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. 2026 Jun 16;65(34):e9060680. doi: 10.1002/anie.9060680

Quadruple Hydrogen‐Bonds Engineering for Intrinsically Stretchable and Healable Semiconducting Polymers

Yuanhe Gu 1,2, Sichun Wang 1, Yiran Liu 3, Zhen Yang 1, Wenhao Li 1, Junyi Lu 2, Jie Xu 2, Guangxin Gu 2, Yunqi Liu 1,3, Zhiyuan Zhao 3,✉, Cheng Zhong 4,✉, Zhengran Yi 1,2,✉, Yan Zhao 1,✉
PMCID: PMC13480495  PMID: 42299074

ABSTRACT

Developing intrinsically stretchable and healable semiconducting polymers with high charge‐carrier mobility is critical for next‐generation flexible electronics; however, integrating these conflicting functionalities remains a formidable challenge. Here, we report a “quadruple‐hydrogen‐bonds end‐capping” strategy to realize high‐performance stretchable and healable semiconducting polymers. By incorporating quadruple hydrogen‐bonds between end‐capping units linked with alkyl spacers into polymer backbone, we engineer a supramolecular architecture that achieves enhanced crystallinity and improved ordered packing with reduced π‐π stacking distance, and also superior stretchabillity with molecular‐ordering retention during stretching. Moreover, enhanced chain mobility together with dynamic and reversible and hydrogen‐bonding sites in the architecture contribute to efficient healing. Consequently, our designed semiconducting polymer exhibits a more than 2‐fold increase in mobility, while demonstrating stable mobility retention under strain, high mobility recovery after healing, and scalability in fully stretchable transistor arrays. This work provides an effective molecular design strategy for achieving simultaneous improvements in electrical performance, mechanical stretchability, and healing ability in organic electronics.

Keywords: healing, intrinsic stretchability, mobility, organic thin film transistors, semiconducting polymers

A quadruple‐hydrogen‐bonds end‐capping strategy constructs a supramolecular architecture in DPPTVT‐UPy that simultaneously and significantly enhances charge‐carrier mobility, stretchability, and healing ability, overcoming the conventional molecular‐packing trade‐off among these properties and enabling high‐performance stretchable and healable transistors.

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1. Introduction

Semiconducting polymers have become promising candidates as charge‐transport materials for the next generation of wearable and biomedical applications, due to their soft nature, low cost, and solution processability [1, 2, 3, 4, 5, 6]. To achieve practical application, semiconducting polymers are required to simultaneously possess excellent electronic performance, mechanical stretchability, and maintainable electronic performance during stretching [7, 8, 9]. Several strategies have been developed to realize stretchable semiconductors with impressive mobility in recent years, such as physical blending of nanofibers with elastomers or other additives [10, 11, 12], which has demonstrated some potential. However, these multicomponent systems require strict and specific means for preparation and manufacture [13], restricting the broad application of such approaches. Therefore, to achieve simple fabrication process and thereby highly‐integrated and flexible devices, developing intrinsically stretchable semiconducting polymers with good electronic performance has been regarded as the most promising approach [14, 15, 16, 17]. In addition, as a practical matter, semiconducting polymers are highly required to sustain their functions under long‐term and repeated deformations. But in most cases, multiple stretching‐releasing cycles will inevitably lead to degraded electronic performance of semiconducting polymers [18]. Therefore, it is also desired to develop healable semiconducting polymers that enable the recovery of their functions. Overall, exploring high‐mobililty, stretchable, and healable semiconducting polymers has emerged as a key future direction in flexible electronics.

For the molecular engineering of intrinsically stretchable high‐performance semiconducting polymers, varieties of effective strategies were explored, including nonconjugated‐spacers incorporation [19, 20, 21, 22], dynamic‐bonding incorporation [23, 24, 25, 26], terpolymerization [14, 16, 27, 28], and side‐chain modification [29, 30, 31, 32, 33], etc. The underlying principle of these design strategies is to accurately reduce the overall crystallinity of polymer thin film while retaining short‐range ordered molecular packing. In the thin film microstructure, amorphous regions facilitate strain energy dissipation under strain through molecular conformation change and alignment, while crystalline domains allow effective charge transport. To reduce and prevent the occurrence of the breakage of crystallites, the distribution of small crystalline domains in amorphous regions is promising to acquire high stretchability and good charge‐transport performance. However, most strategies enhanced the stretchability of semiconducting polymers at the expense of the electrical performance, resulting from disrupted conjugation and weakened ordered molecular packing. Consequently, stretchable semiconducting polymers still lag significantly behind their rigid, highly crystalline counterparts in charge carrier mobility. Basically, it remains challenging to break the shackles of conflicting requirement of stretchability and charge‐transport performance on their aggregation structures. Furthermore, although several design strategies for constructing healable semiconducting polymers were explored in recent years, such as introduction of dynamic bonds, the polymers that concurrently exhibit high mobility and robust healing remain limited [25, 26, 34]. Thus, a design strategy that can simultaneously maximize mobility, stretchability, and healing ability represents a grand challenge in flexible electronics.

Introducing dynamic noncovalent bonds is an effective approach to develop stretchable and healable polymers [35, 36]. For instance, an in situ continuous hydrogen‐bond engineering strategy was employed to develop high‐performance stretchable and healable semiconducting polymers, achieving an 81% mobility recovery after healing and a mobility of 1.08 cm2 V−1 s−1 under 100% strain [37]. Zhen et al. introduced hierarchical hydrogen‐bonded conjugation breakers into the polymer backbone to construct multilevel dynamic interconnected polymer network, enabling the concurrent realization of high charge‐carrier mobility, excellent stretchability and remarkable self‐healing ability. Notably, an outstanding charge‐carrier mobility of 1.01 cm2 V−1 s−1 was achieved in fully stretchable transistors under 150% strain, establishing a new performance benchmark for intrinsically stretchable and healable semiconducting polymers [38]. Owing to their weak binding strength, noncovalent bonds are easily broken and reformed reversibly, which facilitates energy dissipation under strain and recovery of mechanical properties [39, 40, 41]. However, the considerably weak binding strength of noncovalent bonds is not beneficial to achieving excellent mechanical properties [42]. Furthermore, excessive dynamic bonding could induce extreme supramolecular aggregates that are detrimental to effective strain dissipation [43]. Motivated by these considerations, we propose a novel design strategy for developing intrinsically stretchable and healable semiconducting polymers with high mobility, through end‐capping incorporation of high‐binding‐strength noncovalent bonds. The end‐capping incorporation can endow semiconducting polymer with low‐content noncovalent bonds, eliminating excessive dynamic bonding. Moreover, we hypothesized that the incorporation of chain extenders would not affect longer‐range orders and overall crystallinity of the polymer, but rather a probably increasing trend due to the strengthened molecular interaction induced by noncovalent bonds. For demonstration, 2‐Ureido‐4[1H]‐pyrimidinone (UPy) was used as end‐capped units and introduced into the backbone of a common semiconducting polymer thienylene‐vinylene‐thienylene‐diketopyrrolopyrrole (DPPTVT) with high charge transport performance. In this structure system, the specially introduced UPy units enable the formation of high‐binding‐strength quadruple hydrogen bonds between the UPy unit [44, 45], constructing supramolecular polymer DPPTVT‐UPy (Figure 1a). In addition, alkyl spacers are linked with the UPy units to enhance their flexibility. This supramolecular architecture was engineered to impart high stretchability and enable efficient healing (Figure 1b). To better illustrate the superiority of our design strategy, reference polymer DPPTVT without UPy units was also prepared. Contrary to the typical reduction in performance for many strategies, we report that this engineering significantly enhances charge mobility—achieving a more than two‐fold increase over DPPTVT in organic thin‐film transistors (OTFTs). This improvement was accompanied by superior mechanical properties, including higher crack onset strain and lower tensile modulus. Crucially, DPPTVT‐UPy‐based transistors exhibited excellent durability, maintaining stable mobility under strain in both parallel and perpendicular directions. Furthermore, the supramolecular architecture endows the polymer with efficient solvent‐ and thermal‐healing capabilities. Hence, this strategy enables simultaneous and significant improvements in electrical performance, mechanical stretchability, and healing ability (Figure 1c), which provides a promising avenue for the development of high‐performance stretchable and healable semiconducting polymers.

FIGURE 1.

FIGURE 1

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Molecular design and properties of semiconducting polymers. (a) Chemical structures and schematic illustrations of supramolecular polymer DPPTVT‐UPy. (b) Schematic representation of the intrinsically stretchable and healable polymer film. (c) Comparison of key performance metrics between DPPTVT‐UPy and DPPTVT.

2. Results and Discussion

2.1. Material Synthesis and Characterization

The synthetic route to UPy derivative (UPy‐Br) and polymers (DPPTVT‐UPy and DPPTVT) is illustrated in Scheme 1. The synthesis of UPy‐Br started from the α‐alkylation of thiophene (1), giving 2‐(6‐bromo‐hexyl)‐thiophene (2). Then, potassium phthalimide was alkylated with compound 2 to afford amine precursor (3), followed by Ing‐Manske deprotection [46] to give thioalkylamine (4). Compound 4 was then subjected to an amidation reaction using di‐tert‐butyldicarbonate [(Boc)2O] to obtain Boc‐protected thioalkylamine (5). Subsequent reaction with N‐Bromosuccinimide (NBS) yielded Boc‐protected 5‐bromothioalkylamine (6). Trifluoroacetic acid was used to deprotect Boc groups of 6, furnishing 5‐bromothioalkylamine (7). The desired compound UPy‐Br was successfully synthesized by reacting compound 7 with N,N’‐Carbonyldiimidazole‐activated 2‐amino‐4‐hydroxy‐6‐methylpyrimidine (8) [47]. A Stille coupling polymerization between M1 and M2 generated reference polymer DPPTVT. The target polymer DPPTVT‐UPy was prepared using the similar synthetic method, the only difference being the stoichiometric ratio of monomers and additional addition of UPy‐Br in the synthetic process of DPPTVT‐UPy. Specifically, the stoichiometric ratio of M1, M2, and UPy‐Br was strictly controlled at 1.05: 1: 0.1 during the synthesis of DPPTVT‐UPy. This ratio was chosen to satisfy the stoichiometric balance required for effective end‐capping [48, 49], while simultaneously achieving an M n close to that of the reference DPPTVT to ensure a fair comparison of their performance [50]. All the intermediates and polymers were acquired with satisfactory yields, and their detailed synthetic procedures are described in Supporting Information. The molecular structures of the intermediates were verified by NMR spectroscopy and mass spectrometry (Figures S1–S22). The molecular weights of the polymers were evaluated through high‐temperature gel permeation chromatography (HT‐GPC). Both polymers exhibited comparable molecular weight with M n of 26.0 and 29.5 kDa for DPPTVT and DPPTVT‐UPy, respectively (Figures S23 and S24), which facilitates a clear and unambiguous exploration of the contribution of quadruple hydrogen‐bonds engineering to the electrical performance, mechanical stretchability, and healing ability of semiconducting polymer. X‐ray photoelectron spectroscopy (XPS) was used to detect the end‐capped UPy units. Figure S25a,b show clear Br 3d and Sn 3d5/2 peaks for DPPTVT, which are entirely absent in DPPTVT‐UPy. Combined with the higher nitrogen content of DPPTVT‐UPy relative to DPPTVT (Figure S25c,d), these observations indicate that DPPTVT retained residual bromine and tin species at the polymer chain termini, whereas DPPTVT‐UPy was effectively end‐capped with UPy units [51, 52].Moreover, high‐resolution solid‐state 1H NMR spectrum of DPPTVT‐UPy (Figure S26) showed three characteristic peaks (10.0–14.0 ppm) in the downfield region, which are assigned to amide protons in the UPy dimers [42]. The results demonstrate the formation of quadruple hydrogen bonds between the UPy units for DPPTVT‐UPy, thereby constructing supramolecular polymer structure [53].

SCHEME 1.

SCHEME 1

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The synthetic route to UPy derivative and polymers.

Thermal properties of DPPTVT and DPPTVT‐UPy were investigated through differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA), which reveals high decomposition temperature exceeding 400°C and occurrence of no distinct phase transition for both polymers (Figures S27 and S28). Also, cyclic voltammetry (CV) measurements were conducted to explore their electrochemical properties. Both polymers displayed stronger oxidative peaks relative to their reductive peaks (Figure S29). From the onset oxidation potential, the highest occupied molecular orbital (HOMO) energy levels of DPPTVT and DPPTVT‐UPy were determined to be –5.32 and –5.37 eV, respectively. The optical bandgap (E g opt) of the polymer films can be estimated from their absorption onsets in UV–vis absorption spectroscopy (Figure S30, 1.35 and 1.34 eV for DPPTVT and DPPTVT‐UPy, respectively). The E g opt and the HOMO levels gave the lowest unoccupied molecular orbital (LUMO) energy levels of −3.97 eV (DPPTVT) and −4.03 eV (DPPTVT‐UPy).

2.2. Improved Mechanical Properties and Electrical Performance

To understand the contribution of the incorporated quadruple hydrogen bonds to the mechanical property of the polymer, we first investigated the crack onset strain of the polymers, which is one of the most important indicators in stretchable semiconducting polymers. According to the film on elastomer (FOE) method [52], the crack onset strain of films was measured through stretching the film on PDMS substrates under different strains, and then detected micron‐scaled cracks from optical microscopy. As shown in Figure S31, it could be observed that the films gradually began to crack as the increasing of the strain. The cracks emerged within the DPPTVT film at 90% strain, while the DPPTVT‐UPy film exhibited higher crack onset strain of up to 120%, indicating higher stretchability of DPPTVT‐UPy. Moreover, evident wrinkles were observed in DPPTVT films after relaxation from 100% strain. In comparison, only scattered small wrinkles were present in DPPTVT‐UPy films after relaxation from 140% strain, implying better mechanical reversibility of DPPTVT‐UPy relative to DPPTVT. To acquire more insight into the influence of the incorporated quadruple hydrogen bonds on the mechanical property, we also estimated elastic modulus, elongation at break and toughness of polymer films. The strain–stress curves of polymer films were acquired using film on water (FOW) method (Figures 2a and S32) [54], and the corresponding mechanic parameters of polymer films were extracted from strain–stress curves. As shown in Figure 2b, DPPTVT‐UPy possessed lower elastic modulus of 113 MPa, greater elongation at break of 32.2%, and higher toughness of 2.1 MJ/m3 (vs. 209 MPa, 15.7% and 1.2 MJ/m3 for DPPTVT, respectively.), which helps to explain why the DPPTVT‐UPy film has higher crack onset strain relative to the DPPTVT film. The good mechanical properties of DPPTVT‐UPy are favorable for achieving high‐performance stretchable OTFT devices. Note that DPPTVT‐UPy showed a lower elastic modulus compared with its reference polymer DPPTVT. In general, elastic modulus of the polymer film increases when intermolecular hydrogen bonds link the chains of polymer [55]. But for our designed structure system of supramolecular polymer DPPTVT‐UPy, we inferred that the incorporation of alkyl spacers linked with the UPy units could counteract the effect of the introduced low‐content hydrogen bonds on the flexibility of polymer, leading to lower elastic modulus.

FIGURE 2.

FIGURE 2

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Mechanical properties and electrical performance of semiconducting polymers. (a) Schematic depiction of film‐on‐water (FOW) method. (b) The corresponding elongation at break, elastic modulus, and toughness derived from strain–stress curves. (c) Storage modulus (E′), loss modulus (E″), and tanδ of films characterized by dynamic mechanical analysis (DMA) under a temperature‐sweep mode. (d) Schematic diagram of bottom gate/bottom contact (BGBC) OTFT device structure. (e) Representative output and transfer curves of DPPTVT‐UPy film by on‐center spin‐coating process. (f) The extracted mobilities of DPPTVT and DPPTVT‐UPy films characterized in BGBC transistors.

The glass transition temperature (T g) of polymers is another considerable parameter to assess the degree of ductility [56]. In general, polymer with lower backbone T g is more conducive to plastic deformation during stretching. Dynamic mechanical analysis (DMA) was employed to to determine the T g values of DPPTVT and DPPTVT‐UPy. As shown in Figure 2c, DPPTVT‐UPy showed a lower backbone T g than DPPTVT, which is likely attributed to the end‐capping incorporation of alkyl spacers linked with the UPy units. The decrease of backbone T g suggests an increase of free volume of polymer chains, which could explain higher mechanical deformation and ductility for DPPTVT‐UPy.

To evaluate the electrical performance of DPPTVT and DPPTVT‐UPy, bottom‐gate/bottom‐contact (BGBC) OTFT devices were fabricated (Figure 2d). Heavily doped Si and a 300 nm thick SiO2 layer functioned as the gate and dielectric, respectively, with Au source/drain electrodes. Following standard cleaning and surface passivation with octadecyltrichlorosilane (OTS) via vapor deposition, the active layers were deposited using both on‐center and off‐center spin‐coating methods [57]. After a thermal annealing step, the device characteristics were evaluated. Typical transfer and output curves for both fabrication approaches are presented in Figures 2e, S33 and S34, with the corresponding performance summarized in Figure 2f and Table S1. For on‐center spin‐coating method, DPPTVT‐UPy showed hole mobility of 1.7 cm2 V−1 s−1, which is almost twice as high as that of the reference polymer DPPTVT (0.9 cm2 V−1 s−1). When off‐center spin‐coating method was used to cast semiconductor layer, OTFT devices displayed better electrical performance than those using on‐center spin‐coating methods, which can be attributed to the induction of highly aligned polymer chain. The increase of field‐effect mobility for DPPTVT‐UPy was more than that for DPPTVT. In comparison, off‐center spin‐coated DPPTVT‐UPy film exhibited remarkably improved electrical performance with the mobility of up to 3.4 cm2 V−1 s−1, which is over twice that of DPPTVT. Remarkably, this strategy yields one of the most significant improvements in charge‐carrier mobility among reported approaches for stretchable semiconducting polymers (Figure S35).

2.3. Molecular Packing

Basically, molecular packing of semiconducting polymers significantly influences their electrical performance and mechanical properties [14]. We first examined the influence of the introduced quadruple hydrogen bonds on overall molecular ordering of polymer thin films through two‐dimension grazing‐incidence wide‐angle x‐ray scattering (2D GIWAXS). Figure 3a–c display the 2D GIWAXS patterns and the corresponding 1D in‐plane and out‐of‐plane line cuts. Both (h00) alkyl chains stacking signals and (010) π‐π stacking signals were observed in the out‐of‐plane and in‐plane directions, respectively, suggesting edge‐on packing modes for both polymers. Table S2 summarizes the detailed structural parameters of the polymers extracted from the GIWAXS patterns. Compared with DPPTVT, the lamellar packing distance and π‐π stacking distance of DPPTVT‐UPy were both reduced from 23.6 to 23.2 Å, and from 3.50 to 3.45 Å, respectively. This can be ascribed to the incorporation of the quadruple hydrogen bonds, inducing the strengthened molecular interactions and thereby leading to densely packed chains for DPPTVT‐UPy. The better molecular packing may enable higher crystallinity of polymer thin film. The relative degree of crystallinity (rDoC) was used to quantitatively evaluate the crystallinity of the polymers [27]. As shown in Table S2, DPPTVT‐UPy had higher rDoC than DPPTVT, indicating higher crystallinity or longer‐range order for DPPTVT‐UPy [14]. Specially, most of design strategies for achieving low elastic modulus are based on some sacrifice of the overall crystallinity of polymer thin film. By contrast, our design strategy can endow DPPTVT‐UPy with higher crystallinity, which is the key to achieving good electrical performance. Furthermore, based on the Scherrer equation [58], we also estimated the coherence lengths (CCL) of the polymers. The results show that the CCL of DPPTVT‐UPy is longer than that of DPPTVT. The long CCL suggests that DPPTVT‐UPy chains have a tendency to form large crystalline domains.

FIGURE 3.

FIGURE 3

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Molecular packing properties of semiconducting polymers. (a) 2D GIWAXS patterns of polymer thin films. The corresponding in‐plane (b) and out‐of ‐plane (c) line cuts of films from 2D GIWAXS patterns. Normalized UV–vis absorption spectra of polymer solutions (d) and films (e). Relative degree of crystallinity (rDoC) (g) and coherence lengths (CCL) (h) change under strain.

The aggregation states in solutions and films of DPPTVT and DPPTVT‐UPy were evaluated using UV–vis absorption spectroscopy. As shown in Figure 3d,e, both DPPTVT and DPPTVT−UPy exhibited dual absorption bands in the 400−500 nm and 700−800 nm regions, which are attributed to the π‐π* and intramolecular charge transfer (ICT) transitions, respectively. Generally, the ICT absorption results from polymer aggregation or interchain stacking [59]. The 0−0 peak positions are presented in Figure S36. Compared with the reference polymer DPPTVT, the 0−0 peak position of DPPTVT‐UPy was re‐shifted from 782.0 to 783.5 nm in solution, while re‐shifted 0−0 peak was also observed from 786.0 to 789.5 nm for the DPPTVT‐UPy film. Furthermore, both polymers exhibited a strong 0−0 peak accompanied by a shoulder 0−1 peak. The relative degree of aggregation can be evaluated by comparing the absorption intensity ratio between the 0−0 peak (A 0−0) and the 0−1 peak (A 0−1) [60]. As shown in Figure S37, DPPTVT‐UPy had a larger A 0−0/A 0−1 value than DPPTVT both in solution and thin film state. All these results indicate more ordered interchain packing and aggregation for DPPTVT‐UPy [16]. Together with the 2D GIWAXS results, we deduced that the end‐capping incorporation of the quadruple hydrogen bonds into the backbone and thereby the formation of supramolecular structure, could effectively induce the strengthened interchain interactions, which enables more ordered packing with shorter π‐π stacking distance and higher crystallinity in DPPTVT‐UPy film.

Moreover, the polymer chain orientation in films under various strain was investigated by comparing dichroic ratios from polarized UV–vis spectra (Figures S38 and S39). As the strain increases, dichroic ratios of both DPPTVT and DPPTVT‐UPy gradually rose with the increases of strain, which is due to strain‐induced alignment of chains. Specially, the increase in dichroic ratios of DPPTVT‐UPy is more distinct than that of DPPTVT, which may facilitate anisotropic charge transport. The evolution of crystalline domains was also systematically characterized by analyzing three key parameters derived from GIWAXS data: the lamellar stacking distance, rDoC, and CCL (Figures S40 and 3f‒h). The films were transferred onto a SEBS substrate, mechanically stretched to a specific strain, and subsequently exposed to an x‐ray beam for data acquisition. In the direction parallel to stretching, the DPPTVT film showed strain‐dependent increase in lamellar stacking distance, whereas the DPPTVT‐UPy film exhibited increased lamellar stacking distance with applied strain from 0% to 50%, after which it stabilized and remained constant at higher strains, primarily attributable to its high toughness. For rDoC and CCL, DPPTVT exhibited larger reduction under strain compared to DPPTVT‐UPy, suggesting the higher occurrence of breakage of crystalline domains for DPPTVT in the dissipation process. In addition, more distinct anisotropy in rDoC and CCL under strain was observed for DPPTVT‐UPy relative to DPPTVT, which could be attributed to higher degree of strain‐induced chain alignment for DPPTVT‐UPy, as evidenced by polarized UV–vis spectra. Notably, the DPPTVT‐UPy film displayed an increase in rDoC and CCL along the stretching direction, implying the reorientation of crystalline domain under strain [15]. Taken together, the results imply that the two polymers have multiple mechanisms for strain energy dissipation: (i) chain elongation and alignment in amorphous regions; (ii) crystalline reorientation and alignment; (iii) breakage of crystalline domains. For DPPTVT‐UPy, the first two mechanisms dominate the dissipation process, whereas all the three mechanisms are operative for DPPTVT.

2.4. Application in Stretchable and Healable Transistors

To investigate the electrical performance of the polymer films under various strains, we fabricated stretchable transistor devices using a bottom‐gate/top‐contact (BGTC) architecture. In the device fabrication procedures, semiconducting polymer films were fabricated on OTS‐treated Si/SiO2 substrates and subsequently annealed. The films were then peeled using a PDMS elastomer, mechanically stretched to the desired strain, and transferred onto SEBS‐coated Si/SiO2 substrates. The devices were completed by depositing gold source and drain electrodes (Figures 4a and S41). The transfer and output characteristic curves of the devices under various strains are presented in Figures 4b and S42–S45. The electrical performance was summarized in Table S3. Figure 4c displays the relationship between normalized carrier mobility and various strains. For the DPPTVT film, the mobility along the stretching direction was nearly unchanged, while the mobility perpendicular to the stretching direction decreased by more than 60%. In contrast, the charge mobilities of DPPTVT‐UPy were gradually improved as the increasing of strain at parallel direction. Furthermore, the carrier mobilities perpendicular to the stretching direction for DPPTVT‐UPy were basically stable at a level. Taken together, DPPTVT‐UPy exhibited higher mobilities than DPPTVT when subjected to strain at both perpendicular and parallel directions. These results were unified with the dichroic ratios from polarized UV–vis spectra. Compared with DPPTVT, the greater increase in dichroic ratios of DPPTVT‐UPy may be propitious to anisotropic charge transport, thereby leading to improved charge mobilities with the increase of strain at parallel direction [27]. At the perpendicular direction, the greatly decreased mobility of DPPTVT film during stretching, most likely attributed to the formation of cracks under strain and thereby blocking‐up of charge‐transport pathways. In addition, we investigated the mobility evolution of polymer films under various stretching‐releasing cycles at 25% strain to evaluate their durability. Figures 4d and S46–S49 present the normalized mobilities and the corresponding representative transfer and output curves. The results demonstrate that the DPPTVT‐UPy films outperform their DPPTVT counterpart in terms of durability. In particular, after 500 stretching‐releasing cycles under 25% strain, the average charge mobility of the DPPTVT‐UPy film in the perpendicular to the stretching direction remains above 90%.

FIGURE 4.

FIGURE 4

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Electrical performance of semiconducting polymers under various strain and fully stretchable transistor devices. (a) Device configuration of stretchable transistor device based on semiconducting film. (b) Transfer curves of DPPTVT‐UPy film parallel to the stretched direction under various strain. (c) Normalized mobility of semiconducting films under various strain. (d) Normalized mobility of semiconducting films under various stretching‐releasing cycles at 25% strain. (e) Schematic illustration of the fully stretchable transistor architecture. Output (f) transfer curves (g) of a fully stretchable transistor array fabricated with DPPTVT‐UPy (dashed lines, gate current). (h) Optical images of a fully stretchable 5 × 5 transistor array under strain from 0% to 150%. Inset: photograph and device configuration of the transistor array. (i) Mapping statistical distribution of on‐current in the fully stretchable transistor array. Transfer curves (j), mobility and on‐current (k) change of a fully stretchable transistor device under various strain.

Leveraging the excellent mobility retention of DPPTVT‐UPy films under strain, we fabricated fully stretchable transistor arrays (5 × 5) in BGBC configuration. In this architecture, carbon nanotubes (CNTs) were employed as the gate and source/drain electrodes, while SEBS functioned as both the dielectric layer and the elastomeric substrate (Figure 4e,h). The fabrication process achieved a 100% device yield, with the resulting arrays exhibiting high switching performance and exceptional uniformity in charge‐carrier mobility (Figures 4f–i and S50). To evaluate their stretchability, the transistors were subjected to various strain conditions. Notably, the carrier mobility remained virtually unchanged upon stretching to 150% strain parallel to the charge transport direction (Figure 4j,k and Table S4). This represents one of the highest mobility‐retention capabilities under mechanical strain among reported stretchable semiconducting polymers, rivaling even that of some conjugated polymer/elastomer blends (Figure S51). The observed minor reduction in on‐current is primarily attributed to the strain‐induced increase in channel length [10].

Besides intrinsic stretchability, the healing ability of semiconducting polymers in transistor devices was investigated. In this process, polymer thin films were intentionally damaged to generate nano‐cracks by stretching, followed by the healing of the damaged films via post‐treatments. Figure 5a shows the AFM height images for pristine, damaged and healed films of DPPTVT and DPPTVT‐UPy. It was observed that all the initial thin‐film images were evenly distributed to form a uniform network. Also, to reasonably contrast the healing properties of the polymers, we controlled the films by stretched strain to form similar sized cracks, approximately 200 nm wide. DPPTVT and DPPTVT‐UPy films were subjected to strain of approximately 25% and 35%, respectively. From AFM images and height distributions of healed films, nano‐cracks in the damaged DPPTVT‐UPy film were clearly healed after post‐treatment of vapor and heat, whereas the damaged DPPTVT film did not completely heal. Then, we prepared transistor devices to test the charge mobility of these pristine, damaged, and healed films. Representative output and transfer curves for DPPTVT and DPPTVT‐UPy are shown in Figures 5b,c, S52 and S53, respectively, and the saturated mobility was extracted and summarized in Table S5. The mobility was normalized for comparison those of damaged and recovery films in Figure 5d. Consistent with the AFM images, the post‐treated healed DPPTVT‐UPy film demonstrated substantial mobility recovery to 84% of its pristine value, in stark contrast to DPPTVT's 53% recovery. The above results unambiguously demonstrate the improved healing ability of DPPTVT‐UPy in transistor device relative to DPPTVT.

FIGURE 5.

FIGURE 5

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Healing properties of semiconducting polymers. (a) AFM images for pristine, cracked, and self‐healed based on post‐treatments films. Inset parts show the height distribution of the films. (b) The transfer curves for pristine, cracked and healed based on post‐treatments films of DPPTVT. (c) The transfer curves for pristine, cracked, and self‐healed based on post‐treatments films of DPPTVT‐UPy. (d) The mobility of the pristine, cracked, and self‐healed based on post‐treatments films.

2.5. The Proposed Mechanism for Boosted Charge Mobility, Stretchability, and Healing Ability

We proposed that the supramolecular architecture, constructed by end‐capping incorporated quadruple hydrogen bonds in the polymer backbone, was critical for achieving simultaneous and significant improvement in stretchability, electrical performance, and healing ability. The supramolecular architecture could effectively induce the strengthened interchain interactions, which greatly facilitates molecular packing and thereby significantly enhances charge transport properties. Also, the introduced quadruple hydrogen bonds with alkyl spacers in this architecture endow DPPTVT‐UPy with higher crack onset strain, lower elastic modulus, greater elongation at break, higher toughness, and low T g, which contributes to improved stretchability. At the stage of tensile deformation, the dissipation of strain energy is facilitated by mechanisms such as chain stretching, sliding, and alignment. The introduced alkyl spacers for DPPTVT‐UPy enable increased free volume of polymer chains, leading to improved stretchability under strain. And the chain sliding and alignment occurred within the amorphous regions, accompanied by the reorientation and alignment of crystalline domains along the stretching direction. Note that DPPTVT‐UPy had higher chain alignment degree relative to DPPTVT, which could facilitate anisotropic charge transport. Meanwhile, the introduced quadruple hydrogen bonds in the supramolecular architecture render DPPTVT‐UPy high toughness, which enables less decrease of crystalline order under strain compared to DPPTVT, as confirmed by GIXRD. Moreover, the supramolecular architecture of DPPTVT‐UPy enhances chain mobility (evidenced by a lower T g, relative to DPPTVT) and provides dynamic, reversible hydrogen‐bonding sites that enable efficient healing [61].

To investigate our proposed mechanism, we performed molecular dynamics (MD) simulations of the polymers. Figure 6a,b present representative snapshots from the MD trajectories for DPPTVT (Figure 6a, left panel) and DPPTVT‐UPy (Figure 6b, right panel), where the terminal UPy groups are distinctly highlighted in yellow. The square boxes illustrate the initial, unstretched configurations, while the elongated rectangular boxes display the systems under 40% applied strain. From these visualizations, three critical observations can be made: In the unstretched state, the DPPTVT‐UPy film exhibited a noticeably higher degree of parallel polymer chain orientation compared to the DPPTVT film. Upon stretching, the polymer chains in both polymers predominantly aligned parallel to the stretching direction, a behavior that is consistent with the limited change observed in their mobility along this axis. Significantly, distinct voids were observed for DPPTVT after stretching, whereas DPPTVT‐UPy maintained a relatively dense and continuous structure. This improved structural integrity in DPPTVT‐UPy is directly attributed to the robust supramolecular network formed by its quadruple hydrogen bonds. To further elucidate the energy dissipation mechanism and investigate whether the reversible breaking of hydrogen bonds contributes to the enhanced fracture strain, we monitored the evolution of H‐bond counts in the DPPTVT‐UPy system during the stretching process (Figure S54). Interestingly, the number of H‐bonds remains remarkably stable, consistently fluctuating between 180 and 195 without a noticeable downward trend even at high strain levels. This indicates that the quadruple hydrogen bonds are largely preserved during stretching rather than being continuously broken and reformed. Therefore, in this system, the quadruple hydrogen bonds do not primarily act as sacrificial bonds for energy dissipation. Instead, they function as robust physical hydrogen‐bonding sites that maintain the integrity of the polymer network, facilitate efficient stress distribution, and prevent premature chain sliding and localized stress concentration. Consequently, the strain energy is mainly dissipated through chain stretching, sliding, and alignment, while the stable quadruple‐hydrogen‐bonding junctions enable the polymer network to withstand larger deformation before fracture. In addition, we fitted the stress–strain curves of the polymer thin films obtained during the dynamic stretching process (Figure S55). The calculated Young's moduli for DPPTVT and DPPTVT‐UPy were found to be 810 and 670 MPa, respectively. This trend is consistent with our experimental observations, although the calculated values are notably higher than the experimental ones. This discrepancy arises from the significantly faster stretching speeds employed in MD simulations relative to experimental tests. To better isolate the effects of hydrogen bonding and chain flexibility on elastic modulus, we also performed MD simulations on a new control polymer, DPPTVT‐Hexyl, which contains end‐capped alkyl spacers analogous to those in DPPTVT‐UPy but lacks the UPy units (Figure S56a). The fitted stress–strain curve of DPPTVT‐Hexyl obtained during the dynamic stretching process is shown in Figure S56b. The calculated Young's modulus of DPPTVT‐Hexyl is 660 MPa, which is only slightly lower than that of DPPTVT‐UPy but markedly reduced compared to that of DPPTVT. These results indicate that the incorporation of flexible alkyl spacers plays the dominant role in reducing the elastic modulus of the supramolecular polymer DPPTVT‐UPy, while the hydrogen bonding contributes a modest stiffening effect.

FIGURE 6.

FIGURE 6

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Molecular dynamics (MD) simulations of semiconducting polymers. (a) Representative snapshots from the MD trajectories for DPPTVT (left panel). (b) Representative snapshots from the MD trajectories for DPPTVT DPPTVT‐UPy (right panel). (c) Interchain stacking frequency change under strain. (d) Interchain stacking atomicity under strain. (e) Interchain stacking frequency and atomicity after healing. (f) Maximum pore diameter under strain. (g) Maximum pore diameter after healing.

Moreover, we investigated the interchain stacking of the polymers during both dynamic stretching and the subsequent healing process. Figure 6c–e illustrate the changes in two key parameters quantifying interchain stacking strength: interchain stacking frequency and atomicity, respectively. The stacking frequency counts the number of π‐π stacking events between planar residues from different polymer chains, where a stacking event is identified when any pair of nonhydrogen atoms is closer than 3.6 Å and the dihedral angle between the two planes is less than 35°. The atomicity quantifies the total number of atom pairs participating in interchain stacking, counting only atom pairs closer than 3.6 Å within stacked residue pairs; when an atom participates in multiple pairings, only the shortest distance is counted. Given that polymer chains tend to orient along the stretching direction, the strength of these interchain stacking interactions plays a crucial role in influencing the chain mobility perpendicular to the stretching axis. Our analysis reveals: in the unstretched state, DPPTVT‐UPy exhibited stronger interchain stacking compared to DPPTVT. During stretching, the interchain stacking in DPPTVT degraded more rapidly than in DPPTVT‐UPy. After healing, DPPTVT‐UPy also had stronger interchain stacking relative to DPPTVT. This comparative analysis demonstrates that DPPTVT‐UPy possesses superior retention and healing capabilities for its interchain stacking throughout the dynamic stretching and recovery cycles.

Finally, we also examined the evolution of pore size during the stretching process. Initially, the average pore radius appears to decrease as existing pores become elongated. However, further stretching leads to the emergence of larger voids within the system. After the healing process, these voids contracted. Notably, DPPTVT‐UPy consistently showed smaller initial pore sizes, experienced a lesser degree of pore enlargement during stretching, and exhibited a more effective reduction in pore size after healing compared to DPPTVT (Figure 6f,g). This pore size evolution is closely coupled with the degradation of interchain stacking, exhibiting a distinct transition at approximately 20% strain. Below 20% strain, chain alignment compresses the free volume (decreasing maximum pore size, Figure 6f) while keeping π‐π stacking largely intact (Figure 6c,d). Beyond 20% strain, further stretching forcefully pulls adjacent chains apart, disrupting the interchain stacking and simultaneously opening up larger internal voids (Figure 6f). Crucially, the quadruple hydrogen bonds in DPPTVT‐UPy act as a robust dynamic network that resists excessive chain separation, thereby significantly mitigating both interchain stacking degradation and pore enlargement compared to DPPTVT.

3. Conclusion

In summary, we have developed a “quadruple‐hydrogen‐bonds end‐capping” strategy to realize intrinsically stretchable and healable semiconducting polymers with enhanced charge carrier mobility. By incorporating the quadruple hydrogen bonds between end‐capping units linked with alkyl spacers into the backbone of DPPTVT, we constructed the supramolecular polymer DPPTVT‐UPy, which can effectively strengthen interchain interaction, and thus induce higher crystallinity and facilitate more ordered packing with shorter π‐π stacking distance in DPPTVT‐UPy film. Remarkably improved mobility of up to 3.4 cm2 V−1 s−1 was achieved for DPPTVT‐UPy in OTFT devices, more than two‐fold enhancement compared to the nonfunctionalized DPPTVT. Furthermore, this engineering endows DPPTVT‐UPy with higher crack onset strain, lower elastic modulus, greater elongation at break, enhanced toughness, and low T g, which contributes to improved stretchability and the retention of molecular ordering during stretching. Stretchable transistors fabricated with DPPTVT‐UPy exhibited constant even increased charge mobility under strain in both parallel and perpendicular directions, alongside excellent durability under repeated stretching‐releasing cycle. Moreover, DPPTVT‐UPy demonstrated excellent scalability in fully stretchable transistor arrays with robust mobility retention under strain. Concurrently, the supramolecular architecture of DPPTVT‐UPy enhances chain mobility and provides dynamic, reversible hydrogen‐bonding sites that enable efficient healing, resulting in better mobility recovery compared to DPPTVT. MD simulations further reveal that the supramolecular architecture of DPPTVT‐UPy results in lower Young's moduli, smaller initial pore sizes, and also superior retention and healing capabilities in interchain stacking and pore size during dynamic stretching and healing process, which could provide an insight into the mechanism for the improvement in charge mobility, stretchability, and healing ability. These results demonstrate that quadruple hydrogen‐bonds engineering is a facile and effective approach for achieving high‐performance stretchable and healable semiconducting polymers with high mobility.

Author Contributions

Yuanhe Gu: investigation, writing – original draft, visualization. Sichun Wang: investigation, writing – review and editing. Yiran Liu: investigation, writing – review and editing. Zhen Yang: investigation. Wenhao Li: investigation. Junyi Lu: investigation. Jie Xu: investigation. Guangxin Gu: supervision, writing – review and editing. Yunqi Liu: supervision, writing – review and editing. Zhiyuan Zhao: supervision, writing – review and editing. Cheng Zhong: supervision, writing – review and editing. Zhengran Yi: supervision, writing – review and editing, conceptualization. Yan Zhao: conceptualization, supervision, writing – review and editing.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File 1: anie73225‐sup‐0001‐SuppMat.docx.

ANIE-65-e9060680-s001.docx (13.9MB, docx)

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (Grant Numbers: 62461160268, 52473297, and T2441002).

Contributor Information

Zhiyuan Zhao, Email: zhaozhiyuan@iccas.ac.cn.

Cheng Zhong, Email: zhongcheng@whu.edu.cn.

Zhengran Yi, Email: yizr@fudan.edu.cn.

Yan Zhao, Email: zhaoy@fudan.edu.cn.

Data Availability Statement

The data that supports the findings of this study are available in the Supporting Information of this article.

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Associated Data

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Supplementary Materials

Supporting File 1: anie73225‐sup‐0001‐SuppMat.docx.

ANIE-65-e9060680-s001.docx (13.9MB, docx)

Data Availability Statement

The data that supports the findings of this study are available in the Supporting Information of this article.

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