligning with the formation of purer yet more fibrillar phases. Such pronounced morphological optimization can be ascribed to the strong interfacial interactions between ASH and the donor/acceptor components, which facilitate more favorable packing and phase organization. This fibrillar morphology facilitates continuous charge percolation pathways, while the finer yet purer domains promote efficient exciton diffusion and dissociation, jointly underpinning the observed JSC enhancement in devices.

Nitrogen K-edge resonant soft X-ray scattering (NK-RSoXS) was used to assess phase separation and domain formation in the blends. Energy contrast analysis identified 399.2 eV (N 1s→π*) as optimal for visualizing Y5TV-rich domains (Supplementary Fig. 29 a–b). The scattering profiles revealed a distinct hump at q≈0.0107 Å⁻1 in the UVC blend, corresponding to a characteristic domain size of ~58.5 nm (Fig. 4c, Supplementary Fig. 29c). Upon ASH incorporation, this feature became more intense and shifted to lower q (~0.0095 Å⁻1), reflecting the formation of larger and purer domains.

The combined results of GIWAXS, AFM-IR, and NK-RSoXS reveal that structural evolution under ASH incorporation is in line with the enhanced EQE response, where increased molecular ordering and refined, purer phase-separated domains collectively broaden absorption shoulders, extend the continuity for exciton diffusion and carrier transport across multiple length scales, and reinforce JSC and FF gains, consistent with results of TA spectroscopy and device EQEs. A schematic summary of the morphology evolution and the synergistic roles of crosslinker and filler is illustrated in Fig. 4d, highlighting the improved fibrillar domains and networks that underpin the enhanced mechanical and optoelectronic performance of the UVC+ASH system.

Proof-of-concept for applications of IS-OPV in powering
Compared to conventional solar cells on rigid substrates, IS-OPVs offer advantages for powering devices on curved and deformable surfaces. These include applications such as portable energy sources for camping tents, inflatable rescue gear, and wearable electronics for real-time bio-signal or motion tracking in dynamic environments. Such scenarios demand not only high flexibility and stretchability but also robust mechanical resilience to avoid device failure under repetitive or intense mechanical loads. To showcase real-world utility, we integrated the MCFT-based IS-OPV into a wearable, wireless electromyography (EMG) monitoring system (Fig. 5b), where it served as a self-powered energy module. The system successfully recorded muscle electrical signals, distinguishing between contraction and relaxation states (inset, Fig. 5b). Such functionality is directly applicable to biomechanical monitoring, including myoelectric prosthesis control, soft exoskeletons, human-machine interfaces, and fatigue detection via frequency-domain EMG analytics.

In the proof-of-concept demonstration, the IS-OPV wristband integrated three serially connected large area devices with a total effective area of 4.5 cm2 (1.5 cm2 × 3, PCE=14.25%, Supplementary Fig. 30). Under standard solar illumination (1 sun, 100 mW cm⁻²), this corresponds to a theoretical maximum output power of 65 mW (P = Ilight × area × PCE). The standalone wireless EMG monitoring system exhibited an average power load of 360 mW, of which the signal acquisition and processing module accounted for ~100 mW, while the wireless data transmission consumed the majority (~250 mW). With a battery capacity of 0.7 Wh, the system operated for 1.96 hours without external charging. When coupled with the IS-OPV charger, the operational time was extended to 2.30 hours, a 17% increase (Fig. 5c). Extrapolating this concept, enlarging the active area beyond 30 cm² would yield >435 mW under full sunlight (Fig. 5d), a level approaching the continuous power requirement of the device, and thereby potentially enabling near-continuous operation in outdoor conditions. These results underscore the feasibility of MCFT-enhanced IS-OPVs as mechanically robust, energy-autonomous power sources for next-generation wearable and off-grid electronics.

Methods
Materials and synthesis
Monomers and catalysts are purchased from Sigma-Aldrich, Strem, or Volt-Amp Optoelectronic Tech Co. Ltd, Dongguan, China, and used without further purification. All the solvents are purchased from Sigma-Aldrich and used without further purification.

Polymerization of PzA5. TzBI-dF-2Br (396.3 mg, 0.5050 mmol), BDT-Th-Sn (500.0 mg, 0.5316 mmol), TOT-2Br (7.0 mg, 0.0160 mmol), TVT-2Br (3.7 mg, 0.0106 mmol), and 12.3 mg Pd(PPh3)4 were dissolved in 15 mL anhydrous chlorobenzene in a 50 mL pressure tube under N2 protection. The mixture was sealed and stirred at 140 °C for 72 h. The crude product was collected by precipitation in methanol after cooling to room temperature. Then, Soxhlet extraction using a 0.22 μm membrane was carried out with methanol, acetone, n-hexane, dichloromethane, tetrahydrofuran (THF), and chloroform stepwise.

Synthesis of trimethyl(5-(pent-4-en-1-yl)thiophen-2-yl)stannane (TV-Sn). Under −78 °C and nitrogen flow protection, 200.0 mg of 2-(pent-4-en-1-yl)thiophene (1.31 mmol) was dissolved in 20 mL anhydrous tetrahydrofuran in a 3-neck flask. 0.78 mL n-BuLi solution (2.0 M in cyclohexane) was added dropwise and stirred for 1.5 h at 20 °C afterwards. The flask was cooled down to −78 °C again, and the chlorotrimethylstannane solution (1.0 M in THF) was added dropwise. It’s then stirred under 30 °C for another 8 h. The crude product was washed and extracted with saturated potassium fluoride water solution and dichloromethane, respectively. It was then concentrated using vacuum rotary vapour and preliminary purified using recrystallization with DCM:MeOH=1:1. An oily yellow crude product was obtained and used for the next step.

Synthesis of Y5TV. Under the protection of nitrogen, Y5-2Br (200 mg, 0.1067 mmol), TV-Sn (84.0 mg, 0.2668 mmol), and Pd(PPh3)4 (12.3 mg, 0.0107 mmol) are dissolved in 100 mL anhydrous chlorobenzene, heated to 100 °C, and stirred for 10 h and collected by precipitation into methanol. The crude product is purified with a polypropylene gel chromatography column with tetrahydrofuran as the eluent, obtaining a dark green solid with a yield of 32%. 1H NMR (400 MHz, Chloroform-d) δ 9.16 (s, 2H), 8.94 (dd, J=8.4 Hz, 1.5 H), 8.68 (dd, 1.5 H), 8.09 (d, J=5.2 Hz, 1.5 H), 7.94–7.88 (m, 4H), 7.80 (d, J=8.0 Hz, 1.5 H), 7.40 (d, J=10.6 Hz, 3H), 7.12–6.98 (m, 3H), 5.86 (m, 3H), 5.05 (m, 6H), 4.76 (m, 4H), 4.65 (d, 2H), 4.20 (s, 2H), 3.23 (d, 4H), 2.69 (m, 6H), 2.15 (m, 8H), 2.01–1.67 (m, 11H), 1.56 (m, 6H), 1.55–0.69 (m, 88H).

Simulation methods
The DFT calculation is carried out with the Gaussian 16 and Multiwfn software. The structure optimization for the ASH filler cluster and the molecular is conducted based on B3LYP-D3/6-31g**. The dimer conformations are obtained with the forcite molecular dynamics. The exact situation of the dimer interaction is simulated based on B3LYP-D3/6-31+g** and illustrated with the Multiwfn and VMD software, with blue, green, and red colour scaling the electron density corresponding to the bonds, the vdW interaction, and the steric hindrance.

The radius of gyration is obtained from the forcite with the most stable structures going through the following procedures:
For NFA molecular: (1) the crystal structure, including 108 molecules, is built. (2) The cells are subjected to successive dynamics of 100 ps molecular dynamics (NVT, T=300 K), 50 ps molecular dynamics (NPT, T=400 K), and 100 ps molecular dynamics (NVT, T=300 K) to simulate the change in the crystal.
For the blends: (1) An amorphous cell with decamer chains of the polymer donor and NFAs (initial density = 0.5 g cm−3) is built using the structures from the DFT calculation. (2) 5 annealing cycles from 300 to 1000 K (NPT) are conducted to obtain random distributions. (3) The cells are subjected to successive dynamics of 100 ps molecular dynamics (NVT, T=300 K), 200 ps molecular dynamics (NPT, T=400 K), and 100 ps molecular dynamics (NVT, T=300 K) to simulate the change in the cell. The final structures are used for further analysis.

The MD stress–strain simulations were performed using a Perl script. Each system was first equilibrated for 200 ps under zero external stress (300 K, NPT), followed by the application of successive uniaxial stresses of 0.001, 0.05, 0.10, 0.15, 0.25, 0.30, 0.35, and 0.40 GPa. At each stress level, the cell was equilibrated for 100 ps (300 K, NPT) to extract the actual stress and strain values of the cell. Fracture onset was identified from the stress response. In the visualization, PzA5 donor, L8BO acceptor, crosslinked PzA5:Y5TV network, and ASH filler are shown in red, blue, purple, and cyan, respectively. Representative snapshots at selected steps were taken to illustrate the locations where cracks initiate.

Finite element analysis of the mechanical properties
To provide an intuitive visualization of the tensile process, finite element models were constructed in Abaqus. The matrix dimensions were set to 30×40 (units), and spherical fillers with diameters of 8 and 4 (units) were embedded. For computational efficiency, a linear elastic constitutive model was adopted. The Young’s modulus of the matrix was set to 370 MPa, while that of the fillers was either 100 MPa (representing soft elastomeric additives) or 2000 MPa (representing rigid fillers). The unit system was kept consistent across the model.

This simplification does not affect the validity of the observed trends, as the purpose of the simulation was to reproduce the qualitative behavior of uniaxial tensile experiments. A static analysis step was employed: one side of the model was fully fixed, while a constant tensile force of 10 N was applied along the X-axis at the opposite side. The resulting stress distributions are presented in Supplementary Fig. 10.

The comparison between the elastomer and filler blending is using the model in Supplementary Fig. 10, with the balls representing the elastomer or filler particles. A stretching force is applied to the right side with the left side of the model fixed. With moduli much less than or exceeding the matrix, the stress distribution in the matrix differs a lot. The sample using low-modulus blends has a significantly larger displacement. Notably, the stress concentrates at the interface between the low-modulus additive and the matrix, demonstrating heterogeneous intra-stress distribution due to the mismatch of modulus and elongation rate. In the circumstance of high-modulus filler, the filler particles take large shares of the stress due to their high modulus, which could alleviate the stress distribution in the photoactive materials. Besides, the negligible deformation of the filler particle prevents the stress concentration at the interface. These explain the improvement in stress resistance for photoactive layers using the high-modulus filler.

Supplementary Fig. 10d is the bilayer model with two different moduli; the two sides of the bar are fixed, and an upward force is applied to bend it. The model using the modulus of that measure for UVC+ASH film shows the least displacement under the bending force, demonstrating the improvement in stress durability for devices using photoactive layer with higher moduli.

Device fabrication and photovoltaic performance
The solid organic solar cells were prepared with a device structure of ITO/PEDOT:PSS/active layer/PNDIT-F3N-Br/Ag. Patterned ITO substrates were cleaned with tetrahydrofuran, deionized water, and isopropanol in an ultrasonic cleaner for 15 minutes each and 3 cycles. After drying at 75 °C, the ITO substrates were plasma-cleaned for 150 seconds. PEDOT:PSS was then spin-coated on the ITO substrates at 3000 rpm for 30 s and annealed at 150 °C for 15 min. The active layer (donor and acceptors are 5.4 and 6.5 mg mL−1, respectively) was dissolved in chloroform altogether and stirred for 12 h. The thickness of this layer was controlled between 90 and 110 nm, measured by a Bruker Dektak XT profilometer. It’s followed by thermal annealing at 100 °C for 5 min. The PNDIT-F3N-Br layer was then spin-coated onto the active layers in a methanol solution. Finally, Ag electrodes (~100 nm) were deposited on the PNDIT-F3N-Br.

For devices with active layers more than binary compositions, Y5TV takes 20% parts of the acceptors (1.3 mg mL−1) and L8BO takes the rest (5.2 mg mL−1). PETMP is dissolved in the blend solutions with a concentration of 0.065 mg mL−1. ASH is added to the blend solutions with a concentration of 0.3 mg mL−1.

The stretchable organic solar cells were prepared with a device structure of TPU/PH1000/PEDOT:PSS (4083)/active layer/EGaIn. TPU substrates were cleaned with deionized water in an ultrasonic cleaner for 15 minutes and 3 cycles. After drying at 30 °C, the TPU substrates were plasma-cleaned for 180 seconds. PH1000 (with 5 vol% dimethyl sulfoxide, 2 vol% polyethyleneglycol, 0.5 vol% Zonyl fluorosurfactant FS-30 blended to improve the performance) was then spin-coated on the TPU substrates at 2000 rpm for 60 s and annealed at 100 °C for 10 min. PEDOT:PSS (CLEVIOS P VP AI 4083) was then spin-coated at 3000 rpm for 60 s and annealed at 100 °C for 10 min. The active layer materials (donors and acceptors are 5.4 and 6.5 mg mL−1, respectively) are dissolved in chloroform altogether and stirred for 12 h, then spin-coated as the parameters used in solid devices. Methanol was then spin-coated onto the active layers. Finally, EGaIn electrodes were deposited on the active layers by spraying as shown in Supplementary Fig. 19 with the masks of 0.04 cm2 area and 1.5 cm2 area, respectively. The area of the IS-OPV devices is estimated by multiplying the strain by (1+ε).

The hole-only devices were configured by ITO/PEDOT:PSS/active layer/MoO3/Ag, while the electron-only device structure was ITO/ZnO/active layer/PFN-Br/Ag. The fabrication process was similar to the device above. Mobilities are calculated with the Mott-Gurney equation: J=9ε0εRμhV2/8d3, which means μ = 8Jd3 / 9ε0εRV2 (cm2 V−1 s−1). The J–V curves were measured with a Keithley 2400 source meter under 100 mW cm−2, AM 1.5G solar simulator (Enlitech SS–F5). The J–V curves were measured with forward scan mode from −0.1 V to 1.0 V, and the scan step was 0.01 V with a dwell time of 1 ms. The light intensity was calibrated with a standard silicon solar cell certified by NREL. Light-intensity dependences were measured by varying the incident flux via the illumination aperture, followed by calibrating the intensity with the standard silicon solar cell. The EQE spectra were tested on a commercial QE measurement system (Enlitech, QER3011). Hole-only devices were recorded with a Keithley 236 source meter under a dark chamber.

The output power of the IS-OPV module was calculated using the standard relation:
P = Ilight × Aeff × PCE (1)
where Ilight = 100 mW cm−2 (AM 1.5G, 1 sun), Aeff = 4.5 cm2 (three devices in series, each 1.5 cm2), and PCE = 14.4%. This corresponds to a theoretical maximum output power of ~65 mW. For the wearable demonstration, a wireless Canum Venaticorum Arduino UNO-based electromyography (EMG) monitoring prototype was powered by a lithium-polymer battery (nominal voltage 3.7 V, effective capacity ~0.70 Wh), ensuring stable device operation. The average load power of the system was estimated from the measured discharge time (1.96 h) to be ~360 mW, with ~100 mW attributed to biosignal acquisition and processing and ~260 mW consumed by wireless data transmission. Three large-area IS-OPV devices were connected through a 330-type boost converter and interfaced with an SPV1050 solar power management module, supplying the EMG charging port to partially offset power consumption or recharge the battery. The effective contribution of photovoltaic charging was quantified by comparing discharge durations with and without solar input, with the improvement attributed to the continuous ~65 mW output from the IS-OPV module under 1 sun illumination. For scaling estimation, the same formula was applied to larger module areas, demonstrating that >30 cm² of active area could generate >430 mW under full sunlight in the future, approaching the continuous operating power of the EMG system.

Characterization methods
The NMR measurements were conducted on a Bruker AVANCE 400 spectrometer operating at 400 MHz, with tetramethyl silane (TMS) used as the internal standard. UV-vis absorption spectra of the solutions or the films spin-cast on quartz glass were measured using a Shimadzu UV-3600 UV-vis-NIR spectrometer. The FTIR is measured on a Nicolet IS50-Nicolet Continuum, Thermo Fisher Scientific. The XPS is measured on an Axis Supra+, Kratos.

The IR-AFM is measured on an Anasys nanoIR3 system, Bruker. Intensity maps were mean-subtracted and Hann-windowed, followed by a 2D Fourier transform and azimuthal averaging to obtain the PSD.

The fully free-standing tensile test was performed using a HY-0230 horizontal tensile system (Shanghai Hengyi Precision Instrument Co., Ltd), following procedures described in our previous work31. The process includes four steps:
(1) Thin film preparation: A PEDOT:PSS (Clevios P VP AI 4083) layer was spin-coated onto oxygen plasma-treated ITO glass at 2000 rpm, followed by thermal annealing at 150 °C for 10 minutes. The active layer blend solution was then spin-coated on top and annealed at 100 °C for 5 minutes inside a nitrogen glovebox.
(2) Film transfer: The coated substrates were floated on deionized water, allowing the PEDOT:PSS layer to dissolve and the ITO glass to sink. The hydrophobic active layer remained afloat, forming a free standing film on the water surface.
(3) Film mounting: The floating film was gently attached to the tensile stage clamps using van der Waals forces from two frosted aluminium blocks. Water was then slowly removed, and the film was air-dried on the stage.
(4) Tensile testing: The film was stretched at a constant rate of 0.03 mm min⁻1 to record the stress-strain response. For cyclic testing, the film was extended to 5% strain at 3 mm min⁻1, held for 5 seconds, and then returned to its original position at the same rate.

GIWAXS experiments were carried out at beamline 7.3.3 of the Advanced Light Source (ALS), Lawrence Berkeley National Laboratory (LBNL). Thin films were deposited on PEDOT:PSS/silicon wafer substrates under the same processing conditions as those employed for device fabrication. The X-ray energy was fixed at 10 keV. Measurements were performed with the samples mounted on a stage inside a helium chamber, using an incident angle of 0.16°. The sample-to-detector distance was ~280 mm, calibrated with silver behenate. Scattering patterns were recorded with a Pilatus 2M detector, featuring a pixel size of 0.172 mm × 0.172 mm.

RSoXS was performed at beamline 11.0.1.2 of ALS. The neat and blend thin films were prepared on PEDOT:PSS-coated wafer substrates under the same conditions as device fabrication, and then floated on deionized water. After dissolution of the PEDOT:PSS layer, the thin films were transferred onto PDMS membranes with a pin-hole (diameter=2.5 mm) at the center. All samples were loaded onto a plate in the sample chamber and measured under high vacuum. The X-ray beam energy was screened within the N K-edge (390.0–400.4 eV). Scattering patterns were recorded using a Princeton Instrument PI-MTE CCD camera, with a pixel size of 0.027 mm × 0.027 mm and an exposure time of 5 seconds.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This study achieved a synergistic balance between power conversion efficiency (PCE) and stretchability in organic photovoltaics (OPVs) through a strategy termed \"molecular crosslinking and filler toughening (MCFT).\" By combining thiol-ene \"click\" crosslinking and the incorporation of ASH as a filler, the authors observed stress dispersion, which simultaneously led to an enhancement in both PCE and crack-onset strain. This impressive achievement in PCE and maximum strain is subsequently demonstrated through a demonstration.\nWhile the reported photovoltaic performance is indeed impressive, the manuscript suffers from a lack of detailed discussion and thorough explanation, which hinders a comprehensive evaluation of its validity. Substantial revisions are necessary to improve its scientific rigor.\n1. A more appropriate discussion regarding the role of ASH is required. The discussion of the AFM-IR results is largely qualitative and does not sufficiently support the claim that ASH addition significantly enhances all parameters, including JSC. Furthermore, the schematic illustration in Figure 4d lacks clarity regarding the experimental basis for its depiction, and it fails to clearly elucidate the role of ASH. The figure needs revision, possibly by adding more detailed explanations.\n2. In Figure 2c, it is unclear which features represent the ASH particles. It is difficult to confirm how the figure supports the statement \"with ASH particles absorbing and redistributing strain while undergoing minimal deformation.\" A scale bar for each color is essential for interpreting the image. In Supplementary Figure 7, \"Ball 1\" and \"Ball 2\" are indicated (in very small font), but their meaning remains ambiguous.\n3. On page 7, the statement \" This is consistent with the optical imaging under strain (Supplementary Fig. S9) where ASH domains visibly interrupt fracture propagation.\" is not convincing to me. Please specify the direction and magnitude of applied strain for both the upper and lower panels in Fig. 2d. From these images, the differences in voids are not immediately apparent, making it difficult to understand what can be inferred. The voids appear to originate near the ASH particles, so how can this observation be claimed to be consistent with Fig. S9, which suggests visible interruption of fracture propagation? Additionally, please add strain direction on Fig. S9.\n4. Figure 3b shows a significant increase in EQE in the 300-400 nm range. The reason for this notable increase should be elucidated.\n5. Figure 2a suggests that UVC+ASH undergoes elastic deformation up to 5% strain, followed by plastic deformation beyond that point. Does the PCE recover to its original value when a strain greater than 5% (e.g., 20%) is applied and then releasing them to 0%? Additionally, cyclic strain stability measurements should be performed. I recommend testing at strain values comparable to previous literature (e.g., 10% or 20%) for comparison.\n6. On page 11, the statement \"Owing to the elevated modulus imparted by the MCFT strategy, the device retained over 80% of its initial PCE while maintaining structural integrity\" is made. Where is this specific result presented? If it refers to Figure 3, then it is redundant to present it again as Fig. 5, and would be more appropriate to mention this result when discussing Fig. 3, perhaps by moving Fig. 5a to the Supplementary Information.\n7. The demonstration shown in Figure 5b is not an adequate \"showcase real-world utility.\" If the intention is to demonstrate that the OPV performance does not change with muscle contraction/expansion, the range of stretchability demonstrated is entirely different from that of real muscle movement. If this is not the intention, then the current setup does not inherently require stretchability for its functionality. Additional explanations or a re-conceptualization of the demonstration setup are advised.\n8. The numbering of SI figures in the main text appears inconsistent, making it challenging to follow the discussion. This needs to be thoroughly checked and corrected.\n9. Overall, the Supplementary Information figures suffer from a severe lack of adequate descriptions, making it difficult for them to support the arguments in the main text. A thorough revision is needed, including adding captions corresponding to each panel, providing detailed explanations within the captions, and clearly specifying what each figure illustrates.", "reviewer_2": "In this manuscript, Luo et al. demonstrate a MCFT strategy to enhance the stretchability of films by integrating a UV-triggered covalent network between conjugated polymer donors and crosslinkable non-fullerene acceptors with a hydrophobic silica filler. Systematic investigations reveal that ASH fillers function as both crack stoppers and morphological modulators and Y5TV acceptor not only suppresses crack propagation under mechanical stress but also optimizes domain size, crystallinity, and phase separation. Consequently, the PzA5:L8BO:Y5TV:ASH device achieves an impressive COS of ~42% while retaining a high PCE of 16.5% in stretchable configurations. The authors further validate the practical utility of these intrinsically stretchable OPVs by successfully integrating them into a self-powered, wireless electromyography monitoring system, underscoring their potential for energy-autonomous wearable electronics. However, several critical concerns warrant attention, and the authors are advised to revise the manuscript accordingly to address these issues.\n1. In the Abstract, the statement “…with 80% of the initial efficiency retained under nearly 40% strain, which is the highest value reported to date for stretchable OPVs” is inaccurate. According to a recent study (Science, 2025, 387, 381), stretchable organic solar cells based on the PNTB6-Cl:BTP-Si4 system retained 82% of the initial PCE under a much higher tensile strain of 80%, which is the highest value reported to date for stretchable OPVs. Furthermore, Fig. 3F also omits this record value. The authors must correct this claim and include the updated data.\n2. Please provide a clear schematic diagram to illustrate the role of material cross-linking and ASH.\n3. According the contact angel data, please provide the Flory-Huggins interaction parameter to explain the miscibility between the two materials.\n4. In Fig 2a, the crack-onset strain (COS) of the UVC-treated film with ASH shows a slight decrease compared to the UVC-treated film without ASH. This suggests that the incorporation of ASH does not enhance the mechanical performance. In addition, what is the thickness of these films used to test mechanical properties?\n5. Apart from the storge modulus, please provide the other data of DMA, for example: loss modulus, and tan δ behavior. And then provided the Tg and sub-Tg based on the different sample.\n6. While the authors conducted finite element simulations to demonstrate the mechanical properties of the UVC-treated film with ASH, the absence of a reference sample (e.g., PzA5:L8BO, UVC-treated film without ASH) limits the interpretability of the results. It is recommended that the authors supplement their analysis with comparative data to establish a baseline for evaluation. Additionally, the manuscript should include detailed descriptions of the simulation methodology, including key input parameters and modeling assumptions, to ensure reproducibility and facilitate validation by readers.\n7. The mechanical property data obtained from molecular dynamics (MD) simulations appear significantly higher than the experimental test results. Could the authors provide a detailed explanation for this discrepancy? Additional, provide a more comprehensive description of the MD simulation methodology.\n8. ASH is an insulating material, and adding it can significantly increase JSC, which is an incomprehensible phenomenon. The authors should explain this result.\n9. Compared to the PzA5:L8BO and PzA5:L8BO:Y5TV samples, the hole mobility of PzA5:L8BO:Y5TV exhibits a remarkable increase after UVC treatment, while its electron mobility shows a significant decrease (Table S1). Please explain it.\n10. As shown in Table S3, the trends in CCL value and area are inconsistent. Please explain it. Also, the relevant crystal parameters related to (111) and (021) mentioned in the manuscript are missing.\n11. Please provide the following key information: 1) specific performance parameters (e.g., PCE, FF, JSC, stretchability stability, etc.) and detailed fabrication processes of the stretchable organic photovoltaic module; 2) operational parameters of the EMG sensor in the integrated device (including required power and voltage range); 3) the power supply source for the wireless signal transmission module (e.g., photovoltaic module or other power sources). The technical details of each module should be clarified to ensure the integrity and reproducibility of the study.\n12. There are some errors in the manuscript. Please revise them. For example, in Fig S17a, the tensile strain of the PzA5:L8BO blend film should 0 to 2.5% not 0 to 40%."}, "author_response": "Respond to the reviewers’ comments\nReviewer 1\nThis study achieved a synergistic balance between power conversion efficiency (PCE)\nand stretchability in organic photovoltaics (OPVs) through a strategy termed\n\"molecular crosslinking and filler toughening (MCFT).\" By combining thiol-ene\n\"click\" crosslinking and the incorporation of ASH as a filler, the authors observed stress\ndispersion, which simultaneously led to an enhancement in both PCE and crack-onset\nstrain. This impressive achievement in PCE and maximum strain is subsequently\ndemonstrated through a demonstration.\nWhile the reported photovoltaic performance is indeed impressive, the manuscript\nsuffers from a lack of detailed discussion and thorough explanation, which hinders a\ncomprehensive evaluation of its validity. Substantial revisions are necessary to improve\nits scientific rigor.\n1. A more appropriate discussion regarding the role of ASH is required. The discussion\nof the AFM-IR results is largely qualitative and does not sufficiently support the claim\nthat ASH addition significantly enhances all parameters, including JSC. Furthermore,\nthe schematic illustration in Figure 4d lacks clarity regarding the experimental basis for\nits depiction, and it fails to clearly elucidate the role of ASH. The figure needs revision,\npossibly by adding more detailed explanations.\nResponse: We fully agree with the reviewer on the comments. In the revised\nmanuscript, we expanded the discussion of the AFM-IR, RSoXS, and GIWAXS results\nin the “MCFT morphology” section to illustrate the role of ASH and explicitly link the\nobserved morphological changes with the corresponding photovoltaic improvements,\nincluding the J enhancement. Specifically, we now explain how the denser and more\nSC\ninterconnected fibrillar network (in AFM-IR image), the improved crystallinity and π–\nπ stacking (in GIWAXS pattern), and the larger, purer domains with long-range\norganization (in RSoXS fitted results) each contribute to improved exciton dissociation,\nbalanced carrier transport, reduced recombination, and ultimately the increases in J ,\nSC\nFF, and V .\nOC\nIn addition, to provide a quantitative assessment of morphology changes, we\nadjusted the AFM-IR scale and performed power spectral density (PSD) analysis.\nFurthermore, Fig. 4d and its description has been updated to clarify the\nexperimental basis for each schematic element, ensuring that the role of ASH as both a\nmorphological modulator and mechanical reinforcement is clearly supported by the\nexperimental results.\nMain text:\n“Power spectral density (PSD) analysis of the AFM topographies shows a\npronounced peak emerging at k≈0.0025 Å−1 (d=251 nm) upon crosslinking and\nfurtherly shifting to ∼0.0015 Å−1 for UVC+ASH, accompanied by a marked\nsuppression of spectral power at low (<0.001 Å−1) and high (>0.01 Å−1) frequencies\n(Supplementary Fig. 22–23). These features indicate a growth of the characteristic\ndomain size together with the damping of large-scale inhomogeneity and nanoscale\nroughness, aligning with the formation of purer yet more fibrillar phases. Such\npronounced morphological optimization can be ascribed to the strong interfacial\ninteractions between ASH and the donor/acceptor components, which facilitate more\nfavorable packing and phase organization. This fibrillar morphology facilitates\ncontinuous charge percolation pathways, while the finer yet purer domains promote\nefficient exciton diffusion and dissociation, jointly underpinning the observed J\nSC\nenhancement in devices.”\n“The strengthened π–π stacking would increase carrier mobility, benefiting both\nJ and FF, while the improved crystalline order lowers trap-assisted recombination,\nSC\nproviding a modest V gain. This enhancement likely arises from the gel-like\nOC\nrheological behaviour of ASH (Supplementary Fig. 26), which modulates phase\nseparation kinetics during solvent evaporation and thus reshapes the thin film\nmorphology.”\n“Upon ASH incorporation, this feature became more intense and shifted to lower\nq (~0.0095 Å⁻1), reflecting the formation of larger and purer domains.\nThe combined results of GIWAXS, AFM-IR, and NK-RSoXS reveal that structural\nevolution under ASH incorporation is in line with the enhanced EQE response, where\nincreased molecular ordering and refined, purer phase-separated domains collectively\nbroaden absorption shoulders, extend the continuity for exciton diffusion and carrier\ntransport across multiple length scales, and reinforce J and FF gains, consistent with\nSC\nresults of TA spectroscopy and device EQEs. A schematic summary of the morphology\nevolution and the synergistic roles of crosslinker and filler is illustrated in Fig. 4d,\nhighlighting the improved fibrillar domains and networks that underpin the enhanced\nmechanical and optoelectronic performance of the UVC+ASH system.”\na b PzA5:L8BO d PzA5:L8BO\n0\nt y -2\nsi UVC\nn\nt e -4\nn\nI\nUVC\nUVC+ASH\nUVC\n4\n2\n1 2\n-2\nc\nN K-edge at 399.2 eV -4\nPzA5:L8BO UVC+ASH\nUVC\nUVC+ASH\nUVC+ASH\nsit y 2\nn\ne\nnt\n-2\nFiber/domain of L8BO\nPzA5 donor Y5TV\n-4\nPETMP crosslinker\n0.003 0.01 0.03\nASH filler cluster\nQ vector (Å-1)\nFig. 4 | Morphology properties for the photoactive films. a Grazing incidence wide-angle X-ray\nscattering sector-averaged I–q curves in the IP (dashed lines) and OOP (solid lines). b AFM-IR\nmapping of the blend films at the wavenumber of 1426 cm−1. c NK-RSoXS averaged I–q curves for\nPzA5:L8BO, UVC, and UVC+ASH blends. d Schematic of the effect of crosslinker and filler in the\nblend. The purple and cyan translucent strips represent the synergetic interwoven networks in the\nblend film.\nSupplementary information:\na\nb c\nSupplementary Fig. 22 | Power spectral density (PSD) fitting process. PSD fitting\nof AFM-IR mapping of the a PzA5:L8BO, b UVC, and c UVC+ASH blend films at the\nwavenumber of 1426 cm−1.\nPzA5:L8BO\nUVC\nUVC+ASH\nSupplementary Fig. 23 | Power spectral density (PSD) plots derived from AFM-\nIR mapping of the blend films.\n2. In Figure 2c, it is unclear which features represent the ASH particles. It is difficult to\nconfirm how the figure supports the statement \"with ASH particles absorbing and\nredistributing strain while undergoing minimal deformation.\" A scale bar for each color\nis essential for interpreting the image. In Supplementary Figure 7, \"Ball 1\" and \"Ball 2\"\nare indicated (in very small font), but their meaning remains ambiguous.\nResponse: We thank the reviewer for this helpful suggestion. Figure 2c was\nredrawn to better illustrate the particles and the stress-absorbing and redistributing role\nof the high-modulus filler particles, achieved by comparing particles moduli of 100 and\nredistributing strain” was revised to “absorbing and redistributing stress”. The stress\ndistribution color scale and scale bar were adjusted to clearly demarcate particle\nboundaries. In Supplementary Figure 7 (now Supplementary Fig. 8), the “Ball 1” and\n“Ball 2” labels, which referred to particles of identical modulus but different diameters\nto simulate filler size polydispersity, were removed. This ambiguous notation was\nreplaced with a new Supplementary Fig. 8a, where color contrast is used to distinguish\nphotoactive material and particle material domains. Collectively, the revised Fig. 2c and\nSupplementary Fig. 8 reveal the mesoscopic mechanism by which high-modulus fillers\nabsorb and redistribute stress while undergoing negligible deformation, thereby\nestablishing a deformable but toughening filler network that reinforces the low-\nmodulus photovoltaic active layer.\nMain text:\n“Under tensile load, stress is concentrated at the filler–matrix interfaces, with ASH\nparticles absorbing and redistributing stress while undergoing minimal deformation.”\nE =100 MPa\nparticle\na b c Stress (MPa)\nP a)40\nP\nU U\nz\nV V\nA\nC C\n5\n+\n:L\nA\nB\nH\nO\nk P a) 7.5\nN\ness (\nM\n1\n2\n3\n0\n0\n0 1\n1\n3\n0\n9\n6\n9\n5\n9\nM\nM\nP\nP\nP\na\na\na\nΔ\n(\nE'0.2\n%-2 %\n2\n5\n.\n.\n5\n0\n1.9 2.0\nE activelayer (matrix) =200 MPa\n0 0.0\n0 10 20 30 40 50 PzA5 PzA5:L8BO UVC UVC+ASH 0\nStrain (%)\nE =2000 MPa\nparticle\nPzA5:L8BO\nd P U z V A C 5 + :L A 8 S B H O( ( σ σ a a p p p p = = 0 0 . . 1 2 5 5 G G σ P P a a , , P L C A ε ε 8 z r S = = o A B H s O 5 2 9 s f l i 5 1 p a i l n l o c e . . k 7 0 c l r y e e % % c m d p l u t e m ) ) o s r t r a e d t r o e n ri o al r s e e str ess ( M P a) 2 3 4 0 0 0 P U U z V V A C C 5 + :L A 8 S B H O C f ra s c t ( k r % - a o i ) n ns 6 e 0 t 40 2 Y 0 ield 0 stra 3 6 9 1 i 2 n (% 30 ) 0 600 90 U U 0 V V 12 C C 0 + m 0 Y A ( o o M S u d H P n u a g lu ) 's s\nσ\nc y\ncli c t e\nnsil\n1\n0\n(\ns\nM\nt x r i e m\nP\ns\na\nu s\n)\n6 m 0 40\n20\n0\n1\n4 8\n2\n0 0\n0\n0 0\n10 20 30 40\n(\ns Y\nM\nt i r e e\nP\nl s\na\nd s\n)\n0 5 10\nFig. 2 | Mechanical performance of the blends. a Stress-strain curves of neat and blend films. b\nDynamic modulus of the blend films. c Stress distribution in finite element analysis on a matrix with\nfiller particles of different modulus. d Snapshot of the stopped-extended crack in the UVC+ASH\nblend film (with molecular colours as the legend). e Stress-strain curves at the 100th cycle below\nyielding strain (3.1% for PzA5:L8BO, 8.8% for UVC, and 10.2% for UVC+ASH blend films,\nrespectively). f Summary of various mechanical parameters for blend films.\nSupplementary information:\na d\nE=203MPa\nb e\nStress (MPa) Stress (MPa)\n0.5 E=369MPa 1\nE=1090MPa\n0 0\nc f\n0\nlow modulus elastomer\nStretc hi n g\ndis\nplace\nme nt ( m m)\n0\n1\n.\n.\n5\n0 high modulus filler\nBe\nn di n g\ndis\nplace\nme nt ( m m)\n−\n−\n−\n−\n4\n3\n2\n1\n1\n3\n2\n0\n6\n0\n9\n9\n3\nM\nM\nP\nP\nP\na\na\na\n0.0\n0.0 0.5 1.0 0.0 0.5 1.0\nTime increment Time increment\nSupplementary Fig. 8 | Finite element analysis of filler and bilayer models. a\nSchematic of particles embedded in the photoactive matrix, with different diameters\nused to simulate size polydispersity. b Stress distribution in blends under 10 N tensile\nload with identical matrix modulus (200 MPa) but different particle moduli: 100 MPa\n(soft elastomer, top) and 2000 MPa (hard filler, bottom). c Elongation under 10 N\ntensile load with varying particle moduli. d Bilayer model consisting of a photoactive\nlayer on a substrate (Young’s modulus of the substrate is set as 500 MPa). e Stress\ndistribution in the bilayer under 10 N mid-span load as the modulus of the photoactive\nlayer increases. f Bending deformation under mid-span load with photoactive layers of\ndifferent moduli.\n3. On page 7, the statement \" This is consistent with the optical imaging under strain\n(Supplementary Fig. S9) where ASH domains visibly interrupt fracture propagation.\"\nis not convincing to me. Please specify the direction and magnitude of applied strain\nfor both the upper and lower panels in Fig. 2d. From these images, the differences in\nvoids are not immediately apparent, making it difficult to understand what can be\ninferred. The voids appear to originate near the ASH particles, so how can this\nobservation be claimed to be consistent with Fig. S9, which suggests visible\ninterruption of fracture propagation? Additionally, please add strain direction on Fig.\nS9.\nResponse: In the revised version of the manuscript, we specified the direction and\nmagnitude of the applied stress, as well as the corresponding strain of the illustrated\nsnapshot before crack in Fig. 2d.\nTo provide a clearer comparison, we added a new figure (Supplementary Fig. 10)\nshowing the sequential fracture evolution in the three models. In the PzA5:L8BO model,\ncracks nucleate at the L8BO interfaces (L8BO molecules are colored blue), consistent\nwith the limited stretchability of conventional blends using only small molecular\nacceptors. In the PzA5:L8BO:Y5TV (crosslinked with PETMP) model, cracks still\ninitiate at L8BO interfaces but the presence of tie chains delays fracture, and\ndisentanglement of crosslinked chains (crosslinked donor and acceptor chains are\ncolored purple) leads to a pronounced yielding process. In the crosslinked\nPzA5:L8BO:Y5TV model with ASH fillers, this disentanglement process is further\narrested when it reaches the filler interfaces (ASH filler clusters are colored cyan),\npreventing cracks from growing further.\nImportantly, the successive snapshots demonstrate that voids originate between\nL8BO molecules in all three models rather than at ASH particles; on the contrary, ASH\ndomains act as barriers that stop or deflect crack propagation. We apologize for the\nearlier presentation of only the final snapshot, which may have caused some\nmisunderstanding. With this clarification, the fracture-arresting role of ASH fillers\nobserved in the simulations (Fig. 2d and Supplementary Fig. 10) is consistent with the\noptical imaging in Fig. S9 (now Supplementary Fig. 11), where crack paths are visibly\ninterrupted by ASH domains. The strain direction has also been added.\nMain text.\nE =100 MPa\nparticle\na b c Stress (MPa)\nP a)40\nP\nU U\nz\nV V\nA\nC C\n5\n+\n:L\nA\nB\nH\nO\nk P a) 7.5\nN\ness (\nM\n1\n2\n3\n0\n0\n0 1\n1\n3\n0\n9\n6\n9\n5\n9\nM\nM\nP\nP\nP\na\na\na\nΔ\n(\nE'0.2\n%-2 %\n2\n5\n.\n.\n5\n0\n1.9 2.0\nE activelayer (matrix) =200 MPa\n0 0.0\n0 10 20 30 40 50 PzA5 PzA5:L8BO UVC UVC+ASH 0\nStrain (%)\nE =2000 MPa\nparticle\nPzA5:L8BO\nd P U z V A C 5 + :L A 8 S B H O( ( σ σ a a p p p p = = 0 0 . . 1 2 5 5 G G σ P P a a , , P L C A ε ε 8 z r S = = o A B H s O 5 2 9 s f l i 5 1 p a i l n l o c e . . k 7 0 c l r y e e % % c m d p l u t e m ) ) o s r t r a e d t r o e n ri o al r s e e str ess ( M P a) 2 3 4 0 0 0 P U U z V V A C C 5 + :L A 8 S B H O C f ra s c t ( k r % - a o i ) n ns 6 e 0 t 40 2 Y 0 ield 0 stra 3 6 9 1 i 2 n (% 30 ) 0 600 90 U U 0 V V 12 C C 0 + m 0 Y A ( o o M S u d H P n u a g lu ) 's s\nnsil\n20\n0\n0 0\nσ\nc y\ncli c t e\n1\n0\n(\ns\nM\nt x r i e m\nP\ns\na\nu s\n)\n6 m 0 40\n1\n4 8\n2\n0 0\n0\n10 20 30 40\n(\ns Y\nM\nt i r e e\nP\nl s\na\nd s\n)\n0 5 10\nFig. 2 | Mechanical performance of the blends. a Stress-strain curves of neat and blend films. b\nDynamic modulus of the blend films. c Stress distribution in finite element analysis on a matrix with\nfiller particles of different modulus. d Snapshot of the stopped-extended crack in the UVC+ASH\nblend film (with molecular colours as the legend). e Stress-strain curves at the 100th cycle below\nyielding strain (3.1% for PzA5:L8BO, 8.8% for UVC, and 10.2% for UVC+ASH blend films,\nrespectively). f Summary of various mechanical parameters for blend films.\nSupplementary information:\nApplied\nPzA5:L8BO:Y5TV PzA5:L8BO:Y5TV\nStress PzA5:L8BO\n(crosslinked by PETMP) (crosslinked by PETMP)+ASH\n(horizontal\ndirection)\n0.15 GPa\n0.20 GPa\n0.25 GPa\nSupplementary Fig. 10 | Representative MD snapshots and stress values during\ncrack initiation and growth under tensile loading. Key trajectories are shown for\ndifferent molecular dynamics models, highlighting the evolution of crack formation and\npropagation.\npristine 40% strain\ntensile\n10 μm\nSupplementary Fig. 11 | Optical microscopy images of the UVC+ASH film on a\nTPU substrate at 0% and 40% strain.\n4. Figure 3b shows a significant increase in EQE in the 300-400 nm range. The reason\nfor this notable increase should be elucidated.\nResponse: The EQE increase in the 300–400 nm region originates from ASH-\ninduced morphological optimization: GIWAXS data show improved local molecular\nordering, which strengthens donor absorption shoulders, while RSoXS data reveal purer\ndomains that facilitate exciton dissociation. These structural improvements directly\ntranslate into faster exciton diffusion, more efficient dissociation, and sustained carrier\ntransport, as confirmed by transient spectra and device measurements. Such\nmorphology-driven enhancements in charge generation and transport underpin the\nobserved EQE rise in the short-wavelength region.\nMain text:\n“The combined results of GIWAXS, AFM-IR, and NK-RSoXS reveal that structural\nevolution under ASH incorporation is in line with the enhanced EQE response, where\nincreased molecular ordering and refined, purer phase-separated domains collectively\nbroaden absorption shoulders, extend the continuity for exciton diffusion and carrier\ntransport across multiple length scales, and reinforce JSC and FF gains, consistent\nwith results of TA spectroscopy and device EQEs. A schematic summary of the\nmorphology evolution and the synergistic roles of crosslinker and filler is illustrated in\nFig. 4d, highlighting the improved fibrillar domains and networks that underpin the\nenhanced mechanical and optoelectronic performance of the UVC+ASH system.”\n5. Figure 2a suggests that UVC+ASH undergoes elastic deformation up to 5% strain,\nfollowed by plastic deformation beyond that point. Does the PCE recover to its original\nvalue when a strain greater than 5% (e.g., 20%) is applied and then releasing them to\n0%? Additionally, cyclic strain stability measurements should be performed. I\nrecommend testing at strain values comparable to previous literature (e.g., 10% or 20%)\nfor comparison.\nResponse: We thank the reviewer for this important suggestion. Regarding the first\nquestion, the cyclic strain stability test itself directly addresses whether PCE can\nrecover after stretching to a specific strain and releasing it back to 0%. As shown in\nSupplementary Fig. 18, when UVC+ASH devices are stretched beyond 5% strain,\npartial recovery is possible at moderate levels, but PCE cannot fully recover after 20%\nstrain cycles. Specifically, after repeated 20% cyclic stretching and release, the PCE\ndrops to 7.4% and 44.8% of the initial value (Supplementary Fig. 18 g–i), indicating\nirreversible degradation. It is also important to note that strain is a manifestation of\napplied stress. The synergistic crosslinking–filler design substantially increases the\nactive-layer modulus, so that a given external force translates into much smaller strains\nthan in softer films. Thus, while UVC+ASH devices are not optimized for repeated 20%\nstrain cycles, they are inherently less likely to encounter such extreme deformations\nunder realistic mechanical loading, as even 10% strain already corresponds to a stress\napproaching 40 MPa. This highlights that modulus enhancement is as critical as\nelongation for sustaining device durability.\nWe added cyclic strain stability tests for both devices using PzA5:L8BO and the\nUVC+ASH active layers (Supplementary Fig. 18 g–i): the devices maintain reasonable\nstability under repeated 10% strain cycles, whereas significant and irreversible\ndegradation occurs under 20% cyclic stretching, consistent with previous reports.\nMain text:\n“In addition, cyclic stretching tests at defined strains (5%, 10%, or 20%) over 100–\n200 cycles further underline the advantage of UVC+ASH, which exhibits significantly\nhigher performance retention than PzA5:L8BO throughout repeated loading–\nunloading processes.”\nSupplementary information:\na 5 b 5 c\nPzA5:L8BO UVC+ASH 100 PzA5:L8BO\n0 0 UVC+ASH\n-2\nC\n(\nm )\nu\nA\nr r\nc\ne\nm\nnt\nDensity\n- -\n-\n2 1\n1\n- 0 5\n0\n1 2 3 0\n(%\n0 0 0\n)\n-2\nC\n(\nm )\nu\nA\nr r\nc\ne\nm\nnt\nDensity\n- -\n-\n2 1\n1\n- 0 5\n0\n2\n3 4\n0 (%\n0\n0\n0 0\n)\nNor\nmalized\nPCE (%)\n2 5\n7\n5 0\n5\n-25 40 -25\n-30 -30 0\n0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 0 15 30 45\nVoltage (V) Voltage (V) Stretching strain (%)\nd 5 e 5 f\nPzA5:L8BO UVC+ASH 100 PzA5:L8BO\n0 0 UVC+ASH\n-2\nC\n( m\n)\nu\nA\nr r\nc\ne\nm\nnt\nDensity\n- -\n-\n-\n2 2\n1\n1 -\n5 0\n5\n0 5\nTensil e 1 0 0 c 0 ycles\n-2\nC\n( m\n)\nu\nA\nr r\nc\ne\nm\nnt\nDensity\n- -\n-\n-\n2 2\n1\n1 -\n5 0\n5\n0 5 Tensil e 0 1\n2\n0\n0\nc 0\n0\nycles\nNor\nmalized PCE (%)\n2\n5\n7\n5\n0\n5\n200\n-30 -30\n0\n0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 0 50 100 150 200 250 300\nVoltage (V) Voltage (V) Tensile cycle\ng 5 h 5 i 120\nTensile cycles Tensile cycles PzA5:L8BO ε=10%\n-2\nC\n(\nm )\nu\nA\nr r\nc\ne\nm\nnt\nDensity\n− −\n−\n−\n2 1\n1\n0 5\n0 5\nU\nz\nV\nA\nC\n5\n+\n: L\nA\nε ε\n0\nε\nε\n0 8 = =\n=\n=\nP\nB 1 1\n1\nO 0 0\n0\n0\n% %\n%\n%\n1 2\n1\n2\n0 0\n0\n0\n0 0\n0\n0\nc c\nc\nc\ny y\ny\ny\nc c\nc\nc\nl l\nl\nl\ne e\ne\ne\n-2\nC\n(\nm )\nu\nA\nr r\nc\ne\nm\nnt\nDensity\n− −\n−\n−\n2 1\n1\n0 5\n0 5\n0 ε\nε\n0 8 =\n=\n=\nP B 2\n2\n0\n0\n%\n%\n%\n1\n1\n2\n0\n0\n0\n0\n0\n0\nc\nc\nc\ny\ny\ny\nc\nc\nc\nl\nl\nl\ne\ne\ne\nNor\nmalized\nPCE 10\n4\n6\n8 0\n0\n0\n=\n= = 0 1 4\n0\n1 2 . 0 0 1\n.\n% .\n1\n5 % % %\n20\n−25 −25\n0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 0 0 50 100 150 200\nVoltage (V) Voltage (V) Tensile cycle\nSupplementary Fig. 18 | The photovoltaic performance of the intrinsically\nstretchable organic solar cells under stretching strain and cyclic tensile tests. a–c\nSustained strain=0, 10, 20, 30, and 40%. d–f Cyclic strain=5% for 0, 100, and 200\ncycles. g–i Cyclic strain=0, 10, and 20% for 0, 100, and 200 cycles.\n6. On page 11, the statement \"Owing to the elevated modulus imparted by the MCFT\nstrategy, the device retained over 80% of its initial PCE while maintaining structural\nintegrity\" is made. Where is this specific result presented? If it refers to Figure 3,\nthen it is redundant to present it again as Fig. 5, and would be more appropriate to\nmention this result when discussing Fig. 3, perhaps by moving Fig. 5a to the\nSupplementary Information.\nResponse: In the revised manuscript, we have moved Fig. 5a to the Supplementary\nInformation and now mention the corresponding result in the discussion of Fig. 3 to\navoid redundancy.\nMain text:\n“Crucially, these devices maintained 80% of their initial efficiency under a stretching\nloading of 1 kg resulting in tensile strain of approximately 40% (Fig. 3e, Supplementary\nFig. 17–18), representing one of the highest reported levels of strain endurance among\nIS-OPVs.”\n7. The demonstration shown in Figure 5b is not an adequate \"showcase real-world\nutility.\" If the intention is to demonstrate that the OPV performance does not change\nwith muscle contraction/expansion, the range of stretchability demonstrated is\nentirely different from that of real muscle movement. If this is not the intention, then the current\nsetup does not inherently require stretchability for its functionality. Additional\nexplanations or a re-conceptualization of the demonstration setup are advised.\nResponse: We fully agree with the reviewer that actual muscle contraction and\nrelaxation involve complex deformations are not fully captured by our demonstration.\nOur intention here was to simplify this scenario by considering the expansion of arm\ncircumference during muscle tightening, which naturally imposes tensile strain on the\nattached device. This provides a straightforward rationale for requiring stretchability in\nsuch wearable applications.\nTo clarify the issues, we revised the description in the main text and added an\nillustrative schematic of arm expansion in Fig. 5a. Besides, the core contribution of this\nwork lies in the MCFT-based active layer design, and the simplified demonstration is\nintended only to validate its effectiveness. Translation of this strategy into practical\napplication scenarios will require further efforts and exploration in future studies.\nAdditional explanations are also supplemented:\nMain text:\n“To showcase real-world utility, we integrated the MCFT-based IS-OPV into a\nwearable, wireless electromyography (EMG) monitoring system (Fig. 5b), where it\nserved as a self-powered energy module. The system successfully recorded muscle\nelectrical signals, distinguishing between contraction and relaxation states (inset, Fig.\n5b). Such functionality is directly applicable to biomechanical monitoring, including\nmyoelectric prosthesis control, soft exoskeletons, human-machine interfaces, and\nfatigue detection via frequency-domain EMG analytics.”\na\nmuscle contraction\nmuscle relaxation\nIntrinsically stretchable OPV powered\nwearable wireless EMG monitoring\nPower (mW)\nb c\n2000 1200\nwithout IS-OPV\nal 1.96 h\nn 1600\nsi\ng\n-\n2)\nm 900\no g r a\np\nhi c\nwith IS-OPV e n sit y\n(\nW\n1200\n600\no\nm\ny\n2.30 h\nht\ni\nnt 800\n400\ne\nc t r Li g\n400\n250 300\nEl\n100\n59\n0 0\n4000 6000 8000 0 5 10 15 20 25 30 35 40\nEMG operation time (s) Device area (cm2)\nFig. 5 | Proof-of-concept for IS-OPV. a Mechanical resistance of an intrinsically stretchable OPV\n(IS-OPV) device to a 1 kg load. b A wearable wireless electromyography (EMG) monitoring\nprototype charged by IS-OPVs fitted and packed as part of the elastic wristband. c The collected\nsignals and the maximum operation time of the wireless EMG monitoring prototype in outdoor\nsunlight sports scenes (with and without IS-OPV power supply devices). d The power output of IS-\nOPVs with different areas under different light intensities estimated based on a PCE of 16.5%.\n8. The numbering of SI figures in the main text appears inconsistent, making it\nchallenging to follow the discussion. This needs to be thoroughly checked and corrected.\nResponse: We thank the reviewer for pointing this out. We have thoroughly audited\nand corrected the Supplementary Information figure numbering. All SI figure numbers\nin the main text, captions, and SI have been renumbered to follow the order of first\ncitation, and the corresponding in-text references have been updated for consistency.\n9. Overall, the Supplementary Information figures suffer from a severe lack of adequate\ndescriptions, making it difficult for them to support the arguments in the main text.\nA thorough revision is needed, including adding captions corresponding to each panel,\nproviding detailed explanations within the captions, and clearly specifying what each\nfigure illustrates.\nResponse: The Supplementary Figure captions have been thoroughly revised to\ninclude panel-by-panel descriptions and detailed explanations, ensuring that each figure\nis clearly interpretable and directly supports the main text.\nReviewer #2\nIn this manuscript, Luo et al. demonstrate a MCFT strategy to enhance the stretchability\nof films by integrating a UV-triggered covalent network between conjugated polymer\ndonors and crosslinkable non-fullerene acceptors with a hydrophobic silica filler.\nSystematic investigations reveal that ASH fillers function as both crack stoppers and\nmorphological modulators and Y5TV acceptor not only suppresses crack propagation\nunder mechanical stress but also optimizes domain size, crystallinity, and phase\nseparation. Consequently, the PzA5:L8BO:Y5TV:ASH device achieves an impressive\nCOS of ~42% while retaining a high PCE of 16.5% in stretchable configurations. The\nauthors further validate the practical utility of these intrinsically stretchable OPVs by\nsuccessfully integrating them into a self-powered, wireless electromyography\nmonitoring system, underscoring their potential for energy-autonomous wearable\nelectronics. However, several critical concerns warrant attention, and the authors are\nadvised to revise the manuscript accordingly to address these issues.\n1. In the Abstract, the statement “…with 80% of the initial efficiency retained under\nnearly 40% strain, which is the highest value reported to date for stretchable OPVs” is\ninaccurate. According to a recent study (Science, 2025, 387, 381), stretchable organic\nsolar cells based on the PNTB6-Cl:BTP-Si4 system retained 82% of the initial PCE\nunder a much higher tensile strain of 80%, which is the highest value reported to date\nfor stretchable OPVs. Furthermore, Fig. 3F also omits this record value. The authors\nmust correct this claim and include the updated data.\nResponse: We thank the reviewer for pointing this out. The record reported in\nScience 2025, 387, 381 (PNTB6-Cl:BTP-Si4) has been included in Fig. 3f of our\nmanuscript, where the data point at PCE=14.6% with ε ≈ 80% at 80% PCE corresponds\nto this literature.\nTogether with two other representative reports (PCE = 16.36% with ε = 21% at 80%\nPCE, and PCE = 15.2% with ε = 40% at 80% PCE), these benchmarks delineate the\ncurrent trade-off between PCE and stretchability (e.g. strain, ε) in state-of-the-art\nstretchable OSCs. Our device here (PCE = 16.5%, ε ≈ 39.4% at 80% retention) lies\nabove the guideline connecting these three benchmark points, highlighting another\npathway to balance efficiency and strain tolerance. We further note that while some\nstrategies emphasize achieving extreme strain endurance, other scenarios prioritize\nmaximizing efficiency with moderate but sufficient strain tolerance. As the practical\nrequirements remain diverse and application-specific, it is essential to pursue record\nperformance across different strain regimes.\nWe have also corrected the Abstract into “which is one of the highest values\nreported to date for stretchable OPVs”.\nMain text:\n(Abstract) “The silica filler promotes enhanced aggregation and molecular\nordering in both donor and acceptor materials, enabling a power conversion efficiency\nexceeding 16% for intrinsically stretchable devices, with 80% of the initial efficiency\nretained under nearly 40% strain, which is one of the highest values reported to date\nfor stretchable OPVs.”\n“Depending on the target scenario, either higher strain tolerance or maximized\nefficiency at moderate strain may be more relevant, suggesting the need for\ncomplementary design strategies across different strain regimes.”\na b c\n-2\nC\n( m\n)\nu\nA\nr r\nc\ne\nm\nnt\ndensity\n−\n−\n−\n3\n2\n1 0\n0\n0\nE\nQE (%)\n1\n2\n4\n6 8 0\n0\n0\n0\n0 0 0\nP U\nU\nz V\nV\nA C\nC\n5\n+\n:L\nA\nB\nH\nO\n0\n5\n1\n1 2 2 3\n0\n5 0 5 0 -2\nI\n)\nntegrated\ncurrent density ( m A c m\nNor\nmalized\ndecay at 649 n m\n−\n−\n− − −\n1\n0\n0 0 0\n.\n.\n. . .\n0\n8\n6 4 2 P (h z o A l 5 e G t U U P r S a z V V A n B C C s 5 + f : e L A r 8 S B c H h O annel)\n(×1\n2 4\n5\n0\nk\n. .\n.\nH\n-\n8 7\n1\n8 0\n9\n2s-1)\n0.0 0.2 0.4 0.6 0.8 1.0 300 400 500 600 700 800 900 1000 0.1 1 10 100 1000\nVoltage (V) Wavelength (nm) Delay time (ps)\nd e f\n0 0\nPzA5:L8BO Strain (%) this work\n2) m UVC+ASH -2) m 0 16 (39.4%, 16.5%)\nm\nA/c\n−10 m A\nc\n-10\n1\n2\n0\ndensity (\n−20\nC\n(\nurrent\ndensity\n-20\n3\n4\n0\nCE\n(%)\n1\n1\n2\n4\n−30 -30 10\n0.0 0.2 0.4 0.6 0.8 1.0 0.0 0.2 0.4 0.6 0.8 1.0 10 20 30 40 50 60 70\nVoltage (V) Voltage (V) Stretching strain at PCE80% (%)\ng\n0-0/0-1\npeak rati o\n1\n1\n1 1\n.\n.\n.\n.\n0\n1\n3 4\n0.1\nP P\nU U\nz z\nV V\nA A\nC C\n5 5\n+ +\n: : L L\nA A\n8 8\nS S\nB B\nH H\nO O\nD A\n10\nh\n-2\nC\n)\nurrent\ndensity ( m A c m\n1\n2 3\n0\n0\n0 0\n0 2\nP\nP P U U U\nz\nz z V V V\nA\nA A C C C\n5\n5 5 + + +\n:\n: :\nL\nL L A A A\n8\n8 8\nB\nB B H H H\nO\nO O ε ε ε\nε\nε ε = = =\n=\n= = 0 1 4\n0\n1 2 . 0 0 1\n.\n% .\n1\n5 % % %\ni\n- 2 -\n-\n5 1\n)\nV\ns\nc\na\nm\nrrier\nm o\nbility (\n×10 1 1 1 1\n1 2\n6\n8\n0 2 4 6\n8 0\n0.1 1 1 0\nP\nU\nz\nV\nA\nC\n5\n+\n:L\nA\nB\nH\nO\nStrain (%) Time (μs) Strain (%)\nFig. 3 | Photovoltaic performance of the devices. a J–V curves, b EQE response, and c transient\nabsorption at donor GSB for rigid devices using different blend films including PzA5:L8BO, UVC,\nand UVC+ASH. d J–V curves for stretchable devices fabricated on TPU substrate. (inset: the\nschematic architecture of stretchable devices). e J–V curves under stretching strains for UVC+ASH\nstretchable devices fabricated on TPU substrate. f Summary of photovoltaic performance and\nstretching strain of 80% PCE retention of IS-OPVs reported in the literature (dashed lines represent\nthe upper limit and the average level of the performances). g Evolution of the 0-0/0-1 absorption\npeak ratio under increasing strains (D and A correspond to the peak of donor and acceptor,\nrespectively) for PzA5:L8BO and UVC+ASH blend films. h Charge extraction by linearly\nincreasing voltage (photo-CELIV) results of PzA5:L8BO and UVC+ASH IS-OPV devices under\nvarying strains. i Dependence of the carrier mobility calculated from photo-CELIV on the stretching\nstrain.\n2. Please provide a clear schematic diagram to illustrate the role of material cross-\nlinking and ASH.\nResponse: In the revised manuscript, we have replaced Fig. 4d with a clearer\nschematic diagram to illustrate the respective roles of molecular cross-linking and ASH\nfillers. The schematic highlights how cross-linking forms covalently tied donor–\nacceptor networks that enhance phase continuity and toughness, while ASH particles\nact as high-modulus fillers, redistribute stress, and promote fibrillar morphology.\nMain text:\na b PzA5:L8BO d PzA5:L8BO\n0\nt y -2\nsi UVC\nn\nt e -4\nn\nI\nUVC\nUVC+ASH\nUVC\n4\n2\n1 2\n-2\nc\nN K-edge at 399.2 eV -4\nPzA5:L8BO UVC+ASH\nUVC\nUVC+ASH\nUVC+ASH\nsit y 2\nn\ne\nnt\n-2\nFiber/domain of L8BO\nPzA5 donor Y5TV\n-4\nPETMP crosslinker\n0.003 0.01 0.03\nASH filler cluster\nQ vector (Å-1)\nFig. 4 | Morphology properties for the photoactive films. a Grazing incidence wide-angle X-ray\nscattering sector-averaged I–q curves in the IP (dashed lines) and OOP (solid lines). b AFM-IR\nmapping of the blend films at the wavenumber of 1426 cm−1. c NK-RSoXS averaged I–q curves for\nPzA5:L8BO, UVC, and UVC+ASH blends. d Schematic of the effect of crosslinker and filler in the\nblend. The purple and cyan translucent strips represent the synergetic interwoven networks in the\nblend film.\n3. According the contact angel data, please provide the Flory-Huggins interaction\nparameter to explain the miscibility between the two materials.\nResponse: Based on the contact angle data with water and diiodomethane\n(Supplementary Fig. 2), we calculated the surface energy components of the neat donor,\nneat acceptor, and their blends using the Owens–Wendt–Rabel–Kaelble (OWRK)\nmethod. By further converting the total surface energies to solubility parameters, the\nFlory–Huggins interaction parameter (χ) between the donor/acceptor and the ASH filler\nwas estimated according to χ =V /RT(δ −δ )2. The obtained χ values are 0.008 for\nAB ref A B AB\nPzA5 and 0.003 for Y5TV, indicating good miscibility with the ASH filler. These\nquantitative results are now included in the Supplementary Information to strengthen\nour discussion of compatibility.\nMain text:\n“The Flory–Huggins interaction parameters (χ) derived from contact angle data\nwere below 0.04 for both donor and acceptor with ASH, further supporting their\nfavourable miscibility and stable dispersion (Supplementary Fig. 4).”\nSupplementary information:\na b\nSi 2p ASH χ of ASH with\nY\nP\n5\nzA\nTV\n5:\n+\nY\nA\n5T\nS\nV\nH\n+ASH o\nn,\nχ 0.03 P\nY5\nzA\nTV\n5\naliz e d I\nnt\ne\nnsit y PzA5+ASH\ngi\nn s i\nnt\ner\na\ncti\n0.02\nm g\nN or - H u 0.01\ny\nor\nFl\n0.00\n95 100 105 110 PzA5 Y5TV\nBinding Energy (eV) Photoactive material\nSupplementary Fig. 4 | Flory–Huggins interaction parameters and XPS analysis\nof ASH blends. a X-ray photoelectron spectroscopy (XPS) of Si 2p for ASH blended\nwith different photoactive materials. b Calculated Flory–Huggins interaction\nparameters (χ) between ASH filler and donor/acceptor components. Contact angles\nwith water and diiodomethane were analyzed using the Owens–Wendt–Rabel–Kaelble\n(OWRK) method to obtain surface energy components. From these, the solubility\nparameters (δ) of the donor (PzA5), acceptor (Y5TV), and ASH filler were derived, and\n𝑉𝑟𝑒𝑓 2\n𝜒 = (δ −δ ) , where 𝑉𝑟𝑒𝑓 is the reference molar volume, R is the gas constant,\n𝑖𝑗 𝑖 𝑗\n𝑅𝑇\nand T is the absolute temperature.\n4. In Fig 2a, the crack-onset strain (COS) of the UVC-treated film with ASH shows a\nslight decrease compared to the UVC-treated film without ASH. This suggests that the\nincorporation of ASH does not enhance the mechanical performance. In addition, what\nis the thickness of these films used to test mechanical properties?\nResponse: The COS shows only a negligible change from 43.8% (UVC) to 42.9%\n(UVC+ASH). It is important to note that toughening is generally defined as the ability\nof a material to resist crack initiation and propagation by improving stress transfer and\nenergy dissipation, which goes beyond COS alone. From this perspective, mechanical\nperformance enhancement can also be evaluated by increased modulus and toughness.\nIn our case, the incorporation of ASH maintains the COS essentially unchanged while\nsubstantially reinforcing the modulus and cyclic energy dissipation, which together\nrepresent a clear toughening reinforcement effect.\nThe thickness of the free-standing films used for mechanical testing was ~100 nm,\nconsistent with that of the device active layers.\nWe also added the following sentence to emphasize that the UVC+ASH blend\nmaintains a comparable COS while markedly reinforcing the modulus and energy\ndissipation, thereby highlighting its toughening reinforcement effect.\nMain text:\n“This combination of nearly unchanged COS with markedly reinforced modulus\nhighlights the role of ASH in providing a clear toughening reinforcement effect.”\n5. Apart from the storge modulus, please provide the other data of DMA, for example:\nloss modulus, and tan δ behavior. And then provided the Tg and sub-Tg based on the\ndifferent sample.\nResponse: We thank the reviewer for this helpful suggestion. The dynamic\nmodulus–strain data originally presented were obtained using a mechanical testing\nsetup equipped with a dynamic loading module, which reports the overall dynamic\nmodulus (equivalent to the storage modulus, E′). This simplified configuration does not\nmeasure the phase lag between stress and strain, and therefore cannot provide loss\nmodulus (E″), tan δ, or T /sub-T values.\ng g\nTo address the reviewer’s request, we have additionally performed standard\ntemperature-sweep DMA measurements. These results reveal that T decreases slightly\ng\nfrom 285 °C to 280 °C upon UVC+ASH incorporation, and the UVC+ASH films also\nexhibit a pronounced rise in both E′ and E″ around 100 °C compared to the pristine\nblend, consistent with additional structural reinforcement introduced by the\ncrosslinker–filler synergy.\nSupplementary information:\na b\n0.025\nPzA5 M\nPa) PzA5: L\nS\n8\nt\nB\no\nO\nra g e m o d u l u s\nUVC+ A\nS\nt\nS\no\nH\nrage modulus 0.25\no d ul\nus\nE' ( M P a)\n0\n0\n.\n.\n0\n0\n1\n2\n5\nSt\nora ge\nm o d ul us ( 100 L T o an ss δ m odulus L T o an ss δ modulus\n0\n0\n.\n.\n1\n2\n5\nm\nD y\nn a mi\nc\n0\n0\n.\n.\n0\n0\n0\n1\n5\n0\nd ul\nus ( M\nPa)10\n0.10\no\nm 0.05\n0.000 L\noss 1\n0.1 0.2 1 2 50 100 150 200 250 300 350\nStrain (%) Temperature (°C)\nSupplementary Fig. 7 | Dynamic modulus for the blend films. a Dynamic modulus\nmeasured with tensile dynamic loading. b Dynamic mechanical analysis (DMA) results\nof PzA5:L8BO and UVC+ASH film under 1 Hz and 0.15% oscillation.\n6. While the authors conducted finite element simulations to demonstrate the\nmechanical properties of the UVC-treated film with ASH, the absence of a reference\nsample (e.g., PzA5:L8BO, UVC-treated film without ASH) limits the interpretability\nof the results. It is recommended that the authors supplement their analysis with\ncomparative data to establish a baseline for evaluation. Additionally, the manuscript\nshould include detailed descriptions of the simulation methodology, including key\ninput parameters and modeling assumptions, to ensure reproducibility and facilitate\nvalidation by readers.\nResponse: We thank the reviewer for this constructive suggestion. Finite element\nanalysis (FEA) cannot realistically simulate the PzA5:L8BO and UVC films without\nASH because of the disparity in scales among the components. Imaging and scattering\nmeasurements reveal that the donor–acceptor phase separation domains in the\nUVC+ASH films are ~70 nm, and even smaller (<30 nm) in the PzA5:L8BO reference.\nIn contrast, the ASH filler particles are orders of magnitude larger. This size mismatch\nprecludes direct modeling of individual donor, acceptor, crosslinked networks, and\nfiller particles within a single FEA framework. Therefore, the bulk heterojunction was\nsimplified as a homogeneous elastic matrix, and the mechanical effect of embedded\nfiller particles was examined.\nAs shown in Fig. 2c, the matrix properties were parameterized using the measured\nmodulus of the UVC elastic film. The embedded particles were assigned either a low\nmodulus of 100 MPa (representing soft elastomeric additives) or a high modulus of\nthese two cases highlights the contrast between blending with soft elastomers and with\nrigid fillers, capturing the essential reinforcement mechanism while maintaining\ncomputational feasibility.\nThe key input parameters, boundary conditions, and assumptions of the FEA are\nnow described in detail in the revised Methods and Supplementary Information to\nensure reproducibility and allow validation by readers.\nMain text:\n(Method part) “To provide an intuitive visualization of the tensile process, finite\nelement models were constructed in Abaqus. The matrix dimensions were set to 30 ×\n40 (units), and spherical fillers with diameters of 8 and 4 (units) were embedded. For\ncomputational efficiency, a linear elastic constitutive model was adopted. The Young’s\nmodulus of the matrix was set to 370 MPa, while that of the fillers was either 100 MPa\n(representing soft elastomeric additives) or 2000 MPa (representing rigid fillers). The\nunit system was kept consistent across the model. This simplification does not affect the\nvalidity of the observed trends, as the purpose of the simulation was to reproduce the\nqualitative behavior of uniaxial tensile experiments. A static analysis step was\nemployed: one side of the model was fully fixed, while a constant tensile force of 10 N\nwas applied along the X-axis at the opposite side. The resulting stress distributions are\npresented in Supplementary Fig. 8.”\nE =100 MPa\nparticle\na b c Stress (MPa)\nP a)40\nP\nU U\nz\nV V\nA\nC C\n5\n+\n:L\nA\nB\nH\nO\nk P a) 7.5\nN\ness (\nM\n1\n2\n3\n0\n0\n0 1\n1\n3\n0\n9\n6\n9\n5\n9\nM\nM\nP\nP\nP\na\na\na\nΔ\n(\nE'0.2\n%-2 %\n2\n5\n.\n.\n5\n0\n1.9 2.0\nE activelayer (matrix) =200 MPa\n0 0.0\n0 10 20 30 40 50 PzA5 PzA5:L8BO UVC UVC+ASH 0\nStrain (%)\nE =2000 MPa\nparticle\nPzA5:L8BO\nd P U z V A C 5 + :L A 8 S B H O( ( σ σ a a p p p p = = 0 0 . . 1 2 5 5 G G σ P P a a , , P L C A ε ε 8 z r S = = o A B H s O 5 2 9 s f l i 5 1 p a i l n l o c e . . k 7 0 c l r y e e % % c m d p l u t e m ) ) o s r t r a e d t r o e n ri o al r s e e str ess ( M P a) 2 3 4 0 0 0 P U U z V V A C C 5 + :L A 8 S B H O C f ra s c t ( k r % - a o i ) n ns 6 e 0 t 40 2 Y 0 ield 0 stra 3 6 9 1 i 2 n (% 30 ) 0 600 90 U U 0 V V 12 C C 0 + m 0 Y A ( o o M S u d H P n u a g lu ) 's s\nnsil\n20\n0\n0 0\nσ\nc y\ncli c t e\n1\n0\n(\ns\nM\nt x r i e m\nP\ns\na\nu s\n)\n6 m 0 40\n20\n0\n1\n4 8\n2\n0 0\n0\n0 0\n10 20 30 40\n(\ns Y\nM\nt i r e e\nP\nl s\na\nd s\n)\n0 5 10\nFig. 2 | Mechanical performance of the blends. a Stress-strain curves of neat and blend films. b\nDynamic modulus of the blend films. c Stress distribution in finite element analysis on a matrix with\nfiller particles of different modulus. d Snapshot of the stopped-extended crack in the UVC+ASH\nblend film (with molecular colours as the legend). e Stress-strain curves at the 100th cycle below\nyielding strain (3.1% for PzA5:L8BO, 8.8% for UVC, and 10.2% for UVC+ASH blend films,\nrespectively). f Summary of various mechanical parameters for blend films.\n7. The mechanical property data obtained from molecular dynamics (MD) simulations\nappear significantly higher than the experimental test results. Could the authors provide\na detailed explanation for this discrepancy? Additional, provide a more comprehensive\ndescription of the MD simulation methodology.\nResponse: We thank the reviewer for this important comment. The mechanical\nproperties obtained from MD simulations are indeed higher than the experimental\nresults. This overestimation arises from the intrinsic nature of the MD cell: the\nsimulated box has an extremely small cross-sectional area and is constructed with\nperiodic boundary conditions in all three directions. As a result, the system has no free\nsurfaces or edges, and therefore lacks edge-induced stress concentration and crack\ninitiation, which are unavoidable in real films. The absence of such failure pathways\nleads to higher apparent modulus and strength in the simulations compared with\nexperimental tests. Explanations for this discrepancy have now been incorporated into\nthe discussion, and the MD methodology has been described in greater detail to ensure\nreproducibility.\nMain text:\n“To investigate fracture dynamics at the molecular level, we performed molecular\ndynamics (MD) simulations under periodic boundary conditions, which suppress edge\neffects and emphasize intrinsic tensile responses, with snapshots taken near the failure\npoint (Fig. 2d, Supplementary Fig. 9–10).”\n(Method part) “The MD stress–strain simulations were performed using a Perl\nscript. Each system was first equilibrated for 200 ps under zero external stress (300 K,\nNPT), followed by the application of successive uniaxial stresses of 0.001, 0.05, 0.10,\n0.15, 0.25, 0.30, 0.35, and 0.40 GPa. At each stress level, the cell was equilibrated for\n100 ps (300 K, NPT) to extract the actual stress and strain values of the cell. Fracture\nonset was identified from the stress response. In the visualization, PzA5 donor,\nL8BO\nacceptor, crosslinked PzA5:Y5TV network, and ASH filler are shown in red, blue, purple,\nand cyan, respectively. Representative snapshots at selected steps were taken to\nillustrate the locations where cracks initiate.”\n8. ASH is an insulating material, and adding it can significantly increase Jsc, which is\nan incomprehensible phenomenon. The authors should explain this result.\nResponse: The increase in J does not originate from any direct contribution of\nSC\nASH itself, since ASH is essentially an inert silica aerogel cluster without photovoltaic\nactivity, electrical conductivity, or doping effects. Rather, the improvement arises from\nits role as a morphology modulator, similar to the function of solvent additives or solid\nadditives commonly used in organic photovoltaics. In bulk heterojunction OPVs,\nefficient exciton dissociation and charge extraction require a delicate balance: domains\nmust be sufficiently pure and continuous to enable high mobility and long carrier\nlifetimes, while remaining small enough to ensure effective exciton harvesting without\nexcessive recombination. Incorporating ASH helps to tune this balance by promoting\nmore ordered molecular packing and more favorable phase separation, while avoiding\nthe formation of trap centers owing to its insulating nature. This results in enhanced\nexciton dissociation and more efficient charge transport, ultimately yielding the\nobserved increase in J . Similar strategies employing high-boiling-point solvent\nSC\nadditives, solid-state additives, or some inorganic clusters have been widely reported in\nthe OPV field to improve device performance through morphology control.\nWe added the following discussion to clarify this point.\nMain text:\n“Such pronounced morphological optimization can be ascribed to the strong\ninterfacial interactions between ASH and the donor/acceptor components, which\nfacilitate more favorable packing and phase organization. This fibrillar morphology\nfacilitates continuous charge percolation pathways, while the finer yet purer domains\npromote efficient exciton diffusion and dissociation, jointly underpinning the observed\nJ enhancement in devices.”\nSC\n9. Compared to the PzA5:L8BO and PzA5:L8BO:Y5TV samples, the hole mobility of\nPzA5:L8BO:Y5TV exhibits a remarkable increase after UVC treatment, while its\nelectron mobility shows a significant decrease (Table S1). Please explain it.\nResponse: As shown in Table S1, the electron and hole mobilities of PzA5:L8BO\n(3.04×10−4 and 2.50×10−4 cm² V⁻¹ s⁻¹) are both lower than those of PzA5:L8BO:Y5TV\n(4.84×10−4 and 5.16×10−4 cm² V⁻¹ s⁻¹). After UV-crosslinking (UVC), both electron\nand hole mobilities decrease markedly (7.20×10−5 and 1.11×10−4 cm² V⁻¹ s⁻¹). By\ncontrast, when ASH is present in the PzA5:L8BO:Y5TV blend during UVC treatment,\nthe mobilities recover to 4.55×10−4 (holes) and 4.92×10−4 (electrons). In summary, the\noverall trend is consistent for both carriers, following UVC < PzA5:L8BO <\nUVC+ASH < PzA5:L8BO:Y5TV.\nThis behavior can be rationalized as follows: during crosslinking in the\nPzA5:L8BO:Y5TV film, the formation of UV-induced covalent networks is\naccompanied by structural disruption and continuous evolution of interfacial domains\nunder high-energy irradiation, which introduces localized degradation and impedes\ncharge transport. When ASH is incorporated in the blend, it mitigates these adverse\neffects by stabilizing domain structures and promoting more favorable molecular\norganization. Consequently, UVC+ASH films retain mobilities comparable to those of\nPzA5:L8BO:Y5TV, while simultaneously establishing a reinforcing network that\nenhances stretchability.\n10. As shown in Table S3, the trends in CCL value and area are inconsistent. Please\nexplain it. Also, the relevant crystal parameters related to (111) and (021) mentioned in\nthe manuscript are missing.\nResponse: The area of a diffraction peak reflects the scattering intensity, which\ndepends on the volume fraction, orientation distribution, and degree of crystallinity of\nthe ordered domains. In contrast, the crystalline coherence length (CCL) is derived from\nthe FWHM of the peak, which quantifies the average size of coherently ordered regions.\nAs these two parameters describe different aspects of molecular ordering, their trends\nare not necessarily correlated, and variations in peak area do not directly translate into\nchanges in CCL.\nBesides, the (11̅1) and (021) parameters are now supplemented.\nSupplementary information:\nSupplementary Table S3 | Fitting results for π-π and lamellar stacking peaks of the\nneat and blend films.\nLocation FWHM d-spacing CCL\nSample Sector Type Reflection Area\n(Å−1) (Å−1) (Å) (Å)\nPzA5 OOP Lorentz 010 1.76 215.36 0.28 1.79 20.27\nIP Lorentz 100 0.29 220.33 0.08 10.66 72.57\nOOP Lorentz 010 1.34 128.84 0.73 2.35 7.66\nL8BO\nIP Lorentz 100 0.42 110.50 0.16 7.41 35.03\nOOP Lorentz 010 1.30 88.42 0.50 2.41 11.21\nY5TV\nIP Lorentz 100 0.34 88.49 0.19 9.18 29.39\nOOP Lorentz 010 1.75 956.80 0.40 1.79 13.94\nIP Lorentz 100 0.32 479.89 0.14 9.77 40.86\nPzA5:L8BO\noff-axis Lorentz 021 0.49 87.66 0.25 6.36 22.20\nIP Lorentz 11̅1 0.44 138.62 0.12 7.19 45.03\nOOP Lorentz 010 1.73 993.18 0.43 1.82 13.08\nIP Lorentz 100 0.32 510.53 0.12 9.91 48.46\nUVC\noff-axis Lorentz 021 0.46 217.66 0.30 6.85 18.93\nIP Lorentz 11̅1 0.42 127.35 0.19 7.40 29.42\nOOP Lorentz 010 1.75 1082.21 0.38 1.79 14.53\nIP Lorentz 100 0.31 348.15 0.10 10.10 54.24\nUVC+ASH\noff-axis Lorentz 021 0.56 132.11 0.22 5.61 25.99\nIP Lorentz 11̅1 0.44 222.31 0.12 7.17 46.97\n11. Please provide the following key information: 1) specific performance parameters\n(e.g., PCE, FF, JSC, stretchability stability, etc.) and detailed fabrication processes of\nthe stretchable organic photovoltaic module; 2) operational parameters of the EMG\nsensor in the integrated device (including required power and voltage range); 3) the\npower supply source for the wireless signal transmission module (e.g., photovoltaic\nmodule or other power sources). The technical details of each module should be\nclarified to ensure the integrity and reproducibility of the study.\nResponse: The key information is summarized as follows:\n1) Stretchable OPV device and module performance: For a 1.5 cm2 device, the\nphotovoltaic parameters are V = 0.887 V, J = 24.89 mA cm−2, FF = 64.58%, and\nOC SC\nPCE = 14.25%. Under 40% tensile strain, the corresponding values are V = 0.886\nOC\nV, J = 23.83 mA cm−2, FF = 54.03%, and PCE = 11.40%. For integration, we\nSC\narranged three 1.5 cm2 units in series, as illustrated in the updated Supplementary\nFig. S14. The resulting effective active area was 4.5 cm2, yielding a theoretical\nmaximum power of ~65 mW. Due to interconnection and power management losses,\nthe realized efficiency is lower, and we plan to develop monolithic modules to\nminimize these losses in future work.\n2) EMG sensor operational parameters: The wireless Arduino UNO–based\nelectromyography (EMG) monitoring prototype was powered by a lithium-polymer\nbattery (nominal voltage 3.7 V, effective capacity ~0.70 Wh). The average load\npower of the EMG system was ~360 mW, including ~100 mW for biosignal\nacquisition and processing and ~260 mW for wireless data transmission. These\nvalues were extracted from the measured discharge time (1.96 h without solar input).\n3) Power supply source for wireless transmission: Three IS-OPV devices were\nconnected through a 330-type boost converter and interfaced with an SPV1050\nsolar power management module, which supplied the EMG charging port. In this\nconfiguration, the IS-OPV module provided continuous output under illumination,\nwhich partially recharged the battery and offset the system load, extending the\noperation time from 1.96 h to 2.30 h. For scaling estimation, an active area\nexceeding 30 cm² could deliver >430 mW under sunlight, approaching the\ncontinuous operating power of the EMG system.\nThese details have now been incorporated into the revised manuscript to ensure\nreproducibility.\nMain text:\n(Method part) “The output power of the IS-OPV module was calculated using the\nstandard relation:\n𝑃 = 𝐼 ×𝐴 ×𝑃𝐶𝐸\n𝑂𝑃𝑉 𝑙𝑖𝑔ℎ𝑡 𝑒𝑓𝑓\nwhere 𝐼 100 mW cm−2 (AM 1.5G, 1 sun), 𝐴 4.5 cm2 (three devices in\n𝑙𝑖𝑔ℎ𝑡 𝑒𝑓𝑓\nseries, each 1.5 cm2), and PCE 14.4%. This corresponds to a theoretical maximum\noutput power of ~65 mW. For the wearable demonstration, a wireless Canum\nVenaticorum Arduino UNO-based electromyography (EMG) monitoring prototype was\npowered by a lithium-polymer battery (nominal voltage 3.7 V, effective capacity ~0.70\nWh), ensuring stable device operation. The average load power of the system was\nestimated from the measured discharge time (1.96 h) to be ~360 mW, with ~100 mW\nattributed to biosignal acquisition and processing and ~260 mW consumed by wireless\ndata transmission. Three large-area IS-OPV devices were connected through a 330-\ntype boost converter and interfaced with an SPV1050 solar power management module,\nsupplying the EMG charging port to partially offset power consumption or recharge\nthe battery. The effective contribution of photovoltaic charging was quantified by\ncomparing discharge durations with and without solar input, with the improvement\nattributed to the continuous ~65 mW output from the IS-OPV module under 1 sun\nillumination. For scaling estimation, the same formula was applied to larger module\nareas, demonstrating that >30 cm² of active area could generate >430 mW under full\nsunlight in the future, approaching the continuous operating power of the EMG system.”\n12. There are some errors in the manuscript. Please revise them. For example, in Fig\nS17a, the tensile strain of the PzA5:L8BO blend film should 0 to 2.5% not 0 to 40%.\nResponse: We have corrected the mistaken in Fig S17 (now Supplementary Fig.\n19). We also reviewed all figures for consistency in axis ranges, units, and panel labels\nand corrected minor typos accordingly.\na b 2.5\n2.0 2.0 ε=0\nε=0\nn n\no o\npti\n1.5\npti1.5\nor or\nb s b s\nA A1.0\nε=2.5%\n1.0 ε=40%\n0.5\n500 600 700 800 900 500 600 700 800 900\nWavelength (nm) Wavelength (nm)\nSupplementary Fig. 19 | In-situ absorption spectra during fully free-standing\nstretching test. a Absorption spectra of the PzA5:L8BO blend film from 0 to 2.5%\ntensile strain. b Absorption spectra of the UVC+ASH blend film from 0 to 40%\ntensile strain.\nRespond to the reviewers’ comments\nReviewer 1\nI'm pleased to see that the authors addressed all the comments raised by the reviewer\nappropriately. This paper is now ready for publication.\nResponse: We thank the reviewer for the positive feedback and the encouraging\nrecommendation for publication. We truly appreciate the reviewer’s insightful\ncomments and constructive suggestions throughout the review process, which helped\nus strengthen the manuscript substantially.\nReviewer 2\nThe manuscript has improved significantly. However, a few critical issues still need to\nbe resolved before it can be accepted.\n1. Please provide a diagram illustrating the cross-linking reaction mechanism for Y5TV,\nPETMP and ASH, rather than the morphological schematic shown in Figure 4d.\nResponse: We fully agree with the reviewer that a schematic diagram illustrating\nthe cross-linking reaction mechanism would be helpful for readers to understand the\nchemical process. Accordingly, we have added Supplementary Scheme 2, which depicts\nthe “click” cross-linking mechanism among the vinyl groups within the donor (PzA5)\nand acceptor (Y5TV) molecules via the thiol groups of the PETMP crosslinker, forming\nan intermolecularly linked network. ASH, as an inert filler, does not participate in the\nchemical cross-linking reaction. Its role arises mainly from non-covalent interactions\n(Fig. 1f), which allow ASH to function as both a mechanical reinforcing filler and a\nmorphology modulator.\nTo clarify this, we added the following sentence in the main text and\nSupplementary Scheme 2.\nMain text:\n“These alkene moieties enable efficient thiol-ene “click” crosslinking. We employed\npentaerythritol tetra(3-mercaptopropionate) (PETMP) as the tetra-thiol crosslinker\n(Supplementary Scheme 2).”\na\nPETMP bridge\nb\nPzA5 Y5TV L8BO\nSupplementary Scheme 2 | Cross-linking reaction mechanism. a General thiol-ene\n“click” reaction of PETMP: thiol groups add across vinyl group to form saturated\nthioether linkages under UV initiation. b Resulting intermolecular cross-linked network\nformed between donors (PzA5, red) and acceptors (Y5TV, purple) via PETMP bridges.\n2. Why does the introduction of ASH lead to a slight decrease in COS while significantly\nincreasing the toughness, and what is the underlying mechanism?\nResponse: The slight decrease in crack-onset strain (COS) but marked increase in\ntoughness with ASH arises because crack initiation and pre-fracture energy dissipation\nare controlled by different processes.\ninterfacial heterogeneities that act as weak stress concentrators, promoting earlier strain\nlocalization and thus a slightly lower COS. However, before catastrophic fracture, a\nfiller could enable multiple pre-fracture dissipation mechanisms, including enhanced\nviscoelastic hysteresis, matrix shear yielding/plastic flow, micro-voiding/crazing, and\ninterfacial debonding/pull-out, stabilizing stress redistribution in the cross-linked\nnetwork. These processes raise the dissipation before fracture, thereby substantially\nincreasing toughness even though the COS is slightly reduced. To clarify this, we added\nthe following sentence in the manuscript.\nMain text:\n“The further introduction of ASH resulted in a slightly reduced COS of 43%, likely due\nto local stiffness contrast from the rigid filler that promotes early strain localization.\nNevertheless, interfacial debonding and viscoelastic dissipation around the filler enable\nefficient energy absorption, resulting in markedly enhanced toughness and modulus\n(1090 MPa for UVC+ASH).”\n3. Please provide the standard deviation of the mechanical properties for the four\nmaterials in Figure 2, along with the corresponding curves.\nResponse: We included three independent tensile curves for each sample in Fig.\n2a, showing good reproducibility across all measurements. The corresponding standard\ndeviations of the mechanical parameters (modulus and crack-onset strain) are\nsummarized in Supplementary Table 1. These additions confirm the statistical\nreliability of the mechanical data. To clarify this, we revised following sentence in the\nmanuscript.\nMain text:\n“The mechanical enhancement conferred by the MCFT strategy was first assessed\nthrough uniaxial tensile testing of free-standing thin films (Fig. 2a), showing consistent\n3\nresults across three independent specimens for each sample, with the corresponding\ntensile parameters and standard deviations summarized in Supplementary Table 1.”\nE =100 MPa\nparticle\na b c Stress (MPa)\nP a)4\n5\n0\nU U\nz z\nV V\nA A\nC C\n5 5\n+\n:L\nA\nB\nH\nO\nk P a)\n10\n7.\n.\n5\n0 9.5 N\n01\n0.5\nStr e s s ( M 2 3 0 0 3 1 9 0 2 90 M M Pa Pa ( E'0.2 %-2 % 5.0 E activelayer (matrix) =370 MPa\n10 229 MPa Δ 2.5 1.9 2.0\n0\n0.0\n0.9 N 01\n0 10 20 30 40 50 PzA5 PzA5:L8BO UVC UVC+ASH 0\nStrain (%) E =2000 MPa\nparticle\nPzA5:L8BO\nd P U z V A C 5 + :L A 8 S B H O( ( σ σ a a p p p p = = 0 0 . . 1 2 5 5 G G σ P P a a , , P L C A ε ε 8 z r S = = o A B H s O 5 2 9 s f l i 5 1 p a i l n l o c e . . k 7 0 c l r y e e % % c m d p l u t e m ) ) o s r t r a e d t r o e n ri o al r s e e str ess ( M P a) 2 3 4 0 0 0 P U U z V V A C C 5 + :L A 8 S B H O C f ra s c t ( k r % - a o i ) n ns 6 e 0 t 40 2 Y 0 ield 0 stra 3 6 9 1 i 2 n (% 30 ) 0 600 90 U U 0 V V 12 C C 0 + m 0 Y A ( o o M S u d H P n u a g lu ) 's s\nnsil\n20\n0\n0 0\nσ\nc y\ncli c t e\n1\n0\n(\ns\nM\nt x r i e m\nP\ns\na\nu s\n)\n6 m 0 40\n1\n4 8\n2\n0 0\n0\n10 20 30 40\n(\ns Y\nM\nt i r e e\nP\nl s\na\nd s\n)\n0 5 10\nFig. 2 | Mechanical performance of the blends. a Stress-strain curves of neat and blend films\n(three independent specimens per blend, with the maximum modulus marked). b Dynamic modulus\nof the blend films. c Stress distribution in finite element analysis on a matrix with\nfiller particles of different modulus. d Snapshot of the stopped-extended crack in the UVC+ASH\nblend film (with molecular colours as the legend). e Stress-strain curves at the 100th cycle below\nyielding strain (3.1% for PzA5:L8BO, 8.8% for UVC, and 10.2% for UVC+ASH blend films,\nrespectively). f Summary of various mechanical parameters for blend films.\nSupplementary information:\nSupplementary Table 1 | Summary of mechanical parameters for PzA5,\nPzA5:L8BO, UVC, and UVC+ASH films.\nCrack-onset strain Young's modulus\nfilms SD a SD a\n(%) (MPa)\nPzA5 23.5 12.9 14.7 4.6 195.0 170.0 229.4 24.4\nPzA5:L8BO 4.0 4.5 2.7 0.8 203.0 103.0 133.4 41.9\nUVC 43.8 40.4 35.5 3.4 369.3 392.0 286.2 45.5\nUVC+ASH 42.2 39.3 33.7 3.5 1090.1 761.8 725.3 164.0\na standard variation based on 3 tensile tests.\n4\n4. The authors could employ femtosecond transient absorption spectroscopy to explore\nhow ASH incorporation influences carrier dynamics.\nResponse: ASH improves carrier dynamics, as it accelerates interfacial hole\ntransfer and speeds the build-up of charge-separated populations, thereby mitigating\ntrap-related losses.\nTo quantify this, we performed femtosecond transient absorption on the donor GSB\nand fitted the growth with a biexponential model (zero at 0.2 ps). The rise lifetimes (τ ,\n1\nτ ) are 1.18 ps, 8.88 ps for PzA5:L8BO, 2.82 ps, 27.77 ps for UVC, and 0.63 ps, 12.89\n2\nps for UVC+ASH. Crosslinking slows both components (consistent with UV-induced\ntraps), whereas ASH mitigates this degradation and accelerates the dynamics: τ\n1\nshortens to 0.63 ps (faster hole-transfer build-up) and τ drops from 27.77 ps (UVC) to\n2\n12.89 ps, indicating quicker formation of long-lived charge-separated populations.\nThese kinetic gains align with the enhanced carrier collection and device metrics of the\nUVC+ASH films (Supplementary Fig. 13–15).\nMain text:\n“To assess the effect of ASH on interfacial charge transfer, we analyzed the donor\nground-state bleach (GSB) growth using a biexponential model, referenced at 0.2 ps.\nThe extracted rise lifetimes (τ , τ ) are 1.18 ps, 8.88 ps (PzA5:L8BO), 2.82 ps, 27.77 ps\n1 2\n(UVC), and 0.63 ps, 12.89 ps (UVC+ASH). Crosslinking slows both components,\nconsistent with UV-induced traps, whereas adding ASH mitigates this degradation,\nrestoring and accelerating the dynamics: τ shortens to 0.63 ps (faster hole-transfer\n1\nbuild-up) and τ decreases from 27.77 ps (UVC) to 12.89 ps, indicating a quicker\n2\nestablishment of the charge-separated population. These kinetic improvements align\nwith the enhanced carrier collection and device metrics of the UVC+ASH films\n(Supplementary Fig. 13–15).”\n5\na ΔA (mOD)\n3.31\nPzA5:L8BO\nGSB CT\nESA\n100\nps) −1.24\ne\n(\nm\nTi 10 −3.51\n1 −5.78\nAcceptor\nGSB Acceptor\nPIA\n0.1 −8.05\n600 700 800 900\nWavelength (nm)\nΔA (mOD)\nc ΔA (mOD) d\n3.31\n3.31 UVC+ASH\nUVC\n1000\nESA 100\ne\n(\nps) 100 −1.24\nm\ne\n(\nps) −1.24\nTi m 10 −3.51 Ti 10 −3.51\n1 −5.78 1 Acceptor −5.78\nAc G ce S p B tor Acceptor GSB Acc P e IA ptor\nPIA 0.1 −8.05\n0.1 −8.05 600 700 800 900\n600 700 800 900\nWavelength (nm)\nWavelength (nm)\nSupplementary Fig. 13 | Contour images of transient absorption spectra for blend\nfilms. Blend films of PzA5:L8BO, UVC, and UVC+ASH were measured under\nphotoexcitation at 780 nm.\nSupplementary Fig. 14 | Transient absorption spectra of PzA5:L8BO, UVC, and\nUVC+ASH blend films at varying delay times.\n6\n1.0 donor GSB 1.0 Acceptor exciton\nPzA5:L8BO (hole transfer channel)\nD UVC D 0.8 PzA5:L8BO\nO0.8 UVC+ASH O UVC\nd\nΔ ExpDecay2 Fit\nd\nΔ 0.6 UVC+ASH\ne0.6 e\nz z 0.4\nm\nali\n0.4 1. τ 1 1 8 8. τ 8 2 8 m\nali\n0.2 0. τ 3 1 5 19 τ . 2 38\no r 2.82 27.77 o r 0.25 18.18\nN0.2 0.63 12.89 N 0.0 0.19 11.90\n0.0 −0.2\n1 10 100 1000 0.1 1 10 100 1000\nDelay time (ps) Delay time (ps)\nSupplementary Fig. 15 | Summarized signals and fitted results from the transient\nabsorption.\n5. As shown in Table S1, the electron and hole mobilities of PzA5:L8BO (3.04×10-4 and\n2.50×10-4 cm2 V-1 s-1) are both lower than those of PzA5:L8BO:Y5TV (4.84×10-4 and\n5.16×10-4 cm2 V-1 s-1). After UV-crosslinking (UVC), both electron and hole mobilities\ndecrease markedly (7.20×10-5 and 1.11×10-4 cm2 V-1 s-1). Nevertheless, the current\ndensity and PCE increase. Typically, low mobilities result from severe recombination\ndue to poor morphology. The authors are requested to explain this phenomenon.\nResponse: The apparent incosistency higher J despite lower SCLC mobilities\nSC\narise from the difference beween dark single carrier measurements and actual change-\ntransport regime under illumination. The SCLC mobilities are extracted from injection-\nlimited diodes under dark condition, which are particularly sensitive to traps introduced\nby UV crosslinking.\nTo examine the recombination behavior under operating conditions, we analyzed\nthe light-intensity (P ) dependence of open-circuit voltage (V ) and short-circuit\nlight OC\ncurrent density (J ), which were described by V ∝𝑛kT/q×ln(P ) and J ∝P α,\nSC OC lignt SC lignt\nwhere k is the Boltzmann constant, T is the absolute temperature, and q is the\nelementary charge. The UVC film shows a slightly higher 𝑛 value of 1.36 for P -\nlignt\ndependent V , indicating increased trap-assisted recombination. However, the P -\nOC lignt\ndependent J reveal that α increases from 0.886 to 0.956, signifying reduced\nSC\nbimolecular recombination, which is likely the major origin of an overall increase in\nJ despite the moderate reduction in the dark SCLC mobility. To clarify this, we\nSC\nrevised the following analysis in the manuscript.\nMain text:\n“The crosslinked UVC device exhibited a similar PCE of 16.7% but with an increased\nJ and a decreased V . The lower V is consistent with the presence of traps\nSC OC OC\nintroduced by UV crosslinking, resulting in higher non-radiative recombination and\nlower SCLC mobilities (Supplementary Fig. 12 a-b and Supplementary Table 2). This\nis reflected by the light-intensity dependences of V and J (Supplementary Fig. 12\nOC SC\nc-d): 𝑛 increases from 1.33 to 1.36, indicating higher trap-assisted recombination in\nUVC, whereas 𝛼 increases from 0.886 to 0.956, indicating improved charge collection\nwith reduced bimolecular recombination.”\nSupporting information:\n8\na b\n1.0\nHole-only device Electron-only device\nPzA5:L8BO 1.4 PzA5:L8BO\nUVC UVC\n- 1 ) m0.8 U Li V ne C a + r A fi S t H - 1 ) m U Li V ne C a + r A fi S t H\n2 c 2 c 1.2\n1 / 1 /\nA A\n2\n(\n0.6 2\n(\n1 / J 1 / J 1.0\n0.4 0.8\n1.5 2.0 2.5 1.5 2.0 2.5 3.0\nVappl (V)\n2\n) Vappl (V)\n-m\nc dc\nA\nV )0.90 PzA5:L8BO n=1.17 m PzA5:L8BO\ne\n( UVC\ny\n(\nolt a\ng\n0.87 L\nU\ni\nV\nne\nC\na\n+\nr\nA\nf\nS\nit\nH\ne n\nsit\nL\nU\ni\nV\nne\nC\na\n+\nr\nA\nf\nS\nit\nH\ne n\nsit 0.84 n=1.33\ne\nnt d 10 α=0.989\nc r α=0.956\nr r\nci u\ne n -0.81 n=1.36 uit c 5 α=0.886\np c\nO r\n0.78 -\nci\n10 100 rt 10 100\no\nh\nLight intensity (mW cm-2) S Light intensity (mW cm-2)\nSupplementary Fig. 12 | Charge transport and light-intensity dependences of the\nblends. a,b The J1/2–V curves for holy-only and electron-only devices using\nPzA5:L8BO, UVC, and UVC+ASH blends. c,d The light-intensity (P ) dependence\nlight\nof open-circuit voltage (V ) and short-circuit current density (J ) for organic solar\nOC SC\ncells based on PzA5:L8BO, UVC, and UVC+ASH with the α and n factors marked.\nThe P dependence of V and J are described by V ∝  𝑛kT/q×ln(P ) and\nlight OC SC OC lignt\nJ ∝P α, where k is the Boltzmann constant, T is the absolute temperature, and q is\nSC lignt\nthe elementary charge.\n9"}, {"round": 2, "reviewer_comments": {"reviewer_1": "I'm pleased to see that the authors addressed all the comments raised by the reviewer appropriately. This paper is now ready for publication.", "reviewer_2": "The manuscript has improved significantly. However, a few critical issues still need to be resolved before it can be accepted.\n1. Please provide a diagram illustrating the cross-linking reaction mechanism for Y5TV, PETMP and ASH, rather than the morphological schematic shown in Figure 4d.\n2. Why does the introduction of ASH lead to a slight decrease in COS while significantly increasing the toughness, and what is the underlying mechanism?\n3. Please provide the standard deviation of the mechanical properties for the four materials in Figure 2, along with the corresponding curves.\n4. The authors could employ femtosecond transient absorption spectroscopy to explore how ASH incorporation influences carrier dynamics.\n5. As shown in Table S1, the electron and hole mobilities of PzA5:L8BO (3.04×10-4 and 2.50×10-4 cm² V-1 s-1) are both lower than those of PzA5:L8BO:Y5TV (4.84×10-4 and 5.16×10-4 cm2 V-1 s-1). After UV-crosslinking (UVC), both electron and hole mobilities decrease markedly (7.20×10-5 and 1.11×10-4 cm² V-1 s-1). Nevertheless, the current density and PCE increase. Typically, low mobilities result from severe recombination due to poor morphology. The authors are requested to explain this phenomenon."}, "author_response": "Respond to the reviewers’ comments\nReviewer 1\nI'm pleased to see that the authors addressed all the comments raised by the reviewer\nappropriately. This paper is now ready for publication.\nResponse: We thank the reviewer for the positive feedback and the encouraging\nrecommendation for publication. We truly appreciate the reviewer’s insightful\ncomments and constructive suggestions throughout the review process, which helped\nus strengthen the manuscript substantially.\nReviewer 2\nThe manuscript has improved significantly. However, a few critical issues still need to\nbe resolved before it can be accepted.\n1. Please provide a diagram illustrating the cross-linking reaction mechanism for Y5TV,\nPETMP and ASH, rather than the morphological schematic shown in Figure 4d.\nResponse: We fully agree with the reviewer that a schematic diagram illustrating\nthe cross-linking reaction mechanism would be helpful for readers to understand the\nchemical process. Accordingly, we have added Supplementary Scheme 2, which depicts\nthe “click” cross-linking mechanism among the vinyl groups within the donor (PzA5)\nand acceptor (Y5TV) molecules via the thiol groups of the PETMP crosslinker, forming\nan intermolecularly linked network. ASH, as an inert filler, does not participate in the\nchemical cross-linking reaction. Its role arises mainly from non-covalent interactions\n(Fig. 1f), which allow ASH to function as both a mechanical reinforcing filler and a\nmorphology modulator.\nTo clarify this, we added the following sentence in the main text and\nSupplementary Scheme 2.\nMain text:\n“These alkene moieties enable efficient thiol-ene “click” crosslinking. We employed\npentaerythritol tetra(3-mercaptopropionate) (PETMP) as the tetra-thiol crosslinker\n(Supplementary Scheme 2).”\na\nPETMP bridge\nb\nPzA5 Y5TV L8BO\nSupplementary Scheme 2 | Cross-linking reaction mechanism. a General thiol-ene\n“click” reaction of PETMP: thiol groups add across vinyl group to form saturated\nthioether linkages under UV initiation. b Resulting intermolecular cross-linked network\nformed between donors (PzA5, red) and acceptors (Y5TV, purple) via PETMP bridges.\n2. Why does the introduction of ASH lead to a slight decrease in COS while significantly\nincreasing the toughness, and what is the underlying mechanism?\nResponse: The slight decrease in crack-onset strain (COS) but marked increase in\ntoughness with ASH arises because crack initiation and pre-fracture energy dissipation\nare controlled by different processes.\ninterfacial heterogeneities that act as weak stress concentrators, promoting earlier strain\nlocalization and thus a slightly lower COS. However, before catastrophic fracture, a\nfiller could enable multiple pre-fracture dissipation mechanisms, including enhanced\nviscoelastic hysteresis, matrix shear yielding/plastic flow, micro-voiding/crazing, and\ninterfacial debonding/pull-out, stabilizing stress redistribution in the cross-linked\nnetwork. These processes raise the dissipation before fracture, thereby substantially\nincreasing toughness even though the COS is slightly reduced. To clarify this, we added\nthe following sentence in the manuscript.\nMain text:\n“The further introduction of ASH resulted in a slightly reduced COS of 43%, likely due\nto local stiffness contrast from the rigid filler that promotes early strain localization.\nNevertheless, interfacial debonding and viscoelastic dissipation around the filler enable\nefficient energy absorption, resulting in markedly enhanced toughness and modulus\n(1090 MPa for UVC+ASH).”\n3. Please provide the standard deviation of the mechanical properties for the four\nmaterials in Figure 2, along with the corresponding curves.\nResponse: We included three independent tensile curves for each sample in Fig.\n2a, showing good reproducibility across all measurements. The corresponding standard\ndeviations of the mechanical parameters (modulus and crack-onset strain) are\nsummarized in Supplementary Table 1. These additions confirm the statistical\nreliability of the mechanical data. To clarify this, we revised following sentence in the\nmanuscript.\nMain text:\n“The mechanical enhancement conferred by the MCFT strategy was first assessed\nthrough uniaxial tensile testing of free-standing thin films (Fig. 2a), showing consistent\n3\nresults across three independent specimens for each sample, with the corresponding\ntensile parameters and standard deviations summarized in Supplementary Table 1.”\nE =100 MPa\nparticle\na b c Stress (MPa)\nP a)4\n5\n0\nU U\nz z\nV V\nA A\nC C\n5 5\n+\n:L\nA\nB\nH\nO\nk P a)\n10\n7.\n.\n5\n0 9.5 N\n01\n0.5\nStr e s s ( M 2 3 0 0 3 1 9 0 2 90 M M Pa Pa ( E'0.2 %-2 % 5.0 E activelayer (matrix) =370 MPa\n10 229 MPa Δ 2.5 1.9 2.0\n0\n0.0\n0.9 N 01\n0 10 20 30 40 50 PzA5 PzA5:L8BO UVC UVC+ASH 0\nStrain (%) E =2000 MPa\nparticle\nPzA5:L8BO\nd P U z V A C 5 + :L A 8 S B H O( ( σ σ a a p p p p = = 0 0 . . 1 2 5 5 G G σ P P a a , , P L C A ε ε 8 z r S = = o A B H s O 5 2 9 s f l i 5 1 p a i l n l o c e . . k 7 0 c l r y e e % % c m d p l u t e m ) ) o s r t r a e d t r o e n ri o al r s e e str ess ( M P a) 2 3 4 0 0 0 P U U z V V A C C 5 + :L A 8 S B H O C f ra s c t ( k r % - a o i ) n ns 6 e 0 t 40 2 Y 0 ield 0 stra 3 6 9 1 i 2 n (% 30 ) 0 600 90 U U 0 V V 12 C C 0 + m 0 Y A ( o o M S u d H P n u a g lu ) 's s\nnsil\n20\n0\n0 0\nσ\nc y\ncli c t e\n1\n0\n(\ns\nM\nt x r i e m\nP\ns\na\nu s\n)\n6 m 0 40\n1\n4 8\n2\n0 0\n0\n10 20 30 40\n(\ns Y\nM\nt i r e e\nP\nl s\na\nd s\n)\n0 5 10\nFig. 2 | Mechanical performance of the blends. a Stress-strain curves of neat and blend films\n(three independent specimens per blend, with the maximum modulus marked). b Dynamic modulus\nof the blend films. c Stress distribution in finite element analysis on a matrix with\nfiller particles of different modulus. d Snapshot of the stopped-extended crack in the UVC+ASH\nblend film (with molecular colours as the legend). e Stress-strain curves at the 100th cycle below\nyielding strain (3.1% for PzA5:L8BO, 8.8% for UVC, and 10.2% for UVC+ASH blend films,\nrespectively). f Summary of various mechanical parameters for blend films.\nSupplementary information:\nSupplementary Table 1 | Summary of mechanical parameters for PzA5,\nPzA5:L8BO, UVC, and UVC+ASH films.\nCrack-onset strain Young's modulus\nfilms SD a SD a\n(%) (MPa)\nPzA5 23.5 12.9 14.7 4.6 195.0 170.0 229.4 24.4\nPzA5:L8BO 4.0 4.5 2.7 0.8 203.0 103.0 133.4 41.9\nUVC 43.8 40.4 35.5 3.4 369.3 392.0 286.2 45.5\nUVC+ASH 42.2 39.3 33.7 3.5 1090.1 761.8 725.3 164.0\na standard variation based on 3 tensile tests.\n4\n4. The authors could employ femtosecond transient absorption spectroscopy to explore\nhow ASH incorporation influences carrier dynamics.\nResponse: ASH improves carrier dynamics, as it accelerates interfacial hole\ntransfer and speeds the build-up of charge-separated populations, thereby mitigating\ntrap-related losses.\nTo quantify this, we performed femtosecond transient absorption on the donor GSB\nand fitted the growth with a biexponential model (zero at 0.2 ps). The rise lifetimes (τ ,\n1\nτ ) are 1.18 ps, 8.88 ps for PzA5:L8BO, 2.82 ps, 27.77 ps for UVC, and 0.63 ps, 12.89\n2\nps for UVC+ASH. Crosslinking slows both components (consistent with UV-induced\ntraps), whereas ASH mitigates this degradation and accelerates the dynamics: τ\n1\nshortens to 0.63 ps (faster hole-transfer build-up) and τ drops from 27.77 ps (UVC) to\n2\n12.89 ps, indicating quicker formation of long-lived charge-separated populations.\nThese kinetic gains align with the enhanced carrier collection and device metrics of the\nUVC+ASH films (Supplementary Fig. 13–15).\nMain text:\n“To assess the effect of ASH on interfacial charge transfer, we analyzed the donor\nground-state bleach (GSB) growth using a biexponential model, referenced at 0.2 ps.\nThe extracted rise lifetimes (τ , τ ) are 1.18 ps, 8.88 ps (PzA5:L8BO), 2.82 ps, 27.77 ps\n1 2\n(UVC), and 0.63 ps, 12.89 ps (UVC+ASH). Crosslinking slows both components,\nconsistent with UV-induced traps, whereas adding ASH mitigates this degradation,\nrestoring and accelerating the dynamics: τ shortens to 0.63 ps (faster hole-transfer\n1\nbuild-up) and τ decreases from 27.77 ps (UVC) to 12.89 ps, indicating a quicker\n2\nestablishment of the charge-separated population. These kinetic improvements align\nwith the enhanced carrier collection and device metrics of the UVC+ASH films\n(Supplementary Fig. 13–15).”\n5\na ΔA (mOD)\n3.31\nPzA5:L8BO\nGSB CT\nESA\n100\nps) −1.24\ne\n(\nm\nTi 10 −3.51\n1 −5.78\nAcceptor\nGSB Acceptor\nPIA\n0.1 −8.05\n600 700 800 900\nWavelength (nm)\nΔA (mOD)\nc ΔA (mOD) d\n3.31\n3.31 UVC+ASH\nUVC\n1000\nESA 100\ne\n(\nps) 100 −1.24\nm\ne\n(\nps) −1.24\nTi m 10 −3.51 Ti 10 −3.51\n1 −5.78 1 Acceptor −5.78\nAc G ce S p B tor Acceptor GSB Acc P e IA ptor\nPIA 0.1 −8.05\n0.1 −8.05 600 700 800 900\n600 700 800 900\nWavelength (nm)\nWavelength (nm)\nSupplementary Fig. 13 | Contour images of transient absorption spectra for blend\nfilms. Blend films of PzA5:L8BO, UVC, and UVC+ASH were measured under\nphotoexcitation at 780 nm.\nSupplementary Fig. 14 | Transient absorption spectra of PzA5:L8BO, UVC, and\nUVC+ASH blend films at varying delay times.\n6\n1.0 donor GSB 1.0 Acceptor exciton\nPzA5:L8BO (hole transfer channel)\nD UVC D 0.8 PzA5:L8BO\nO0.8 UVC+ASH O UVC\nd\nΔ ExpDecay2 Fit\nd\nΔ 0.6 UVC+ASH\ne0.6 e\nz z 0.4\nm\nali\n0.4 1. τ 1 1 8 8. τ 8 2 8 m\nali\n0.2 0. τ 3 1 5 19 τ . 2 38\no r 2.82 27.77 o r 0.25 18.18\nN0.2 0.63 12.89 N 0.0 0.19 11.90\n0.0 −0.2\n1 10 100 1000 0.1 1 10 100 1000\nDelay time (ps) Delay time (ps)\nSupplementary Fig. 15 | Summarized signals and fitted results from the transient\nabsorption.\n5. As shown in Table S1, the electron and hole mobilities of PzA5:L8BO (3.04×10-4 and\n2.50×10-4 cm2 V-1 s-1) are both lower than those of PzA5:L8BO:Y5TV (4.84×10-4 and\n5.16×10-4 cm2 V-1 s-1). After UV-crosslinking (UVC), both electron and hole mobilities\ndecrease markedly (7.20×10-5 and 1.11×10-4 cm2 V-1 s-1). Nevertheless, the current\ndensity and PCE increase. Typically, low mobilities result from severe recombination\ndue to poor morphology. The authors are requested to explain this phenomenon.\nResponse: The apparent incosistency higher J despite lower SCLC mobilities\nSC\narise from the difference beween dark single carrier measurements and actual change-\ntransport regime under illumination. The SCLC mobilities are extracted from injection-\nlimited diodes under dark condition, which are particularly sensitive to traps introduced\nby UV crosslinking.\nTo examine the recombination behavior under operating conditions, we analyzed\nthe light-intensity (P ) dependence of open-circuit voltage (V ) and short-circuit\nlight OC\ncurrent density (J ), which were described by V ∝𝑛kT/q×ln(P ) and J ∝P α,\nSC OC lignt SC lignt\nwhere k is the Boltzmann constant, T is the absolute temperature, and q is the\nelementary charge. The UVC film shows a slightly higher 𝑛 value of 1.36 for P -\nlignt\ndependent V , indicating increased trap-assisted recombination. However, the P -\nOC lignt\ndependent J reveal that α increases from 0.886 to 0.956, signifying reduced\nSC\nbimolecular recombination, which is likely the major origin of an overall increase in\nJ despite the moderate reduction in the dark SCLC mobility. To clarify this, we\nSC\nrevised the following analysis in the manuscript.\nMain text:\n“The crosslinked UVC device exhibited a similar PCE of 16.7% but with an increased\nJ and a decreased V . The lower V is consistent with the presence of traps\nSC OC OC\nintroduced by UV crosslinking, resulting in higher non-radiative recombination and\nlower SCLC mobilities (Supplementary Fig. 12 a-b and Supplementary Table 2). This\nis reflected by the light-intensity dependences of V and J (Supplementary Fig. 12\nOC SC\nc-d): 𝑛 increases from 1.33 to 1.36, indicating higher trap-assisted recombination in\nUVC, whereas 𝛼 increases from 0.886 to 0.956, indicating improved charge collection\nwith reduced bimolecular recombination.”\nSupporting information:\n8\na b\n1.0\nHole-only device Electron-only device\nPzA5:L8BO 1.4 PzA5:L8BO\nUVC UVC\n- 1 ) m0.8 U Li V ne C a + r A fi S t H - 1 ) m U Li V ne C a + r A fi S t H\n2 c 2 c 1.2\n1 / 1 /\nA A\n2\n(\n0.6 2\n(\n1 / J 1 / J 1.0\n0.4 0.8\n1.5 2.0 2.5 1.5 2.0 2.5 3.0\nVappl (V)\n2\n) Vappl (V)\n-m\nc dc\nA\nV )0.90 PzA5:L8BO n=1.17 m PzA5:L8BO\ne\n( UVC\ny\n(\nolt a\ng\n0.87 L\nU\ni\nV\nne\nC\na\n+\nr\nA\nf\nS\nit\nH\ne n\nsit\nL\nU\ni\nV\nne\nC\na\n+\nr\nA\nf\nS\nit\nH\ne n\nsit 0.84 n=1.33\ne\nnt d 10 α=0.989\nc r α=0.956\nr r\nci u\ne n -0.81 n=1.36 uit c 5 α=0.886\np c\nO r\n0.78 -\nci\n10 100 rt 10 100\no\nh\nLight intensity (mW cm-2) S Light intensity (mW cm-2)\nSupplementary Fig. 12 | Charge transport and light-intensity dependences of the\nblends. a,b The J1/2–V curves for holy-only and electron-only devices using\nPzA5:L8BO, UVC, and UVC+ASH blends. c,d The light-intensity (P ) dependence\nlight\nof open-circuit voltage (V ) and short-circuit current density (J ) for organic solar\nOC SC\ncells based on PzA5:L8BO, UVC, and UVC+ASH with the α and n factors marked.\nThe P dependence of V and J are described by V ∝  𝑛kT/q×ln(P ) and\nlight OC SC OC lignt\nJ ∝P α, where k is the Boltzmann constant, T is the absolute temperature, and q is\nSC lignt\nthe elementary charge.\n9"}, {"round": 3, "reviewer_comments": {"reviewer_2": "The authors addressed all my concerns and the manuscript can be accepted."}, "author_response": ""}]

Decision letter