Exploring Wide‐Range Alkyl Bridge Length Variations in Polymer Semiconductors: From Pristine to Blend Films for High‐Mobility Stretchable TFTs

TL;DR

This paper introduces a new method to design stretchable polymer transistors with high performance by adjusting alkyl bridge lengths, enabling durable and flexible electronics.

Contribution

A systematic study of alkyl bridge length variations in polymer semiconductors to enhance stretchability and charge transport in transistors.

Findings

Optimized polymer achieved 6.4 cm² V⁻¹ s⁻¹ mobility at 0% strain.

Maintained 0.6 cm² V⁻¹ s⁻¹ mobility at 100% strain.

Fabricated a 38-device stretchable TFT array with 5.5 cm² V⁻¹ s⁻¹ average mobility.

Abstract

The development of intrinsically stretchable thin‐film transistors (TFTs) with high mobility is essential for next‐generation deformable electronics, including wearable displays and bio‐integrated systems. However, most approaches to improve stretchability in polymer semiconductors compromise charge transport due to disrupted molecular ordering. Here, we report a systematic exploration of wide‐range of alkyl bridge length variations of donor–acceptor‐type conjugated polymers to control crystallinity and morphology without altering the polymer backbone. We also propose a method to quantify the relative degree of crystallinity, enabling comparison across different polymer systems. When blended with an elastomer and aligned via solution shearing, the optimized polymer exhibited a maximum mobility of 6.4 cm2 V−1 s−1 at 0% strain (V
DS = −40 V). The polymer stretchable device maintained measurable mobility (0.6 cm2 V−1 s−1) at 100% strain under the perpendicular to the channel direction under a low V
DS of −10 V. Furthermore, wafer‐scale photopatterning enabled fabrication of a 38‐device intrinsically stretchable TFT array with high uniformity and an average mobility of 5.5 cm2 V−1 s−1 at 0% strain. This work establishes a molecular design framework that elucidates the link between structure, mechanical resilience, and electrical performance, offering generalizable principles and a scalable platform for high‐performance deformable electronics.

A side‐chain engineering strategy enables polymer semiconductors to achieve record‐high mobility and mechanical resilience in intrinsically stretchable transistors. By correlating crystallinity with charge transport and stretchability, this work establishes a predictive design rule. The resulting polymer/elastomer composites support wafer‐scale device fabrication, offering a scalable platform for next‐generation deformable electronics.

Figures

Molecular design strategy for polymer semiconductors (PSCs). (A) Molecular structures of synthesized conjugated polymers with systematically varied branched alkyl bridge length (C2 to C16). (B) UV–vis/NIR absorption spectra of neat PSC films, highlighting the redshift and vibronic peak evolution with increasing alkyl bridge length. (C) Vibronic peak ratio (A0‐0/A0‐1) and absorption maximum (λ max) at different alkyl bridge lengths, indicating increased interchain aggregation at longer side chains. (D) Charge carrier mobilities of neat PSCs in rigid thin‐film transistors fabricated on an ODTMS‐treated SiO2/n+‐Si substrate. Error bars indicate the SD for four devices. (E) Crack onset strains of neat PSC films. Error bars indicate the SD for two to four devices.

Molecular ordering analysis of polymer semiconductors (PSCs). (A) 2D GIWAXS patterns of neat PSC films with alkyl bridge lengths of C2, C8, and C16. (B) 1D in‐plane GIWAXS profile of a neat C8 PSC film with peak deconvolution into p1 (amorphous halo) and p2 (π–π stacking). (C) Plots of the relative crystalline population (CP) and molecular ordering quality (Q) as functions of the alkyl bridge length, revealing a non‐monotonic trend and maximum at C8. CP is the area ratio of p2 to p1 (p2/p1), and Q is the ratio of the second‐ to the first‐order crystal coherence length [Lc(200)/Lc(100)]. (D) Correlation between the charge carrier mobility and alkyl bridge length based on an empirical model incorporating the two structural metrics CP and Q. (E) Schematic illustrations of molecular packing in representative PSCs (C2, C8, and C16), reflecting CP and Q. (F) Structural parameter analysis showing increased lamellar d‐spacing, Lc(100), and π–π distance at longer side chains. (G,H) 3D visualization of molecular packing in C8 and C16 PSCs, incorporating data from (F).

Solution‐sheared polymer/SEBS composites for intrinsically stretchable thin‐film transistors (TFTs). (A) Process schematic of blade‐assisted solution shearing of a polymer/SEBS composite film, followed by PDMS‐mediated transfer onto a cross‐linked SEBS substrate. (B) AFM phase and conducting‐AFM images of a C8 composite film, showing nanofiber alignment and directional current flow along the shearing direction (scale bar: 1 µm). (C) Charge carrier mobilities of rigid TFTs fabricated using spin‐coated (SC) and solution‐sheared (SS) films with different alkyl bridge lengths. (D) Transfer characteristics of rigid TFTs across the polymer series at V DS = −10 V. (E) Comparison of the craze and crack onset strains of composite films. (F) Optical images of a C8 composite film under strain applied parallel and perpendicular to the shear direction. The right‐side panels show magnified views of the damage morphology (scale bar: 40 µm). (G and H) Mobility retention (µ/µ 0) of stretchable TFTs under increasing uniaxial strain applied parallel (G) and perpendicular (H) to the channel direction. Error bars indicate the SD for five different devices. The insets show the mobility retention at 50% strain (µ 50/µ 0). (I) Charge carrier mobilities of stretchable TFTs at V DS = −10 and −40 V and corresponding mobility retention (µs tretchable/µ rigid).

Wafer‐scale integration and device characteristics of an intrinsically stretchable thin‐film transistor (TFT) array. (A) Photograph of an 8‐inch wafer‐scale intrinsically stretchable TFT array based on a SEBS‐blended C6 polymer (scale bar: 200 µm). (B) Charge carrier mobility distribution across 38 devices in the array, with an average of 5.5 cm2 V−1 s−1 and maximum of 6.4 cm2 V−1 s−1. (C) Threshold voltage (V TH) distribution, centered at approximately −0.4 V. (D–F) Output (D) and transfer (E) characteristics and gate‐voltage‐dependent mobility (F) of a representative device.

Full text

Introduction

1

Stretchable electronics, capable of dynamically transforming their shape through multidirectional deformation, have emerged as a promising platform for next‐generation technologies that have the potential to surpass conventional flexible electronics [1, 2]. Their exceptional mechanical adaptability enables new opportunities across a range of applications, including displays [3, 4, 5, 6], health monitoring devices [7, 8, 9], bio‐integrated electronics [10, 11], and soft robotics [12]. To support the development of these applications, substantial research efforts have focused on stretchable thin‐film transistors (TFTs), which serve as essential circuit components in fully deformable electronic systems [13, 14, 15, 16, 17]. Compared with geometrically engineered deformable circuits with flexible or rigid components [18, 19, 20], intrinsically stretchable TFTs fabricated from inherently stretchable materials offer enhanced device density, mechanical robustness, and mechanical compliance [17, 21, 22, 23].

Polymer semiconductors (PSCs) have long attracted attention as active channel materials owing to their inherent mechanical flexibility, which is superior to those of conventional inorganic semiconductors [24, 25]. Significant progress has been made in designing new polymers by tuning backbone rigidity through the introduction of conjugation‐breaking spacers [26], dynamic bonding units [27], and other structural modifications [28]. More recently, researchers have developed advanced polymer systems, such as terpolymers [23] and fluorine‐functionalized donor–acceptor copolymers [29], to simultaneously control the crystalline domain size and interconnectivity within thin films, both of which are key parameters for enhancing mechanical stretchability. While these approaches have successfully achieved stretchability exceeding 100% and partially mitigated the trade‐off between charge carrier mobility and mechanical deformability, the mobility remains significantly limited (typically <1 cm^2^ V^−1^ s^−1^). Incorporating an elastomer into the polymer as a mechanical buffer is a robust materials engineering strategy for enhancing stretchability, as its inherent elasticity facilitates efficient stress dissipation and suppresses crack propagation, while preserving the electrical performance of the PSC [30, 31]. This approach further improves charge carrier mobility when combined with a coating method that induces molecular alignment, such as solution shearing [21, 32]. However, despite these and other materials engineering strategies, stretchable PSCs still exhibit charge transport characteristics comparable to those of amorphous silicon, underscoring the need for further breakthroughs to meet the performance requirements of high‐end stretchable electronics. To this end, it is essential to move beyond the current focus on polymer/elastomer blends—which largely remains at the level of identifying material combinations and optimizing blending conditions—toward more integrated approaches that incorporate molecular design principles through molecular engineering.

Side‐chain engineering, one of the most established molecular design strategies for controlling molecular packing in thin films, has been employed in numerous studies to investigate the influence of alkyl chains on the performance of PSCs [33, 34]. However, these efforts have been largely confined to rigid device configurations that focus exclusively on electrical characteristics [35, 36, 37, 38, 39]. Only recently have a few studies begun to explore the correlation between polymer chain dynamics and stretchability—primarily emphasizing mechanical properties—in PSCs tailored for stretchable electronics [40]. Notably, most prior studies have focused on a narrow range of short side‐chain lengths, lacking a comprehensive comparison across a broader spectrum from very short to long bridging units [41]. Critically, systematic investigations of how molecular structural variations affect the thin‐film morphology of both polymer/elastomer blends and neat polymer films—and how morphological changes, in turn, influence the performance of intrinsically stretchable TFTs—remain exceedingly limited.

In this study, we systematically explored the wide‐range side‐chain engineering of stretchable PSCs by varying the length of the branched alkyl bridge (from C2 to C16) while preserving the polymer backbone. Fundamental investigations were conducted to understand the effects of the alkyl bridge length on solid‐state aggregation, polymer chain alignment, and the resulting mechanical and electrical properties of neat PSC films and polymer/elastomer blends. By examining a wide range of alkyl bridge lengths, we identified a polymer that achieved excellent electrical performance while maintaining adequate mechanical stretchability within the trade‐off range between the two properties. Notably, blending the newly synthesized polymers with an elastomer and applying solution‐shearing alignment allowed us to systematically examine how variations in side‐chain length influence molecular ordering and the resulting device behavior under stain. Building on this advancement, we believe that this technology is poised to become a key circuit component in next‐generation stretchable electronic systems.

Results

2

Material Properties of PSC Films

2.1

A series of 3,6‐bis(5‐bromothiophen‐2‐yl)‐2,5‐bis(4‐decylhexadecyl)‐2,5‐dihydropyrrolo[3,4‐c]pyrrole‐1,4‐dione (DPPT)‐based copolymers was synthesized by incorporating thienothiophene (TT) as an electron‐donating unit. The alkyl bridge length was systematically varied from a very short (C2) to a very long chain (C16) across seven derivatives (C2, C4, C6, C8, C10, C12, and C16) while maintaining the DPPT‐TT backbone (Figure 1A; Schemes S1–S3). These polymers were designated as DPPT‐C2‐TT to DPPT‐C16‐TT according to the number of carbon atoms in the branched alkyl chain. Notably, only even‐numbered alkyl bridge lengths were selected to account for the even–odd effects of alkyl side chains [36]. Moreover, the number‐average molecular weights (M n) of all polymers used in this study were adjusted to approximately 80–100 kg mol^−1^, as the molecular weight significantly influences electrical and mechanical properties [42, 43, 44]. The Mn and polydispersity indices of the polymers are summarized in Table S1.

Molecular design strategy for polymer semiconductors (PSCs). (A) Molecular structures of synthesized conjugated polymers with systematically varied branched alkyl bridge length (C2 to C16). (B) UV–vis/NIR absorption spectra of neat PSC films, highlighting the redshift and vibronic peak evolution with increasing alkyl bridge length. (C) Vibronic peak ratio (A0‐0/A0‐1) and absorption maximum (λ max) at different alkyl bridge lengths, indicating increased interchain aggregation at longer side chains. (D) Charge carrier mobilities of neat PSCs in rigid thin‐film transistors fabricated on an ODTMS‐treated SiO2/n+‐Si substrate. Error bars indicate the SD for four devices. (E) Crack onset strains of neat PSC films. Error bars indicate the SD for two to four devices.

First, the intrinsic electrical and mechanical properties of the polymers in their pristine state (neat PSC films) were investigated. Ultraviolet‐visible/near‐infrared (UV–vis/NIR) spectroscopy was performed in both solution and thin‐film states to analyze the vibronic peak features of the polymers and gain insight into their intermolecular aggregation behavior. Conjugated PSCs tend to form aggregates, and the intensity ratio of the 0–1 and 0–0 vibronic peaks serves as a reliable indicator of their relative degree of aggregation [45, 46, 47, 48]. As shown in Figure 1B, increasing the alkyl bridge length from C2 to C16 enhanced the absorption intensity of the 0–0 peak, reduced that of the 0–1 peak, and induced a redshift in the absorption maximum (λ max). λ max and the vibronic intensity ratio exhibited a linear correlation (Figure 1C), indicating that polymers with longer alkyl bridges exhibit stronger polymer aggregation (i.e., interchain packing order) than those with shorter ones. A similar trend was also observed in the solution phase (Figure S1 and Table S1).

Subsequently, we evaluated the intrinsic charge transport characteristics of the polymers in bottom‐gate, top‐contact (BG‐TC) TFTs fabricated by spin‐coating the neat PSC films onto an n‐octadecyltrimethoxysilane (ODTMS)‐modified SiO_2_/n^+^‐Si substrate. The charge carrier mobility increased with the alkyl bridge length up to DPPT‐C8‐TT, which exhibited the highest mobility, and then decreased with further elongation (Figure 1D; Figure S2). The average mobilities were 0.59, 0.93, 1.01, and 0.61 cm^2^ V^−1^ s^−1^ for DPPT‐C2‐TT, DPPT‐C6‐TT, DPPT‐C8‐TT, and DPPT‐C16‐TT, respectively. Interestingly, although both the vibronic intensity ratio and λ max increased monotonically with increasing alkyl bridge length, the charge carrier mobility exhibited a non‐monotonic trend. This highlights the need for detailed structural analysis to identify the key morphological factors governing charge transport behavior. These factors were elucidated through subsequent structural investigations, which will be discussed in the next section.

To understand the mechanical performance of the neat PSC films, the crack onset strain was determined as a measure of stretchability. The trend in mechanical behavior was opposite to that in electrical performance (Figure 1E; Figure S3). The crack onset strain decreased with increasing alkyl bridge length up to C8, then partially recovered at longer alkyl bridge lengths. Specifically, the average values were 43%, 33%, 20%, and 33% for DPPT‐C2‐TT, DPPT‐C6‐TT, DPPT‐C8‐TT, and DPPT‐C16‐TT, respectively. This inverse correlation underscores the trade‐off between charge transport efficiency and mechanical compliance (Figure S4), highlighting the necessity of optimizing the alkyl bridge length to balance electrical functionality and mechanical durability.

Morphological Differences across Various Alkyl Bridge Lengths

2.2

To gain structural insight into molecular packing behavior, grazing‐incidence wide‐angle X‐ray scattering (GIWAXS) measurements were performed on neat PSC films (Figure 2A; Figure S5). In particular, we quantified the crystalline population (CP) represented by the volume fraction of the π–π stacking region—which is a key parameter describing the quantity of ordered domains, along with the quality of molecular ordering within those domains [49]. The relative CP, which is the CP of one film relative to that of another with the same material and thickness, can be quantified [50, 51] by separating the scattering intensity arising from amorphous domains from that arising from ordered domains [52]. However, in our case—where films with different alkyl bridge lengths represent distinct materials—such direct comparison is not feasible because of the limited reliability of the absolute GIWAXS signal intensity across the samples. To address this, we employed a normalization‐based approach, wherein the π–π stacking peak area was normalized using the amorphous halo peak area as an internal reference.

Molecular ordering analysis of polymer semiconductors (PSCs). (A) 2D GIWAXS patterns of neat PSC films with alkyl bridge lengths of C2, C8, and C16. (B) 1D in‐plane GIWAXS profile of a neat C8 PSC film with peak deconvolution into p1 (amorphous halo) and p2 (π–π stacking). (C) Plots of the relative crystalline population (CP) and molecular ordering quality (Q) as functions of the alkyl bridge length, revealing a non‐monotonic trend and maximum at C8. CP is the area ratio of p2 to p1 (p2/p1), and Q is the ratio of the second‐ to the first‐order crystal coherence length [Lc(200)/Lc(100)]. (D) Correlation between the charge carrier mobility and alkyl bridge length based on an empirical model incorporating the two structural metrics CP and Q. (E) Schematic illustrations of molecular packing in representative PSCs (C2, C8, and C16), reflecting CP and Q. (F) Structural parameter analysis showing increased lamellar d‐spacing, Lc(100), and π–π distance at longer side chains. (G,H) 3D visualization of molecular packing in C8 and C16 PSCs, incorporating data from (F).

Figure 2B displays the amorphous halo (p1) and π–π stacking peaks (p2) extracted from the 1D in‐plane GIWAXS profiles of the neat PSC films (Figure S6). As shown in Figure 2C, the relative CP, defined as the area ratio of p2 to p1, increased with the alkyl bridge length, reaching a maximum at C8, and subsequently decreasing upon further alkyl elongation. Analysis as a function of the polar angle shows that the in‐plane and out‐of‐plane directions follow the same overall trend (Figure S7). Notably, this trend closely parallels the charge carrier mobility behavior shown in Figure 1D. Moreover, to assess the quality of molecular ordering within the crystalline domains [21, 53], the ratio of the second‐ to the first‐order crystal coherence length [Lc(200)/Lc(100)] in the out‐of‐plane direction was analyzed. This ratio markedly decreased for polymers with an alkyl bridge length above C8, indicating a deterioration in internal structural regularity. Notably, the trend was observed in the ratio of the third‐ to the first‐order and third‐ to the second‐order crystal coherence length (Figure S8). By combining this ordering metric with the relative CP, we developed an empirical correlation model that quantitatively reproduced the observed mobility trend across alkyl bridge lengths: µ = −0.42 + 0.983∙CP + 0.835∙Q, where Q is the quality of molecular ordering derived from the Lc(200)/Lc(100) ratio (Figure 2D). Although polymers with a longer alkyl bridge (C12 and C16) exhibited a higher relative CP than those with a shorter one (C2 and C4), their reduced Q resulted in comparable mobilities at both extremes of the alkyl bridge length spectrum. In addition, the relative CP exhibited a strong inverse correlation with the crack onset strain (Figure S9), suggesting that the mechanical properties of thin films were strongly associated with the volume fraction of ordered regions.

To provide a visual understanding of the morphological differences across various alkyl bridge lengths, Figure 2E presents schematics of molecular packing for representative cases (C2, C8, and C16), which incorporate other structural parameters obtained from the GIWAXS measurements, including the lamellar d‐spacing of crystalline domains, π–π stacking distance, and Lc(100) (Figure 2F; Table S2). These schematics qualitatively reflect the relative CP and molecular ordering characteristics discussed above. The representative molecular packing models of C8 and C16 depicted in Figure 2G,H, respectively, further illustrate how differences in structural parameters translate into variations in molecular organization.

All three structural parameters—lamellar d‐spacing, π–π stacking distance, and Lc(100)—monotonically changed with the alkyl bridge length, showing no strong correlation with the charge carrier mobility. This indicates that conventional structural descriptors typically used for conjugated PSCs cannot fully account for the electrical performance of the alkyl‐bridge‐modulated polymer system. Instead, as demonstrated in Figure 2D, a combined consideration of the relative CP and Q offers a more accurate and predictive framework for understanding charge transport behavior, as both parameters serve as proxies for molecular ordering characteristics [49].

TFTs with Stretchable Composite Films

2.3

To explore the application of PSCs in stretchable TFTs, we employed a composite strategy aimed at enhancing mechanical stretchability [30]. Specifically, the conjugated polymers were blended with polystyrene‐block‐poly(ethylene‐ran‐butylene)‐block‐polystyrene (SEBS), a representative soft elastomer chosen for its favorable compatibility. The SEBS elastomer and DPPT‐based polymers exhibit similar surface energies, which helps suppress excessive phase separation and mitigate the morphological disruption typically observed in polymer blends [21, 30]. The absorption redshift and spectral linearity observed for neat PSC films (Figure 1B,C) were retained after SEBS blending (Figure S10 and Table S3). Similarly, GIWAXS analysis revealed that both trends in relative CP and Q (Figures S11–S13) were nearly identical to those observed for the neat PSC counterparts. These results indicate that blending with SEBS does not significantly alter the molecular ordering characteristics of PSCs with different alkyl bridge lengths. To induce molecular alignment in the composite films comprising SEBS and a DPPT‐based polymer, solution shearing was performed using a microtrench‐patterned blade on a nanogrooved SiO_2_/Si substrate [21], followed by transfer to a stretchable substrate (Figure 3A). The atomic force microscopy (AFM) phase image revealed well‐aligned nanofibers within the SEBS‐blended C8 PSC layer along the shearing direction (Figure 3B). This directional morphology is further supported by conducting‐AFM measurements, which showed local current flow following the nanofiber orientation.

Solution‐sheared polymer/SEBS composites for intrinsically stretchable thin‐film transistors (TFTs). (A) Process schematic of blade‐assisted solution shearing of a polymer/SEBS composite film, followed by PDMS‐mediated transfer onto a cross‐linked SEBS substrate. (B) AFM phase and conducting‐AFM images of a C8 composite film, showing nanofiber alignment and directional current flow along the shearing direction (scale bar: 1 µm). (C) Charge carrier mobilities of rigid TFTs fabricated using spin‐coated (SC) and solution‐sheared (SS) films with different alkyl bridge lengths. (D) Transfer characteristics of rigid TFTs across the polymer series at V DS = −10 V. (E) Comparison of the craze and crack onset strains of composite films. (F) Optical images of a C8 composite film under strain applied parallel and perpendicular to the shear direction. The right‐side panels show magnified views of the damage morphology (scale bar: 40 µm). (G and H) Mobility retention (µ/µ 0) of stretchable TFTs under increasing uniaxial strain applied parallel (G) and perpendicular (H) to the channel direction. Error bars indicate the SD for five different devices. The insets show the mobility retention at 50% strain (µ 50/µ 0). (I) Charge carrier mobilities of stretchable TFTs at V DS = −10 and −40 V and corresponding mobility retention (µs tretchable/µ rigid).

The electrical performance of the SEBS‐blended PSCs was evaluated using TFTs fabricated on a SiO_2_/Si substrate (Figure 3C; Figures S14–S16). Compared with the neat PSC devices, the spin‐coated SEBS‐blended PSC devices exhibited a 20%–60% increase in mobility, while maintaining its dependence on the alkyl bridge length. Although the overall trend was preserved, particularly from C2 to C8, longer chains (>C8) showed increased sensitivity to processing parameters such as SEBS blending and deposition method (Figure S17). Compared with the devices fabricated via spin‐coating, those fabricated via solution shearing on a nanogrooved substrate demonstrated significantly enhanced charge carrier mobility (Figure 3C,D; Figures S18 and S19). Notably, the solution‐sheared SEBS‐blended C8 PSC device exhibited the best performance, with the mobility reaching 4.9 cm^2^ V^−1^ s^−1^ and on/off current ratio exceeding 10^5^.

Figure 3E shows the mechanical properties of the solution‐sheared SEBS‐blended PSC films. The overall trend in mechanical behavior was preserved after blending with SEBS. However, the crack onset strain of the blended films increased markedly, exceeding 100%. Additionally, the crack onset strain measured perpendicular to the shearing direction was 10%–20% higher than that measured parallel to it. This anisotropy is likely due to the structural vulnerability of the aligned nanofibers under strain applied along the direction of molecular orientation.

A critical observation in this study is the onset of surface defect formation—commonly referred to as craze initiation—prior to the appearance of visible cracks (Figure 3F; Figure S20), which is also observed in AFM images (Figure S21). For the SEBS‐blended C8 PSC, which exhibited the highest charge carrier mobility, the crack onset strain was 80%–100%, whereas the craze onset strain was significantly lower, ranging from 30% to 40% (Figure 3E). Similar behavior was observed at other alkyl bridge lengths, although the gap between the craze and crack onset strain became more pronounced at longer alkyl bridges. Consequently, the craze onset strain of the SEBS‐blended PSC films tended to correspond more closely to the crack onset strain of the neat PSC films (Figure S22). This indicates that the craze behavior of the blended system more accurately reflects the inherent mechanical properties of the PSC.

Figure 3G,H shows the electrical performance of the stretchable TFTs based on the SEBS‐blended PSCs (C4, C6, C8, and C16) (Figure S23). The direction of solution shearing was aligned with the TFT channel, and uniaxial strain ranging from 0% to 50% was applied either parallel or perpendicular to the channel direction. As the strain increased, the charge carrier mobility gradually decreased. The mobility was fully recovered after applying and releasing strain in the direction perpendicular to the shearing direction, while noticeable degradation occurred under parallel strain. Notably, the rate of mobility degradation was strongly correlated with the alkyl bridge length dependence of the craze onset strain (insets of Figure 3G,H). At 50% strain, C4 exhibited the lowest mobility degradation, followed by C16, C6, and C8. This reveals a clear correlation between the electrical performance under mechanical strain and microscopic morphological response of the film. These findings indicate that the craze onset strain is a useful and predictive metric for the performance degradation of stretchable TFTs under mechanical deformation (Figure S24). In addition, the contact resistance of the TFTs was observed to be lower than the channel resistance even under strain (Figure S25).

During the transfer of the SEBS‐blended PSCs to fabricate stretchable TFTs, the charge carrier mobility decreased (Figure 3I). Although the extent of degradation varied with the alkyl bridge length, the C8‐based device exhibited the most pronounced loss, with the mobility reduced to approximately 60% of that of the corresponding rigid device. The C4‐ and C6‐based devices retained 80% and 76% of their rigid‐device mobility, respectively. The observed mobility degradation appears to be correlated with the craze onset strain trend, suggesting that mechanical damage incurred during transfer may be a contributing factor. Notably, the SEBS‐blended C8 PSC seemed to experience stress exceeding the critical threshold during transfer. Among the stretchable TFTs evaluated in this work, the SEBS‐blended C6 PSC device exhibited the highest mobility at 0% strain, showing 3.4 and 6.1 cm^2^ V^−1^ s^−1^ at V DS of −10 and −40 V, respectively. During repeated cycling between 0% and 100% strain in the direction perpendicular to the channel, the TFTs maintained a stable performance with minimal degradation, showing a mobility of 0.6 cm^2^ V^−1^ s^−1^ at the initial 100% strain and 0.3 cm^2^ V^−1^ s^−1^ after 10 000 cycles at V DS of −10 V (Figure S26).

Stretchable TFT Arrays

2.4

Intrinsically stretchable TFTs based on the SEBS‐blended C6 PSC were fabricated on an 8‐inch wafer substrate, employing cross‐linked SEBS as both the stretchable substrate and gate insulator and photopatterned microcracked gold as the stretchable gate, source, and drain electrodes. Each TFT device featured a channel width and length of 100 and 10 µm, respectively, and was evaluated on a test unit measuring 3 cm × 4 cm (Figure 4A; Figures S27 and S28). The detailed fabrication procedure is described in the Methods section and in our previous work [21]. Figure 4B,C presents the distributions of the charge carrier mobility and threshold voltage (V TH), respectively, obtained from measurements of 38 individual TFT devices. Under 0% strain (V DS = −40 V), the array exhibited an average mobility of 5.5 cm^2^ V^−1^ s^−1^ and average V TH of −0.43 V. The highest mobility obtained among the measured devices was 6.4 cm^2^ V^−1^ s^−1^ (Table S4). Figure 4D–F shows the representative output curves, transfer characteristics, and extracted mobilities of the intrinsically stretchable TFTs, respectively. In Figure 4F, the extracted field‐effect mobility exhibited slight overestimation due to indicative to non‐ideal saturation behavior [54]. The reliability factor was approximately 70%, which is slightly lower than that of rigid TFTs (Figures S16 and S19, Supporting Information), likely due to the incorporation of stretchable components. Nevertheless, the corresponding effective mobility remained high at 4.5 cm^2^ V^−1^ s^−1^, representing a resonably accurate and reliable estimation of charge transport in stretchable systems.

Wafer‐scale integration and device characteristics of an intrinsically stretchable thin‐film transistor (TFT) array. (A) Photograph of an 8‐inch wafer‐scale intrinsically stretchable TFT array based on a SEBS‐blended C6 polymer (scale bar: 200 µm). (B) Charge carrier mobility distribution across 38 devices in the array, with an average of 5.5 cm2 V−1 s−1 and maximum of 6.4 cm2 V−1 s−1. (C) Threshold voltage (V TH) distribution, centered at approximately −0.4 V. (D–F) Output (D) and transfer (E) characteristics and gate‐voltage‐dependent mobility (F) of a representative device.

Conclusions

3

We developed intrinsically stretchable TFTs with DPPT‐based conjugated PSCs by systematically engineering the alkyl side chain. This molecular design enabled precise control over the degree of crystallinity, which represents the volume fraction of ordered domains. This allowed for tunable performance within the trade‐off range between charge transport and mechanical stretchability despite their clear inverse relationship. When blended with SEBS and aligned via solution shearing, the optimized polymer exhibited a maximum mobility of 6.4 cm^2^ V^−1^ s^−1^ at V DS = −40 V under 0% strain, and the resulting stretchable devices maintained measurable mobility under large deformation in the perpendicular direction. Furthermore, we demonstrated wafer‐scale integration of photopatterned stretchable TFT arrays, highlighting the scalability and uniformity of the fabrication process. This performance level approaches those of inorganic TFTs, which are currently deployed in commercial stretchable display backplanes, underscoring the strong potential of intrinsically stretchable polymer‐based TFTs for real‐world integration. These results highlight alkyl side‐chain engineering as a practical and versatile strategy for simultaneously optimizing charge transport and mechanical resilience in stretchable semiconductors. Building on this advancement, we envision stretchable polymer‐based TFTs to play a key role in next‐generation deformable electronics, including integrated stretchable displays, bio‐interface systems, and wearable computing platforms.

Experimental Section/Methods

4

Materials

4.1

For the synthesis of the new conjugated polymers, tetrahydrofuran, allylmagnesium bromide, diglyme anhydrous, sodium borohydride, boron trifluoride diethyl etherate (purified by redistillation, ≥46.5% BF_3_ basis), triphenylphosphine, N‐bromosuccinimide, 2,5‐dihydro‐3,6‐di‐2‐thienyl‐pyrrolo[3,4‐c]pyrrole‐1,4‐dione,11‐(bromomethyl)tricosane, tris(dibenzylideneacetone)dipalladium(0), tri(o‐tolyl)phosphine, and chlorobenzene were purchased from Sigma‐Aldrich Co. (Seoul, Republic of Korea) and Tokyo Chemical Industry Co. (Seoul, Republic of Korea). 11‐(2‐Bromoethyl)tricosane, 11‐(3‐bromopropyl)tricosane, and 2,5‐bis(trimethylstannyl)thieno[3,2‐b]thiophene were purchased from Good Chem (Gangwon‐do, Republic of Korea), Derthon (Shenzhen, China), and Lumtec (Taipei, Taiwan), respectively.

For the fabrication and characterization of the devices, acetone, isopropyl alcohol (IPA), hexane, toluene, and chloroform were obtained from Samchun Pure Chemical. Anhydrous hexane, anhydrous chlorobenzene, decane, ammonium hydroxide solution (NH_4_OH, 28%), and ODTMS were obtained from Sigma‐Aldrich. Octadecyltrichlorosilane (ODTS) and Novec 7200 were obtained from Acros Organics and 3 m, respectively.1‐Bromooctane (1‐BO) and pentafluorobenzenethiol (PFBT) were purchased from Tokyo Chemical Industry. For SEBS, Tuftec H1221 (SEBS H1221, 12 wt% styrene), used as the encapsulation layer, and Tuftec H1052 (SEBS H1052, 18 wt% styrene), used as the stretchable substrate and gate insulator, were purchased from Asahi Kasei. The Kraton G1652 M polymer (SEBS G1652, 30 wt% styrene) used as the elastic semiconducting layer was purchased from Trimex Co., Ltd. Polydimethylsiloxane (PDMS; Sylgard 184), used for the crack onset strain test and transfer process, was purchased from Dow Corning. Bis(6‐((4‐azido‐2,3,5,6‐tetrafluorobenzoyl)oxy)hexyl)decanedioate, used as the azide crosslinker, was purchased from Medigen. All materials were used without further purification.

Synthesis of New Conjugated Polymers

4.2

The conjugated polymers with controlled alkyl bridge lengths (C2, C4, C6, C8, C10, C12, and C16) were synthesized as follows. The DPPT monomer (0.2 mmol, 1.0 equiv), 2,5‐bis(trimethylstannyl)thieno[3,2‐b]thiophene (0.2 mmol, 1.0 equiv), tris(dibenzylideneacetone)dipalladium(0) (0.006 mmol, 0.03 equiv), and tri(o‐tolyl)phosphine (0.024 mmol, 0.12 equiv) were dissolved in chlorobenzene (10 mL). The solution was flushed with nitrogen for 10 min and then sealed. The mixture was stirred at 120°C for 24 h and then cooled to room temperature and precipitated in methanol. The precipitated solid was collected by filtration and purified by Soxhlet extraction sequentially using methanol, acetone, hexane, CH_2_Cl_2_, and CHCl_3_. The CHCl_3_ solution was collected and then concentrated under reduced pressure. The concentrated solution was added dropwise to excess methanol. The precipitate was collected by filtration and dried under vacuum at room temperature for 12 h to afford the DPPT‐TT polymer. The molecular weights and polydispersity indices of the polymers are given in Table S1.

Fabrication of Rigid TFTs

4.3

BG‐TC‐type rigid TFTs were fabricated by spin‐coating a neat PSC or SEBS‐blended PSC film onto a ODTMS‐treated SiO_2_ (300 nm)/n^+^‐Si (ODTMS‐SiO_2_) substrate, which was prepared according to our previously reported method (Figure S14) [21]. For the neat PSC film, a PSC solution was prepared by stirring a chloroform solution containing DPPT‐TT (5 mg mL^−1^) at 45°C for 6 h. Before spin‐coating, the solution was filtered through a 0.45‐µm polyvinylidene fluoride (PVDF) filter. For the SEBS‐blended PSC film, a chlorobenzene solution containing DPPT‐TT and SEBS at a weight ratio of 4:6 (7.5 mg mL^−1^) was stirred at 140°C for 2 h and then filtered through a 0.45‐µm PVDF filter. Subsequently, the solution was spin‐coated onto the ODTMS‐SiO_2_ substrate at 1000 rpm. After annealing the thin film at 190°C for 1 h in a glove box, a stacked layer comprising MoO_x_ (10 nm) and Au (100 nm) was thermally deposited as source (S)/drain (D) electrodes through a metal shadow mask. The channel length and width were 100 and 1000 µm, respectively.

For the BG‐TC‐type rigid TFTs based on solution‐sheared films, microtrench patterns on a silicon blade for solution shearing were prepared using a standard photolithography technique. After photopatterning, the blade was immersed in an ODTS solution (4 mM, anhydrous hexane) for 1.5 h, annealed at 120°C for 30 min, and cleaned with chloroform. The SEBS‐blended PSC solution (DPPT‐TT:SEBS = 4:6, 20 mg mL^−1^ in chlorobenzene) was sheared onto the ODTMS‐SiO_2_ substrate at a speed within the range of 1–2 mm s^−1^. The tilt angle of the ODTS‐treated blade relative to the substrate and the gap between the blade and substrate were 8° and 20 µm, respectively. The substrate temperature during solution shearing was set to 70°C (Figure S29). The sheared film was annealed at 190°C for 1 h in a glove box. BG‐TC TFTs were fabricated by thermally depositing S/D electrodes (MoO_x_ (10 nm)/Au (100 nm)) onto the sheared film on the ODTMS‐SiO_2_ substrate. The channel length and width were 100 and 1000 µm, respectively.

Fabrication of Intrinsically Stretchable TFTs

4.4

The devices were fabricated on an 8‐inch glass substrate following our previous report [21]. Briefly, the layer stack consisted of dextran (sacrificial layer), hard crosslinked SEBS (stretchable substrate), microcracked gold (gate electrode), crosslinked SEBS (gate insulator, capacitance: 3.9 nF/cm^2^), and microcracked gold (S/D electrode) [21]. A solution‐sheared SEBS‐blended PSC film formed on an ODTMS‐modified SiO_2_ substrate was then transferred onto the substrate surface treated with PFBT and 1‐BO [21]. The PSC film was patterned by mechanical scribing using a probe tip mounted on a precision positioner. Finally, an SEBS encapsulation layer was formed, the dextran layer was removed, and an SEBS support film was attached to the backside of the device, followed by drying under vacuum at room temperature for >30 min [21].

Characterization

4.5

^1^H and ^13^C nuclear magnetic resonance (NMR) spectra were measured at room temperature using a Bruker AVANCE III HD 500 MHz spectrometer with CDCl_3_ as the solvent and tetramethylsilane as the internal standard (Data S1–S14). The molecular weights and polydispersity indices of the polymers relative to those of a polystyrene standard were measured by gel permeation chromatography using an Agilent PL‐GPC 220 system equipped with a TSKgel GMHHR‐H(S) HT2 column (Tosoh Bioscience LLC). The polymers were eluted with 1,2,4‐trichlorobenzene heated to 160°C at a flow rate of 1 mL min^−1^. UV–vis absorption spectra were measured using a Shimadzu UV‐3600 Plus UV–Vis–NIR spectrophotometer. The thickness of the thin films was measured using a KLA Tencor P‐17 profiler. GIWAXS measurements were conducted at the 3C SAXS beamline of the Pohang Accelerator Laboratory. An incidence angle of approximately 0.12° was selected to ensure complete X‐ray penetration of the film. The exposure time was 5 s for all samples. All electrical properties of the TFTs were measured using a probe station connected to a Keithley 4200A‐SCS parameter analyzer. The mechanical properties of the thin films were determined using a film‐on‐elastomer method. Each film on an ODTMS‐SiO_2_ substrate was first transferred onto a PDMS substrate (precursor:crosslinker = 12.5:1, w/w). The films were subjected to strain to observe crack formation and determine the crack onset strain by optical microscopy (Olympus DSX1000). For the stretching tests of the TFTs, a handmade jig was used to fix the film in a stretched state. A TFT substrate cut to a size of 3 cm × 4 cm was placed in the jig, and the electrical properties were measured in the stretched state. The capacitance of the gate insulator in the stretched state was calculated by assuming a constant dielectric constant and estimating the thickness from the microscopically measured area under the assumption of a uniform volume.

In stretchable TFTs, the reliability factor (r) and effective mobility (µ eff) were calculated from the device parameters and the TFT characteristics [54], as follows: r=|IDSmax|−|IDS0||VGSmax|2/∂|IDS|∂VGS2, µ_ eff _ =  r × µ. Here, |IDSmax| is the experimental maximum source‐drain current at the maximum gate voltage |VGSmax|, and |IDS0| is source‐drain current at V GS = 0 V.

Author Contributions

H.K., H.K., Y.K., and G.H.L. conceived and designed the experiments and interpreted the results. H.K. and Y.K. conducted the fabrication and characterization. H.K. and B.‐L.L. synthesized the polymers. H.K. and G.H.L analyzed the experimental data and wrote the manuscript. E.K., A.C., K.K., S.‐G.K, J‐Y.K, and J.Y.J. contributed to the fabrication and characterization of materials and devices. S.G., J.S. Y.L., and S.‐J.L. contributed to the synthesis of the polymers. Y.Y. and G.H.L. oversaw the project, revised the manuscript, and led the effort to completion.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: smll72251‐sup‐0001‐SuppMat.docx