Localized Potential Regulation on Polymer Donor Backbone Suppresses Energetic Disorder for Efficient, Stable and Scalable Organic Solar Cells

ABSTRACT

High energetic disorder and short exciton diffusion lengths in low‐crystallinity polymer donors (P _Ds) critically limit exciton dissociation, charge transport, efficiency, stability, and scalability of organic solar cells (OSCs). Here we report a localized potential regulation strategy to intrinsically suppress energetic disorder of P _D through precise backbone engineering. By introducing a strongly electron‐deficient unit into the structurally disordered backbone of the DP1, we developed a P _D DP10 that exhibits enhanced backbone rigidity, optimized local electronic polarization, and favorable miscibility with the acceptor L8‐BO. These features synergistically reduce the exciton binding energy and suppress exciton–phonon coupling, thereby extending the exciton diffusion length from 15.5 to 19.9 nm. Consequently, DP10 enables balanced, trap‐tolerant charge transport and reduced non‐radiative recombination in devices. The DP10:L8‐BO binary system achieves a device efficiency of 19.51%, together with exceptional thermal stability and outstanding thickness‐tolerant performance. Additionally, a DP10:L8‐BO:BTP‐eC9 ternary device reaches an efficiency of 20.57% and a 15.1 cm² solar module delivers 17.20% efficiency, with impressive fill factors of 81.04% and 77.39%, respectively. This work establishes localized potential regulation as a powerful molecular design principle of P _Ds for simultaneously achieving high efficiency, thermal stability, and processing robustness in next‐generation OSCs.

A localized potential strategy yields modified polymer donor DP10 with improved molecular packing and local charge polarization. This extends the exciton diffusion length and stabilizes charge transport, enabling organic solar cells to achieve 20.57% efficiency and a 15.1 cm² module to reach 17.20%, while also offering good thermal stability and thickness tolerance.

1. Introduction

Organic solar cells (OSCs) have emerged as one of the most promising photovoltaic technologies, owing to their intrinsic advantages such as light weight, mechanical flexibility, semitransparency, and compatibility with low‐temperature, solution‐based manufacturing processes [1, 2, 3]. Over the past decade, sustained progress in molecular design [4, 5], morphology control [6, 7], and device engineering [8] has driven the power conversion efficiencies (PCEs) of single‐junction OSCs to exceed 20% [9, 10]. In parallel, advances in large‐area coating techniques and module fabrication have highlighted the potential of OSCs for building‐integrated photovoltaics, portable power sources, and emerging applications where conventional inorganic photovoltaics are less suitable [1, 2, 11]. Despite these impressive achievements, the translation of laboratory‐scale efficiencies into stable, process‐tolerant, and industrially viable devices remains a central challenge for the field [2, 12, 13, 14, 15, 16]. Realizing this goal requires active layers capable of sustaining thermal stress while maintaining high efficiency at increased thicknesses [17]. From a device‐physics perspective, this necessitates efficient three‐dimensional exciton and carrier transport to counteract two dominant degradation pathways: (i) thermally‐induced phase coarsening, which accelerates Frenkel exciton quenching [17, 18, 19] and (ii) thickness‐dependent space‐charge accumulation, which weakens the internal electric field and exacerbates recombination losses [20, 21, 22, 23]

To mitigate these effects, extending exciton diffusion lengths and enhancing charge carrier mobility within the active layer have emerged as primary solutions. For example, deuteration of the vinyl linker in Y‑series small molecule acceptors (SMAs) has been shown to effectively reduce vibrational relaxation [24], resulting in an approximately 50 % enhancement in exciton lifetime [25]. In parallel, ternary blending strategies incorporating highly crystalline small‑molecule donors or SMAs have been developed to prolong carrier diffusion lengths in the host system, thereby alleviating efficiency losses in thick active layers [4, 22, 26, 27]. Solid‑additive strategies have also emerged as effective approaches to coordinately regulate donor/acceptor crystallization kinetics, resulting in concurrent improvements in device efficiency, thermal stability, and thickness tolerance [28, 29, 30, 31]. Although these approaches have proven effective at the device level, they largely circumvent the intrinsic limitations of polymer donors. Consequently, the exciton diffusion length (L _D) of state‐of‐the‐art polymer donors still lags significantly behind that of their non‐fullerene acceptor counterparts [32, 33, 34, 35, 36]. Strategies that intrinsically enhance exciton and charge‑carrier dynamics within polymer donors themselves are still lacking, constituting a critical bottleneck in both performance optimization and mechanistic understanding of OSCs.

Polymer donors inherently suffer from several limitations, including disordered molecular packing, restricted charge transport, high exciton binding energy (E _b), and short exciton lifetimes [35]. These deficiencies largely stem from the torsional flexibility and conformational randomness of conjugated polymer backbones [37]. Although the donor–acceptor (D–A) architecture promotes intramolecular charge transfer (ICT), the use of relatively weak electron‐withdrawing units often leads to significant spatial overlap between hole and electron wavefunctions along the conjugated backbone [38]. While such localization can suppress extensive geometric relaxation upon excitation and thereby reduce the internal reorganization energy (λ _int), it simultaneously hampers charge separation and significantly elevates E _b [39]. These intrinsic constraints severely limit the contribution of low‐crystallinity polymer donors to efficient energy and charge transport in bulk heterojunction systems. Therefore, beyond ensuring energy‐level alignment and spectral complementarity with acceptor materials [40], rational molecular design must focus on enhancing electronic polarizability, reducing E _b, and stabilizing molecular packing [41, 42]. Such an approach, targeted at mitigating exciton confinement and suppressing energetic disorder, is essential for establishing efficient exciton diffusion pathways and robust charge–carrier extraction.

To address the challenge of balancing stability and scalable fabrication in low‐crystallinity polymer donors, we exploit the unique attributes of the low‐cost building block TTzCl (Figure 1a), previously developed by our group [43, 44]. Its asymmetric molecular structure and solubilizing side chains induce quasi‐random backbone conformations [45], endowing the resulting polymers with excellent solution processability [44]. However, the disordered molecular packing of relevant polymers limits thermal stability. Notably, by implementing homologous backbone modifications that strengthen and balance intra‐ and intermolecular interactions, unfavorable chain vibrations and traps (or defects) can be effectively suppressed [46, 47], thereby enhancing structural order and stability while preserving favorable processing characteristics. Guided by this rationale, herein we developed a modified polymer donor, DP10, derived from DP1 via a localized potential regulation strategy (Figure 1a). This involved copolymerizing 10 mol% of a strongly electron‐deficient ester‐substituted thiazole (TClTzE) unit, which shares a similar structural skeleton with TTzCl but features deeper energy levels and multiple non‐covalent interactions with adjacent benzo[1,2‐b:4,5‐b']dithiophene (BDT) units (Figure 1b and Figure S1, Table S1). This targeted modification reinforces backbone rigidity and induces localized electronic polarization along the polymer chain. As a result, DP10 exhibits reduced E _b and weakened exciton‐phonon coupling in the thin film, leading to a pronounced extension of L _D from 15.5 nm (DP1) to 19.9 nm (DP10). Crucially, the reduced energetic disorder in DP10 is well‐preserved upon blending with the SMA L8‐BO [48] (Figure S2), enabling efficient charge generation and the formation of interconnected stacking networks that facilitate exciton diffusion and charge‐carrier transport. The DP10:L8‐BO system yields an excellent PCE of 19.51%, along with outstanding thermal stability and robust thickness tolerance. Furthermore, by incorporating the SMA BTP‐eC9 [49] (Figure S2) to extend near‐infrared absorption and further optimize the blend morphology, a ternary system with a PCE of 20.57% was attained. This high efficiency, combined with the demonstration of a scalable large‐area (15.1 cm²) module (PCE = 17.20%, FF = 77.39%), places DP10 on par with state‐of‐the‐art polymer donors such as PM6 [50] and D18 [51]. These results demonstrate that the localized potential regulation strategy effectively minimizes energetic disorder and provides a robust design principle for developing efficient polymer donors in multi‐material systems.

FIGURE 1.

Material properties of polymer donors DP1 and DP10. (a) Molecular structures and design strategies. (b) NBO‐based orbital overlap diagrams of the TClTzE unit. (c) Cyclic voltammograms and energy level diagrams (referenced to Fc/Fc⁺). (d) The ultraviolet‐visible‐near infrared (UV‐vis‐NIR) absorption spectra of donor films. (e) In‐plane and out‐of‐plane CCLs and hole mobilities. (f) Enlarged views of the hole‐electron distribution diagrams of donor blocks (blue and red areas represent hole and electron distributions, respectively). (g) Dielectric properties of pure donor film devices. (h) Temperature‐dependent PL spectra of DP1 film. (i) Temperature‐dependent PL spectra of DP10 film. (j) Exciton lifetimes and diffusion lengths of films.

2. Results and Discussion

2.1. Polymer Design, Characterization, and Excited‐State Dynamics

The polymer donor DP10 was synthesized via Stille coupling copolymerization, employing a BDT derivative as the electron‐donating unit and a combination of two brominated electron‐accepting monomers, TTzCl‐Br and TClTzE‐Br, introduced at a feed molar ratio of 90:10. It should be noted that this feed ratio was selected after optimization because the DP10‐based devices exhibited the highest PCE among the modified polymer donors (DP5, DP10 and DP15) with different feed ratios (Figure S3 and Table S2). Detailed synthetic procedures for the TClTzE‐Br monomer and DP10 are provided in the Supporting Information (Figure S4), together with the corresponding nuclear magnetic resonance spectra presented in Figures S5–S11. High‐temperature gel permeation chromatography, conducted at 160°C using 1,2,4‐trichlorobenzene as the eluent, yielded a number‐average molecular weight (M _n) of 52.8 kg mol⁻¹ and a dispersity (Đ) of 2.32 for DP10 (Figure S12). For comparison, the reference polymer DP1 [44], composed exclusively of TTzCl acceptor units, exhibited a similar M _n of 46.3 kg mol⁻¹ and a slightly broader Đ of 2.76 (Figure S12). The solubility of DP1 and DP10 in chloroform at room temperature was measured to be 34.5 mg mL⁻¹ and 30.9 mg mL⁻¹, respectively (Figure S13). The lower solubility of DP10 indicates stronger pre‑aggregation in solution, which likely results in a less hybridized donor–acceptor interfacial morphology in the active layer.

To validate the design strategy of localized potential regulation on the polymer backbone, density functional theory (DFT) calculations were integrated with experimental characterizations to systematically investigate the optoelectronic disparities between DP1 and DP10. Geometry optimizations were performed on representative oligomeric fragments, with long alkyl side chains truncated to methyl groups to reduce computational complexity. All calculations were carried out at the PBE0/def2‐SVP level, including GD3BJ dispersion corrections, while a polarizable continuum model with diphenyl ether (dielectric constant ε _r = 3.7) was employed to approximate the dielectric environment of the solid state [52]. As illustrated in Figure S14, frontier molecular orbital analysis indicates that the incorporation of the strongly electron‐deficient TClTzE unit induces only a negligible shift in the highest occupied molecular orbital (HOMO) level, from −5.34 eV for DP1 to −5.37 eV for DP10. In contrast, the lowest unoccupied molecular orbital (LUMO) level undergoes a more pronounced downshift, decreasing from −2.81 to −2.93 eV. This minor HOMO variation is further supported by ionization potentials (E _IP) derived from cyclic voltammetry (CV) measurements (Figure 1c), which yield values of −5.52 eV for DP1 and −5.56 eV for DP10. The marginal difference in E _IP suggests that the introduction of TClTzE has a limited impact on the donor‐acceptor energy offset and is therefore unlikely to be the primary factor responsible for the observed variations in V _OC [53].

Natural bond orbital (NBO) analysis [54] (Table S1) reveals that multiple non‐covalent orbital interactions associated with the TClTzE unit cooperatively stabilize the planar backbone conformation of DP10 (Figure 1b), in contrast to DP1 (Figure S1). In DP1, backbone planarity is primarily governed by a single Cl···S interaction on the thiazole ring (second‐order perturbation energy, E ⁽²⁾ = 2.15 kcal mol^{− 1}; orbital overlap integral, S = 0.389), while the S···C–H interaction involving the alkyl side chain is negligible, indicating a solubility‐driven steric role. Conversely, DP10 features two types of synergistic stabilizing interactions: the O···S contact involving the ester oxygen (E ⁽²⁾ = 1.57 kcal mol⁻¹, S = 0.371) and two Cl···S contacts involving the chlorine atom on the thiophene ring (E ⁽²⁾ ≈ 1.14 kcal mol⁻¹, S ≈ 0.169). The increased multiplicity of effective orbital overlaps in the DP10 segment facilitates stronger electronic delocalization along the backbone [55]. The stronger intramolecular interactions help suppress high‐frequency molecular vibrations along the DP10 polymer chain, thereby promoting intrachain exciton diffusion and charge transport [25]. This stable chain conformation is further supported by potential energy surface scans, which show an increased rotational barrier for DP10 (Figure S15). Experimentally, the enhanced rigidity manifests as a slight red shift in solution absorption (Figure S13) and strengthened π‐π stacking in the solid state (Figure 1d). Consistent with this trend, two‐dimensional grazing‐incidence wide‐angle X‐ray scattering (2D GIWAXS) measurements (Figures S16, S17) were performed, which reveal an increased out‐of‐plane (010) π‐π stacking coherence length for DP10 (crystal coherence length, CCL = 25.4 Å) relative to DP1 (CCL = 22.3 Å) (Figure 1e and Table S3). The resulting improvement in molecular ordering suppresses dynamic disorder during charge transport, leading to a higher space‐charge‐limited current (SCLC) hole mobility for DP10 (9.15  ×  10⁻⁴ cm² V⁻¹ s⁻¹), compared with DP1 (5.23  ×  10⁻⁴ cm² V⁻¹ s⁻¹) (Figure 1e and Figure S18).

To elucidate the impact of localized potential regulation on ICT dynamics, the excited‐state geometries of DP1 and DP10 segments were optimized using time‐dependent DFT [56]. As summarized in Table S4, the S₁ state contributes dominantly to the oscillator strength in both segments, identifying it as the primary optically allowed transition, followed by S₃. Hole–electron distribution analysis (Figure 1f and Figure S19) reveals that although the S₁ state corresponds to a largely localized excitation in both segments, DP10 displays a larger hole–electron centroid distance (D index) than DP1, indicating enhanced spatial charge separation. Notably, for the S₀ → S₃ transition, the excited‐state electron density in DP10 is predominantly localized on the TClTzE unit, accompanied by a substantially increased D index (15.22 Å) and a reduced hole–electron overlap integral (S_r = 0.62), compared with DP1 (D = 4.15, S_r = 0.75). These results demonstrate that the incorporation of the TClTzE unit effectively facilitates the formation of a pronounced ICT state. This trend is consistent with electrostatic potential (ESP) maps (Figure S20), which reveal strongly negative regions localized around the ester oxygen and chlorine substituents, underscoring the electron‐withdrawing contribution of nonconjugated lone‐pair electrons. Based on the optimized ground‐state geometries, total energies of the neutral (E _N), anionic (E _N+1), and cationic (E _N‐1) states were calculated to derive the vertical ionization potential (VIP  = E _{N − 1}  − E_N ), vertical electron affinity (VEA  = E_N   − E _{N + 1}), and exciton binding energy ($E_b=VIP-VEA-E_g^{opt}$, where $E_g^{opt}$ is the optical gap of the S₁ state) [57]. As summarized in Table S4 and Figure 1f, the E _b of the DP1 segment is 0.54 eV, whereas that of the DP10 segment drops to 0.48 eV, indicating an enhanced potential for exciton dissociation.

To further assess the stability of the separated charges, the internal hole reorganization energy (λ_h) was calculated using the standard four‐point method (Table S4, and Figure 1f and Figure S19) [58]. To elucidate the role of the TClTzE unit as a modifying unit, a model polymer segment designated as DP100 consisting of alternating TClTzE and BDT units was simulated. The properties of the modified polymer DP10 were benchmarked against the binary copolymers DP1 and DP100. Notably, DP10 exhibits a higher λ _h (0.17 eV) than DP1 (0.13 eV) and DP100 (0.12 eV), which indicates more robust local electron phonon coupling. This effect is driven by the significant electrostatic potential disparity between the TClTzE and TTzCl moieties within the Holstein polaron framework [39]. The resulting deeper polaron binding energy (E _pol ≈ λ_h/2) stabilizes hole carriers via local lattice deformations. Although increased reorganization energy can impede charge transport, in the highly ordered DP10 films, it serves to suppress rapid carrier recombination by dynamically localizing holes away from recombination‐prone disordered regions. This balance between efficient exciton dissociation and carrier stabilization is consistent with the prolonged carrier lifetimes observed in transient absorption (TA) spectroscopy (Figure S21). Collectively, the reduced E _b and increased E _pol in DP10 synergistically promote efficient photovoltaic performance as discussed below.

The theoretically predicted reduction in E _b is experimentally substantiated by the dielectric properties of the polymer films. DP10 exhibits a higher relative dielectric constant (ε _r = 5.1) than DP1 (ε _r = 4.6), suggesting stronger dielectric screening that effectively weakens the Coulombic attraction between photogenerated hole‐electron pairs. To quantitatively assess this effect, temperature‐dependent photoluminescence (PL) spectra (Figure 1h–i) were fitted using a thermal quenching model [26]:

$$I_T=\frac{I_0}{1+A{e}^{-E_a/k_{B}T}}$$

where I _T is the PL intensity at temperature T, I ₀ is the intrinsic intensity at 0 K, A is a pre‐exponential factor related to the ratio of radiative to non‐radiative rates, E _a is the activation energy (taken as the experimental E _b), k_B is the Boltzmann constant, and T is the absolute temperature. The extracted E _a (Figure S22) are 0.076 eV for DP1 and 0.063 eV for DP10, respectively, confirming that the incorporation of the TClTzE unit effectively reduces the Frenkel exciton binding energy.

To further elucidate the impact of the TClTzE unit on exciton relaxation dynamics, temperature‐dependent PL spectra were deconvoluted into contributions from ICT, π‐π stacking, and tail states (Figures S23, S24) [29]. Low‐temperature PL spectra at 80 K reveal a significantly narrower π‐π emission peak for DP10 compared to DP1, indicative of suppressed excited‐state relaxation. Additionally, the reduced ICT/π‐π intensity ratio in DP10 suggests that localized potential regulation mitigates vibrational relaxation pathways. Quantitative analysis of the temperature‐dependent full width at half maximum (FWHM) for both ICT and π‐π emissions (Figure S25) was fitted using the model [59]:

$$\mathrm{\Gamma }\left(T\right)=\mathrm{\Gamma }\left(0\right)+\sigma T+\frac{{\mathrm{\Gamma }}_{\mathrm{op}}}{e^{h{\omega }_{\mathrm{op}}/k_{B}T}-1}$$

In this expression, Γ(Τ) denotes the temperature‐dependent FWHM, while Γ(0) represents the inhomogeneous broadening. The term σ accounts for exciton‐acoustic phonon coupling (typically negligible), hω _op is the optical phonon energy, and Γ _op represents the exciton‐optical phonon coupling strength. Analysis of the emission peaks reveals that DP10 exhibits smaller Γ(0) values (0.128 eV for ICT and 0.114 eV for π‐π stacking) compared to DP1 (0.135 eV and 0.138 eV, respectively), indicating reduced static structural disorder. Regarding the π‐π stacking emission, DP10 displays a reduced Γ _op of 0.594 eV (vs. 0.839 eV for DP1). This suppressed dynamic coupling is consistent with the enhanced structural order evidenced by the larger CCL in GIWAXS measurements. Equally important, the ICT state in DP10 exhibits a drastically reduced Γ _op (0.714 eV) compared to DP1 (4.195 eV), implying significantly weaker exciton‐vibration coupling along the polymer backbone. These findings further corroborate the reduced energetic disorder and improved molecular rigidity in the DP10 system.

The suppression of non‐radiative decay pathways directly translates into prolonged exciton lifetimes in DP10. Time‐resolved PL (TRPL) measurements (Figure 1j and Figure S26) yield a τ _TRPL of 175.9 ps for DP10, compared with 134.2 ps for DP1. Furthermore, pump‐fluence‐dependent transient absorption spectroscopy was employed to extract the L _D via exciton–exciton annihilation method [60]. The exciton population dynamics were globally fitted using the following integrated rate equation, which accounts for concurrent first‐order recombination and bimolecular annihilation:

$$N\left(t\right)=\frac{N\left(0\right)e^{\left(-k_{1}t\right)}}{1+\frac{k_2}{2k_1}N\left(0\right)\left[1-{\mathrm{e}}^{\left(-k_{1}t\right)}\right]}$$

Here, N(t) represents the singlet exciton density, k ₁ is the intrinsic monomolecular decay rate constant (inverse of the natural lifetime t in the absence of annihilation), and k ₂ denotes the bimolecular annihilation rate constant. As shown in Figure 1j and Figure S27, the L _Ds derived from excitation fluences of 2.5 – 71.0 µJ cm⁻² (at 595 nm) are 15.5 nm for DP1 and 19.9 nm for DP10, respectively. This improvement benefits from the stronger intrachain conformational stability and improved intermolecular packing of DP10 (Figure 1e and Figure S16) [24, 37]. The extended L _D in DP10 is particularly advantageous for efficient exciton diffusion within active layer morphologies featuring pronounced vertical phase separation and relatively large domain sizes.

The impact of the polymer donor on exciton dissociation and relaxation dynamics in L8‐BO‐based blend systems was systematically examined through comprehensive exciton kinetics analysis. As shown in Figure 2a, the DP10:L8‐BO blend displays a higher ε _r (ε _r = 3.5) than the DP1:L8‐BO counterpart (ε _r = 3.3), indicating that the enhanced exciton dissociation capability associated with DP10 is preserved upon blending. TRPL measurements (Figure 2b and Figure S28) further reveal that, despite significant quenching relative to pristine L8‐BO, the DP10:L8‐BO blend maintains a longer exciton lifetime (97.1 ps) than the DP1:L8‐BO system, implying that DP10 effectively mitigates vibrational relaxation within the L8‐BO phase. These observations are supported by temperature‐dependent PL analyses of the blend films (Figure 2c, d and Figures S29–S31). Specifically, the DP10:L8‐BO film yields lower E _a (0.061 eV) and F _op (0.383 eV) values than the DP1:L8‐BO system (0.073 eV and 0.449 eV, respectively). The diminished coupling suggests attenuated exciton‐phonon interactions in the acceptor phase, likely arising from the more ordered molecular packing induced by DP10.

FIGURE 2.

Excited‐state dynamics analysis for DP1:L8‐BO and DP10:L8‐BO systems. (a) Dielectric properties of L8‐BO and blend film devices. (b) Normalized TRPL decay kinetics. (c) Pseudo‐color plots of temperature‐dependent PL spectra for the DP1:L8‐BO film. (d) Pseudo‐color plots of temperature‐dependent PL spectra for the DP10:L8‐BO film. (e) Pseudo‐color TA plots of the blend films under 830 nm excitation. (f) Hole transfer kinetics of the ICT state derived from global analysis of the GSB and corresponding ESA signals. (g) Hole transfer kinetics of the π‐π stacking state derived from global analysis of the GSB and corresponding ESA signals. (h) Comparison of hole transfer rates, lifetimes, and yields in the blend films; the bubble size represents the relative charge yield (A₂ + A₃).

TA spectroscopy under selective acceptor excitation (830 nm) resolves distinct GSB features associated with ICT (549 nm) and π‐π stacking (595 nm) states in the DP10 blend, alongside an ESA feature at 631 nm (Figure 2e and Figure S32). These spectral signatures enable deconvolution of the dominant excitonic pathways (Figure 2f–g). Tri‐exponential global fitting yields time constants assigned to ultrafast delocalization (τ ₁), diffusion‐mediated hole transfer (τ ₂, where exciton diffusion to the donor–acceptor interface is a prerequisite for hole transfer), and bulk carrier recombination (τ ₃). As detailed in Tables S5 and S6, while τ ₁ is comparable between blends, the effective charge yield (A₂ + A₃) visualized in Figure 2h highlights a distinct behavior for DP10. The DP10:L8‐BO system exploits a dual‐channel hole transfer mechanism, combining a moderately slower transfer rate with significantly prolonged carrier lifetimes and enhanced overall yield. Moreover, the DP10:L8‐BO system exhibits longer exciton diffusion times (τ ₂ = 91.04 ps and 45.08 ps) compared to the DP1:L8‐BO system (τ ₂ = 10.93 ps and 2.86 ps). These carrier and exciton behaviors are attributed to the synergistic effect of the polarized TClTzE unit and suppressed phonon coupling, which stabilizes delocalized hole–electron distributions. In contrast, the π‐π stacking pathway in DP1:L8‐BO exhibits a negative charge yield (‐0.10), indicating dominant ground‐state recovery due to excessive donor‐acceptor intermixing, while the ICT pathway shows only a marginal yield of 0.01 due to these losses. Consistently, acceptor GSB kinetics (Figure S33) reveal faster delocalization and diffusion processes in the DP10 blend, confirming enhanced ordering of the L8‐BO phase. Overall, these results demonstrate that localized potential regulation effectively suppresses both static and dynamic disorder, thereby synergistically stabilizing charge generation and transport in the blend system.

2.2. Photovoltaic Properties and Carrier Dynamics

The photovoltaic performances of DP1‐ and DP10‐based devices were systematically compared using conventional OSC devices with an architecture of Glass/ITO/2‐PACz/donor: L8‐BO/PDINN/Ag. The current density–voltage (J–V) characteristics and photovoltaic parameters of the optimal devices are presented in Figure 3a,b and detailed in Table 1. Devices based on the DP1:L8‐BO blend achieved a PCE of 16.50%, with a V _OC of 0.888 V, a J _SC of 25.76 mA cm⁻², and a FF of 72.17%. In contrast, the DP10‐based devices exhibited markedly improved overall performance, achieving a higher PCE of 19.51%. This improvement arises from concurrent increases in all key photovoltaic parameters, most notably V _OC (0.918 V) and FF (79.08%), together with an elevated J _SC of 26.88 mA cm⁻². The reliability of the measured J _SC is corroborated by external quantum efficiency (EQE) spectra (Figure 3c), which show discrepancies of less than 5% between the integrated and experimentally obtained current densities.

FIGURE 3.

Photovoltaic performance and carrier dynamics of the DP1:L8‐BO and DP10:L8‐BO devices. (a) J‐V curves. (b) Box plots of PCE statistics. (c) EQE spectra. (d) J _ph and V _eff characteristics. (e) Charge mobilities derived from Photo‐CELIV measurements. (f) Dependence of V _OC versus P _light. (g) Charge lifetime as a function of charge density, derived from TPV and CE measurements. (h) Normalized TPC curves. (i) Energy loss analysis detailing the ∆E ₁, ∆E ₂, and ∆E ₃ components.

TABLE 1.

Photovoltaic parameters of optimized binary system active layers (thickness ≈ 110 nm) under the standard of AM 1.5G illumination (100 mW cm⁻²).

System	V OC (V)	J SC (mA cm−2)	JEQE SC (mA cm−2)	FF (%)	PCE (PCE a ) (%)
DP1:L8‐BO	0.888	25.76	24.53	72.17	16.50 (15.97 ± 0.42)
DP10:L8‐BO	0.918	26.88	26.14	79.08	19.51 (19.06 ± 0.22)

^a The statistics were obtained from over ten devices.

To elucidate the physical origins underlying the superior device performance of DP10 relative to the reference analogue DP1, a series of device physics characterizations was performed. These analyses validate the critical role of the localized potential regulation strategy in optimizing exciton dissociation and charge carrier dynamics. The exciton dissociation efficiency, η _diss, was evaluated from the dependence of the photogenerated current density (J _ph) on effective voltage (V _eff), where η _diss is defined as J _ph/J _sat (Figure 3d). The DP1:L8‐BO device exhibited an η _diss of 93.6%, whereas the DP10:L8‐BO device achieved a higher value of 98.2%, indicating more efficient exciton dissociation at the donor/acceptor interface. Charge transport properties were further examined via photo‐induced charge carrier extraction by linearly increasing voltage (Photo‐CELIV) measurements (Figure 3e and Figure S34). The DP10:L8‐BO system exhibited a higher carrier mobility of 2.49  ×  10⁻⁴ cm² V⁻¹ s⁻¹, exceeding that of the DP1:L8‐BO system (1.86 × 10⁻⁴ cm² V⁻¹ s⁻¹).

To further elucidate charge recombination dynamics and charge collection processes, steady‐state light‐intensity dependence measurements were integrated with transient kinetics analyses. The dominant recombination mechanism was first probed by analyzing the dependence of V _OC on light intensity (P _light), which follows the relation V _OC∝(nkT/q)lnP _light, where k is Boltzmann's constant, T is temperature, q is the elementary charge, and n denotes the ideality factor. As illustrated in Figure 3f, the DP10:L8‐BO device exhibits a slope of 1.26 kT q⁻¹, corresponding to an n close to unity, whereas the DP1:L8‐BO blend shows a larger slope of 1.34 kT q⁻¹. This reduction in n suggests a suppression of trap‐assisted Shockley–Read–Hall recombination. This conclusion is quantitatively corroborated by recombination order (R) analysis derived from the dependence of carrier lifetime on charge density τ  = τ₀ (n₀/n _c )^λ via transient photovoltage (TPV) and charge extraction (CE) measurements (Figure 3g and Figure S35) [61]. The extracted R value for DP10:L8‐BO is 2.07, which is much closer to the theoretical value of 2 for ideal bimolecular recombination compared to the value of 2.29 for DP1:L8‐BO. The convergence of both steady‐state and transient data confirms that the localized potential regulation strategy effectively minimizes energetic disorder and trap density. Moreover, the relationship $J_{\mathrm{SC}}\propto P_{light}^{\alpha }$ (Figure S36) yielded α values of 0.995 for DP10:L8‐BO and 0.991 for DP1:L8‐BO. The near‐unity α values indicate negligible bimolecular recombination losses under short‐circuit conditions for both systems. The slightly higher α value for DP10:L8‐BO is in agreement with its superior FF, suggesting that enhanced charge transport effectively mitigates space‐charge accumulation even under high photogeneration rates.

The interplay between charge extraction and carrier recombination provides direct insight into the origin of the enhanced device performance observed in the DP10:L8‐BO system. Transient photocurrent (TPC) measurements (Figure 3h) reveal a substantially shorter charge extraction time (τ _TPC) of 0.42 µs for DP10:L8‐BO, compared with 0.85 µs for the DP1:L8‐BO system, indicating more efficient charge collection pathways. Concurrently, TPV decay measurements (Figure S37) demonstrate a significantly prolonged carrier lifetime (τ _TPV) of 9.13 µs for the DP10:L8‐BO system, more than twice that of the DP1:L8‐BO system (4.33 µs). The synergistic effect of faster charge extraction and extended carrier lifetime enables more efficient carrier collection before recombination can occur. Consequently, the DP10:L8‐BO device achieves a superior charge collection efficiency (η _coll) of 88.9%, markedly surpassing the 73.8% obtained for DP1:L8‐BO (Figure 3d), thereby directly linking the optimized recombination dynamics to the superior photovoltaic performance.

Complementing the recombination analysis, energy loss characteristics were quantified using the equation E_loss = E_g   −  qV_OC , yielding total energy losses of 0.537 eV for the DP1:L8‐BO device and 0.499 eV for the DP10:L8‐BO device (Figure 3i and Figure S38). According to the detailed loss decomposition model (Table S7 and Figure 3i), the total loss comprises three components: $\mathrm{\Delta }{E}_1=E_g-q{V}_{OC}^{SQ}$, representing the unavoidable radiative loss above the bandgap; $\mathrm{\Delta }{E}_2=q{V}_{OC}^{SQ}-q{V}_{OC}^{Rad}$, denoting the radiative loss below the bandgap; and $\mathrm{\Delta }{E}_3=q{V}_{OC}^{Rad}-q{V}_{OC}$, corresponding to non‐radiative loss. The radiative‐limit open‐circuit voltage ($V_{OC}^{Rad}$) was obtained by fitting Fourier‐transform photocurrent spectroscopy external quantum efficiency (FTPS‐EQE) and electroluminescence (EL) spectra (Figure S38), giving values of 1.094 V for DP1:L8‐BO and 1.101 V for DP10:L8‐BO. These yield ΔE ₂ values of 0.065 eV and 0.051 eV, respectively. The most pronounced difference lies in ΔE ₃ where DP10:L8‐BO exhibits a reduced non‐radiative loss of 0.183 eV compared to 0.206 eV in DP1:L8‐BO. This confirms that incorporation of the TClTzE unit effectively suppresses non‐radiative recombination, thereby enhancing V _OC. Overall, the DP10:L8‐BO system outperforms the DP1:L8‐BO system across multiple critical processes: more efficient exciton dissociation, balanced and enhanced charge transport, suppressed non‐geminate and non‐radiative recombination, and reduced overall energy loss. These synergistic improvements collectively result in higher J _SC, FF, and V _OC, ultimately delivering superior photovoltaic performance.

2.3. Thick‐Film and Large‐Area Devices

To assess the practical implications of suppressed exciton–phonon coupling, enhanced charge polarization dissociation, and increased molecular ordering in DP10, the thermal stability of the corresponding active layer systems was investigated (Figure 4a). Under prolonged thermal stress at 120 °C, the DP10:L8‐BO system exhibited superior thermal robustness, retaining 80% of its initial efficiency (T ₈₀) for approximately 1000 h, significantly outperforming the thermal lifetime of the DP1:L8‐BO system (T ₈₀ ≈ 340 h). Tracking the photovoltaic parameters (Figure S39) revealed that while photocurrent generation J _SC remained robust in both systems, the DP1:L8‐BO system suffered from severe degradation in V _OC and FF. Conversely, the DP10:L8‐BO system maintained high retention of these parameters, highlighting the importance of the extended exciton diffusion length in preserving device efficiency despite thermally induced energetic disorder and morphological coarsening. Charge transport properties were further probed using SCLC measurements (Figure 4b and Figures S40, S41). The fresh DP10:L8‐BO system displayed higher and more balanced µ _h and electron mobility (µ _e), with µ _h/µ _e = 6.96/6.32  ×  10⁻⁴ cm² V⁻¹ s⁻¹, compared to the DP1 system (µ _h/µ _e = 3.28/4.66  ×  10⁻⁴ cm² V⁻¹ s⁻¹). Crucially, upon thermal aging, the DP10:L8‐BO system retained the majority of their charge transport capability (µ _h/µ _e = 6.61/5.69  ×  10⁻⁴ cm² V⁻¹ s⁻¹), whereas the DP1:L8‐BO system exhibited pronounced mobility degradation (µ _h/µ _e = 2.71/3.10 × 10⁻⁴ cm² V⁻¹ s⁻¹). This superior retention of balanced charge mobilities in DP10‐based devices effectively suppresses space‐charge accumulation, providing a kinetic rationale for the significantly improved FF stability observed under thermal stress.

FIGURE 4.

Thermal stability and morphological evolution of the DP1:L8‐BO and DP10:L8‐BO systems. (a) Normalized PCE retention after aging at 120°C for 1000 h. (b) Stability of charge mobilities (hole and electron) after thermal aging. (c) Out‐of‐plane (010) stacking parameters from GIWAXS analysis. (d) AFM height and phase images of blend films before and after aging. (e) Statistical distribution of domain widths extracted from AFM phase images.

Morphological analyses provide direct structural evidence for the effective suppression of both static and dynamic energetic disorder in the DP10:L8‐BO system. The film‐forming kinetics of the DP1:L8‐BO and DP10:L8‐BO systems during the drying process were characterized by in situ UV‐vis‐NIR spectroscopic measurements (Figure S42). Both systems exhibit a similar kinetic pathway: the acceptor first precipitates and undergoes a relatively long period (t₂), followed by rapid donor precipitation. In the DP1:L8‐BO system, the donor possesses a weaker pre‐aggregated state and thus precipitates more quickly, resulting in a shorter solvent evaporation time (t₁). This reduced time for self‑assembly likely leads to a disordered active layer morphology with increased donor–acceptor miscibility [62]. To probe the evolution of molecular packing under thermal stress, 2D GIWAXS measurements were performed on the blend films before and after thermal aging (Figure 4c and Figures S43, S44, Table S8). In the in‐plane direction, the (100) lamellar stacking diffraction peaks exhibit only marginal changes for both blends upon aging, indicating that the overall lamellar organization is largely preserved. Notably, in the out‐of‐plane direction, the (010) π–π stacking diffraction peak of the DP1:L8‐BO film displays enhanced molecular ordering after thermal aging, manifested by an increased CCL from 16.3 Å to 21.0 Å and a concomitant rise in diffraction intensity, approaching the molecular packing characteristics of the aged DP10:L8‐BO film. Despite this thermally induced increase in long‐range crystallinity, the DP1:L8‐BO system does not exhibit improved thermal stability at the device level. This observation underscores that enhanced crystallinity alone is insufficient to guarantee operational robustness. Instead, these results highlight the critical importance of the localized potential regulation strategy in suppressing energetic disorder, which is essential for maintaining efficient long‐range charge transport and device stability under thermal stress.

Atomic force microscopy (AFM) was employed to further investigate phase separation behavior (Figure 4d). Fresh films of both systems showed comparable root‐mean‐square (RMS) roughness (1.12 nm for DP1:L8‐BO vs. 1.16 nm for DP10:L8‐BO). Phase imaging (Figure 4d,e and Figure S45) revealed that the fresh DP10:L8‐BO film featured a well‐defined nano‐fibrillar network with slightly finer domain spacing (9.4 nm) relative to the DP1:L8‐BO film (8.9 nm). Upon thermal aging, the DP1:L8‐BO film underwent pronounced phase coarsening, resulting in highly heterogeneous, oversized domains (average size: 24.7 nm) and a substantial increase in surface roughness (3.75 nm). In contrast, the DP10:L8‐BO film retained a more uniform nanostructure, with moderately enlarged domains (18.2 nm) and a lower roughness of 3.54 nm. These observations indicate that the enhanced morphological stability of DP10:L8‐BO systems stems from its tailored molecular design. By suppressing phonon vibrations and stabilizing charge‐carrier dynamics, the DP10:L8‐BO active layer maintains efficient charge generation and extraction properties, even in the presence of thermally driven morphological coarsening. This structural resilience underpins the superior thermal stability of devices based on DP10 system.

Building on the suppressed non‐radiative recombination and robust charge transport characteristics, the efficacy of the localized potential regulation strategy was further evaluated in thick‐film active layer architectures (Figure 5a–c and Figure S46, Table 2). Upon increasing the active layer thickness from approximately 110 nm to 310 nm, the DP1:L8‐BO system exhibited pronounced performance degradation, with the PCE decreasing from 16.50% to 12.64% and the FF dropping sharply from 72.17% to 55.18%. In contrast, the DP10:L8‐BO system demonstrated remarkable thickness tolerance: the PCE decreased only slightly from 19.51% to 18.02%, while the FF remained high at 75.26%. This superior robustness originates from the synergistic suppression of both static and dynamic energetic disorder. Dynamically, reduced exciton–phonon coupling and enhanced local polarization promote exciton delocalization and diffusion. Structurally, GIWAXS analyses of thick films confirm that the DP10:L8‐BO system maintains tighter π–π stacking and longer CCL than the DP1:L8‐BO system (Table S8). Transmission electron microscopy (TEM) characterization (Figure S47) further reveals that the DP1:L8‐BO system displays a non‐uniform phase distribution, which originates from complex film‐formation dynamics induced by enhanced pre‐aggregation of DP1 during the drying process (Figure S48). Consistent with the thermal stability results, even in thick films where carrier extraction pathways are substantially extended, the DP10:L8‐BO system effectively suppresses space‐charge accumulation and recombination losses. Conversely, although the DP1‐based blend also exhibits increased packing order in thick films, it remains highly susceptible to energetic disorder‐induced trap formation, leading to severe losses in V _OC and FF.

FIGURE 5.

Thickness tolerance and scalability of the photovoltaic devices. (a) J–V curves of DP1:L8‐BO system with varying active layer thicknesses. (b) J–V curves of DP10:L8‐BO system with varying active layer thicknesses. (c) Dependence of PCE on active layer thickness. (d) J–V curves of small‐area device and large‐area module based on the ternary DP10:L8‐BO:BTP‐eC9 system. (e) Schematic diagram of the large‐area module architecture and the corresponding device photograph.

TABLE 2.

Photovoltaic parameters of optimized thick‐film and ternary system active layers under the standard of AM 1.5G illumination (100 mW cm⁻²).

System	V OC (V)	J SC (mA cm−2)	FF (%)	PCE (PCE a ) (%)
DP1:L8‐BO 206 nm	0.866	24.60	65.04	13.86 (13.50 ± 0.34)
DP1:L8‐BO 307 nm	0.877	26.13	55.18	12.64 (12.09 ± 0.47)
DP10:L8‐BO 192 nm	0.913	26.60	77.85	18.90 (18.64 ± 0.15)
DP10:L8‐BO 322 nm	0.901	26.58	75.26	18.02 (18.02 ± 0.08)
DP10:L8‐BO:BTP‐eC9 0.048 cm2	0.908	27.95	81.04	20.57 (20.32 ± 0.15)
DP10:L8‐BO:BTP‐eC9 15.1 cm2	7.24	3.08	77.39	17.20

^a The statistics were obtained from over ten devices.

The combined insights from thermal stability and thick‐film analyses demonstrate that a localized potential regulation strategy, through effective suppression of energetic disorder and stabilization of active layer morphology, imparts both exceptional thermal robustness and thickness tolerance. To further leverage the robustness of this binary system, a ternary system was constructed by incorporating BTP‐eC9 into the DP10:L8‐BO system (Figure 5d and Figure S49, Table S9). The DP10:L8‐BO:BTP‐eC9 (1:1:0.2, w/w/w) device delivers a boosted PCE of 20.57%, driven by an enhanced J _SC of 27.95 mA cm⁻² (verified as 27.59 mA cm⁻² by EQE integration, Figure S50) and a notably high FF of 81.04%. These performance metrics place DP10 among benchmark polymer donors such as PM6 and D18 (Figure S51, Table S10). Demonstrating the scalability of this system, a large‐area module comprising eight series‐connected sub‐cells with a total active area of 15.1 cm² was fabricated (Figure 5e), achieving an impressive PCE of 17.20% combined with a remarkable FF of 77.39% (Table 2). This result positions the DP10‐based module among the highest‐performing large‐area OSCs reported to date (Figure S52 and Table S11). Taken together, these results highlight localized potential regulation as a pivotal molecular design principle for developing high‐performance, process‐robust polymer donors compatible with industrial‐scale fabrication.

3. Conclusions

In summary, a localized potential regulation strategy is proposed to effectively suppress energetic disorder in polymer donors. By selectively incorporating a strongly electron‐deficient TClTzE unit into the DP1 backbone, the resulting polymer donor DP10 simultaneously enhances backbone rigidity, local electronic polarization, and molecular packing, while preserving favorable miscibility with the acceptor. This molecular‐level regulation leads to synergistic improvements in excitonic and charge‐transport processes, including reduced exciton binding energy, markedly suppressed exciton–phonon coupling, and stabilized intramolecular charge‐transfer states. As a result, DP10 exhibits prolonged exciton lifetimes, an extended exciton diffusion length of 19.9 nm, and balanced trap‐tolerant charge transport. These fundamental advantages translate directly into device‐level gains. Binary DP10:L8‐BO solar cells deliver a high PCE of 19.51%, characterized by simultaneously enhanced V _OC, J _SC, and FF, together with suppressed non‐radiative recombination. Importantly, mitigating energetic disorder confers exceptional thermal‐robustness: DP10‐based systems retain ≈80% of their initial efficiency after approximately 1000 h of thermal aging at 120°C and maintain high efficiency at active layer thicknesses up to 310 nm, demonstrating outstanding thickness tolerance and processing resilience. Beyond binary systems, the practical applicability of this strategy is further validated by a ternary device efficiency of 20.57% and a large‐area module efficiency of 17.20% with an impressive FF of 77.39% (15.1 cm²), underscoring its scalability and manufacturability. Overall, this study reveals that localized potential regulation provides a unifying design principle to reconcile high efficiency, thermal stability, and process tolerance in OSCs. These insights offer a robust molecular blueprint for the development of next‐generation polymer donors.

4. Experimental Section

See experimental details in the Supporting Information.

Author Contributions

L.‐Y.X., Y.M.S., R.S., and J.M. conceived and developed the ideas. Y.M.S. synthesized the polymer donors DP1 and DP10. R.S., Z.C.X., S.H.Z., and Y.G. fabricated the solar cells and performed the device characterization. L.‐Y.X. and Y.M.S. performed the theoretical calculations. Y.H.L. carried out the TEM experiments. X.K.W. and J.W.C. performed the dielectric constant measurements. B.X. carried out the GIWAXS experiments. X.H.W. and X.R.Y. participated in data analysis. C.Z. assisted with the theoretical calculations. L.‐Y.X. conducted the PL, AFM and other measurements. L.‐Y.X., Y.M.S., R.S., and J.M. contributed to the manuscript preparation. All authors have given approval to the final version of the manuscript.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: adma74059‐sup‐0001‐SuppMat.pdf.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.