High-performance n-type stretchable OFETs enabled by molecular engineering of flexible polymers

Abstract

Stretchable organic field-effect transistors (OFETs) have emerged as promising semiconductor devices for flexible electronics, combining mechanical deformability with stable electrical performance. However, developing high-performance n-type stretchable semiconductors remains challenging. In this study, we designed three novel n-type polymers (P1–P3) by incorporating flexible chains into an azo-benzodifurandione-based oligo(p-phenylene vinylene) (azo-BDOPV) backbone, achieving balanced mechanical and electrical properties. Using polydimethylsiloxane substrates, gold and silver nanowire electrodes, and polyvinyl alcohol (PVA) dielectric layers, we fabricated fully stretchable top-gate n-type OFETs. The devices demonstrated excellent initial electron mobilities of 0.44, 0.34, and 0.52 cm² V⁻¹ s⁻¹ for P1–P3 respectively, with P3 showing superior performance. Remarkably, P3 maintained mobilities of 0.48–0.29 cm² V⁻¹ s⁻¹ (strain parallel to the charge transport direction) and 0.42–0.26 cm² V⁻¹ s⁻¹ (strain perpendicular to the charge transport direction) under 15–50% deformation, demonstrating exceptional mechanical–electrical stability. All three polymer films show uniform surface morphology and molecular stacking, with polymer P3 having the most ordered edge-on stacking, which is consistent with its excellent device performance. These results highlight the effectiveness of molecular engineering in developing stretchable n-type semiconductors with mechanical flexibility and efficient charge transport, providing valuable insights for the design and application of high-performance fully stretchable OFETs, advancing the development of next-generation flexible and wearable electronics.

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1. Introduction

As an emerging semiconductor device, organic field-effect transistors (OFETs) have achieved significant advancements, particularly in the research of materials, device structures, and performance optimization.^1–3 By optimizing the molecular structure and preparation process of organic semiconductor materials, researchers have consistently enhanced the electrical performance and stability of the devices.^4–6 However, with the rapid development of flexible electronics technology, organic electronic devices are gradually shifting from the traditional rigid, planarized form to a more flexible and deformable configurations. This trend reflects the growing demand for portability and comfort and has opened up new research directions for the application of organic electronic devices in fields such as biomedicine and artificial intelligence.^7–9 Most traditional OFETs are based on rigid substrates, such as silicon wafers or glass.^10–12 These substrates not only restrict the shape and application scenarios of the devices but also severely limit their use in flexible electronic applications. In contrast, stretchable organic field-effect transistors (SOFETs) offer unique advantages as a semiconductor device capable of maintaining stable electrical properties under mechanical deformation.^13–15 Their core strength lies in the ability to operate efficiently under complex mechanical stresses, such as stretching, bending, or twisting, while preserving robust charge transfer capability. This characteristic endows them with great potential for applications in various fields, including wearable electronic devices,^16,17 electronic skin,^18,19 biomedical implantable devices,²⁰ and human–computer interaction interfaces.²¹

Despite the promising applications of SOFETs, their development still faces many challenges. Firstly, the selection of semiconductor materials is crucial. Ideal stretchable semiconductor materials must possess high carrier mobility, good mechanical flexibility, and excellent chemical stability. In recent years, researchers have made significant progress in developing intrinsic p-type stretchable organic semiconductor materials by introducing flexible chains or units and optimizing polymer molecular structures.^22–25 For example, Zhang et al. successfully developed a flexible polymer semiconductor with high mobility by incorporating a central asymmetric unit (e.g., a spirofluorene unit) into the main chain.²⁶ This design reduces the tensile modulus of the film and significantly improves the charge transport properties. However, research on n-type stretchable polymer semiconductor materials is still in its infancy, and their performance and stability require further improvement. In addition to the design and synthesis of semiconductor materials, the design of device structure is also key to realizing stretchable performance. By optimizing the electrode layout, introducing an elastic substrate, or designing special geometrical structures, the stretchable performance of SOFETs can be significantly enhanced. Architectural innovations like serpentine electrodes,²⁷ pleated geometries,^28,29 and multilayer designs³⁰ synergistically improve mechanical stability while retaining high electrical performance. Nevertheless, although progress has been made in the synthesis of stretchable polymer materials and the fabrication of stretchable transistor devices, achieving high-performance, fully stretchable n-type OFETs remains an urgent problem to be solved.

In this study, three novel stretchable polymers, P1, P2, and P3, were designed and synthesized by introducing flexible chain segments of different lengths into the molecular structure of azo-benzodifurandione-based oligo(p-phenylene vinylene) (azo-BDOPV). These polymers were designed to enhance the mechanical flexibility and stretchability of the materials while maintaining good electrical properties. To achieve this goal, we chose polydimethylsiloxane (PDMS) as the elastomer substrate because of its excellent flexibility and stretchability, and its ability to withstand large deformation without affecting the device performance. During the fabrication of the device, gold (Au) was chosen for the source–drain electrodes, and polyvinyl alcohol (PVA) was used as the insulating layer. PVA, a water-soluble polymer, exhibits both superior insulation characteristics and good adhesion to the PDMS substrate. Finally, we selected solution-processable silver nanowires (AgNWs) as the gate material. Fully stretchable n-type field effect transistors were fabricated with a top-gate-bottom contact structure using polymers P1–P3 as the organic layers by a polymer/dielectric layer composite thin film transfer method and vaporized sprayed metal electrode technique. In the initial undeformed state, the SOFETs based on P1, P2, and P3 exhibit excellent electron mobility of 0.44, 0.34, and 0.52 cm² V⁻¹ s⁻¹, respectively. These values reflect the charge transfer efficiency of these polymers in the unstrained state and are crucial for evaluating their potential applications in flexible electronic devices. Additionally, when these devices are subjected to deformation in different directions and with different magnitudes, the electron mobility decreases to varying degrees due to the flexible chain segments in the polymers. Under identical strain conditions (15%, 30%, and 50%), P3-based SOFETs demonstrated superior performance retention, maintaining the highest electron mobility among the three polymers both before and after mechanical deformation. Concretely, the electron mobility of P3 decreases to 0.48, 0.39, and 0.29 cm² V⁻¹ s⁻¹ (strain parallel to the charge transport direction), and 0.42, 0.35, and 0.26 cm² V⁻¹ s⁻¹ (strain perpendicular to the charge transport direction). These results suggest that polymer P3 has significant potential for application in flexible electronic devices that require a balance of high mechanical stability and electrical performance.

2. Experimental

The synthesis process is shown in the ESI.†

3. Results and discussion

3.1. Materials synthesis and thermal properties

In order to obtain polymer semiconductors with good electronic transport properties, structural units azo-BDOPV and (3,3′-difluoro-[2,2′-bithiophene]-5,5′-diyl)bis(trimethylstannane) with strong electron-withdrawing abilities were selected as monomers, mainly due to their lower lowest unoccupied molecular orbital (LUMO) energy levels, which facilitate electron injection and transport, making them ideal for n-type polymer semiconductor materials. In addition, flexible chain segments M1–M3 were introduced to impart stretchability to the polymers, and the effect of the length of the flexible chain segments on the stretchability and carrier mobility of the polymers was emphasized. The chemical structures and synthetic routes of polymers P1–P3 are shown in Scheme 1. The monomers M1–M3 were obtained with relatively high yields via Friedel–Crafts reaction. In order to guarantee the quality of the final polymers, Stille coupling polycondensation reactions were performed. The molar ratio of monomers azo BDOPV to M1, M2, or M3 was precisely controlled at 90 : 10 to balance charge carrier mobility and polymer flexibility. The experimental results showed that the yields of all three polymers exceeded 90%. The structures of M1–M3 were confirmed by NMR spectra and HRMS measurements, while the structures of P1–P3 were verified by elemental analysis. In addition, the number-average molecular weight (M_n), weight-average molecular weight (M_w), and polydispersity index (PDI) of these polymers were evaluated using GPC. The GPC measurements were performed with 1,2,4-trichlorobenzene as the eluent at 150 °C and calibrated with standard polystyrene (Fig. S1, ESI†). All three polymers exhibited sufficiently high molecular weights (>60 kDa) and moderate PDIs (<2.5) (Table 1). TGA and DSC were conducted to investigate the thermal properties of P1–P3. TGA results indicated that all three polymers possessed good thermal stability, with their 5% weight loss temperatures exceeding 300 °C (Fig. S2, ESI†). Furthermore, DSC curves revealed no glass transition processes between room temperature and 300 °C (Fig. S3, ESI†). These results demonstrate that P1–P3 have sufficient thermal stability during annealing to meet the relevant process requirements.

Scheme 1 The chemical structures and synthetic routes of P1, P2, and P3.

Table 1 Molecular weights, and thermal, and optical properties of P1, P2, and P3

Polymer	M n (kDa)	PDI^a	λ ^absmax (nm)		T dec (°C)
			Solu^b	Film^b
a Determined by GPC at 150 °C. b Determined using UV-vis-NIR absorption spectra in chlorobenzene solution and thin films. c Obtained from TGA traces.
P1	96.2	2.07	847	837	369
P2	69.8	1.42	849	853	415
P3	75.8	1.33	845	849	422

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3.2. Optical properties

The normalized UV-vis-NIR absorption spectra of polymers P1–P3 in dilute chlorobenzene solution and thin films are shown in Fig. 1a and b, respectively. The spectra of all three polymers exhibit a typical dual-band feature, which is manifested by a narrow high-energy absorption band (300–500 nm) and a broad low-energy absorption band (600–1000 nm). The high energy absorption band can be attributed to π–π* transitions in polymer molecules. This type of transition is a typical manifestation of electrons from bonding to antibonding orbitals in conjugated systems.³¹ The low-energy absorption bands mainly originate from the intramolecular charge transfer (ICT) effect, which is closely related to the interaction between the electron donor and acceptor units within the polymer molecules.³² In dilute chlorobenzene solution, the maximum absorption wavelengths (λ^abs_max) of P1, P2, and P3 are 847 nm, 849 nm, and 845 nm, respectively. Compared with P2, the absorption peaks of P1 and P3 exhibit a slight blue shift phenomenon, which may be due to the difference in the electron-absorbing capacity of the flexible chain segments, resulting in the weakening of the ICT effect. The blueshift phenomenon was more significant in the thin-film absorption spectra. The λ^abs_max of P2 was 853 nm, whereas λ^abs_max of P1 and P3 were 837 nm and 849 nm, respectively (Table 1). The extended conjugated structure of azo-BDOPV and the strong donor–acceptor interactions give the three polymers a relatively narrow optical band gap.

Fig. 1 Normalized UV-Vis-NIR absorption spectra of P1, P2 and P3 (a) in dilute chlorobenzene solution and (b) in the thin films.

3.3. Frontier orbital energy levels

The frontier molecular orbital energy levels of the three polymers were accurately characterized by UPS, IPES and cyclic voltammetry (CV) (Fig. 2 and Fig. S4, ESI†). In the UPS measurements of the P3 polymer film, a secondary electron cutoff edge of 16.76 eV was observed, which is higher than that of the other two polymers, indicating that P3 has a relatively shallow HOMO energy level. The E_HOMO of P1, P2 and P3 films was calculated and analyzed to be −5.86, −5.89 and −5.84 eV, respectively. The IPES technique was used to characterize the LUMO energy level (E_LUMO). The E_LUMO measured by IPES was −4.12, −4.07, and −4.14 eV for P1, P2, and P3 films, respectively (Table 2). Among the three polymers, P3 had the deepest E_LUMO energy level, followed by P1, while P2 had the shallowest E_LUMO energy level. For a more comprehensive comparison, the energy levels obtained from CV tests are shown in Table 2, which are generally consistent with the UPS and IPES results. This difference in energy levels has important implications for SOFETs using gold as the source–drain electrode. The deeper E_LUMO facilitates enhanced electron injection and carrier transport performance.³³

Fig. 2 UPS and IPES characterization of three polymer films. (a) P1, (b) P2, and (c) P3.

Table 2 Frontier molecular orbital energy levels and energy gaps of P1, P2, and P3

Polymer	E cutoff [eV]	E H,onset [eV]	E HOMO [eV]	E LUMO [eV]	E g [eV]	E ^CVHOMO ^e [eV]	E ^CVLUMO ^e [eV]	E ^CVg ^f [eV]
a Obtained from UPS spectra. b Calculated from EHOMO = −IP eV, IP = [hν − (Ecutoff − EH,onset)] eV, hv = 21.22 eV. c Determined using IPES spectra. d E g = (ELUMO − EHOMO) eV. e Determined by CV. f E ^CVg = −(E^CVHOMO − E^CVLUMO) eV.
P1	16.72	1.36	−5.86	−4.12	1.74	−5.84	−3.77	2.07
P2	16.62	1.29	−5.89	−4.07	1.82	−5.88	−3.74	2.14
P3	16.76	1.38	−5.84	−4.14	1.70	−5.87	−3.79	2.08

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3.4. Thim film morphology and microstructure

As a key component in stretchable OFETs, the surface properties and microstructure of the polymer semiconductor layer play a crucial role in charge transport performance. The flatness and uniformity of the film as well as the ordering of the molecular arrangement are directly related to the electrical performance of the device.³⁴ In order to explore these properties in depth, AFM was used to characterize the surface morphology of three polymer films. As shown in Fig. 3, the initial film surfaces of the three stretchable polymers have flocculated buildups accompanied by distinct crystalline regions. After thermal annealing at 140 °C, the surface morphology of these films was significantly improved. Thermal annealing resulted in sharper boundaries of crystal fibers and crystalline regions, while increasing the surface roughness of the films. Specifically, the root-mean-square (RMS) roughness values of the pristine and thermally annealed films of polymer P1 were 1.01 and 1.20 nm, respectively, and the corresponding RMS values of polymer P2 were 1.07 and 1.10 nm, respectively, whereas polymer P3 showed an increase in the RMS value from 0.894 nm for the pristine film to 1.12 nm after thermal annealing. Polymer P3 shows the most significant change in roughness before and after thermal annealing, which suggests that the thermal annealing process significantly enhances the inter-chain interactions of the films and promotes the aggregation of crystalline molecules.

Fig. 3 AFM height images of three polymer films. (a), (c) and (e) Pristine films, (b), (d) and (f) films after thermal annealing at 140 °C. (a) and (b) P1, (c) and (d) P2, and (e) and (f) P3.

To deeply investigate the molecular arrangement of the polymers and the stacking pattern of the films, we performed GIWAXS tests on the three polymer films. Fig. 4 illustrates the 2D GIWAXS diffraction patterns of the pristine films of P1, P2, and P3 polymers versus the annealed films. The patterns show that all films exhibit several (h00) Bragg peaks in the out-of-plane direction (q_z-axis), which reveals the existence of an ordered lamellar stacking structure within the films. Meanwhile, the (010) Bragg peaks in the in-plane direction (q_xy axis) indicate that all three polymers adopt an edge-on molecular stacking mode in the films. Compared with the pristine films, all the annealed films exhibit a more ordered molecular stacking pattern, suggesting that the thermal annealing process effectively enhances the microstructural ordering of the films. In addition, the microstructural ordering of the films (especially the annealed films) is gradually enhanced from P2, P1 to P3. In particular, the annealed film of P3 exhibits up to quadruple (h00) diffraction peaks in the q_z-axis direction, and its (010) Bragg peak on the q_xy-axis is the most pronounced among the three polymer films, which further confirms that P3 has the most ordered edge-on molecular stacking pattern. Based on the (100) Bragg peak on the q_z-axis and the (010) Bragg peak on the q_xy-axis, we can deduce the d–d and π–π stacking distances of the annealed films. The specific values are as follows: 27.15/3.34 Å for P1, 27.31/3.39 Å for P2, and 26.88/3.31 Å for P3. This dense and well-organized molecular stacking facilitates the effective hopping of carriers between polymer chains, which in turn achieves a higher carrier transport efficiency.³⁵ Scherrer analysis of the GIWAXS (010) peaks reveals well-developed crystalline domains in all three polymers, with coherence lengths of 5.095 (P1), 5.017 (P2), and 5.305 nm (P3). These substantial domain sizes (>5 nm) account for the similarly high charge carrier mobilities observed across the series. Notably, P3's superior mobility correlates with its largest coherence length (5.305 nm), as extended crystalline domains facilitate more continuous charge transport pathways. Conversely, P2's marginally shorter domain length (5.017 nm) likely introduces additional grain boundaries, explaining its slightly reduced mobility compared to P1 and P3. These findings demonstrate a clear correlation between subtle variations in the crystalline domain size and the observed charge transport characteristics.

Fig. 4 GIWAXS diffraction patterns of three polymer films. (a), (c) and (e) Pristine films, (b), (d) and (f) films after thermal annealing at 140 °C. (a) and (b) P1, (c) and (d) P2 and (e) and (f) P3.

3.5. Charge transport properties and performances of stretchable OFETs

Stretchable transistor devices were prepared using a top-gate bottom-contact structure and PDMS as a flexible substrate for the study of the relationship between strain and charge transport properties of three polymers containing different flexible chain segments. In the device preparation process, Au with excellent electrical conductivity and chemical stability was chosen as the source–drain electrode material; polyvinyl alcohol (PVA) was used for the insulating layer, which is a widely used water-soluble polymer material that not only has excellent insulating properties, but also has good adhesion with the PDMS substrate; and finally, AgNWs with high conductivity and good flexibility were selected as the gate material. To ensure the reliability and generalizability of the test results, all devices were tested under standard environmental conditions. During the testing process, we applied different directions of strain to the stretchable field effect transistors, including parallel and perpendicular to the charge transport direction. The applied deformations were 15%, 30%, and 50%, covering the range from slight deformation to large deformation, and help to comprehensively evaluate the mechanical stability and charge transfer performance of stretchable field effect transistors (Fig. S5–S10, ESI†).

In the initial undeformed state, the stretchable field effect transistors based on the three polymers P1, P2, and P3 all exhibit good electron mobility with values of 0.44, 0.34, and 0.52 cm² V⁻¹ s⁻¹, respectively (Fig. 5). The initial mobility of the three polymers differed at 0% strain. P3 had the highest initial mobility followed by P1 and P2 had the lowest initial mobility. However, due to the different flexible chain segments contained in P1, P2, and P3, these devices show varying degrees of electron mobility degradation when subjected to deformation in different directions and magnitudes. Specifically, when 15%, 30% and 50% deformations are applied, the electron mobility of P1 decreases to 0.37/0.33, 0.31/0.26, and 0.20/0.15 cm² V⁻¹ s⁻¹; the electron mobility of P2 decreases to 0.24/0.26, 0.20/0.19, and 0.17/0.13 cm² V⁻¹ s⁻¹, respectively, under the same deformation conditions. For P3, the electron mobility decreases to 0.48/0.42, 0.39/0.35, and 0.29/0.26 cm² V⁻¹ s⁻¹, respectively (Table 3). Remarkably, comparative analysis reveals that P3-based devices exhibit superior initial electron mobility (0.52 cm² V⁻¹ s⁻¹) and demonstrate the most stable electrical performance post-deformation among the three polymer variants. The decrease in the mobility of P3 is relatively small and it maintains a high mobility at 50% strain. This indicates that P3 better maintains its charge transport properties when subjected to mechanical strain. The mobility of P1 and P2 decreases faster at smaller strains, and the mobility of P1 is always higher than that of P2. From the trend of mobility versus strain, P3 exhibits a better stretchability as it maintains a higher mobility at higher strains. Overall, the P3 polymer exhibits superior charge transport stability and stretchability when subjected to mechanical strain, indicating that the length of the flexible chain has a significant effect on the mechanical and electrical properties of the material. However, the flexible chains also affect the surface morphology and stacking condition of the films, and too long flexible chains may lead to weakened π–π interactions between the molecular chains, thus reducing the charge mobility. Therefore, it is necessary to balance the stretchability and charge mobility by precisely controlling the length of the flexible chains. Here, we compared the performance of intrinsically stretchable OFETs based on P1–P3 prepared in this study with the performance of stretchable OFETs in other research works. It is observed that the device performance in this work has significant advantages in both initial mobility and performance retention after 50% deformation (Table S1, ESI†).^36–41 Additionally, we have subjected the stretchable OFETs based on P1, P2, and P3 to 500 stretching cycles at 30% strain to verify their mechanical stability (Fig. S11 and S12, ESI†). Under cyclic stretching where charge transport is parallel to the strain direction, the mobilities of the three polymers decreased to 0.17, 0.12, and 0.23 cm² V⁻¹ s⁻¹, respectively. When the charge transport is perpendicular to the strain direction, the mobilities dropped to 0.13, 0.09, and 0.19 cm² V⁻¹ s⁻¹, respectively. Despite this, all maintained relatively good electrical characteristics compared to their initial values at 30% strain, demonstrating good mechanical stability and durability. The charge transport properties exhibit a non-monotonic correlation with flexible segment length. While the segment length increases sequentially from P1 < P2 < P3, the corresponding carrier mobility follows an inverted trend (P2 < P1 < P3), indicating that extended chain length alone does not guarantee enhanced charge transport. Specifically, P3 achieves the highest mobility due to an optimal synergy between conformational flexibility (promoting stress dissipation) and ordered molecular packing (facilitating efficient charge hopping). In contrast, P2 displays reduced mobility, likely attributed to compromised packing geometry or dynamic disorder arising from its intermediate chain length. P1, with its rigid backbone, maintains moderate charge transport capability but exhibits inferior mechanical adaptability compared to P3. Similarly, the flexible segment length exerts a significant yet nonlinear influence on optoelectronic behavior. UV-vis absorption spectra reveal that P2 exhibits the strongest redshift, surpassing both P1 and P3, suggesting its intermediate chain length maximizes intramolecular charge transfer (ICT) through balanced backbone dynamics. The pronounced blueshift in P1 implies that shorter segments intensify solid-state intermolecular interactions, whereas P3's retained redshift reflects preserved stacking order despite reduced ICT efficiency. Furthermore, the progressive deepening of LUMO levels from P2 to P3 highlights the critical role of segment length in modulating electron affinity, with longer chains enhancing charge delocalization.

Fig. 5 Charge mobility as a function of strain for three polymer films with charge transport direction (a) parallel or (b) perpendicular to the strain direction.

Table 3 Performance parameters of polymer P1–P3 based stretchable field-effect transistors under different deformations

Strain	P1			P2			P3
	μ e,max (μavg)^a [cm^2 V^−1 s^−1]	V th [V]	I on/Ioff	μ e,max (μavg)^a [cm^2 V^−1 s^−1]	V th [V]	I on/Ioff	μ e,max (μavg)^a [cm^2 V^−1 s^−1]	V th [V]	I on/Ioff
a The average mobility values were calculated from more than 12 devices.
0	0.44 (0.38)	20 (± 4)	∼10^3	0.34 (0.28)	18 (± 5)	10^2–10^3	0.52 (0.45)	20 (± 5)	∼10^3
15%‖	0.37 (0.32)	15 (± 4)	10^2–10^3	0.24 (0.18)	11 (± 4)	10^2–10^3	0.48 (0.39)	20 (± 4)	10^2–10^3
15%⊥	0.33 (0.27)	18 (± 4)	10^2–10^3	0.26 (0.21)	18 (± 4)	10^2–10^3	0.42 (0.35)	19 (± 5)	10^2–10^3
30%‖	0.31 (0.24)	15 (± 5)	10^2–10^3	0.20 (0.15)	12 (± 5)	10^2–10^3	0.39 (0.31)	18 (± 5)	10^2–10^3
30%⊥	0.26 (0.21)	15 (± 4)	10^2–10^3	0.19 (0.12)	14 (± 4)	10^2–10^3	0.35 (0.27)	17 (± 4)	10^2–10^3
50%‖	0.20 (0.14)	14 (± 4)	10^2–10^3	0.17 (0.11)	11 (± 4)	10^2–10^3	0.29 (0.22)	15 (± 4)	10^2–10^3
50%⊥	0.15 (0.09)	12 (± 5)	10^2–10^3	0.13 (0.06)	12 (± 5)	10^2–10^3	0.26 (0.18)	16 (± 5)	10^2–10^3

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These findings show that flexible segment length is a powerful tool for modulating properties, with context-dependent effects. The transition from P1 to P3 shows non-linear improvements in transport due to the competition between packing and flexibility. Optical properties correlate more directly in solution but become complex in films. Electronic structure tuning at the orbital level occurs without disrupting the bandgap. This comprehensive understanding of structure–property relationships should guide future molecular design strategies for optimizing flexible organic electronic materials.

4. Conclusion

In this study, three stretchable polymers, P1, P2, and P3, were successfully prepared by introducing flexible groups with different lengths using azo-BDOPV as the conjugated backbone. In order to obtain stretchable field effect transistor devices with balanced stretchability and electrical properties, we employed PDMS elastomers as the substrate, utilized Au as the source–drain electrode, PVA as the insulating layer, and AgNWs as the gate electrode, combining the advantages of solution processing and vacuum evaporation process with composite film transfer technology and electrode deposition method to realize the fabrication of fully stretchable transistors based on P1–P3 semiconductor materials. Initial tests show that the electron mobility of P1, P2, and P3 films is 0.44, 0.34, and 0.52 cm² V⁻¹ s⁻¹, respectively, demonstrating excellent charge transfer capability. Notably, due to the structural differences in the flexible chain segments, the three types of materials exhibit significantly different performance decay patterns under mechanical strain. Specifically, when 15%, 30%, and 50% strains are applied (strain direction parallel/perpendicular to the charge transport direction), the mobility of P1 decreases to 0.37/0.33, 0.31/0.26, and 0.20/0.15 cm² V⁻¹ s⁻¹, respectively; that of P2 decreases to 0.24/0.26, 0.20/0.19, and 0.17/0.13 cm² V⁻¹ s⁻¹; and that of P3 exhibits optimal electron mobility and mechanical stability, only decreasing to 0.48/0.42, 0.39/0.35, and 0.29/0.26 cm² V⁻¹ s⁻¹. This study not only reveals the variation rules of charge transport properties of polymers with flexible chain segments of different lengths under deformation conditions, but also provides important guidance for the design and optimization of high-performance n-type stretchable electronic devices.

Author contributions

Qian Che: writing – original draft, conceptualization, methodology, investigation, and data curation. Tianhao Zhang: investigation, methodology, and formal analysis. Weifeng Zhang: funding acquisition, resources, supervision, and proofreading the article. Jiadi Chen and Yunchao Zhang: validation, formal analysis, and data curation. Zhihui Chen, Youjia Li and Lei Yang: methodology and validation. Liping Wang: funding acquisition, supervision, and writing – review and editing. Gui Yu: funding acquisition, resources, supervision, writing – review and editing, and proofreading the article.

Data availability

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

Conflicts of interest

The authors declare no conflicts of interest.

Acknowledgements

This work was supported by the Key R&D Program of China (Grant No. 2024YFA1209600), the National Natural Science Foundation of China (Grants No. 22175021, 22475220, 22275194 and 22021002), the Strategic Priority Research Program of the Chinese Academy of Sciences (XDB0520000), and Beijing National Laboratory for Molecular Sciences (BNLMS-CXXM-202101). The authors thank the beamlines BL14B1 and BL02U2 of Shanghai Synchrotron Radiation Facility (SSRF) and beamline 1W1A of Beijing Synchrotron Radiation Facility (BSRF) for providing the beam time and the help from beamline scientists on 2D-GIWAXS measurements. We also appreciate Dr Jianyao Huang for his help in GIXRD analysis.

Notes and references

1. J. Mei, Y. Diao, A. Appleton, L. Fang and Z. Bao, J. Am. Chem. Soc., 2013, 135, 6724–6746.
2. K. Shi, W. Zhang, D. Gao, S. Zhang, Z. Lin, Y. Zou, L. Wang and G. Yu, Adv. Mater., 2018, 30, 1705286.
3. Q. Shi, X. Shi, C. Feng, Y. Wu, N. Zheng, J. Liu, X. Wu, H. Chen, A. Peng, J. Li, L. Jiang, H. Fu, Z. Xie, S. Marder, S. Blakey and H. Huang, Angew. Chem., Int. Ed., 2020, 60, 2924–2928.
4. Q. Che, W. Zhang, X. Wei, Y. Zhou, H. Luo, J. Wei, L. Wang and G. Yu, CCS Chem., 2023, 5, 2603–2616.
5. D. Liu, Y. Zhao, J. Zhang, Z. Wei, Y. Liu and Y. Wang, Angew. Chem., Int. Ed., 2024, 63, e202400061.
6. W. Zhang, K. Shi, J. Lai, Y. Zhou, X. Wei, Q. Che, J. Wei, L. Wang and G. Yu, Adv. Mater., 2023, 35, 2300145.
7. S. Wang, J. Xu, W. Wang, G. N. Wang, R. Rastak, F. M. Lopez, J. W. Chung, S. Niu, V. R. Feig, J. Lopez, T. Lei, S. Kwon, Y. Kim, A. M. Foudeh, A. Ehrlich, A. Gasperini, Y. Yun, B. Murmann, J. B. Tok and Z. Bao, Nature, 2018, 555, 83–88.
8. Z. Luo, B. Peng, J. Zeng, Z. Yu, Y. Zhao, J. Xie, R. Lan, Z. Ma, L. Pan, K. Cao, Y. Lu, D. He, H. Ning, W. Meng, Y. Yang, X. Chen, W. Li, J. Wang, D. Pan, X. Tu, W. Huo, X. Huang, D. Shi, L. Li, M. Liu, Y. Shi, X. Feng, P. K. Chan and X. Wang, Nat. Commun., 2021, 12, 1928.
9. Y. Bian, M. Zhu, C. Wang, K. Liu, W. Shi, Z. Zhu, M. Qin, F. Zhang, Z. Zhao, H. Wang, Y. Liu and Y. Guo, Nat. Commun., 2024, 15, 2624.
10. Z. Chen, X. Wei, J. Huang, Y. Zhou, W. Zhang, Y. Pan and G. Yu, ACS Appl. Mater. Inter., 2019, 11, 34171–34177.
11. S. Duan, T. Wang, B. Geng, X. Gao, C. Li, J. Zhang, Y. Xi, X. Zhang, X. Ren and W. Hu, Adv. Mater., 2020, 32, 1908388.
12. J. Wang, Z. Ren, J. Pan, X. Wu, J. Jie, X. Zhang and X. Zhang, Adv. Mater., 2023, 35, 2301017.
13. J. Mun, Y. Ochiai, W. Wang, Y. Zheng, Y. Zheng, H. Wu, N. Matsuhisa, T. Higashihara, J. B. Tok, Y. Yun and Z. Bao, Nat. Commun., 2021, 12, 3572.
14. Z. Zhou, N. Luo, T. Cui, L. Luo, M. Pu, Y. Wang, F. He, C. Jia, X. Shao, H. Zhang and Z. Liu, Adv. Mater., 2024, 36, 2313312.
15. H. Yue, Y. Wang, S. Luo, J. Guo, J. Jin, G. Li, Z. Meng, L. Zhang, D. Zhou, Y. Zhen and W. Hu, Sci. Adv., 2024, 10, eadq0171.
16. F. Tehrani, H. Teymourian, B. Wuerstle, J. Kavner, R. Patel, A. Furmidge, R. Aghavali, H. Hosseini-Toudeshki, C. Brown, F. Zhang, K. Mahato, Z. Li, A. Barfidokht, L. Yin, P. Warren, N. Huang, Z. Patel, P. P. Mercier and J. Wang, Nat. Biomed. Eng., 2022, 6, 1214.
17. D. Zhong, C. Wu, Y. Jiang, Y. Yuan, M. Kim, Y. Nishio, C. Shih, W. Wang, J. Lai, X. Ji, T. Gao, Y. Wang, C. Xu, Y. Zheng, Z. Yu, H. Gong, N. Matsuhisa, C. Zhao, Y. Lei, D. Liu, S. Zhang, Y. Ochiai, S. Liu, S. Wei, J. B. Tok and Z. Bao, Nature, 2024, 627, 313–320.
18. G. Luo, J. Shi, W. Deng, Z. Chang, Z. Lu, Y. Zhang, R. Pan, J. Jie, X. Zhang and X. Zhang, Adv. Mater., 2023, 35, 2301020.
19. N. Vo, T. Nam, M. Jeong, J. Kim, K. Jung, Y. Lee, G. Ma, X. Gu, J. B. Tok, T. Lee, Z. Bao and J. Oh, Nat. Commun., 2024, 15, 3433.
20. W. Wang, Y. Jiang, D. Zhong, Z. Zhang, S. Choudhury, J. Lai, H. Gong, S. Niu, X. Yan, Y. Zheng, C. Shih, R. Ning, Q. Lin, D. Li, Y. Kim, J. Kim, Y. Wang, C. Zhao, C. Xu, X. Ji, Y. Nishio, H. Lyu, J. B. Tok and Z. Bao, Science, 2023, 380, 735–742.
21. Z. Li, Z. Li, W. Tang, J. Yao, Z. Dou, J. Gong, Y. Li, B. Zhang, Y. Dong, J. Xia, L. Sun, P. Jiang, X. Cao, R. Yang, X. Miao and R. Yang, Nat. Commun., 2024, 15, 7275.
22. J. Y. Oh, S. Rondeau-Gagné, Y.-C. Chiu, A. Chortos, F. Lissel, G.-J. N. Wang, B. C. Schroeder, T. Kurosawa, J. Lopez, T. Katsumata, J. Xu, C. Zhu, X. Gu, W.-G. Bae, Y. Kim, L. Jin, J. W. Chung, J. B.-H. Tok and Z. Bao, Nature, 2016, 539, 411.
23. K. Ditte, J. Perez, S. Chae, M. Hambsch, M. A. Hussein, H. Komber, P. Formanek, S. Mannsfeld, A. Fery, A. Kiriy and F. Lissel, Adv. Mater., 2021, 33, 2005416.
24. J. Mun, Y. Ochiai, W. Wang, Y. Zheng, Y.-Q. Zheng, H.-C. Wu, N. Matsuhisa, T. Higashihara, J. B.-H. Tok, Y. Yun and Z. Bao, Nat. Commun., 2021, 12, 3572.
25. F. Zhang, J. Sun, F. Liu, J. Li, B. Hu, Q. Tang and R. Li, ACS Appl. Mater. Interfaces, 2024, 16, 38324–38333.
26. X. Yu, L. Chen, C. Li, C. Gao, X. Xue, X. Zhang, G. Zhang and D. Zhang, Adv. Mater., 2023, 35, 2209896.
27. K. Li, X. Cheng, F. Zhu, L. Li, Z. Xie, H. Luan, Z. Wang, Z. Ji, H. Wang, F. Liu, Y. Xue, C. Jiang, X. Feng, L. Li, J. A. Rogers, Y. Huang and Y. Zhang, Adv. Funct. Mater., 2019, 29, 1806630.
28. S. Chae, W. Choi, L. Nebel, C. Cho, Q. Besford, A. Knapp, P. Makushko, Y. Zabila, O. Pylypovskyi, M. Jeong, S. Avdoshenko, O. Sander, D. Makarov, Y. Chung, A. Fery, J. Oh and T. Lee, Nat. Commun., 2024, 15, 3071.
29. J. Yang, S. Lee, B. Ma, J. Kim, M. Song, S. Kim, D. Kim, T. Kim and S. Park, Sci. Adv., 2022, 8, eabn3863.
30. M. Khatib, O. Zohar, W. Saliba and H. Haick, Adv. Mater., 2020, 32, 2000246.
31. J. Ha, K. Kim and D. Choi, J. Am. Chem. Soc., 2011, 133, 10364–10367.
32. C. Yang, Z. Pan, K. Zhang, J. Tai, C. Wang, L. Ding, M. Fung and J. Fan, Mater. Horiz., 2023, 10, 945–951.
33. X. Tan, Q. Li, Z. Qin, D. Liu, Y. Liu, P. Wang, Z. Xie, Z. Miao, Y. Lei, Y. Zhang, P. Wang, X. Chen, Z. Liu, C. Gao, W. Hu, H. Zhang and H. Dong, Mater. Horiz., 2025, 12, 1589–1595.
34. S. Fratini, M. Nikolka, A. Salleo, G. Schweicher and H. Sirringhaus, Nat. Mater., 2020, 19, 491–502.
35. Z. Wang, L. D. Virgilio, Z. Yao, Z. Yu, X. Wang, Y. Zhou, Q. Li, Y. Lu, L. Zou, H. I. Wang, X. Wang, J. Wang and J. Pei, Angew. Chem., Int. Ed., 2021, 60, 20483–20488.
36. H. Shim, K. Sim, B. Wang, Y. Zhang, S. Patel, S. Jang, T. J. Marks, A. Facchetti and C. Yu, Nat. Electron., 2023, 6, 349.
37. H. Shim, F. Ershad, S. Patel, Y. Zhang, B. Wang, Z. Chen, T. J. Marks, A. Facchetti and C. Yu, Nat. Electron., 2022, 5, 660.
38. J. Mun, G.-J. N. Wang, J. Y. Oh, T. Katsumata, F. L. Lee, J. Kang, H.-C. Wu, F. Lissel, S. Rondeau-Gagne, J. B.-H. Tok and Z. Bao, Adv. Funct. Mater., 2018, 28, 1804222.
39. M. H. Kim, M. W. Jeong, J. S. Kim, T. U. Nam, N. T. P. Vo, L. Jin, T. I. Lee and J. Y. Oh, Sci. Adv., 2022, 8, eade2988.
40. J. Mun, J. Kang, Y. Zheng, S. Luo, H.-C. Wu, N. Matsuhisa, J. Xu, G.-J. N. Wang, Y. Yun, G. Xue, J. B.-H. Tok and Z. Bao, Adv. Mater., 2019, 31, 1903912.
41. J. Chen, M. Zhu, M. Shao, W. Shi, J. Yang, J. Kuang, C. Wang, W. Gao, C. Zhu, R. Meng, Z. Yang, Z. Shao, Z. Zhao, Y. Guo and Y. Liu, Adv. Mater., 2023, 36, 2305987.

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Footnotes

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5tc01650a

‡ Q. Che and T. H. Zhang contributed equally to this work.

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