Correlative molecular-to-mesoscale evolution in conjugated polymers for intrinsically stretchable organic photovoltaics

Domain: chemistry
Published: 2026-02-20
Editorial outcome: minor_revision

Abstract
Conjugated polymer thin films offer a unique combination of tunable optoelectronic properties and mechanical flexibility, making them as promising materials for intrinsically stretchable optoelectronic devices. However, achieving both mechanical robustness and high device performance remains a key challenge. Addressing this requires a fundamental understanding of how molecular and mesoscale structures evolve under mechanical strain. Here, we employ a comprehensive suite of X-ray spectroscopy and scattering techniques to investigate the multiscale structural evolution of conjugated polymer thin films during uniaxial deformation. We uncover a two-stage morphological response: an initial stage characterized by polymer chain alignment and rapid crystallite disruption, followed by continued chain orientation accompanied by intrachain torsion at higher strains. These correlative structural adaptations govern key material properties, including stress dissipation, optical absorption, and photovoltaic performance. Our findings establish a mechanistic framework for understanding deformation in semiconducting polymers and provide design principles for developing mechanically robust, high-performance stretchable electronics.

Introduction
Intrinsically stretchable organic semiconductor electronics have shown significant interest for applications in wearable sensors, volumetric displays, and energy-harvesting devices1–9. Thin films of π-conjugated polymers are of particular interest in these systems, because their extended π-π stacking and chain statistical flexibility endow them with good optoelectronic performance even under applied strain. However, understanding how deformation drives polymeric structural evolution at various scales remains challenging, since the semi-flexible backbones and high torsional energy barriers of π-conjugated polymers result in mechanical properties far inferior to those of flexible polymers10–14. It is therefore critical to elucidate the correlations among these multiscale structural changes to establish robust structure-property relationships for emerging stretchable devices.

Under mechanical deformation, conjugated polymer thin films are subjected to stress that leads to chain orientation, and even disruption of chemical bonds15–17. For conventional flexible polymers under deformation, strain is relieved by relaxation across multiple length and time scales, a behavior well described by Gaussian-chain models. The rigid backbones and high torsional barriers of π-conjugated polymers, however, severely restrict in-chain conformation coordination. Thus, elastic deformation is minimal, and the large conformational changes are typically accompanied by crystallographic plane dissociation once the yield point is exceeded18. Nevertheless, conjugated polymers have enabled intrinsically stretchable electronics by molecular design and microstructural control. In organic field-effect transistors (OFETs), high-molecular-weight polymers shows retained crystalline domains and reversible chain alignment under >100% strain, thereby preserving electrical performance19. Random terpolymer design can suppress crystallinity and long-range order, while maintain short-range ordered aggregates, resulting in high mobility with improved stretchability20. In organic photovoltaics (OPVs), polymer-polymer blends outperform polymer-small molecule blends because gradual chain alignment and entangled networks dissipates stress while sustaining charge transport pathways21. The addition of high-molecular-weight conjugated polymers offers soft, amorphous regions with entangled and tie chains for stress dissipation, shielding the brittle acceptor crystallites and mitigating efficiency loss17. In organic light-emitting diodes (OLEDs), plasticizer incorporation transforms brittle conjugated polymers into fibrillar morphologies, resolving the conductivity-stretchability trade-off22. Despite these advances, direct in situ probing of the multiscale structural changes in conjugated polymer thin films, and how they collectively govern stress dissipation from the molecular to the mesoscale, remains an outstanding challenge.

Recent advancements in stretchable polymer semiconductors and energy-tunable X-ray characterization tools now offer a path to address this challenge directly23–34. The current work combines high-brightness soft and tender X-ray techniques to unravel the deformation behavior of conjugated polymer thin films from molecular to the mesoscopic scale. Near-edge X-ray absorption fine structure (NEXAFS) spectroscopy reveals strain-induced backbone orientation and intrachain torsion; tender resonant X-ray scattering (TReXS) demonstrates crystallite deformation; and resonant soft X-ray scattering (RSoXS) tracks morphological evolution at larger length scales. We identify a two-step morphological response that modulates mechanical behaviors, which influence absorption and photovoltaic performance. These insights into correlative, multiscale structural adaptation provide a foundation for rational design of durable and high-performance stretchable electronics.

Results
Molecular orientation and torsion
We selected the donor-acceptor copolymer poly[(N,N’-bis(2-octyldodecyl)-naphthalene-1,4,5,8-bis(dicarboximide)-2,6-diyl)-alt-5,5’-(2,2’-bithiophene)] (P(NDI2OD-T2)) as a model polymer for exploring the structural changes of stretchable polymer thin films (Fig. 1a). This polymer comprises alternating bithiophene (BT) and naphthalene diimide (NDI) units, with a dihedral angle of ~45° between NDI and BT (Supplementary Fig. 1). The ~200 nm film shows stretchable without cracks up to 50% strain when supported on a polydimethylsiloxane (PDMS) membrane (Supplementary Fig. 2).

To study the structures at the molecular level, NEXAFS spectroscopy was performed with a film-on-elastomer sample (Supplementary Fig. 3a). The emitted photons from the stretching thin film were collected by a 2D detector, resulting in total fluorescence yield (TFY) spectra with bulk-sensitivity (probing depth: ~100 nm) (Supplementary Fig. 4, 5). The NEXAFS intensity depends on the angle γ between the electrical field ~E of incident X-ray and the electronic transition dipole moment (TDM, O~) within the chemical bonds: I / ~E(cid:2)O~ 2 / (cosγ)2 (Fig. 1b)35. As shown in Fig. 1c, the TFY of polymer shows multiple peaks at 283.6–286.0 eV, corresponding to C 1s→π* transitions with TDMs normal to the conjugated planes. Spectral simulations reveal that the peaks at 283.6–284.4 eV mainly come from transitions localized to the NDIs, while the peaks at 284.4–286.0 eV arise from transitions delocalized across the NDI and BT units (Supplementary Figs. 6, 7)36.

We define the stretch direction (SD) of the film as the x axis, the transverse direction within the film plane as the y axis, and the surface normal as the z axis. Linearly polarized X-rays were incident at a fixed angle θ, with two distinct polarizations: ~E parallel to the x axis (horizontal polarization, H.) and ~E perpendicular to the x axis (vertical polarization, V.). These polarization conditions enable directional NEXAFS measurements to probe chain orientations along the in-plane (IP) and out-of-plane (OOP) directions. Under our experimental geometry, the NEXAFS spectra along the x-, y-, and z-axes were approximated by measurements using the H. beam, the V. beam at an incident angle of 80°, and the V. beam at an incident angle of 20°, respectively. The x-, y-, and z-spectra were used to calculate the refractive indices37, including both the dispersive (δ) and absorptive (β) components (δ is shown in Supplementary Fig. 8). In the pre-stretched film (strain = 0%), the β values at 283.6–286.0 eV show a trend with βz > βx = βy (Fig. 1d). In contrast, for the 50% strain film, the β values at 283.6–286.0 eV exhibit a trend with βz > βy > βx (Fig. 1e). To interpret these results, we simulated the π* transition intensities by aligning the polymer backbone along the x axis, which yield a trend of Iz > Iy > Ix within the 283.6–286.0 eV range (Supplementary Fig. 9). The comparison between experiment and simulation confirms that, upon stretching, chain backbones orient in the SD in the IP direction while remain face-on in the OOP direction.

To assess OOP orientation changes, angle-dependent NEXAFS spectra at incident angles (θ) of 20°, 55°, and 80° were measured (Fig. 1f). At 0% strain, the tilt angle (α) of π* TDMs relative to the surface normal was extracted using the relationship of I = A[1 + 1/2(3cos2θ-1)(3cos2α-1)], where I is the peak area, and A is a constant35,38. This angular dependence of NEXAFS intensity is appropriate for samples with three-fold or higher substrate symmetry, consistent with the as-cast P(NDI2OD-T2) thin films prepared by spin-coating. While some conjugated polymer systems can indeed exhibit IP alignment depending on the film-forming process38, we note that this fitting method should not be considered universally valid for all thin films prior to stretching. Under 50% strain, the film became a two-fold rotational symmetry in the IP direction. The α angle can be fitted from the expression I = A(cos2θcos2α + sin2θsin2αcos2φ), where the φ is the azimuthal angle of π* TDMs in the IP direction. The calculated α angles at Area 1 (π* signals localized in NDI, 283.6–284.4 eV) and Area 2 (total π* signals, 283.6–286.0 eV) are 32.2° and 34.6°, respectively, which are 5.6° and 8.0° decrease compared to the pre-stretched film (Fig. 1g, and Supplementary Fig. 10–13). The larger decrease in intensity for Area 2 compared to Area 1 suggests that the NDI and the entire polymer chain showed different motions in response to strain. Specifically, the BT shows greater intrachain rotation relative to the substrate than the NDI, suggesting a higher torsion angle between BT and NDI.

The intrachain torsion development as a function of strain was evaluated by NEXFAS spectra measured using horizontal X-rays during incremental film strains (Supplementary Fig. 14). With increasing strain, the intensity ratio of Area 1 relative to Area 2 increases and becomes larger at high strains (Fig. 1h). Molecular dynamics (MD) simulations corroborate such observation, showing an exponential rise in the dihedral angle between BT and NDI with increasing strain (Fig. 1i, and Supplementary Fig. 15). Achieving high dihedral angles requires surmounting specific energy barriers between subunits as demonstrated by potential energy surface scans (Supplementary Fig. 16). Together, these results indicate that strain-induced chain orientation is accompanied with intrachain conformational torsion, which plays a critical role in dissipating stress under high strains.

Crystallite orientation, deformation and destruction
To study crystallite orientation and deformation, we employed TReXS in a transmission geometry with horizontal X-rays at 2476 eV (Fig. 2a). The PDMS-supported freestanding configuration provides a practical approach to impose uniaxial strain on thin polymer films during in situ X-ray scattering experiments, as validated by simulated strain and stress maps (Supplementary Figs. 17, 18). The selected photon energy of 2476 eV enables the selective enhancement of the (001) reflection along the SD due to the S 1s→σ* transition localized to the thiophenes (Supplementary Figs. 19–21)39,40. P(NDI2OD-T2) has two main reflections: a (100) reflection at q~0.25 Å−1 and a (001) reflection at q~0.47 Å−1, corresponding to lamellar stacking and in-chain electronic interaction (Fig. 2b)41. As shown in Fig. 2c, we categorize crystallites as either parallel or perpendicular based on the orientation of their chain backbones relative to the SD. Parallel crystallites have chain backbones aligned with the SD, displaying (001) and (100) reflections in the qx (||) and qy (⊥) directions, respectively. Perpendicular crystallites with chain backbones transverse to the SD, show (001) and (100) reflections in the qy and qx directions, respectively.

Under 0–33% increasing strain, the parallel crystallites exhibit a 37% drop in (001) intensity and a 44% drop in (100) intensity (Fig. 2d, and Supplementary Fig. 22–25). In contrast, the perpendicular crystallites show more significant degradation, with (001) intensity decreasing by 58% and (100) by 54%. Such observations are particularly pronounced in the early stage of stretching. Upon releasing the film, the scattering intensities continue to decrease between 30% and 20% strain, likely due to residual stress remaining above the critical threshold for crystallite destruction. Below ~20% strain, no further decrease is observed; however, neither reflection recovers, indicating a permanent loss of crystallinity for both directions. The (001) anisotropic ratio, defined as (I|| - I⊥)/(I|| + I⊥)42, increases from 0.12 to 0.33 during stretch (Supplementary Fig. 26). These results suggest that crystallites are highly strain-sensitive, with parallel crystallites being less susceptible to strain-induced destruction, resulting in enhanced crystallite orientation in the SD.

To further explore crystallite deformation, we calculated the layer packing number for each reflection by taking the ratio of crystalline coherence length (CCL) to d-spacing (Fig. 2e). For parallel crystallites, the packing number for the (100) reflection decreases significantly, while the (001) reflection remained stable, suggesting crystallographic slippage, where the layers slide past one another along the SD (Supplementary Fig. 27a). Such slippage allows the crystallite layers to partially shift without entirely losing their stacked arrangement, thus accommodating strain while keeping some structural integrity. In contrast, perpendicular crystallites showed that packing number for the (001) reflection decreased more noticeably, suggesting that more repeated units are lost from the crystallites, a process we term crystallographic peeling (Supplementary Fig. 27b). Rather than sliding, the crystallite layers pull apart at the edges, resulting in a greater and faster structural disruption, while providing additional amorphous chains that align and undergo intrachain torsion.

To study the π-π stacking developments during stretch, the films were stretched to the desired strains and then placed on silicon wafers for grazing incidence wide angle X-rays scattering (GIWAXS) measurements at the beam energy of 10 keV. This measurement was performed with orthogonal X-ray directions including parallel and perpendicular conditions (Supplementary Fig. 28). The 0% strain thin film shows nearly identical reflections (Supplementary Fig. 29), which support that βx ≈ βy before stretching as demonstrated by NEXAFS. Meanwhile, the (100) peak area in both qx and qy decreases rapidly at small strains (10-20%) and then more gradually at higher strains, with a stronger reduction along qx (Supplementary Fig. 30–32, and Supplementary Table 1-2). This behavior is consistent with the observations at 2476.25 eV in TReXS, indicating that the resonant excitation is not sensitive to the crystalline deformation. The azimuthal intensity of the π-π stacking reflection was used to evaluate crystallite orientation relative to the substrate (Supplementary Fig. 33). With increasing strain, the pole figures show narrowing for the perpendicular condition and broadening for the parallel condition. Therefore, an enhanced face-on orientation of π-π stacking is observed when X-rays probe perpendicular to the SD, in agreement with the angle-resolved NEXAFS results. The opposite trend in pole figure widths suggests reduced twisting of crystallites about the (001) axis but increased tilting about the (100) axis, consistent with the tilted crystallite orientation observed in flow-induced alignment42. Together, these combined mechanisms of crystallite deformation and orientation provide irreversible structural adaptation pathways for stress dissipation.

Morphological evolution under strain
The chain orientation with intrachain torsion and crystallite destruction set the basis for large-scale morphological changes in the film, which we studied using in situ RSoXS in a transmission geometry (Fig. 2a). To optimize the measurements, we estimated the scattering contrast to determine an appropriate soft X-ray energy. As shown in Fig. 3a, atomic force microscopy (AFM) height images reveal the pre-stretched film showing refined fibrillar nanostructures. For the 50% strained film, large-scale fibril aggregates with their long axes align toward the SD are seen, showing a fast Fourier transform (FFT) pattern with an X-shaped anisotropy, which results from the progressive development in both size and orientation of the aggregates (Supplementary Fig. 34). This reflects the formation of a statistically oriented structure induced by strain. We then calculated the orientational contrast in the xy, xz, and yz planes by E4(Δδ2+Δβ2) based on refractive indices δ and β along each axis (Fig. 3b)43,44. This analysis shows a significant intensity increase for the xy- and xz-components at ~284.0 eV, suggesting that chain backbone orientations in both IP and OOP directions strongly influence the scattering signal. Further analysis of the integrated scattering intensity (ISI, ISI/ q2Iq2dq) from the RSoXS I-q curves confirms that at ~284.0 eV, scattering contrast is maximized when the free-standing film is strained to 30% (Fig. 3c, and Supplementary Fig. 35–38).

RSoXS patterns probed at 284.0 eV during a stretch-release process (stretch: 0-30%; release: 30-0%) are shown in Fig. 3d, e. At the pre-stretched state (0%), the film shows a slightly anisotropic scattering pattern that varies with X-ray polarization (Supplementary Fig. 39), suggesting contributions from local molecular orientation or the form factor of core-shell-like fibrillar domains32,45. Under mild strain (5.66%), the scattering pattern changed to an X shape, which elongates transversely with increasing strain (5.66-29.81%). This X-shaped pattern is independent of X-ray polarization (Supplementary Fig. 40), indicating that global orientation occurs upon stretching. The corresponding characteristic domain size is also evaluated using the Debye-Bueche model, which shows a trend of structural reorganization that complements the micro-scale evolution observed by AFM (Supplementary Figs. 41–43). We note that the strained thin-film morphology is complex, as domain and orientational contrasts are intertwined and difficult to decouple. Nevertheless, the consistent X shape observed in AFM FFT images, the enhanced orientational contrast in the xy- and xz-planes at 284.0 eV, and the local maximum ISI at 284.0 eV indicate that the X-shaped scattering arises from strain-induced alignment of polymer chains across the fibrillar morphology. In situ RSoXS measurements at 285.0 eV during stretching show that X-shaped feature disappears (Supplementary Fig. 44), suggesting reduced orientation sensitivity and further supporting the assignment of scattering contrast origin. Thus, the X pattern provides direct information related to the orientation of fibril aggregates. Such results also indicates that although the crystalline scattering exhibits similar strain decay profiles across different X-ray energies, RSoXS contrast variation reveals environment-dependent mechanical behavior. Upon releasing strain, the X shape rotates and intensifies along the SD, likely due to the formation of wrinkles, which introduce vacuum/roughness contrast dominance (Supplementary Fig. 45).

We define the angular positions of the X shape using the azimuthal angles, φ1 and φ2, where Δφ(=φ2 – φ1) represent the degree of orientation relative to the SD (Fig. 3f, Supplementary Figs. 46, 47). As shown in Fig. 3g, h, Δφ was first calculated at 3.54% strain, suggesting that elastic structural changes by minimal molecular rearrangements are limited to very low strain. Then, Δφ shows a two-step growth, with a transition strain at 14.69%. In the first stage, Δφ increases quickly, driven by initial chain orientation and rapid crystallite destruction, an irreversible process that provides an early, foundational orientation to the morphology. In the second stage, Δφ displays another rapid increase, facilitated by the continued orientation and more chains freed by crystallite destruction, coupled with significant intrachain torsion at higher strains. Together, this sequential orientation process highlights a unique morphological adaptation in semiflexible conjugated polymers under mechanical deformation.

Upon release, Δφ shows a linear decrease from 30% to 20.27% strain, after which it becomes relatively stable. The linear decrease is due to partial recovery of φ1 and φ2, indicating that chain orientations relax away from the SD. Notably, the Δφ at 20.27% aligns closely with that at 14.69% during stretch, suggesting that the structural relaxation involves the recovery of intrachain torsion, which is significantly enhanced during the later stage of stretching. However, beyond 20.27% strain, the stable Δφ suggests limited further relaxation, likely due to the formation of wrinkles that dissipate part of the applied strain energy (Supplementary Fig. 48)46. These results highlight the role of aligned chains with intrachain torsion across morphological scales for effective reversible stress dissipation.

Correlations between structural evolution and properties
We further sought to connect the strain-induced multiscale structural changes with the mechanical, optical, and electronic properties of the thin film. As shown in Fig. 4a, the stress-strain (σ-ε) curve of the P(NDI2OD-T2) thin film displays an elastic region (~3%), a yield point (~10%), strain hardening (~14%), and ultimate fracture (~35%). Note that the crack-onset strain from optical microscopy (supported film) is not directly comparable to the fracture strain from tensile testing (free-standing film), as the former reflects delayed surface cracking under strain relaxation by the substrate, whereas the latter represents complete mechanical failure of the film. The yielding and hardening behaviors correspond well with the two-step Δφ growth observed in morphological adaptation under strain. In the initial stage, rapid crystallite destruction and partial chain orientation contribute to morphological plasticity, distributing stress evenly across the film and preventing localized stress concentration, and thus leading to ductile yielding. In the subsequent stage, further chain alignment and the development of intrachain torsion over a large morphological scale increase resistance to additional strain, while effectively dissipating stress, thus contributing to the toughening and strain-hardening response. This connection underscores how multiscale structural changes in conjugated polymers translate directly to enhanced macroscopic mechanical properties, which helps understand why incorporating non-fused units and increasing molecular weight are effective chemical design methods for achieving stretchable conjugated polymers47–51.

Next, we studied the UV-vis-NIR absorption spectra of the thin film during tensile stretching with a film-on-elastomer sample (Fig. 4b). Upon stretching to 50% strain, the main absorption peak at 709 nm exhibits a 10 nm blue shift, while the shoulder around 790 nm diminishes (Fig. 4c). This shift is attributed to the combined effects of crystallite destruction and intrachain torsion under strain, leading to a local order-disorder transition. Time-dependent density functional theory calculations using strain-dependent dimer structures confirm that increasing the dihedral angle between BT and NDI units from ~30° to 90° induces a blue shift in absorption (Supplementary Fig. 49), consistent with reduced π-conjugation and exciton delocalization. The absorption peak also narrows, which we attribute to increasing chain alignment across the morphological scale, limiting the distribution of vertical transition energy levels. These findings highlight how multiscale structural adaptations influence light-harvesting properties, which are critical for devices, such as organic photovoltaics (OPVs) and organic photodetectors (OPDs).

Intrinsically stretchable OPV devices with the structure of TPU/PEDOT:PSS (M-PH1000)/PEDOT:PSS (4083)/active layer/PNDIT-F3N/EGaIn@Ag were fabricated52. The active layer comprised the two conjugated polymers, PO10 (donor) and P(NDI2OD-T2) (acceptor), with a 2:1 weight blend ratio (Supplementary Fig. 50)43. This blend film shows ductile deformation, with a fracture strain of 21% observed from its σ-ε curve and crack-onset strain at 30% on PDMS substrate observed by optical microscopy (Fig. 4d, Supplementary Fig. 51). Similar to the neat thin film, the P(NDI2OD-T2) component in the blend shows blue-shifted absorption, during stretch from 0% to 30% strain, while PO10 contributes only minor changes within its own absorption range (Fig. 4e and Supplementary Fig. 52). The pre-stretched device exhibits a PCE of 6.92%, with a VOC of 0.854 V, a JSC of 13.97 mA cm-2, and an FF of 58.00% (Supplementary Fig. 53, Supplementary Table 3). Upon stretching to 30%, the PCE drops to 5.78% (VOC = 0.858 V, JSC = 12.48 mA cm-2, FF = 53.97%), indicating a loss of 84% of the initial PCE. Such decline is primarily due to the decrease in both JSC and FF (Fig. 4f), which correlate with charge generation, transport, and recombination in the devices.

The exciton dissociation and charge collection efficiency, P(E,T), extracted from photocurrent density-effective voltage (Jph-Veff) curves, decreases from 98.86% for pre-stretched condition to 97.93%, 96.85%, and 95.53% at 10%, 20%, and 30% strain, respectively, indicating reduced charge generation (Supplementary Fig. 54). Photo-CELIV analysis reveals a concurrent decline in charge mobility from 6.59×10-5 to 4.13×10-5 cm2 V-1 s-1 (Supplementary Fig. 55). The light intensity (Plight) dependence of VOC, described by VOC/nkTln(Plight)/q, yields n values of 1.25, 1.42, 1.46, and 1.58 as strain increased, signifying enhanced trap-assisted recombination (Supplementary Fig. 56). To further elucidate this, morphological changes upon thin film stretched by nano-infrared spectroscopy atomic force microscopy (AFM-IR). As shown in Fig. 4g–i, P(NDI2OD-T2) domains probed with 843 cm−1 exhibit nanofibrillar network morphology in the as-cast blend film (Supplementary Fig. 57), which is retained when 10% strain was applied, but transformed into large aggregates at 30% strain. This morphological transformation arises from strain-induced chain orientation, intrachain torsion, and crystalline destruction, leading to local order-disorder transitions. These combined absorption, electrical, and morphological results demonstrate that strain-induced structural evolution directly deteriorates charge generation and transport while promoting recombination losses, ultimately accounting for the observed performance degradation. To assess the effect of cyclic strain, PO10:P(NDI2OD-T2) films on PDMS were stretched at 20% for 300 cycles. AFM images show that the fibrillar morphology gradually blurred with repeated cycling (Supplementary Fig. 58), causing some decline in device performance; however, the device retained over 80% of its initial PCE (Supplementary Fig. 59 and Supplementary Table 4). These results indicate that the overall nanostructure and device performance remain largely preserved under moderate morphological relaxation.

Therefore, these results underscore that strain-induced multiscale structural changes inevitably compromise some electronic properties, a critical consideration for stretchable organic electronics. Recent studies have identified entropic chain alignment and enthalpic crystallite breakage as multimodal energy dissipation pathways for mechanical robustness19,48,52–57. Both effects are directly observed in our system, with a strain-dependent temporal evolution further revealed. Consistent with a recent review highlighting the role of chain alignment and torsional rotation in mechanical-electronic coupling, the in situ and simulation data show increased torsional freedom under strain, likely governed by backbone architecture53. Our findings suggest that incorporating additional stress dissipation mechanisms, such as dynamic bonding and elastomer networks19,48,52–57, can mitigate structural deformation and improve device performance.

Methods
Sample preparation and general characterizations
P(NDI2OD-T2) was synthesized according to a reported procedure58. Its molecular weight was characterized by high-temperature gel permeation chromatography (GPC) on a Tosoh EcoSEC HT system equipped with three TSKgel GMHhr-H(S) HT columns in series, using 1,2,4-trichlorobenzene as the eluent at 135 °C. Calibration was carried out against narrow-dispersity (Đ) polystyrene standards. The number-average molecular weight (Mn) and dispersity of P(NDI2OD-T2) were determined to be 131 kDa and 1.356, respectively. The corresponding molecular-weight distribution curve is shown in Supplementary Fig. 60. The polymer PO10 was synthesized following a reported protocol50, with Mn = 24.4 kDa and Đ = 2.15. The Mn of PNDIT-F3N was determined to be 12.3 kDa by GPC (Waters 2410) using tetrahydrofuran (THF) as the eluent and linear polystyrene standards for calibration. Preparation of thin films of P(NDI2OD-T2) involved several steps: (1) the PEDOT:PSS (Clevios PVPAI 4083, Heraeus) was spin-coated on a wafer at 4000 rpm; (2) P(NDI2OD-T2) thin films (~200 nm thick) were prepared by spin-coating on the PEDOT:PSS/silicon wafer substrates from a chloroform solution (10 mg mL−1). The thicknesses of films were measured by a Tencor Alpha-step 500 Surface Profilometer. The film was then treated with thermal annealing at 100 oC on a hot plate for 10 min. The PDMS membrane was prepared by mixing 6 g base and 0.5 g cross-linker in a 20 mL vial with weight ratio of 12:1 (Sylgard 184, Dow Corning). The mixture was placed in a vacuum box for 10 min to remove the bubbles and then transferred on a 4-inch silicon wafer, and baked in a 70 oC oven for 1 h. Both film-on-elastomer and free-standing samples were prepared following the method as illustrated in Supplementary Fig. 3.

UV-vis-NIR absorption spectra of thin films were measured by a SHIMADZU UV-3600 spectrophotometer. The film-on-elastomer sample was placed with the clamps of a tensile stage and then subjected to the desired strains for the measurements. A pristine PDMS membrane with the same thickness as the film sample was used as a reference to subtract the absorption background. The resulting spectra were normalized to compare their differences in wavelength shift. AFM height images of the PNDI2OD-T2 thin films were captured by a Bruker Multimode 8 Atomic Force Microscope. AFM-infrared (AFM-IR) spectroscopy was performed using an Anasys nanoIR3 system (Bruker). The thin films were stretched to specific strains and transferred onto silicon wafers for imaging.

NEXAFS simulations
The NEXAFS simulations of P(NDI2OD-T2) were done with its simplified structural unit BT-NDI, whose geometry was optimized with density function theory (DFT) at the B3LYP/6-31G* level using the Gaussian 09 package. The resulting molecular model was then used as input for NEXAFS simulations based on the excited electron and core hole (XCH) calculations. The resulting spectra were shifted by 284.5 eV for C K-edge. However, the PBE-GGA functional that is employed can underestimate the band gap, leading to compressed spectra at the energy axis. To alleviate this issue, a dilation factor of 1.1 was applied to C K-edge simulations. The resulting transitions were converted to continuous spectra, which were more relevant to experimental data by using a Gaussian function with peak widths of 0.2 eV. We note that despite these corrections, small residual discrepancy of ~0.2–0.3 eV remains, as the simplified molecular model cannot fully capture the long-range conjugation, and the vacuum-based calculations cannot account for solid-state interactions. These deviations are commonly accepted in DFT-based NEXAFS simulations and do not affect the reliability of spectral assignments33,36,59.

MD simulations
The MD simulation to study the dihedral angle changes of P(NDI2OD-T2) under stretching were performed as follows: (1) The structural geometry of a P(NDI2OD-T2) dimer was optimized using density functional theory (DFT) calculations at the B3LYP/6-31g** level, implemented in Gaussian 16 and analyzed with Multiwfn software61. The optimized dimer structure was then used as the repeat unit to construct a decamer chain of the polymer. (2) An amorphous cell containing the decamer chains was built with an initial density of 0.3 g cm-3 to ensure dense packing. (3) The cell was subjected to incremental applied stress along the x axis (0.001, 0.05, 0.10, 0.15, 0.16, 0.17, 0.18, 0.19, 0.20, 0.21, 0.22, and 0.23 GPa). For each stress, 100 ps of MD simulations were performed under the NVT ensemble at 300 K to capture the evolution of dihedral angles between specific units. The absorption spectra were calculated using a Gaussian 16 software based on cam-B3LYP-D3/TZVP, upon a tetramer structure extracted from the stress-strain MD simulation with rising torsion angle.

TReXS
The TReXS measurements were performed at the Soft Matter Interfaces (SMI) beamline (Beamline 12-ID) of the National Synchrotron Light Source II, Brookhaven National Laboratory. The P(NDI2OD-T2) thin films were transferred onto PDMS substrates with a central circular hole (4 mm in diameter), as illustrated in Supplementary Fig. 3b. The X-ray beam size was 20 μm × 200 μm, ensuring that the beam fully passed through the unsupported film region. The sample was mounted in a tensile stage (Linkam TST350) in vacuum. Both static and in situ scattering images were collected by a Pilatus 300K-W detector under a photon energy range of 2445–2500 eV, with a refined step size of 0.25 eV for the data acquisition through S K-edge (2470–2480 eV). Under TReXS, the reflection intensity is proportional to the square of the structure factor by I(q) / |F(q)|2. The structure factor, F(q), is calculated using the equation F(q) = Σ atoms fj exp(2πiq(cid:2)rj), where rj represents the location of the jth atom in the crystal unit cell and fj represents the energy-dependent atomic scattering factor of the jth atom. Changes in reflection intensity at resonance are caused by variation the fj, which has a real component and an imaginary component, denoted as f = f0 + f’ + if”. The imagery part, f”, is determined using NEXAFS spectra, while the real part, f’, is obtained through a Kramers-Kronig transformation. To ensure consistent comparison across all strain levels, the scattering intensity was normalized to the value at q = 0.37 Å−1 at 0% strain. No diffraction contribution from PDMS was observed; PDMS typically shows a reflection at q~0.40 Å−1 (Supplementary Fig. 61), whereas our TReXS data remain flat in this range, confirming that the measured scattering originates solely from the polymer thin film.

GIWAXS
The GIWAXS measurements of static stretched films were performed at beamline 7.3.3 of ALS, LBNL. The beam energy was 10 keV. The samples were placed in a Helium box with a sample-to-detector distance of ~280 mm, which was calibrated by a silver behenate. The incident angle was 0.16o, which was normalized by a photodiode. The scattering images were collected by a CCD detector (Pilatus 2M) with a pixel size of 0.172 mm × 0.172 mm and an exposure time of 2 s.

NEXAFS and RSoXS
The uniaxial tensile tests of the thin film samples were performed by a tensile stage (200 N, Deben UK Ltd). The elongation and movement rate of clamps were controlled and monitored by a controller that was coupled with Microtest v6.3.30, a computer-based software. NEXAFS (TFY mode) and RSoXS were performed at beamline 11.0.1.2 of ALS, LBNL. The CCD detector (Princeton Instrument PI-MTE) has a pixel size of 0.027 mm × 0.027 mm. For NEXAFS, the tensile rate was set at 1.5 mm min−1. Post-edge normalization of NEXAFS spectra was done by Athena 0.9.26. For the calculation of refractive indices, the absorptive component β was determined from NEXAFS spectra by β = μ/2k, where μ denotes the attenuation coefficient and k is the wave vector. The dispersive component δ was then obtained via Kramers-Kronig transformation with β. For in situ RSoXS, the tensile rate was 0.5 mm min−1. The sample preparation followed the same procedure as shown in Supplementary Fig. 3b. The beam size is 200 μm × 200 μm, ensuring that the beam fully passed through the unsupported film region. The sample-to-detector distance was set at 150 mm. During the tensile test, the scattering images were captured consecutively by the CCD with 5 s exposure time per image.

Stress-strain (σ-ε) curves of thin films
Tensile tests were conducted using a free-standing sample following the reported method50. The polymer solution was spin-coated onto a PEDOT:PSS/silicon wafer substrate. The sample was then floated on deionized water, allowing the PEDOT:PSS layer to dissolve and leaving the active layer free-standing on the water surface. Frosted aluminum blocks attached to the clamps of a tensile stage were carefully adjusted to capture the floating film via van der Waals interactions. After securing the film, the deionized water was gradually removed, leaving the active layer suspended on the tensile stage. The film was stretched at a constant rate of 0.03 mm/min. Stress (σ) and strain (ε) were calculated using the following equations: σ = F/(A×B) and ε = Δl/l0, where F is the applied force measured by a high-sensitivity force sensor, A is the film width (10 mm), B is the film thickness, Δl is the elongation of the film, and l0 is the initial distance between the clamps.

Fabrication of intrinsically stretchable OPV devices
The stretchable OPV devices were fabricated with the structure TPU/PEDOT:PSS (M-PH1000)/PEDOT:PSS (4083)/active layer/PNDIT-F3N/EGaIn@Ag. The stretchable, transparent electrode was prepared by modifying PEDOT:PSS (PH1000, Heraeus) with the addition of 5 vol% dimethyl sulfoxide, 2 vol% polyethylene glycol, 0.5 vol% FS-30, and 0.1% GOPS. The resulting solution was spin-coated onto a TPU substrate at 1500 rpm, followed by baking at 100 °C for 20 min. Next, PEDOT:PSS (4083, Heraeus) was spin-coated onto the PH1000/TPU layer at 3000 rpm and dried at 100 °C for 20 min. The PO10:P(NDI2OD-T2) blend film (2:1 weight ratio) was then spin-coated with its chloroform solution (6 mg/mL for PO10, with 0.5 vol% dibenzyl ether as solvent additive) onto the 4083 layer as the active layer. Post-treatment to the active layer with thermal annealing of 100 °C for 10 min. Then, PNDIT-F3N (methanol solution, 1 mg/mL) was spin-coated at 3000 rpm onto the active layers. Finally, a 100 nm thick Ag layer was thermally evaporated onto the interface through a shadow mask in a vacuum chamber at a pressure of 3×10−7 torr. EGaIn liquid metal was then sprayed onto the Ag layer using a deposition mask, yielding the EGaIn@Ag stretchable electrodes. The initial active area (0.04 cm2) was defined by the overlap of the patterned M-PH1000 bottom electrode and the patterned EGaIn@Ag top electrode in a cross-electrode geometry. During mechanical stretching, the active area deformed geometrically. To account for this, the actual active area was corrected based on dimensional changes measured by optical microscopy. The average area variation was extracted from three independent devices under identical strain conditions and fitted with a three-phase exponential decay function (Supplementary Fig. 62), which was then applied to correct the active area of the tested devices at each strain. J-V curves were measured using a Keithley 2400 source meter with illumination (100 mW cm-2) provided by an AM 1.5G solar simulator (SS-F5-3A, Enlitech) in a nitrogen-filled glovebox. Photo-CELIV measurements were performed using the PAIOS platform (FLUXiM, Switzerland).

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The manuscript by Zhong et al. combined NEXAFS, TReXS, and in situ RSoXS to study the microscale to mesoscale structural evolution of a model conjugated polymer P(NDI2OD-T2) under mechanical strain. The study reveals a two-step response mechanism that finally causes the crystal peeling effect under large strain: first, the stage of polymer chain orientation and rapid crystallite fracture, and secondly, the stage of chain alignment and intrachain torsional deformation. Further, the authors studied P(NDI2OD-T2):PO10 blend systems and makes a correlation between device behavior and morphological disruption resulting from strain-induced meso- and microscale crystalline destruction investigated from resonant X-ray scatterings. With minor revisions and more experiments, this paper is well-suited for publication in Nature Communications:\n1. In lines 117 to 122, the authors used different mathematical equations to fit the tilt angle of the transition dipole moment under 0% and 50% strain. It seems the expression used for the 0% strain condition (I = A/3[1+12(3cos2θ−1)(3cos2(cid:1)−1)]) is assuming a uniaxial orientation distribution where the stacking is face-on predominantly in the z-direction but in-plane orientation in the x and y directions is random. While this may be a reasonable assumption for an unstressed system, there are some conjugated polymers that have some in-plane alignment before stretching depending on the film forming process. The authors should clarify whether this assumption is supported by earlier statements in the manuscript, specifically the claim that βₓ equals βᵧ at 0% strain, and whether such a model is universally valid for the polymer system under study.\n2. In line 132, the authors refer to “Area 1” in the NEXAFS analysis, but the methodology used to extract this value is not described in detail. I assumed that the area under the curve was calculated directly from a fixed energy window. However, the authors should discuss in SI or method part whether a peak fitting approach was considered, and whether this method might offer a more accurate representation of the electronic transitions. Authors should also mention, if no fitting is performed, if using such a peak fitting would change the observed trend or only shift the numerical values.\n3. The display of figure 2 is confusing. First, the italicized terms “para.” and “perp.” overlap with the labels for crystal directions, which create confusion for the reader. Replacing these overlapping labels with clearly marked axes or direction arrows would enhance readability. Second, the data shown in Figures 2c to e would be more interpretable if the corresponding scattering features were marked using the same color to represent. For example, the author can use four different colors (grouped inro two main group, like deep and light shades of blue and red) and denote the corresponding scattering feature and also use the corresponding color in figure 2c (instead of plain yellow).\n4. The authors have discussed the scattering pattern at 2450 eV in TReXS but had not offered further discussion. Given that this energy also excites the (100) peak, it would be useful to know whether the (100) scattering shows the same decay profile under strain as excited by other energy. If the behavior differs, this may suggest distinct mechanical responses in the polymer chain in different environments. The authors should include it in either SI or the body.\n5. In line 225, the authors argue that the RSoXS results align with AFM data; however, AFM images are only provided for samples strained to 50%, which is significantly larger than that of the RSoXS experiment (30%). It remains unclear at what strain level the “X-shaped” anisotropy emerges in AFM. Including an FFT analysis of AFM images at intermediate strains, such as around 30%, would help confirm the structural evolution pathway supported by the X-ray scattering data. Additionally, a comparison of domain sizes derived from AFM and RSoXS would be valuable. Although the feature is not seemingly a peak, the authors still may select some turning points in the small angle range if plot a figure on log-log scale.\n6. The UV-Vis spectrum of the blend film is unclear. The spectral features that appear in the blend film is significantly shifted compared to neat P(NDI2OD-T2). It is unclear whether this shift originates from the P(NDI2OD-T2) component or from PO10. The authors should include a strain-dependent UV-Vis spectrum of neat PO10 films to determine the individual contributions of each material in the blend. This would allow them to confidently attribute the observed blue shift in the blend to a specific molecular origin.\n7. While the manuscript connects morphological evolution to device performance, the authors could elaborate further on how the relatively subtle strain-dependent optical changes (only a few nm shifts in absorption peaks) influence charge generation, transport, or recombination in OPV devices.\n8. The authors could expand the discussion of prior works on strain-induced morphology changes in conjugated polymers. For example, Wu et al. (https://www.nature.com/articles/s41467-023-44099-w) introduced the “relative stretchability” (rS, the ratio of DR and rDoC) to quantitatively analyze the strain induced morphological evolution and crystal breakage. Comments on whether your system follows similar entropic (chain alignment, from DR) and enthalpic (crystallite amorphization/breakage, from rDoC) dissipation pathways would be helpful for the readers. In addition, Zheng et al. (https://pubs.acs.org/doi/full/10.1021/jacs.2c00072) summarized the molecular and morphological processes expected under strain, for example, chain alignment and strain‑induced backbone planarization through monomer torsional rotation and then outlined how such processes map onto mobility retention. Since this manuscript shows direct evidence on monomer rotation to planarize the chain under strain, highlighting how it supports or refines the assumptions in this review paper would be helpful.", "reviewer_2": "The manuscript presents a fundamental investigation into the structural evolution of a conjugated polymer under tensile stress, utilizing a combination of NEXAFS, TReXS, and RSoXS techniques. The studied material is further demonstrated in a fully stretchable photovoltaic device. Overall, the manuscript is well written, and the comprehensive suite of characterizations are clearly correlated to provide a coherent mesoscale picture that advances the fundamental understanding of intrinsically stretchable organic semiconductors during dynamic mechanical deformation. I recommend this work for publication, provided that the following concerns and questions are satisfactorily addressed:\nQ1: In comparing the simulated NEXAFS spectra of P(NDI2OD-T2) with the experimental results in Figure 1c, there remain noticeable discrepancies, particularly around 285.1 eV. What might be the origin of these mismatches? Additionally, the major and minor tick marks on the x-axis of Figure 1c are not clearly visible and should be improved for clarity.\nQ2: Regarding the free-standing experimental setup, as the majority of the film remains supported by the PDMS substrate, could this result in differences between the local strain experienced near the PDMS opening and the overall applied strain? How is the strain distributed across the film, and is there any quantification or mapping of potential strain gradients?\nQ3: For the TReXS and RSoXS experimental setups, which employ a transmission geometry, how is it ensured that the X-ray beam does not also penetrate and contribute signal from the PDMS substrate? What is the beam size relative to the size of the unsupported region (hole)? Additionally, since the scattering intensity is correlated with film thickness—and considering that the film thickness is likely reduced during stretching—is this variation reflected or corrected for in the analysis (for example, in Figure 2d)?\nQ4: While optical microscopy determined the crack-onset of P(NDI2OD-T2) to occur at strains above 50% (Fig. S2), the macroscopic stress-strain curve exhibits mechanical failure before reaching 35% strain (Fig. 4a). Can the authors clarify the origin of this apparent discrepancy?\nQ5: The observation of a two-stage strain response corresponding to distinct stress dissipation mechanisms in P(NDI2OD-T2) is intriguing. To what extent is this behavior generalizable to other conjugated polymers, such as PO10, already studied in this work?\nQ6: The demonstration of robust photovoltaic device performance under up to 50% applied strain is impressive. Is it possible to directly characterize the charge carrier mobility of P(NDI2OD-T2) under strain? Such measurements would provide important insight into the relationship between mechanical deformation, structural evolution, and charge transport properties.\nQ7: For real-world applications, stretchable semiconductors will undergo repeated cyclic stretching rather than a single extension. How do the authors anticipate that cyclic mechanical fatigue will influence the nanostructure and electrical performance over time? Are there any preliminary results or plans to study such effects?\nQ8: The molecular weight of the P(NDI2OD-T2) is very high, 131k; providing the P(NDI2OD-T2) GPC traces and the calibration curves will be helpful. Besides, material information (synthetic procedure, molecular weight) on PO10 and PNDIT-F3N is missing.", "reviewer_3": "This manuscript presents a novel and comprehensive approach to study the multiscale structural evolution of conjugated polymer thin films under strain, using a powerful combination of advanced in situ X-ray techniques (NEXAFS, TReXS, RSoXS). The study provides important mechanistic insights into stress-dissipation pathways, including chain orientation changes, intrachain torsion, crystallite disruption (via “peeling” and “slippage” mechanisms), and a “two-step” morphological reorganization. Such structural changes are successfully connected to mechanical and optoelectronic properties. The experimental design is rigorous, the correlative multimodal methodology is state-of-the-art, and the results offer valuable guidance for the rational design of durable stretchable electronics. Given the novelty of the characterization method and the depth of the mechanistic insights, I strongly recommend publication after minor revisions.\n(1) While the title emphasizes the study of structural evolution on organic photovoltaic, the Introduction does not clearly articulate how key morphological changes under strain impact photovoltaic metrics. A more explicit discussion connecting the multiscale structural changes to performance-limiting mechanisms in organic solar cells would enhance the impact of the work.\n(2) The angle-resolved NEXAFS shows increased α-angles of polymer chains under strain. Are similar changes observed within the crystalline domains? Since face-on π-π stacking is crucial for vertical charge transport in OSCs, clarification on whether the crystallite orientation is similarly affected would be helpful.\n(3) The observed decrease in crystallinity during stretching is reasonable due to the rigid backbone of conjugated polymers. However, it remains unclear why the scattering intensity continues to decrease upon strain release.\n(4) The use of Δφ to quantify morphological orientation is well-executed and convincingly correlates with mechanical strain (σ–ε). However, the physical origin of the “X”-shaped contrast in the RSoXS images should be clarified due to the structural complexity of the stained thin film. Meanwhile, how are the φ positions determined?\n(5) The manuscript suggests that crystallite disruption and increased torsion reduce exciton delocalization, affecting optical absorption. Do the authors have supporting evidence or citations for this claim?\n(6) Regarding the J-V curves of stretchable solar cells, how is the device’s active area defined and measured during strain application? This point should be clarified in the experimental section.", "reviewer_4": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts.", "reviewer_5": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts."}, "author_response": "Response to Reviewers’ Comments\nReviewer #1 (Remarks to the Author):\nThe manuscript by Zhong et al. combined NEXAFS, TReXS, and in situ RSoXS to study\nthe microscale to mesoscale structural evolution of a model conjugated polymer\nP(NDI2OD-T2) under mechanical strain. The study reveals a two-step\nresponse mechanism that finally causes the crystal peeling effect under large strain: first, the\nstage of polymer chain orientation and rapid crystallite fracture, and secondly, the\nstage of chain alignment and intrachain torsional deformation. Further, the authors\nstudied P(NDI2OD-T2):PO10 blend systems and makes a correlation between device\nbehavior and morphological disruption resulting from strain-induced meso- and\nmicroscale crystalline destruction investigated from resonant X-ray scatterings. With\nminor revisions and more experiments, this paper is well-suited for publication in\nNature Communications:\n1. In lines 117 to 122, the authors used different mathematical equations to fit the tilt\nangle of the transition dipole moment under 0% and 50% strain. It seems the expression\nused for the 0% strain condition (I = A/3[1+1/2(3cos2θ−1)(3cos2𝛼−1)]) is assuming\na uniaxial orientation distribution where the stacking is face-on predominantly in the\nz-direction but in-plane orientation in the x and y directions is random. While this may\nbe a reasonable assumption for an unstressed system, there are some conjugated\npolymers that have some in-plane alignment before stretching depending on the film\nforming process. The authors should clarify whether this assumption is supported by\nearlier statements in the manuscript, specifically the claim that βₓ equals β at 0% strain,\ny\nand whether such a model is universally valid for the polymer system under study.\nOur response: Thanks for the comment. At 0% strain, we adopted the fitting equation\ncommonly used for samples with three-fold or higher substrate symmetry. The as-cast\nP(NDI2OD-T2) thin films were prepared by spin-coating without any external\nalignment field, and thus in-plane isotropy is expected. We agree with the reviewer that\nthis assumption is not universally valid for all conjugated polymer systems, since\n1\ncertain systems can indeed exhibit in-plane alignment depending on the film-forming\nprocess. To support our assumption in the present case, we added GIWAXS\nmeasurements of the as-cast P(NDI2OD-T2) thin film probed with orthogonal X-ray\ndirections. The results show nearly identical reflections with similar peak positions,\nintensities, and widths, which confirms that βₓ ≈ β before stretching. In addition, we\ny\nrevised the manuscript to clarify that the fitting model at 0% strain is specifically\njustified for the current P(NDI2OD-T2) system and should not be considered a\nuniversal assumption for all conjugated polymers. The corresponding revisions are\nincluded in the revised manuscript.\nMain text (Page 9, line 7):\n“To study the π-π stacking developments during stretch, the films were stretched to the\ndesired strains and then placed on silicon wafers for grazing incidence wide angle X-\nray scattering (GIWAXS) measurements at the beam energy of 10 keV. This\nmeasurement was performed with orthogonal X-ray directions including parallel and\nperpendicular conditions (Fig. S28). The 0% strain thin film shows nearly identical\nreflections (Fig. S29), which support that βₓ ≈ β before stretching as demonstrated by\ny\nNEXAFS.”\nMain text (Page 8, line 28):\n“This angular dependence of NEXAFS intensity is appropriate for samples with three-\nfold or higher substrate symmetry, consistent with the as-cast P(NDI2OD-T2) thin films\nprepared by spin-coating. While some conjugated polymer systems can indeed exhibit\nIP alignment depending on the film-forming process Error! Bookmark not defined., we note that\nthis fitting method should not be considered universally valid for all thin films prior to\nstretching.”\nReference:\n[39] Schuettfort, T., Thomsen, L. & McNeill, C.R. Observation of a Distinct Surface\nMolecular Orientation in Films of a High Mobility Conjugated Polymer. J. Am. Chem.\nSoc. 135, 1092-1101 (2013).\nFig. S29. GIWAXS (a,b) 2D images and (c) averaged I-q curves in IP and OOP\ndirections of P(NDI2OD-T2) thin film under pre-stretched condition.\n2. In line 132, the authors refer to “Area 1” in the NEXAFS analysis, but the\nmethodology used to extract this value is not described in detail. I assumed that the\narea under the curve was calculated directly from a fixed energy window. However, the\nauthors should discuss in SI or method part whether a peak fitting approach was\nconsidered, and whether this method might offer a more accurate representation of the\nelectronic transitions. Authors should also mention, if no fitting is performed, if using\nsuch a peak fitting would change the observed trend or only shift the numerical values.\nOur response: Thanks for the comment. In addition to calculating the integrated\nintensity from a fixed energy window, we also performed peak fitting of the\ncorresponding NEXAFS features based on Gaussian functions. The results from the\npeak fitting analysis are consistent with those obtained by direct integration, showing\nthe same trend in tilt angle changes under strain, with only minor numerical differences\nin the extracted values. To make this clearer, we have added a description of the fitting\n3\ndetails and results in the Supplementary Information and clarified that both approaches\nlead to the same conclusions.\nSupplementary Information:\nFig. S11. Gaussian multi-peak fitting of the π* region in angle-dependent NEXAFS\nspectra of P(NDI2OD-T2) thin films: (a-c) under 0% strain; (d-f) under 50% strain.\nFig. S12. Linear fitting of peak area of angle-dependent TFY NEXAFS spectra,\n4\nobtained from Gaussian multi-peak fitting of P(NDI2OD-T2) thin films at 0% and 50%\nstrain. The parameters a and b correspond to the intercept and slope, respectively.\nFig. S13. Fitted tilt angles (α) derived from Gaussian multi-peak fitting for the peak at\n284.2 eV and the total π* region. The extracted α values for the 50% strain thin film are\n34.6° and 35.2°, representing decreases of 6.4° and 9.5°, respectively, compared to the\npre-stretched film. These results are consistent with those obtained from direct\nintegration (Fig. S10), confirming the same trend in tilt angle reduction under strain,\nwith only minor numerical differences.\n3. The display of figure 2 is confusing. First, the italicized terms “para.” and “perp.”\noverlap with the labels for crystal directions, which create confusion for the reader.\nReplacing these overlapping labels with clearly marked axes or direction arrows would\nenhance readability. Second, the data shown in Figures 2c to e would be more\ninterpretable if the corresponding scattering features were marked using the same color\nto represent. For example, the author can use four different colors (grouped inro two\nmain group, like deep and light shades of blue and red) and denote the corresponding\nscattering feature and also use the corresponding color in figure 2c (instead of plain\nyellow).\nOur response: Thanks for the comment. To avoid overlap between the crystallite\norientation labels and the q-vector directions, we have replaced “para.” and “perp.” with\n“q ” and “q ” to denote the respective q vectors. In addition, to improve the readability\nx y\nof Figures 2c-e, we now use two color groups: (i) deep and light shades of blue to\n5\nindicate the (001) and (100) reflections of crystallites aligned parallel to the stretch\ndirection, and (ii) deep and light shades of red to indicate the (001) and (100) reflections\nof crystallites aligned perpendicular to the stretch direction. The revised Figure 2 is\nshown below.\nFig. 2. Crystallite deformation and destruction. (a) Experimental setup for TReXS\nand RSoXS in transmission geometry, where TReXS probes crystalline packing and\nRSoXS reveals the morphology of the P(NDI2OD-T2) thin films. (b) TReXS I-q curves\naveraged along the q and q directions at 2450 eV and 2476.26 eV. (c) Schematic\nx y\nillustrations of (100) and (001) reflections for parallel and perpendicular crystallites. (d)\nIntensity and packing number (CCL/d-spacing) of the (100) and (001) reflections as\nfunctions of strain for parallel and perpendicular crystals.\n4. The authors have discussed the scattering pattern at 2450 eV in TReXS but had not\noffered further discussion. Given that this energy also excites the (100) peak, it would\nbe useful to know whether the (100) scattering shows the same decay profile under\nstrain as excited by other energy. If the behavior differs, this may suggest distinct\nmechanical responses in the polymer chain in different environments. The authors\nshould include it in either SI or the body.\nOur response: Thanks for the comment. The (100) scattering decay profile under strain\ndoes not show significant differences between 2450 eV and other X-ray energies.\npeak area in both q and q decreases rapidly at small strains and then more gradually\nx y\nat higher strains, with a stronger reduction along q . This behavior is consistent with\nx\nthe observations at 2476.26 eV in TReXS, indicating that the resonant excitation is not\nsensitive to the crystalline deformation. The data at 2450 eV were included to\ndemonstrate that tender X-rays reproduce the same (100) and (001) reflection\ndistributions as hard X-rays (10 keV), thereby validating the use of TReXS for probing\nstrain-induced crystallite deformation. However, we did not perform in situ TReXS\nmeasurements at 2450 eV itself.\nTo further address the reviewer’s point about environment-specific responses, we\ncompared in situ RSoXS measurements at 284.2 eV and 285.0 eV, where the scattering\ncontrast selectively probes different chemical environments. At 284.2 eV, stretching\ninduces a clear “X”-shaped scattering pattern, reflecting fibril aggregate reorientation,\nwhereas at 285.0 eV this feature is absent, suggesting reduced orientation sensitivity\nand a distinct response under strain. These results confirm that while the (100)\ncrystalline scattering exhibits similar strain decay profiles across different X-ray\nenergies, RSoXS contrast variation reveals environment-dependent mechanical\nbehavior. We have added clarifying sentences to the manuscript accordingly.\nMain text (Page 9, line 7):\n“To study the π-π stacking developments during stretch, the films were stretched to the\ndesired strains and then placed on silicon wafers for grazing incidence wide angle X-\nray scattering (GIWAXS) measurements at the beam energy of 10 keV. This\nmeasurement was performed with orthogonal X-ray directions including parallel and\nperpendicular conditions (Fig. S28). The 0% strain thin film shows nearly identical\nreflections (Fig. S29), which support that βₓ ≈ β before stretching as demonstrated by\ny\nNEXAFS. Meanwhile, the (100) peak area in both q and q decreases rapidly at small\nx y\nstrains (10-20%) and then more gradually at higher strains, with a stronger reduction\nalong q (Fig. S30-32, Table S1-S2). This behavior is consistent with the observations\nx\nat 2476.26 eV in TReXS, indicating that the resonant excitation is not sensitive to the\n7\ncrystalline deformation”\nMain text (Page 11, line 19):\n“In situ RSoXS measurements at 285.0 eV during stretching show that “X”-shaped\nfeature disappears (Fig. S44), suggesting reduced orientation sensitivity and further\nsupporting the assignment of scattering contrast origin. Thus, the “X” pattern provides\ndirect information related to the orientation of fibril aggregates. Such results also\nindicates that although the crystalline scattering exhibits similar strain decay profiles\nacross different X-ray energies, RSoXS contrast variation reveals environment-\ndependent mechanical behavior.”\nSupplementary Information:\nFig. S28. (a) Schematic representation of GIWAXS geometries for strained\nP(NDI2OD-T2) thin film, where the scattering images are collected with the incident\nX-ray roughly perpendicular (perp.) or parallel (para.) to the SD. (b) GIWAXS patterns\nof P(NDI2OD-T2) thin films with various strains measured with the incident X-ray\nbeam (b) perpendicular (q ) and (c) parallel (q ) to the SD.\nx y\nx\nbeam perpendicular to the stretch direction) of P(NDI2OD-T2) thin films under\ndifferent strains: (a) 0%, (b) 10%, (c) 20%, (d) 30%, (e) 40%, and (f) 50%.\nFig. S31. Fitting illustration of GIWAXS (100) reflections in the q direction (X-ray\ny\nbeam parallel to the stretch direction) of P(NDI2OD-T2) thin films under different\nstrains: (a) 0%, (b) 10%, (c) 20%, (d) 30%, (e) 40%, and (f) 50%.\nx y\nfor P(NDI2OD-T2) thin films\nFig. S44 In situ RSoXS images of a P(NDI2OD-T2) thin film at different strains during\n(a) stretching and (b) releasing, which are probed at 285.0 eV.\nTable S1. Fitted results of the GIWAXS (100) reflections in the q direction of\nx\nP(NDI2OD-T2) thin film under different strains. FWHM represents full width at half\nmaximum.\nStrain [%] Type Location [Å-1] Amplitude Area FWHM[Å-1]\n10\nTable S2. Fitted results of the GIWAXS (100) reflections in the q direction of\ny\nP(NDI2OD-T2) thin film under different strains. FWHM represents full width at half\nmaximum.\nStrain [%] Type Location [Å-1] Amplitude Area FWHM[Å-1]\n5. In line 225, the authors argue that the RSoXS results align with AFM data; however,\nAFM images are only provided for samples strained to 50%, which is significantly\nlarger than that of the RSoXS experiment (30%). It remains unclear at what strain level\nthe “X-shaped” anisotropy emerges in AFM. Including an FFT analysis of AFM images\nat intermediate strains, such as around 30%, would help confirm the structural\nevolution pathway supported by the X-ray scattering data. Additionally, a comparison\nof domain sizes derived from AFM and RSoXS would be valuable. Although the feature\nis not seemingly a peak, the authors still may select some turning points in the small\nangle range if plot a figure on log-log scale.\nOur response: Thanks for the comment. To clarify the structural evolution with strain,\nwe have now included additional AFM images at 0%, 10%, 20%, and 30% strain levels,\nalong with their corresponding FFT analyses. In the as-cast film (0% strain), N2200\nexhibits a fibrillar morphology that assembles into clusters, with slight anisotropy\nobservable in the FFT. As strain increases, the fibrillar features become less distinct,\nand bundle-like domains grow in size. An “X”-shaped anisotropic signature begins to\n11\nemerge in the FFT around 20% strain, which becomes more pronounced by 30% strain.\nThis progression aligns temporally with the development of the “X”-shaped scattering\npattern observed in RSoXS, though the two techniques probe different structural levels\nwith different sample configuration.\nTo quantify the morphology evolution, we extracted the power spectral density\n(PSD) from each AFM height image. The turning point in the log-log PSD plots\nprovides an estimate of the characteristic domain size, which increased from 2.2 μm to\n3.9 μm, 4.5 μm, and 4.8 μm for films strained at 0%, 10%, 20%, and 30%, respectively.\nThese results reflect the micrometer-scale aggregation process captured by AFM.\nRSoXS probes the nanometer sized features, which is the structure within the\nbundle domain. There is no turning point in I-q curves in the log-log scale. Thus, we\nanalyzed the characteristic size with Debye-Bueche model, (𝐼(𝑞))−1/2 =\n𝐾(𝑎3𝑄)−1/2(1+𝑎2𝑞2), where K is a constant, Q the scattering invariant, and a the\ncorrelation length (CL). The CL initially decreased by about 50 nm at low strain (<5%),\nlikely due to transverse compression, and then increased from ~150 nm to 185 nm as\nstrain further increased. This trend reflects reorganization and alignment at the\nnanoscale, which complements the micro-scale aggregation observed by AFM.\nAlthough AFM and RSoXS probe different length scales (micrometer vs.\nnanometer), both consistently reveal a strain-induced transition from fibrillar networks\nto globally oriented, aggregated domains. These complementary results support our\nconclusion that the anisotropic “X”-shaped scattering in RSoXS arises from the same\nfibrillar alignment observed by AFM. To clarify this, we added the following sentences\nin the manuscript.\nMain text (page 10, line 15):\n“For the 50% strained film, large-scale fibril aggregates with their long axes aligned\ntoward the SD are seen, showing a fast Fourier transform (FFT) pattern with an “X”\nshaped anisotropy, which resulted from the progressive development in both size and\norientation of the aggregates (Fig. S34). This reflects the formation of a statistically\noriented structure induced by strain.”\n“The corresponding characteristic domain size was also evaluated using the Debye-\nBueche model, which shows a trend of structural reorganization that complements the\nmicro-scale evolution observed by AFM (Fig. S41-S43).”\nSupplementary Information:\nFig. S34. AFM height images and corresponding FFT images of P(NDI2OD-T2) thin\nfilms under different stretch strains: (a) 0%, (b) 10%, (c) 20%, and (d) 30%. The stretch\ndirection is in the horizontal direction of the images.\nshown in Fig. S34) of P(NDI2OD-T2) thin films under different stretch strains: (a) 0%,\n(b) 10%, (c) 20%, and (d) 30%. The dotted lines highlight the turning points giving an\nestimate of the characteristic domain size, which increased from 2.2 μm to 3.9 μm, 4.5\nμm, and 4.8 μm for films strained at 0%, 10%, 20%, and 30%, respectively. These\nresults reflect the micrometer-scale aggregation process captured by AFM.\nFig. S42. (a-b) The I-q curves, averaged at two azimuthal angles of the X-shaped signal\nin RSoXS, i.e., φ (Sector 1) and φ (Sector 2). Inset shows the sector averaging details\n1 2\nof Sector 1 and Sector 2. (c-d) The I-1/2-q2 curves plotted following the Debye-Beuche\nmodel, (𝐼(𝑞))−1/2 = 𝐾(𝑎3𝑄)−1/2(1+𝑎2𝑞2) , where K is a constant, Q is the\nscattering invariant, and a is the correlation length, which described the average domain\nsize. The correlation lengths were fitted from the linear region at low-q range (0.0012-\n0.003 Å-1). (e) Representative fit region (0.001-0.003 Å-1) of the I-1/2-q2 curves. (f)\nRepresentative linear fits of the I-1/2-q2 curves, where the resultant correlation lengths\n14\nequal to the values of (slope/intercept)-1/2 [Å].\nFig. S43. Correlation length as a function of film strain of the P(NDI2OD-T2) thin film.\nThe correlation length in both section 1 and section 2 first decreased by ~50 nm during\nthe initial strain (<5%), likely due to transverse compression, and subsequently\nincreased from ~150 nm to ~185 nm as strain progressed. This trend complements the\nmicro-scale aggregation observed by AFM (Fig. S41).\n6. The UV-Vis spectrum of the blend film is unclear. The spectral features that appear\nin the blend film is significantly shifted compared to neat P(NDI2OD-T2). It is unclear\nwhether this shift originates from the P(NDI2OD-T2) component or from PO10.\nThe authors should include a strain-dependent UV-Vis spectrum of neat PO10 films to\ndetermine the individual contributions of each material in the blend. This would allow\nthem to confidently attribute the observed blue shift in the blend to a specific molecular\norigin.\nOur response: Thanks for the comment. To clarify the origin of the absorption shift in\nthe blend film under strain, we measured the strain-dependent UV-vis absorption\nspectra of neat PO10 films. The result indicates that PO10 exhibits a main absorption\npeak at 578 nm within 400-700 nm, which undergoes only a slight blue shift and\nnarrowing upon stretching. In contrast, the blue shift observed in the blend films\n15\nprimarily occurs in the 720-800 nm region, consistent with the absorption shift of neat\nP(NDI2OD-T2). These results confirm that the strain-induced spectral shift in the\nblends arises from the P(NDI2OD-T2) component. To clarify this point, we revised the\nfollowing sentence in the manuscript.\nMain text (Page 14, line 26):\n“Similar to the neat thin film, the P(NDI2OD-T2) component in the blend showed blue-\nshifted absorption, during stretch from 0% to 30% strain, while PO10 contributes only\nminor changes within its own absorption range (Fig. 4e and Fig. S52).”\nFig. S52. Normalized absorption spectra of the neat PO10 thin film under different\nstretch strains with a film-on-elastomer sample.\n7. While the manuscript connects morphological evolution to device performance,\nthe authors could elaborate further on how the relatively subtle strain-dependent optical\nchanges (only a few nm shifts in absorption peaks) influence charge generation,\ntransport, or recombination in OPV devices.\nOur response: Thanks for the comment. Although the strain-induced absorption\nchanges are relatively small (~10 nm peak shift and a decreased shoulder intensity at\n~790 nm), they signify meaningful variations in electronic coupling and molecular\npacking order. These arise from chain conformational adjustment, orientation, and\npartial crystallite disruption, which collectively reshape the morphology and thus\n16\ninfluence charge generation, transport, and recombination in the devices. To clarify\nthese correlations, we performed measurements of photocurrent density-effective\nvoltage (J -V ) curves, photo-CELIV, and light intensity dependent J-V curves for\nph eff\nstretchable OPV devices under different strains (0%, 10%, 20%, and 30%).\nFrom the J -V analysis, the exciton dissociation and charge collection efficiency,\nph eff\nP(E,T), decreased slightly from 98.86% for pre-stretched condition to 97.93%, 96.85%,\nand 95.53% at 10%, 20%, and 30% strain, respectively, indicating reduced charge\ngeneration under strain. The photo-CELIV results further revealed a gradual decline in\ncharge mobility from 6.59×10-5 to 4.13×10-5 cm2 V-1 s-1 as the strain increased from 0%\nto 30%. Moreover, the light intensity (P ) dependence of V can be described by\nlight OC\n𝐾𝑇\n𝑉 ∝ 𝑛 ln(𝑃 ) , yielding n values of 1.25, 1.42, 1.46, and 1.58 for devices\nOC light\n𝑞\nstrained at 0%, 10%, 20%, and 30%, respectively, suggesting progressively enhanced\ntrap-assisted recombination. Coupled with the AFM-IR images, the results demonstrate\nthat even the subtle strain-dependent optical shifts, originating from local order-\ndisorder transitions to morphological aggregation, translate into measurable\ndeterioration in charge generation and transport, along with increased recombination\nlosses. To clarify this, we have revised the following section in the revised manuscript.\nMain text (Page 15, line 1):\n“Such decline was primarily due to the decrease in both J and FF (Fig. 4f), which\nSC\ncorrelate with charge generation, transport, and recombination in the devices.\nThe exciton dissociation and charge collection efficiency, P(E,T), extracted from\nphotocurrent density-effective voltage (J -V )curves, decreased from 98.86% for pre-\nph eff\nstretched condition to 97.93%, 96.85%, and 95.53% at 10%, 20%, and 30% strain,\nrespectively, indicating reduced charge generation (Fig. S54). Photo-CELIV analysis\nrevealed a concurrent decline in charge mobility from 6.59×10-5 to 4.13×10-5 cm2 V-1\ns-1 (Fig. S55). The light intensity (P ) dependence of V , described by 𝑉 ∝\nlight OC OC\n𝐾𝑇\n𝑛 ln(𝑃 ) , yields n values of 1.25, 1.42, 1.46, and 1.58 as strain increased,\nlight\n𝑞\nsignifying enhanced trap-assisted recombination (Fig. S56). To further elucidate this,\nmorphological changes upon thin film stretched by nano-infrared spectroscopy atomic\nforce microscopy (AFM-IR). As shown in Fig. 4g-i, P(NDI2OD-T2) domains probed\nwith 843 cm-1 exhibit nanofibrillar network morphology in the as-cast blend film (Fig.\nS56), which is retained when 10% strain was applied, but transformed into large\naggregates at 30% strain. This morphological transformation arises from strain-induced\nchain orientation, intrachain torsion, and crystalline destruction, leading to local order-\ndisorder transitions. These combined absorption, electrical, and morphological results\ndemonstrate that strain-induced structural evolution directly deteriorates charge\ngeneration and transport while promoting recombination losses, ultimately accounting\nfor the observed performance degradation. To assess the effect of cyclic strain,\nPO10:P(NDI2OD-T2) films on PDMS were stretched at 20% for 300 cycles. AFM\nimages show that the fibrillar morphology gradually blurred with repeated cycling (Fig.\nS58), causing some decline in device performance; however, the device retained over\n80% of its initial PCE (Fig. S59 and Table S4). These results indicate that the overall\nnanostructure and device performance remain largely preserved under moderate\nmorphological relaxation.”\nSupplementary Information:\nFig. S54. J -V curves of PO10:P(NDI2OD-T2) based stretchable OPV device under\nph eff\ndifferent strains.\nPO10:P(NDI2OD-T2) based stretchable OPV device under different strains.\nlight OC\nstretchable OPV device under different strains.\n8. The authors could expand the discussion of prior works on strain-induced\nmorphology changes in conjugated polymers. For example, Wu et al.\n(https://www.nature.com/articles/s41467-023-44099-w) introduced the “relative\n19\nstretchability” (rS, the ratio of DR and rDoC) to quantitatively analyze the strain\ninduced morphological evolution and crystal breakage. Comments on whether your\nsystem follows similar entropic (chain alignment, from DR) and enthalpic (crystallite\namorphization/breakage, from rDoC) dissipation pathways would be helpful for the\nreaders. In addition, Zheng et al. (https://pubs.acs.org/doi/full/10.1021/jacs.2c00072)\nsummarized the molecular and morphological processes expected under strain, for\nexample, chain alignment and strain-induced backbone planarization through\nmonomer torsional rotation and then outlined how such processes map onto mobility\nretention. Since this manuscript shows direct evidence on monomer rotation to\nplanarize the chain under strain, highlighting how it supports or refines the\nassumptions in this review paper would be helpful.\nOur response: Thanks for the comment. We have now expanded the discussion of\nprior works to better contextualize our findings. Wu et al. (Nat. Commun. 2023, 14,\n8382) proposed the concept of relative stretchability (rS), which integrates the effects\nof chain alignment (DR) and crystallite breakage (rDoC) to describe dual entropic and\nenthalpic dissipation pathways during strain. In our system, both effects were directly\ncaptured, where chain alignment and crystallite destruction were evidenced by\nNEXAFS and TReXS/GIWAXS, respectively. This observation aligns with the dual-\ndissipation framework of Wu et al. but further refines it by revealing a strain-dependent\ntemporal evolution. In relation to the review by Zheng et al. (J. Am. Chem. Soc. 2022,\n144, 4699), we note that their discussion of strain-induced backbone behavior focused\non orientation and rotation. Our in situ and MD simulation data complement these\nfindings: we observe increased dihedral torsion angles under strain, implying enhanced\nrotational freedom. These results clarify that strain-induced molecular responses are\nstrongly dependent on side-chain architecture and packing constraints. To clarify this,\nwe added the following points in the Discussion section of the manuscript.\nMain text (Page 15, line 28):\n“Recent studies have identified entropic chain alignment and enthalpic crystallite\nbreakage as multimodal energy dissipation pathways for mechanical robustness [55].\nevolution further revealed [56]. Consistent with a recent review highlighting the role of\nchain alignment and torsional rotation in mechanical-electronic coupling, the in situ\nand simulation data show increased torsional freedom under strain, likely governed by\nbackbone architecture.”\nReferences:\nmodal energy dissipation and high relative stretchability. Nat. Comm. 14, 8382 (2023).\npolymer electronics. Nat. Comm. 12, 5701 (2021).\nThe manuscript presents a fundamental investigation into the structural evolution of a\nconjugated polymer under tensile stress, utilizing a combination of NEXAFS, TReXS,\nand RSoXS techniques. The studied material is further demonstrated in a fully\nstretchable photovoltaic device. Overall, the manuscript is well written, and the\ncomprehensive suite of characterizations are clearly correlated to provide a coherent\nmesoscale picture that advances the fundamental understanding of intrinsically\nstretchable organic semiconductors during dynamic mechanical deformation. I\nrecommend this work for publication, provided that the following concerns and\nquestions are satisfactorily addressed:\nQ1: In comparing the simulated NEXAFS spectra of P(NDI2OD-T2) with the\nexperimental results in Figure 1c, there remain noticeable discrepancies, particularly\naround 285.1 eV. What might be the origin of these mismatches? Additionally, the major\nand minor tick marks on the x-axis of Figure 1c are not clearly visible and should be\nimproved for clarity.\nOur response: Thanks for the comment. The slight mismatch in peak positions between\nthe simulated and experimental spectra (e.g., 285.0 eV in experiment vs. 284.8 eV in\nsimulation) mainly arises from intrinsic limitations of first-principles NEXAFS\ncalculations. In our case, the simulations were based on the simplified BT-NDI\nstructural unit rather than the full P(NDI2OD-T2) polymer, so long-range conjugation\nand solid-state effects are not fully captured. Moreover, the PBE-GGA functional is\nknown to underestimate bandgaps, which compresses the spectra along the energy axis\neven after applying a dilation factor, while the lack of environmental interactions such\nas interchain packing and substrate effects in the vacuum-based model can also shift\nresonance energies. Finally, the applied alignment (284.5 eV for the C K-edge) and\nGaussian broadening inevitably introduce small offsets. Overall, the deviation of ~0.2-\n0.3 eV is within the typical range reported for DFT-based NEXAFS simulations and\ndoes not affect the spectral assignments or conclusions of NEXAFS analysis. To clarify\nthis, we added the following sentences in the Method section.\n“We note that despite these corrections, small residual discrepancy of ~0.2-0.3 eV\nremains, as the simplified molecular model cannot fully capture the long-range\nconjugation, and the vacuum-based calculations cannot account for solid-state\ninteractions. These deviations are commonly accepted in DFT-based NEXAFS\nsimulations and do not affect the reliability of spectral assignments [34,37,62]”\nWe also revised the Figure 1c to clearly show the major and minor ticks as follows.\nFig. 1. (c) Experimental (TFY) and simulated NEXAFS spectra of P(NDI2OD-T2) in\nthe C 1s→π* transitions range (283.6-286.0 eV), and the contributions from the NDI\nand BT units.\nQ2: Regarding the free-standing experimental setup, as the majority of the film remains\nsupported by the PDMS substrate, could this result in differences between the local\nstrain experienced near the PDMS opening and the overall applied strain? How is the\nstrain distributed across the film, and is there any quantification or mapping of\npotential strain gradients?\nOur response: Thanks for the comment. Finite element analysis (FEA) of the PDMS-\nsupported freestanding geometry (a 4 mm diameter circular hole at the center of the\nPDMS) show that, under a perfect-bonding assumption, the unsupported thin film can\nexperience high local strain amplification near the hole edge: for a 30% applied overall\n23\nstrain of the sample, the model predicts a local strain of 79% at the edge, due to stress\nconcentration. The simulation also shows that the central region of the film exhibits\nuniaxial, spatially uniform in-plane strain, which is over an area much larger than the\nX-ray beam footprint, confirming that the RSoXS/TReXS measurements probe a\nuniaxial strain state.\nIn the FEA, no slip between the film and PDMS was assumed; however,\nexperimental observations indicate that no macroscopic bending occurs in practice,\nsuggesting that strain amplification is partially relaxed and the actual film strain is\ncloser to the nominal applied strain. This setup therefore provides the most practical\napproach for realizing uniaxial stretching of thin polymer films during in situ X-ray\nscattering experiments. We report the nominal applied strain in the manuscript while\nnoting this limitation and include the FEA maps in the Supplementary Information for\ntransparency.\nMain text (Page 7, line 17):\n“The PDMS-supported freestanding configuration provides a practical approach to\nimpose uniaxial strain on thin polymer films during in situ X-ray scattering experiments,\nas validated by simulated strain and stress maps (Fig. S17 and S18).”\nSupplementary Information:\nFig. S17. Finite element analysis (FEA): (a) Simulated models of the P(NDI2OD-T2)\n24\nthin film on a PDMS substrate with a central circular opening. (b) Simulated strain\nmap, showing that the overall strain applied to the PDMS substrate (30%) leads to local\nstrain amplification (79%) near the hole edge. The PDMS substrate and central hole\nhave initial lengths of 15 cm and 0.4 cm, respectively. The inset numbers are in units\nof μm. (c) Side-view strain map, showing significant macroscopic bending under the\nperfect bonding assumption between the PDMS and P(NDI2OD-T2) film. (d)\nSimulated stress map, showing the stress concentration at the central hole, where the\ncentral 0.2 mm region of the unsupported film exhibits a nearly uniform uniaxial strain.\nNotes: FEA simulations of the samples for TReXS and RSoXS were performed using\nAbaqus. The PDMS substrate was modeled as a rectangular block with dimensions of\n10 cm × 15 cm × 0.5 cm, containing a circular opening of 4 mm in diameter at the center.\nA thin P(NDI2OD-T2) film (thickness: 0.2 μm) was conformally placed on the top\nsurface of the PDMS, fully covering the hole region. A perfect bonding condition (no\nslip) was assumed between the P(NDI2OD-T2) film and PDMS, and the two layers\nwere meshed using shared nodes. To simplify the computation while retaining the\nstrain-distribution trend, both materials were described by linear elastic constitutive\nrelations. The elastic moduli were set as 1.74 MPa for PDMS and 63.72 MPa for\nP(NDI2OD-T2). A static step analysis was used, where one side of the model was fully\nfixed while a 30% tensile strain was applied along the x-axis to the opposite side.\nFig. S18. The photos showing a P(NDI2OD-T2) thin film with 0% and 30% strain on\nthe tensile tester. Such photos indicate no macroscopic bending occurs in practice,\nwhich reduces strain amplification and brings the actual film strain closer to the applied\nnominal value.\ngeometry, how is it ensured that the X-ray beam does not also penetrate and contribute\nsignal from the PDMS substrate? What is the beam size relative to the size of the\nunsupported region (hole)? Additionally, since the scattering intensity is correlated\nwith film thickness—and considering that the film thickness is likely reduced during\nstretching—is this variation reflected or corrected for in the analysis (for example, in\nFigure 2d)?\nOur response: Thanks for the comment. The PDMS substrate contains a 4 mm\ndiameter hole, which is substantially larger than the incident X-ray beam size (20 × 200\nμm2 for TReXS and 200 × 200 μm2 for RSoXS), ensuring that the beam fully passes\nthrough the unsupported film region. In addition, no diffraction contribution from\nPDMS is observed: PDMS typically shows a reflection near q~0.40 Å-1, whereas our\nTReXS data remain flat in this range, confirming that the measured scattering originates\nsolely from the active layer.\nFor RSoXS measurements, the PDMS substrate thickness (0.5-0.8 mm) effectively\nblocks the C K-edge soft X-rays, preventing beam penetration (hundreds of nanometers\nfor attenuation length at ~285 eV) and background scattering from the substrate.\nTherefore, the signals collected in both TReXS and RSoXS arise exclusively from the\npolymer thin films.\nRegarding the scattering intensity in TReXS analysis, we applied normalization to\naccount for possible variations in film thickness upon stretching. Specifically, the\nscattering intensity was normalized to the intensity at q position of 0.37 Å-1 under 0%\nstrain, ensuring consistent comparison across all strain levels. To clarify this, we added\nthe following notes in the Method section.\nMethod:\nTReXS\n“The P(NDI2OD-T2) thin films were transferred onto PDMS substrates with a central\ncircular hole (4 mm in diameter), as illustrated in Fig. S3b. The X-ray beam size was\n20 μm × 200 μm, ensuring that the beam fully passed through the unsupported film\n26\nregion.”\n“To ensure consistent comparison across all strain levels, the scattering intensity was\nnormalized to the value at q = 0.37 Å-1 at 0% strain. No diffraction contribution from\nPDMS was observed; PDMS typically shows a reflection at q~0.40 Å-1 (Fig. S61),\nwhereas our TReXS data remain flat in this range, confirming that the measured scattering\noriginates solely from the polymer thin film.”\nNEXAFS and RSoXS\n“The sample preparation followed the same procedure as shown in Fig. S3b. The beam\nsize is 200 μm × 200 μm, ensuring that the beam fully passed through the unsupported\nfilm region.”\nFig. S61. WAXS (a) 2D image and (b) I-q curves averaged in the parallel (para.) and\nperpendicular (perp.) directions.\nQ4: While optical microscopy determined the crack-onset of P(NDI2OD-T2) to occur\nat strains above 50% (Fig. S2), the macroscopic stress-strain curve exhibits mechanical\nfailure before reaching 35% strain (Fig. 4a). Can the authors clarify the origin of this\napparent discrepancy?\nOur response: Thanks for the comment. The apparent discrepancy arises because the\ntwo measurements probe different mechanical conditions and failure criteria. The\noptical microscopy test characterizes the crack-onset strain (COS) of a thin film\nsupported on a compliant PDMS substrate, where strain is partially relaxed through the\nsubstrate, and visible surface cracks appear only after local stress concentration has\n27\naccumulated. In contrast, the macroscopic stress-strain curve in Fig. 4a was measured\nfrom a free-standing film, where mechanical failure is defined by a sudden stress drop\ndue to complete fracture. Therefore, the supported film can appear to sustain higher\nnominal strain before visible cracking, even though local strain in the film may already\nbe high. Such differences between supported and free-standing geometries have also\nbeen reported in stretchable semiconducting thin films. To clarify this distinction, we\nadded the following sentence in the manuscript.\nMain text (Page 13, line 23):\n“Note that the crack-onset strain from optical microscopy (supported film) is not\ndirectly comparable to the fracture strain from tensile testing (free-standing film), as\nthe former reflects delayed surface cracking under strain relaxation by the substrate,\nwhereas the latter represents complete mechanical failure of the film.”\nQ5: The observation of a two-stage strain response corresponding to distinct stress\ndissipation mechanisms in P(NDI2OD-T2) is intriguing. To what extent is this behavior\ngeneralizable to other conjugated polymers, such as PO10, already studied in this work?\nOur response: Thanks for the comment. The observed two-stage strain response in\nP(NDI2OD-T2) is closely related to its mechanical behavior, including yielding and\nhardening, as reflected in the stress-strain curves. This behavior arises from how stress\nis dissipated within conjugated polymers, which varies with molecular structure and\nsolid-state morphology. In this work, P(NDI2OD-T2) is a highly crystalline polymer,\nwhereas PO10 exhibits much lower crystallinity. Therefore, it is reasonable to expect\ndistinct strain-response characteristics for PO10. Meanwhile, to observe a similar two-\nstage morphological evolution using in situ RSoXS, it is necessary to identify an\nappropriate soft X-ray energy that can enhance the scattering contrast, particularly the\norientational contrast. However, such an energy is not always experimentally accessible.\nThe use of P(NDI2OD-T2) here thus serves as an example to demonstrate how a\ncombination of X-ray scattering and spectroscopy techniques can reveal multi-scale\nstructural evolution in conjugated thin films and its correlation with electronic\n28\nperformance.\nIndeed, we are currently extending this study to polymers with lower or near-\namorphous crystallinity using in situ RSoXS and molecular dynamics simulations for\nstretchable organic solar cells and organic light-emitting diodes. Preliminary results\nindicate that strain-induced orientation occurs across molecular to mesoscale structures,\naccompanied by torsional conformational changes, which are consistent with the\nobserved in this work. We are continuing this analysis and plan to report these findings\nin a future publication. To clarify this point, we have revised the following sentences in\nthe revised manuscript.\nMain text (Page 17, line 11):\n“Although the two-step morphological response is demonstrated using the highly\ncrystalline P(NDI2OD-T2), its occurrence is expected to depend on the molecular\nstructure and thin-film morphology. Ongoing studies on polymers with lower\ncrystallinity or near-amorphous states aim to extend these insights and examine the\ngenerality of the observed behavior. Nevertheless, our findings establish a mechanistic\nframework for understanding how conjugated polymer systems maintain device\nfunctionality under large deformations.”\nQ6: The demonstration of robust photovoltaic device performance under up to 50%\napplied strain is impressive. Is it possible to directly characterize the charge carrier\nmobility of P(NDI2OD-T2) under strain? Such measurements would provide important\ninsight into the relationship between mechanical deformation, structural evolution, and\ncharge transport properties.\nOur response: Thanks for the comment. To further elucidate the relationship between\nmechanical deformation and charge transport, we have evaluated the charge carrier\nmobility of the stretchable OPV device under strain using photo-CELIV measurements.\nThe mobility decreased from 6.59×10-5 at 0% strain to 4.13×10-5 cm2 V-1 s-1 at 30%\nstrain. We have also evaluated the charge generation and recombination by photocurrent\ndensity-effective voltage (J -V ) curves and light intensity dependence of V ,\nph eff OC\n29\nrespectively. Coupled with the AFM-IR images, the results collectively indicate that the\nstrain-induced absorption shifts and structural evolution, from local order-disorder\ntransitions to morphological aggregation, translate into measurable deterioration in\ncharge generation and transport, accompanied by enhanced recombination losses. To\nclarify this, we revised the following section in the manuscript.\nMain text (Page 15, line 3):\n“The exciton dissociation and charge collection efficiency, P(E,T), extracted from\nphotocurrent density-effective voltage (J -V )curves, decreased from 98.86% for pre-\nph eff\nstretched condition to 97.93%, 96.85%, and 95.53% at 10%, 20%, and 30% strain,\nrespectively, indicating reduced charge generation (Fig. S54). Photo-CELIV analysis\nrevealed a concurrent decline in charge mobility from 6.59×10-5 to 4.13×10-5 cm2 V-1\ns-1 (Fig. S55). The light intensity (P ) dependence of V , described by 𝑉 ∝\nlight OC OC\n𝐾𝑇\n𝑛 ln(𝑃 ) , yields n values of 1.25, 1.42, 1.46, and 1.58 as strain increased,\nlight\n𝑞\nsignifying enhanced trap-assisted recombination (Fig. S56). To further elucidate this,\nmorphological changes upon thin film stretched by nano-infrared spectroscopy atomic\nforce microscopy (AFM-IR). As shown in Fig. 4g-i, P(NDI2OD-T2) domains probed\nwith 843 cm-1 exhibit nanofibrillar network morphology in the as-cast blend film (Fig.\nS57), which is retained when 10% strain was applied, but transformed into large\naggregates at 30% strain. This morphological transformation arises from strain-induced\nchain orientation, intrachain torsion, and crystalline destruction, leading to local order-\ndisorder transitions. These combined absorption, electrical, and morphological results\ndemonstrate that strain-induced structural evolution directly deteriorates charge\ngeneration and transport while promoting recombination losses, ultimately accounting\nfor the observed performance degradation.”\nSupplementary Information:\nph eff\ndifferent strains.\nPO10:P(NDI2OD-T2) based stretchable OPV device under different strains.\nlight OC\nstretchable OPV device under different strains.\nQ7: For real-world applications, stretchable semiconductors will undergo repeated\ncyclic stretching rather than a single extension. How do the authors anticipate that\ncyclic mechanical fatigue will influence the nanostructure and electrical performance\nover time? Are there any preliminary results or plans to study such effects?\nOur response: Thanks for the comment. To assess the influence of cyclic mechanical\ndeformation on nanostructure and device stability, PO10:P(NDI2OD-T2) blend films\nwere transferred onto PDMS substrates and subjected to cyclic stretching at 20% strain\nfor 100 and 300 cycles. The initially well-defined fibrillar morphology became\nprogressively blurred with increased surface roughness after repeated stretching,\nindicating partial disruption of the nanofibrillar network. Correspondingly, the\nstretchable OPV devices exhibited a gradual decline in performance but retained over\n80% of its initial PCE after 300 strain cycles, demonstrating good mechanical durability.\nThese results suggest that while cyclic strain introduces moderate morphological\nrelaxation and roughening, the overall nanostructure integrity and device function\nremain largely preserved. Further studies are planned to systematically probe the\nfatigue-induced morphological evolution and charge transport behavior under more\n33\nprolonged cyclic deformation. To clarify this, we added the following sentences in the\nmanuscript.\nMain text (Page 15, line 19):\n“To assess the effect of cyclic strain, PO10:P(NDI2OD-T2) films on PDMS were\nstretched at 20% for 300 cycles. AFM images show that the fibrillar morphology\ngradually blurred with repeated cycling (Fig. S58), causing some decline in device\nperformance; however, the device retained over 80% of its initial PCE (Fig. S59 and\nTable S4). These results indicate that the overall nanostructure and device performance\nremain largely preserved under moderate morphological relaxation.”\nSupplementary Information:\nFig. S58. AFM height images of PO10:P(NDI2OD-T2) thin films on PDMS substrates\nunder (a) pre-stretched condition, (b) after 100 cycles at 20% strain, and (c) after 300\ncycles at 20% strain.\nPO10:P(NDI2OD-T2) thin films under pre-stretched condition, after 100 cycles at 20%\nstrain, and after 300 cycles at 20% strain.\nTable S4. Device parameters of stretchable OPV device based on PO10:P(NDI2OD-\nT2) thin films under cyclic stretch at 20% strain.\nCycles V [V] J [mA cm-2] FF [%] PCE [%] Normalized PCE\nOC SC\n0 0.855 13.91 55.67 6.62 1\n100 0.813 14.47 51.73 6.09 0.92\n300 0.797 13.94 50.50 5.61 0.85\nQ8: The molecular weight of the P(NDI2OD-T2) is very high, 131k; providing the\nP(NDI2OD-T2) GPC traces and the calibration curves will be helpful. Besides,\nmaterial information (synthetic procedure, molecular weight) on PO10 and PNDIT-\nF3N is missing.\nOur response: We thank the reviewer for this valuable suggestion. We have included\nthe GPC molecular-weight distribution of P(NDI2OD-T2) in the Supplementary\nInformation (Fig. S60). The key GPC parameters are provided (Retention time = 24.287\nmin; M = 131 kDa; M = 178 kDa; Dispersity = 1.356). Molecular weights are reported\nn w\nas polystyrene (PS) equivalents. The GPC experimental details in the Method section\nhave also been revised.\nFig. S60. Molecular weight distribution and cumulative curves of P(NDI2OD-T2).\naveraged molecular weight of 24.4 kDa and dispersity of 2.15 [Macromolecules 56,\n8928-8938 (2023)]. The number-average molecular weight of PNDIT-F3N was\ndetermined as 12.3 kDa by a Waters GPC 2410 with tetrahydrofuran (THF) as the eluent\nusing linear polystyrene standards. Such key information has been added in the Method\nsection of the manuscript as follows.\nMethod:\n“P(NDI2OD-T2) was synthesized according to a reported procedure [61]. Its molecular\nweight was characterized by high-temperature gel permeation chromatography (GPC)\non a Tosoh EcoSEC HT system equipped with three TSKgel GMHhr-H(S) HT columns\nin series, using 1,2,4-trichlorobenzene as the eluent at 135 °C. Calibration was carried\nout against narrow-dispersity (Đ) polystyrene standards. The number-average\nmolecular weight (M ) and dispersity of P(NDI2OD-T2) were determined to be 131\nn\nkDa and 1.356, respectively. The corresponding molecular-weight distribution curve is\nshown in Fig. S36. The polymer PO10 was synthesized as described in our previous\nwork [52], with M = 24.4 kDa and Đ = 2.15. The M and Đ of PNDIT-F3N were\nn n\ndetermined to be 6.25 kDa and 1.34 by GPC (Waters 2410) using tetrahydrofuran (THF)\nas the eluent and linear polystyrene standards for calibration.”\n[52] Li, K. et al. 8.0% Efficient all-polymer solar cells based on novel starburst polymer\nacceptors. Sci. China Chem. 61, 576-583 (2018).\n[61] Luo, X. et al. Intrinsically Stretchable Organic Photovoltaic Thin Films Enabled\nby Optimized Donor–Acceptor Pairing. Macromolecules 56, 8928-8938 (2023).\nThis manuscript presents a novel and comprehensive approach to study the multiscale\nstructural evolution of conjugated polymer thin films under strain, using a powerful\ncombination of advanced in situ X-ray techniques (NEXAFS, TReXS, RSoXS). The study\nprovides important mechanistic insights into stress-dissipation pathways, including\nchain orientation changes, intrachain torsion, crystallite disruption (via “peeling” and\n“slippage” mechanisms), and a “two-step” morphological reorganization. Such\nstructural changes are successfully connected to mechanical and optoelectronic\nproperties. The experimental design is rigorous, the correlative multimodal\nmethodology is state-of-the-art, and the results offer valuable guidance for the rational\ndesign of durable stretchable electronics. Given the novelty of the characterization\nmethod and the depth of the mechanistic insights, I strongly recommend publication\nafter minor revisions.\n(1) While the title emphasizes the study of structural evolution on organic photovoltaic,\nthe Introduction does not clearly articulate how key morphological changes under\nstrain impact photovoltaic metrics. A more explicit discussion connecting the multiscale\nstructural changes to performance-limiting mechanisms in organic solar cells would\nenhance the impact of the work.\nOur response: Thanks for the comment. In the revised manuscript, we have expanded\nthe Introduction to explicitly connect strain-induced morphological changes with\nperformance-limiting mechanisms in organic solar cells. In particular, we now highlight\nhow chain alignment and entanglement dissipate stress and how ductile polymer\nnetworks can shield brittle acceptor crystallites, thereby mitigating efficiency losses\nunder high strain. To provide broader context, we also referenced reported mechanisms\nof strain-induced structural evolution in stretchable OFETs and OLEDs, illustrating\nhow multiscale morphological changes affect charge transport and optoelectronic\nproperties across different device types. These additions aim to present a more\ncomprehensive picture of structure-property relationships under mechanical\ndeformation. The following sentences were added to the Introduction section:\n“Nevertheless, conjugated polymers have enabled intrinsically stretchable electronics\nby molecular design and microstructural control. In organic field-effect transistors\n(OFETs), high-molecular-weight polymers shows retained crystalline domains and\nreversible chain alignment under >100% strain, thereby preserving electrical\nperformance [19]. Random terpolymer design can suppress crystallinity and long-range\norder, while maintaining short-range ordered aggregates, resulting in high mobility with\nimproved stretchability [20]. In organic photovoltaics (OPVs), polymer-polymer blends\noutperform polymer-small molecule blends because gradual chain alignment and\nentangled networks dissipate stress while sustaining charge transport pathways [21].\nThe introduction of high-molecular-weight conjugated polymers offers soft, amorphous\nregions with entangled and tie chains for stress dissipation, shielding the brittle acceptor\ncrystallites and mitigating efficiency loss [22]. In organic light-emitting diodes\n(OLEDs), plasticizer incorporation transforms brittle conjugated polymers into fibrillar\nmorphologies, resolving the conductivity-stretchability trade-off [23]. Despite these\nadvances, direct in situ probing of the multiscale structural changes in conjugated\npolymer thin films, and how they collectively govern stress dissipation from the\nmolecular to the mesoscale remains an outstanding challenge.”\nReferences:\n19. Wu, H.-C. et al. Highly stretchable polymer semiconductor thin films with multi-\nmodal energy dissipation and high relative stretchability. Nat. Comm. 14, 8382 (2023).\n20. Mun, J. et al. A design strategy for high mobility stretchable polymer\nsemiconductors. Nat. Comm. 12, 3572 (2021).\n21. Wang, Z. et al. Revealing the Strain-Induced Morphological, Mechanical, and\nPhotovoltaic Evolution of Self-Encapsulated and Semitransparent Intrinsically\nStretchable Organic Solar Cells. ACS Mater. Lett. 6, 1811-1819 (2024).\n22. Peng, Z. et al. Unraveling the Stretch-Induced Microstructural Evolution and\nMorphology–Stretchability Relationships of High-Performance Ternary Organic\nPhotovoltaic Blends. Adv. Mater. 35, 2207884 (2023).\n38\n23. Jin-Hoon Kim, Jin-Woo Park, Intrinsically stretchable organic light-emitting diodes.\nSci. Adv. 2021, 7, eabd9715.\n(2) The angle-resolved NEXAFS shows increased α-angles of polymer chains under\nstrain. Are similar changes observed within the crystalline domains? Since face-on π-π\nstacking is crucial for vertical charge transport in OSCs, clarification on whether the\ncrystallite orientation is similarly affected would be helpful.\nOur response: Thanks for the comment. The increased α-angles observed in angle-\nresolved NEXAFS reflect an overall enhancement of face-on chain orientation,\nincluding both crystalline and amorphous regions. To explicitly probe crystalline\ndomains, we performed GIWAXS measurements on strained films transferred to silicon\nwafers. The azimuthal intensity distribution of the π-π stacking reflection showed\nnarrowing when probed perpendicular to the strain direction and broadening when\nprobed parallel. This indicates enhanced face-on π-π stacking orientation using X-ray\nto probe the perpendicular direction, consistent with the NEXAFS results. Moreover,\nthe contrasting changes in pole figure width suggest reduced twisting of crystallites\nabout the (001) axis but increased tilting about the (100) axis. Together, these findings\ndemonstrate that crystallite deformation and orientation contribute to irreversible\nstructural adaptation for stress dissipation under strain. To clarify this point, we have\nadded the following sentences in the manuscript.\nMain Text (Insertion, Page 9, Line 7):\n“To study the π-π stacking developments during stretch, the films were stretched\nto the desired strains and then placed on silicon wafers for grazing incidence wide angle\nX-ray scattering (GIWAXS) measurements at the beam energy of 10 keV. This\nmeasurement was performed with orthogonal X-ray directions including parallel and\nperpendicular conditions (Fig. S28). The 0% strain thin film shows nearly identical\nreflections (Fig. S29), which support that βₓ ≈ β before stretching as demonstrated by\ny\nNEXAFS. Meanwhile, the (100) peak area in both q and q decreases rapidly at small\nx y\nstrains (10-20%) and then more gradually at higher strains, with a stronger reduction\n39\nalong q (Fig. S30-32, Table S1-S2). This behavior is consistent with the observations\nx\nat 2476.26 eV in TReXS, indicating that the resonant excitation is not sensitive to the\ncrystalline deformation. The azimuthal intensity of the π-π stacking reflection was used\nto evaluate crystallite orientation relative to the substrate (Fig. S33). With increasing\nstrain, the pole figures show narrowing for the perpendicular condition and broadening\nfor the parallel condition. Therefore, an enhanced face-on orientation of π-π stacking is\nobserved when X-rays probe perpendicular to the SD, in agreement with the angle-\nresolved NEXAFS results. The opposite trend in pole figure widths suggests reduced\ntwisting of crystallites about the (001) axis but increased tilting about the (100) axis,\nconsistent with the tilted crystallite orientation observed in flow-induced alignment [43].\nTogether, these combined mechanisms of crystallite deformation and orientation\nprovide irreversible structural adaptation pathways for stress dissipation.”\nReference:\n[43] Persson, N.E., Engmann, S., Richter, L.J. & DeLongchamp, D.M. In Situ\nObservation of Alignment Templating by Seed Crystals in Highly Anisotropic Polymer\nTransistors. Chem. Mater. 31, 4133-4147 (2019).\nSupplementary Information:\nP(NDI2OD-T2) thin film, where the scattering images are collected with the incident\nX-ray roughly perpendicular (perp.) or parallel (para.) to the SD. (b) GIWAXS patterns\nof P(NDI2OD-T2) thin films with various strains measured with the incident X-ray\nbeam (b) perpendicular (q ) and (c) parallel (q ) to the SD.\nx y\nFig. S33. (a) Illustration of the circular average for the generation of (010) pole figures\nand (b) fitting example for the pole figures using Gaussian function. (c) (010) pole\nfigure full width at half maximum (FWHM) as functions of strain.\n(3) The observed decrease in crystallinity during stretching is reasonable due to the\nrigid backbone of conjugated polymers. However, it remains unclear why the scattering\nintensity continues to decrease upon strain release.\nOur response: Thanks for the comment. The continued decrease in scattering intensity\nduring strain release can be rationalized as follows: (i) in the initial stage of release (30%\n→ ~20% strain), residual stress remains above the critical threshold for crystal stability,\ncausing further disruption of crystallites even as the macroscopic strain decreases;\n(ii) below ~20% strain, the applied stress falls below this threshold, so no additional crystal\ndestruction occurs, but the previously disrupted structures cannot self-reorganize due\nto irreversible chain slip and peeling. As a result, the scattering intensity does not\nrecover upon full release. To clarify this, we added the following sentences in the\nmanuscript.\nMain text (Page 8, line 18):\n“Upon releasing the film, the scattering intensities continued to decrease between 30%\nand 20% strain, likely due to residual stress remaining above the critical threshold for\ncrystallite destruction. Below ~20% strain, no further decrease was observed; however,\nneither reflection recovered, indicating a permanent loss of crystallinity for both\ndirections.”\n(4) The use of Δφ to quantify morphological orientation is well-executed and\nconvincingly correlates with mechanical strain (σ–ε). However, the physical origin of\nthe “X”-shaped contrast in the RSoXS images should be clarified due to the structural\ncomplexity of the stained thin film. Meanwhile, how are the φ positions determined?\nOur response: Thanks for the comment. The “X”-shaped contrast in the RSoXS\npatterns originates from oriented fibrillar structures, as supported by the consistent “X”-\nlike features observed in the FFT analysis of AFM images. We acknowledge that the\nmorphology of strained thin films is complex, where domain and orientational contrasts\nare intertwined, making it challenging to decouple their individual contributions.\nNevertheless, orientational contrast, which can be quantitatively analyzed, provides a\nreasonable approximation, as the collective alignment of polymer chains forms oriented\nfibrils. Although this approach may not fully capture the complete complexity of the\nsystem, it effectively accounts for the main source of contrast. In particular, we\nobserved a pronounced increase in orientational contrast in the xy- and xz-planes at\n284.0 eV under strain, consistent with the experimental ISI results. To further clarify\nthis point, we have added in situ RSoXS images of the P(NDI2OD-T2) thin film at\nvarious strains during stretching and releasing, which show that the “X”-shaped feature\nis absent when probed at 285.0 eV. These experimental observations and calculations\nstrongly support that the “X”-shaped scattering feature originates from the strain-\ninduced orientation of polymer chains within the fibrillar morphology, with its visibility\nbeing highly sensitive to the selected X-ray energy. To address this, we added the\nfollowing sentence in the manuscript.\nMain text (Page 10, line 27):\n“We note that the strained thin-film morphology is complex, as domain and\norientational contrasts are intertwined and difficult to decouple. Nevertheless, the\nconsistent “X” shape observed in AFM FFT images, the enhanced orientational contrast\nin the xy- and xz-planes at 284.0 eV, and the local maximum ISI at 284.0 eV indicate\nthat the “X”-shaped scattering arises from strain-induced alignment of polymer chains\nacross the fibrillar morphology. In situ RSoXS measurements at 285.0 eV during\nstretching show that “X”-shaped feature disappears (Fig. S44), suggesting reduced\norientation sensitivity and further supporting the assignment of scattering contrast\norigin. Thus, the “X” pattern provides direct information related to the orientation of\nfibril aggregates.”\nSupplementary Information:\nFig. S44. In situ RSoXS images of a P(NDI2OD-T2) thin film at different strains during\n(a) stretching and (b) releasing, which are probed at 285.0 eV.\nTo determine the φ positions, the azimuthal intensity profiles were fitted using\nGaussian functions. For clarity, an example of the peak-fitting is provided in Fig. S47.\n42\n(a)\n6\n-6\n350\n300\n250\n200\n150\n0\n100\n40\n0 50 100 150 200 250 300 350\n(b) 6\n-6\n350\n300\n250\n200\n150\n0\n160\n120\n0 50 100 150 200 250 300 350\nFig. S47. Example of Gaussian-based peak fitting used to extract (a) φ and (b) φ from\n1 2\nazimuthal intensity profiles obtained at 284.0 eV.\n(5) The manuscript suggests that crystallite disruption and increased torsion reduce\nexciton delocalization, affecting optical absorption. Do the authors have supporting\nevidence or citations for this claim?\nOur response: Thanks for the comment. To support our interpretation, we performed\ntime-dependent density functional theory (TD-DFT) calculations using dimer structures\nextracted from the MD simulations of the amorphous polymer under different strain\nconditions. The results reveal that when the intra-chain dihedral angle increases from\n30° to 90° under high strain, the calculated absorption peak exhibits a noticeable blue\nshift. This trend indicates reduced conjugation and exciton delocalization along the\npolymer backbone, as well as weakens electronic coupling. To clarify this, we added\nthe following sentence in the manuscript.\nMain text (Page 14, line 12):\n43\n“Time-dependent density functional theory calculations using strain-dependent dimer\nstructures confirm that increasing the dihedral angle between BT and NDI units from\n~30° to 90° induces a blue shift in absorption (Fig. S49), consistent with reduced π-\nconjugation and exciton delocalization.”\nSupplementary Information:\nFig. S49. Time-dependent density functional theory (TD-DFT) calculated absorption\nspectra of strain-dependent dimer structures with different the dihedral angle between\nBT and NDI units.\n(6) Regarding the J-V curves of stretchable solar cells, how is the device’s active area\ndefined and measured during strain application? This point should be clarified in the\nexperimental section.\nOur response: Thanks for the comment. The initial device area (0.04 cm2) was defined\nby the overlap between the patterned M-PH1000 bottom electrode and the patterned\nEGaIn@Ag top electrode in a cross-electrode geometry. During mechanical stretching,\nthe active area deformed geometrically. To account for this, we corrected the active area\nby quantifying dimensional changes using optical microscopy. Specifically, the average\narea change was obtained from three independent devices under identical strain\nconditions, and the results were fitted with a three-phase exponential decay function\n(ExpDec3). This fitting curve was then used to correct the active area of the tested\ndevices at each applied strain. To clarify this, we added the following sentences to the\nMethod:\n“The initial active area (0.04 cm2) was defined by the overlap of the patterned M-\nPH1000 bottom electrode and the patterned EGaIn@Ag top electrode in a cross-\nelectrode geometry. During mechanical stretching, the active area deformed\ngeometrically. To account for this, the actual active area was corrected based on\ndimensional changes measured by optical microscopy. The average area variation was\nextracted from three independent devices under identical strain conditions and fitted\nwith a three-phase exponential decay function (Fig. S62), which was then applied to\ncorrect the active area of the tested devices at each strain.”\nFig. S62. The correction curve for the active area of intrinsically stretchable OPV\ndevices. Black squares denote the averaged active areas of three independent devices,\nwhile the red line represents the corresponding fit using a three-phase exponential decay\nfunction.\nI co-reviewed this manuscript with one of the reviewers who provided the listed reports.\nThis is part of the Nature Communications initiative to facilitate training in peer review\nand to provide appropriate recognition for Early Career Researchers who co-review\nmanuscripts.\nReviewer #5 (Remarks to the Author):\nI co-reviewed this manuscript with one of the reviewers who provided the listed reports.\nThis is part of the Nature Communications initiative to facilitate training in peer review\nand to provide appropriate recognition for Early Career Researchers who co-review\nmanuscripts.\n46"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have adequately addressed the manuscript revisions, and this version has improved significantly. Most of my comments have been addressed with additional experiments and explanations in the text where needed. The only part that remains slightly confusing is the additional AFM analysis for comment 5. The authors introduced power spectral density (PSD) curve analysis of the AFM results, but the results shown in Fig. S41 differ from their description. From the figure, the order of domain size from small to large is 0% (black), 10% (red), 30% (blue), and 20% (yellow), while in the figure caption it was listed as 0%, 10%, 20%, and 30%. Additionally, the authors did not apply the same PSD analysis to their original AFM data from the 50% strain film for comparison. While the manuscript can be published in Nature Communication, I would recommend some additional clarification of the newly added AFM analysis.", "reviewer_2": "The revision has addressed all the comments. I now recommend accepting the paper.", "reviewer_3": "The authors have thoroughly addressed my comments, and the manuscript has been revised accordingly.Thus, I would recommend it for publication.", "reviewer_4": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts.", "reviewer_5": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts."}, "author_response": "Response to Reviewers’ Comments\nReviewer #1 (Remarks to the Author):\nThe authors have adequately addressed the manuscript revisions, and this version has\nimproved significantly. Most of my comments have been addressed with additional\nexperiments and explanations in the text where needed. The only part that remains\nslightly confusing is the additional AFM analysis for comment 5. The authors\nintroduced power spectral density (PSD) curve analysis of the AFM results, but the\nresults shown in Fig. S41 differ from their description. From the figure, the order of\ndomain size from small to large is 0% (black), 10% (red), 30% (blue), and 20% (yellow),\nwhile in the figure caption it was listed as 0%, 10%, 20%, and 30%. Additionally, the\nauthors did not apply the same PSD analysis to their original AFM data from the 50%\nstrain film for comparison. While the manuscript can be published in Nature\nCommunication, I would recommend some additional clarification of the newly added\nAFM analysis.\nOur response: Thanks for the comment. We apologize for the color mismatch in the\npreviously submitted Fig. S41. We have now corrected the dotted-line colors to match\nthe order in the caption: 0% (black), 10% (red), 20% (yellow), and 30% (blue).\nTo address the second comment, we have also performed the same PSD analysis\non the AFM height images of the 50% strained film. The 50% PSD curve shows an\noverall higher PSD amplitude and a peak shifted to a lower spatial-frequency range,\nconsistent with the formation of larger aggregated domains under high strain. We\nfurther refined the calculation of characteristic domain size using the inverse of the\nspatial frequency k corresponding to the turning or peak positions in the PSD plots.\nThe extracted characteristic domain size increases from 3.6 μm to 6.2 μm, 7.1 μm, 7.7 μm,\nand 8.3 μm for films strained at 0%, 10%, 20%, 30%, and 50%, respectively. Both the\nfigure and the caption have been updated for clarity as show below.\nshown in Fig. S34 for 0-30% strains and Fig. 3a for 50% strain) of P(NDI2OD-T2) thin\nfilms under different stretch strains. The dotted lines highlight the turning or peak\npositions used to estimate the characteristic domain size, calculated as the inverse of\nthe corresponding spatial frequency k. The characteristic domain size increases from\n3.6 μm to 6.2 μm, 7.1 μm, 7.7 μm, and 8.3 μm for films strained at 0%, 10%, 20%, 30%,\nand 50%, respectively. These results reflect the micrometer-scale aggregation process\ncaptured by AFM.\nThe revision has addressed all the comments. I now recommend accepting the paper.\nReviewer #3 (Remarks to the Author):\nThe authors have thoroughly addressed my comments, and the manuscript has been\nrevised accordingly. Thus, I would recommend it for publication.\nReviewer #4 (Remarks to the Author):\nI co-reviewed this manuscript with one of the reviewers who provided the listed reports.\nThis is part of the Nature Communications initiative to facilitate training in peer review\nand to provide appropriate recognition for Early Career Researchers who co-review\nmanuscripts.\nReviewer #5 (Remarks to the Author):\nI co-reviewed this manuscript with one of the reviewers who provided the listed reports.\nThis is part of the Nature Communications initiative to facilitate training in peer review\nand to provide appropriate recognition for Early Career Researchers who co-review\nmanuscripts.\n3"}]

Decision letter