Optically transparent semiconducting polymer nanonetwork for flexible and transparent electronics

Significance

When various electronic appliances used in everyday life become deformable and transparent, they will provide tremendous versatility in the design and use of see-through, smart mobile applications, exceeding the limitations of the best developed conventional silicon technologies, which are available only in rigid, opaque forms. However, even recently discovered innovative semiconducting components have failed to simultaneously achieve such flexibility and transparency. Thus, the existing options still comprise only hard, planar, or opaque materials, and obtaining a “key” material for creating truly flexible and transparent electronics has presented a formidable challenge. We report an effective means of creating a “truly flexible, perfectly transparent” and high-mobility semiconducting material and demonstrate several high-end flexible and transparent applications based on a polymeric semiconductor system.

Abstract

Simultaneously achieving high optical transparency and excellent charge mobility in semiconducting polymers has presented a challenge for the application of these materials in future “flexible” and “transparent” electronics (FTEs). Here, by blending only a small amount (∼15 wt %) of a diketopyrrolopyrrole-based semiconducting polymer (DPP2T) into an inert polystyrene (PS) matrix, we introduce a polymer blend system that demonstrates both high field-effect transistor (FET) mobility and excellent optical transparency that approaches 100%. We discover that in a PS matrix, DPP2T forms a web-like, continuously connected nanonetwork that spreads throughout the thin film and provides highly efficient 2D charge pathways through extended intrachain conjugation. The remarkable physical properties achieved using our approach enable us to develop prototype high-performance FTE devices, including colorless all-polymer FET arrays and fully transparent FET-integrated polymer light-emitting diodes.

Optically transparent and mechanically flexible circuitries have long been desired for next-generation electronics requiring unprecedented features, such as “see-through” visibility, deformability, and even skin-attachable functionality for health care systems (1–3). This new paradigm for electronic applications has motivated researchers to eagerly pursue new innovative semiconducting materials, and one promising candidate is the class of materials called semiconducting conjugated polymers (4). Their unique benefits, including mechanical flexibility, light weight, and processing advantages based on high-throughput fabrication processes using solution-printing technologies, have accelerated the development of these materials as key building blocks for next-generation ubiquitous systems (2, 5, 6). Nevertheless, these materials still cannot fulfill the ultimate requirements for future “flexible” and “transparent” electronics (FTEs). Together with their inferior charge-carrier mobility because of conformational and energetic disorder (7), their high light absorption in the visible range, which is inherent to this class of materials (absorption coefficient ∼10⁵ cm⁻¹) (8), makes it difficult to apply these materials in FTEs. Indeed, despite extensive investigations seeking a suitable model system for FTEs by varying the polymer-structure design and the processing techniques used, the simultaneous achievement of optical transparency and high mobility in semiconducting polymers remains a formidable challenge (9, 10).

Among the various types of semiconducting polymers, low-bandgap polymers using the donor–acceptor (D-A) copolymerization scheme are promising candidate materials for FTE applications. These semiconducting copolymers usually exhibit much less absorption in the visible range compared with other typical midbandgap polymers because of their red-shifted π–π* absorption spectrum, which exhibits strong absorption in the near-infrared (IR) region (11). Several D-A copolymers have recently been found to show exceptionally high mobility (exceeding 1 cm² V⁻¹⋅s⁻¹), despite their relatively low crystalline order (12–14). However, because of their high optical density, even for ultrathin films (thickness t < 100 nm), it remains difficult to obtain fully transparent and colorless thin films using copolymers of this type. Moreover, obtaining high mobility typically requires undesirable processing techniques, such as high-temperature annealing (12) and macroscopic alignment processes (15), which are not readily compatible with flexible electronics. Therefore, the realization of truly colorless semiconducting layers with high mobility for FTEs remains to be achieved.

Multicomponent systems consisting of various polymer blends have recently attracted particular attention because of the tunability of their material properties (16). Recent reports have shown that blends containing a relatively small amount of semiconducting polymer in an inert polymer matrix exhibit charge-transport characteristics that are comparable or even superior to those of the pristine forms (17–22). However, a comprehensive understanding of the underlying mechanism of this intriguing phenomenon has not yet been achieved. We note that this polymer-blending approach can provide new opportunities for the development of innovative polymer systems for FTEs. Here, by introducing a diketopyrrolopyrrole (DPP)-based semiconducting copolymer (DPP2T) into an inert polystyrene (PS) matrix, we create a polymer blend system that demonstrates both high mobility and high transparency approaching 100% without developing any color. We discover that a small amount of DPP2T in an amorphous PS matrix forms a web-like, continuously connected nanonetwork that spreads throughout the thin film formed during solution deposition while remaining confined in a thin fibrous structure. Detailed study reveals that this network structure of DPP2T provides highly efficient charge pathways with substantially reduced structural and energetic disorder through its extended intrachain conjugation. This approach therefore enables us to fabricate prototype high-performance FTE devices.

Results and Discussion

The molecular structures of DPP2T and PS are depicted in Fig. 1A. DPP2T consists of a strong electron-deficient unit, namely, DPP and an electron-rich segment of 2,5-di-2-thienylthieno[3,2-b]thiophene. In general, combining strongly fused moieties causes the resulting polymer backbone to be planar and rigid (12). We used PS as the inert polymer (E_g ∼4 eV, SI Appendix, Fig. S1) to prepare a fully transparent DPP2T/PS blend system. Pure DPP2T and DPP2T/PS (15/85 wt % ratio) composite powders were separately dissolved in organic solvents, and the solutions were subsequently spin cast onto substrates to obtain thin films (thickness t ∼10 nm).

Fig. 1.

Optical and structural studies of DPP2T and DPP2T/PS films. (A) Chemical structures of DPP2T and PS. (B) Optical absorption spectra of DPP2T and DPP2T/PS (15/85 wt % ratio) in solution and as solid-state films. (C) Optical transmittance (T_r) spectra of DPP2T and DPP2T/PS films. (Inset) Optical image of a pure DPP2T film and a DPP2T/PS film on glass substrates. The t of both films is ∼10 nm. (D) TEM images of DPP2T and DPP2T/PS films. (Scale bar, 200 nm.) (E) Normalized 2D GIWAXS patterns of DPP2T and DPP2T/PS films. (F and G) Schematic illustrations of the structural morphologies of (F) DPP2T and (G) DPP2T/PS films.

DPP2T exhibits strong near-IR absorption (onset at a wavelength of λ = 920 nm; energy gap of E_g ∼1.35 eV) arising from the alternating structure of repeating units of DPP and 2,5-di-2-thienylthieno[3,2-b]thiophene (23) (Fig. 1B). This characteristic absorption spectrum makes DPP2T far more transparent in the visible range than other midbandgap polymers such as poly(3-hexylthiophene), but it still exhibits considerable absorption in this spectral region (SI Appendix, Fig. S2) and has a greenish tint (23), as shown in Fig. 1C. By contrast, a blend system of DPP2T/PS that contains a small proportion of DPP2T (15 wt % with respect to PS) is almost perfectly transparent throughout the visible range (average transmittance of T_a ∼99% for 380–700 nm; SI Appendix, Fig. S3) and develops no color (Fig. 1C, Inset).

Comparing the nanomorphology of the films using transmission electron microscopy (TEM) reveals a striking feature of the DPP2T/PS blend system (Fig. 1D). Whereas the TEM image of pure DPP2T is typical of this class of materials, showing stacked polymer aggregates with random orientations, DPP2T/PS shows a phase-separated heterostructure with a web-like polymer network that consists of linked fibrils with widths of a few tens of nanometers. The continuously connected 2D nanonetwork structure consists solely of DPP2T, as confirmed through the elemental mapping of sulfur (S) via energy-dispersive X-ray spectroscopy line-scan analysis in high-angle annular dark-field scanning-transmission electron microscopy mode (SI Appendix, Fig. S4).

A more detailed structural analysis using grazing-incidence wide-angle X-ray scattering (GIWAXS) measurements also reveals a large difference in the nanomorphology of the polymer chains between the two films (Fig. 1E). The 2D GIWAXS patterns of DPP2T exhibit a dominant edge-on orientation, as inferred by the existence of two patterns: one consisting of (h00) Bragg peaks in the out-of-plane direction with a lamellar spacing of 23.2 Å and the other consisting of a (010) peak in the in-plane direction with a π–π-spacing of 3.85 Å. However, DPP2T is estimated to possess rather randomly oriented crystallites and considerably disordered regions, as indicated by the broad (h00) peaks and the wide arc pattern around q ∼1.32 Å⁻¹, which arises from the amorphous polymer phase (7) (Fig. 1F). By contrast, DPP2T/PS exhibits no characteristic peaks but rather a wide halo attributed to amorphous PS (and probably also to DPP2T) (SI Appendix, Fig. S5), indicating no crystalline-ordered structures of the polymer chains (detailed 1D profile studies are presented in SI Appendix, Fig. S6). The thin, fiber-like DPP2T bundles in the nanonetwork structure of DPP2T/PS, which seem to be composed of only a few polymer chains, as shown in Fig. 1G, produce quite different GIWAXS patterns (see SI Appendix, Note S1 for additional structural studies based on optical absorption spectra).

Based on the remarkable contrast in the nanostructures of the films, we can expect to observe fundamentally different charge transport between the two systems. Therefore, we analyzed the charge-transport characteristics of the two films by fabricating top-gate, bottom-contact (TGBC) field-effect transistors (FETs) using the DPP2T and DPP2T/PS layers (SI Appendix, Fig. S7). Fig. 2A compares the devices’ transfer characteristics at room temperature. Both devices exhibit clear p-channel FET characteristics with negligible hysteresis and low contact resistance (SI Appendix, Fig. S8). However, whereas the saturation- and linear-regime field-effect mobilities (μ) of the DPP2T FET are as high as 0.80 cm² V⁻¹⋅s⁻¹ and 0.52 cm² V⁻¹⋅s⁻¹, respectively (with average μ-values of μ_asat = 0.67 cm² V⁻¹⋅s⁻¹ and μ_alin = 0.40 cm² V⁻¹⋅s⁻¹), the DPP2T/PS FET exhibits much higher maximum saturation- and linear-regime μ-values of 3.1 cm² V⁻¹⋅s⁻¹ and 1.6 cm² V⁻¹⋅s⁻¹, respectively (with μ_asat = 1.6 cm² V⁻¹⋅s⁻¹ and μ_alin = 1.0 cm² V⁻¹⋅s⁻¹) (SI Appendix, Figs. S9 and S10). Therefore, we can clearly observe that the charge-transport characteristics of the DPP2T/PS devices are substantially enhanced compared with those of the pure DPP2T devices. Considering that the content of semiconducting DPP2T in these DPP2T/PS devices is only ∼15% (corresponding to an insulating PS content of ∼85%), which is much less than that in pure DPP2T devices, this observation is certainly worth further investigation.

Fig. 2.

Correlation between the transport characteristics and morphological properties of DPP2T/PS blend films. (A) Representative room-temperature transfer characteristics of pure DPP2T and DPP2T/PS FETs under a V_DS of −60 V. The channel length and width of the devices are 40 μm and 1 mm, respectively. (B) Hole μ-values of FETs fabricated with DPP2T/PS blend films of various concentration ratios. The curved line indicates the trend in the hole μ-value observed with the variation of the DPP2T concentration. The vertical lines (whiskers) indicate the range from the 10th to the 90th percentile. The minimum and maximum values are indicated by asterisks. (Inset) Schematic illustration of the device structure. (C) TEM images of DPP2T/PS films of various concentration ratios. These images show the evolutionary phase change of the DPP2T in the PS matrix as the concentration of DPP2T increases. (Scale bar, 200 nm.)

We hypothesize that the improved transport properties of the DPP2T/PS devices are associated with the 2D nanonetwork formation in the DPP2T/PS blend films. Therefore, to more directly observe the correlation between the nanomorphologies and the charge-transport properties of DPP2T/PS films, we investigated the room-temperature FET characteristics of DPP2T/PS blend films with various concentration ratios (Fig. 2B). Here, the corresponding TEM images of DPP2T/PS films were found to show quite different nanomorphologies for different blending ratios, as shown in Fig. 2C. The morphological phase formation of different concentration of DPP2T in a PS matrix can be roughly divided into three regimes: (i) at low DPP2T concentrations, the DPP2T forms long, thin fibrils, which become thicker and more highly interconnected as the DPP2T concentration increases; (ii) at a moderate DPP2T concentration, the DPP2T forms completely percolated 2D nanonetworks; and (iii) at high DPP2T concentrations, the fibrillar bundles are aggregated and stacked, and the morphology thus approaches that of the pure DPP2T phase. These morphological changes are expected to affect the charge-transport properties of the blend films. Fig. 2B shows the dependence of the saturation-regime μ on the DPP2T/PS blending ratio (SI Appendix, Table S1, Fig. S11, and Note S2). With an increasing DPP2T content, the DPP2T/PS devices exhibit rapidly increasing μ beginning at 3 wt % DPP2T (with respect to the total solid) and peak μ at 15 wt %. In this range, the best charge pathways seem to form as the fibrous structures become percolated. Further increasing the DPP2T content above 15 wt % results in a gradual decrease in μ as the characteristics of DPP2T/PS approach those of pure DPP2T. Therefore, we can confirm that the trend in μ observed for the devices directly parallels the morphological trend observed in the DPP2T/PS.

To obtain further insight into the charge-transport regimes and the detailed transport mechanisms, we investigated the temperature (T)-dependent FET characteristics of DPP2T and DPP2T/PS. Here, the μ-values were measured in the linear regime under various low drain-source voltages (V_DS), which were assumed to yield a uniform electric field strength (F) along the channel (SI Appendix, Figs. S12 and S13). Arrhenius plots of the T-dependent μ-values are presented in Fig. 3A, from which we extracted the activation energies (E_A) using the Arrhenius relation, μ ∝ exp(−E_A/kT), where k is the Boltzmann constant. Interestingly, the plots show two different activation regimes, with a transition occurring at T ∼190 K (24, 25). In the low-T regime, DPP2T shows an E_A of ∼19 meV, which is fairly low compared with the values reported for other high-mobility polymers (26, 27). By contrast, a surprisingly low E_A of ∼5 meV is observed for DPP2T/PS (28). Similarly, a lower E_A value for DPP2T/PS (compared with that for pure DPP2T) is observed even in the high-T regime. The transition temperature observed in Fig. 3A is expected to correspond to the thermal energy necessary to overcome the local hole barrier between aggregated and amorphous DPP2T regions, which causes the charge transport to be spatially confined within these ordered regions at low T (SI Appendix, Fig. S14). This explains why the same transition occurs regardless of the level of PS blending. From a structural perspective, the large, randomly oriented aggregates surrounded by amorphous regions that are present in the pure DPP2T film cause the paracrystallinity-dominated π–π transport to act as the rate-limiting factor (7), whereas efficient chain-backbone transport becomes the most relevant contribution in the DPP2T/PS system (29, 30) (Fig. 1 F and G and SI Appendix, Fig. S15). Consequently, the carrier drift in the DPP2T/PS film exhibits exceptionally low thermal coefficients, reflecting single-chain conduction with occasional hopping at interchain connections. In addition, the decrease in the E_A of DPP2T as F increases in the high-T regime implies the occurrence of F-assisted tunneling and hopping at pervasive localized states including deep traps (31), whereas DPP2T/PS shows no noticeable F dependence (Fig. 3B). Furthermore, in semilogarithmic plots of μ as a function of F^1/2 (Fig. 3C), the linear increase in μ with F^1/2 indicates that DPP2T exhibits Poole–Frenkel-like behavior, which is characteristic of disorder-limited charge transport (32, 33). By contrast, DPP2T/PS shows an almost negligible F dependence of μ, which indicates the formation of nearly undisturbed charge pathways (see SI Appendix, Note S3 and Figs. S16 and S17 for further theoretical analyses of the charge transport in DPP2T/PS).

Fig. 3.

Measurements and modeling of T-dependent device characteristics. (A) Arrhenius plots of the T-dependent linear μ-values for DPP2T and DPP2T/PS FETs under various V_DS values: −2 V (squares), −4 V (circles), −6 V (triangles), and −8 V (inverted triangles). (B) E_A as a function of the channel F obtained in the high-T (>190 K) and low-T (<190 K) regimes for DPP2T and DPP2T/PS. (C) Semilogarithmic plots showing the linear variation in the linear μ as a function of F^1/2 for DPP2T and DPP2T/PS FETs at T = 280 K.

It is interesting that the difference in the morphological state modulated by PS blending can result in such large variations in the charge-transport characteristics of DPP2T. DPP2T has a rigid backbone structure because of its strongly fused moieties, which are responsible for its strong intrachain and interchain charge transfer and relatively high intermolecular stacking. However, its crystallization is disturbed because of the complex chain entanglement that unavoidably occurs during solution processing because of the limited space available for the movement of the crowded DPP2T chains and fibrillar bundles (34) (SI Appendix, Fig. S18). Therefore, a pure DPP2T film is composed of not only well-ordered interchain stacks but also disordered amorphous regions, as shown in Fig. 1F.

By contrast, the DPP2T chains in DPP2T/PS are diluted by the PS matrix and confined within a narrow but continuously connected web-like structure, in which large-scale π–π and lamellar stacking are inhibited. The separation behavior of DPP2T and PS is attributed to the immiscibility of these two polymers (35). However, the flexible PS matrix seems to prevent the entanglement of the rigid DPP2T chains by providing a more flexible surrounding environment (SI Appendix, Note S4 and Fig. S19). Therefore, the unique structure of DPP2T/PS is assumed to consist of a series of DPP2T units extending along the long axis of each fibril, in which the structural discontinuity is significantly reduced as the interchain aggregation decreases (see the schematic illustration in Fig. 1G). Therefore, DPP2T/PS is expected to support efficient intramolecular charge transport through continuously connected “clean” pathways without severe structural or energetic disorder because the extended DPP2T units promote a greater extent of intramolecular charge delocalization (26, 36), and thus, its near-intrinsic molecular performance manifests at the length scale of the device channels (see SI Appendix, Notes S5 and S6 and Figs. S20–S24 for a more detailed discussion).

The excellent carrier mobility and high transparency of DPP2T/PS allow us to directly use this material in FTE applications. To this end, we prepared flexible and transparent FETs (FT-FETs) with a TGBC configuration on poly(ethylene-2,6-naphthalate) (PEN) films (t ∼125 μm), in which all layers were fabricated from polymeric materials via solution processing. Fig. 4A shows a photograph and an illustration of the device structure of our large-area FT-FET array (10 cm × 10 cm, 1,650 FETs). To fabricate the transparent source/drain and top-gate electrodes, poly(3,4-ethylenedioxythiophene):poly(styrenesulfonate) (PEDOT:PSS) was inkjet-printed onto the PEN substrate and the gate insulating layer (SI Appendix, Fig. S25). Fig. 4B shows the optical transmittance spectra of the FT-FETs measured under different layering conditions (SI Appendix, Table S2). Because all of the layers on the PEN substrate are thin and highly transparent, the transmittance of the FET device is quite high throughout the visible range (T_a ∼86% for 380–700 nm), almost equal to that of the pure PEN substrate (SI Appendix, Fig. S26). The transfer characteristics of the FT-FETs are presented in Fig. 4C. Although a slight hysteresis and wavy curves are observed in the output characteristics (SI Appendix, Fig. S27), because of the relatively low conductivity of the PEDOT:PSS electrodes (σ < 700 S cm⁻¹), overall, FT-FETs were uniformly fabricated over the entire device area and showed stable and clear field-effect characteristics, with a high saturation-regime μ of 0.80 cm² V⁻¹⋅s⁻¹ (μ_a = 0.64 cm² V⁻¹⋅s⁻¹) and a linear-regime μ of 0.23 cm² V⁻¹⋅s⁻¹ (μ_a = 0.15 cm² V⁻¹⋅s⁻¹). Our FT-FETs are also durable against bias stress and mechanical deformation (Fig. 4D); after 2,000 switching cycles and up to bending 1,000 cycles (at a bending radius of R ∼5 mm), we detect no severe performance degradation (Fig. 4E and SI Appendix, Fig. S28).

Fig. 4.

Large-area FT-FET device. (A) A photograph and an illustration of the FET device structure of a large-area FT-FET device. In the magnified optical microscopy image, inkjet-printed PEDOT:PSS electrodes are shown. (Scale bar, 500 μm.) (B) T_r spectra of FT-FETs measured under different layering conditions. (C) Transfer characteristics of FT-FETs. The channel length and width are 100 μm and 1,000 μm, respectively. (D) Cycle stability test of FT-FETs with an alternating on/off gate-voltage (V_GS) pulse (1 Hz) under a constant V_DS of −30 V before and after 1,000 bending cycles at a bending radius of R ∼5 mm. (E) Transfer characteristics of the FT-FETs before and after 1,000 bending cycles and 2,000 electrical switching cycles (1 Hz) with alternating gate pulses of −60 V (on state) and 0 V (off state) under a V_DS of −30 V.

To take full advantage of the advantageous characteristics of DPP2T/PS and to demonstrate a potential electronic application requiring high-performance transparent driving circuits, we also fabricated prototypes of flexible and transparent FET-integrated polymer light-emitting diodes (FT-FET-PLEDs). Fig. 5 A and B shows a schematic illustration and photographs of our FT-FET-PLED devices, which were prepared by fabricating transparent DPP2T/PS FETs directly on top of PLEDs to drive them (one transistor, one diode architecture) (SI Appendix, Fig. S29). A thin Au (t ∼15 nm) drain electrode on the FET acts as the semitransparent anode of the PLED, and the generated light is emitted through the transparent FET layers on top (SI Appendix, Fig. S30). The molybdenum oxide (MoO_x)/PEDOT:PSS hole-injecting layer offers a supporting surface for sequential solution processing for the fabrication of the top FET. The top FET operates by supplying a gate voltage (V_GS)-modulated drain-source current (I_DS) to the PLED, which has an aryl-substituted poly(p-phenylene vinylene) derivative (PDY-132) emissive layer, resulting in the stable modulation of yellow light with a maximum luminance (L) of ∼252 cd m⁻² (Fig. 5C). The corresponding electroluminescence (EL) spectrum does not show any significant differences from that of the PLED without an FET because of the high transparency of the FET components (Fig. 5D). Additionally, FT-FET-PLED devices with different emission colors were prepared using various emissive layers consisting of the white-light-emitting polymer SPW-111, poly[2-methoxy-5-(2-ethylhexyloxy)-1,4-phenylenevinylene] (MEH-PPV) and poly(9,9-di-n-octylfluorene-alt-benzothiadiazole) (F8BT), and the resulting devices successfully demonstrate the modulation of the emitted light by the top FET even during bending (Fig. 5E and SI Appendix, Figs. S31 and S32 and Movie S1). Thus, our system represents a major step toward achieving transparent and deformable all-polymer active-matrix displays.

Fig. 5.

FT-FET-PLED integrated devices. (A) Schematic illustration of the FT-FET-PLED device structure. (B) Photographs of FT-FET-PLED devices with different emissive layers: PDY-132 (yellow), SW-111 (white), MEH-PPV (red), and F8BT (green). The photographs show the PLEDs being driven by the integrated FETs during bending. (C) The L of the yellow FT-FET-PLED device and the I_DS supplied to the PLED by the integrated FET as a function of V_GS. The driving voltage (V_DS) is fixed at −60 V. For the integrated FET, the channel length and width are 50 μm and 38 mm, respectively. (D) EL spectra of normal yellow PLED and integrated yellow FT-FET-PLED devices. (E) L-V_GS characteristics of FT-FET-PLED devices with various emissive layers. V_DS is fixed at −60 V.

In summary, our results describe a 2D copolymer network with both ultratransparency and high μ prepared using the DPP2T/PS blend system. The dimensionally confined 2D nanonetwork structure of the DPP2T in the PS enables much more efficient charge transport through the extended intrachain conjugation of disorder-reduced clean pathways. Simultaneously, the DPP2T/PS blend system also offers exceptional transparency that approaches 100% because of its low visible-light absorption and thin net-like structure. We demonstrated the fabrication of FTE devices using our completely colorless, high-μ DPP2T/PS system. Our results provide scientific insights into the physics of charge transport in conjugated polymer systems and a pathway toward unprecedented technological applications in organic electronics.

Materials and Methods

DPP2T/PS Blend Solution.

DPP2T (M_r ∼73 kg mol⁻¹, 1-Material) and PS (M_r ∼95 kg mol⁻¹, Sigma-Aldrich) were precisely weighed and mixed at various weight ratios. Pure DPP2T and DPP2T/PS were separately dissolved in 1,2,4-trichlorobenzene/chloroform (80/20 volume % ratio) at a fixed total concentration of 2 mg mL⁻¹.

TGBC FET Fabrication.

Glass slides (Eagle XG^TM, Corning) were cleaned via sequential ultrasonication in water, acetone, and isopropyl alcohol. Thermally evaporated Au (15 nm) source/drain electrodes were patterned using shadow masks. The blend solution was spin cast onto the substrate in an inert nitrogen atmosphere, and the films were subsequently annealed at 80 °C for 10 min to remove residual solvent. The thickness of the semiconducting layers was controlled to be ∼10 nm. For the gate-insulating layer, CYTOP (CTL-809M, Asahi Glass Co., Ltd.) diluted with CT-Solv.180 solvent (4:1 volume ratio) was used. The insulating materials were spin cast onto the semiconducting layer, resulting in a thickness of ∼550 nm, and the films were subsequently annealed at 80 °C for 30 min. The measured capacitance of the CYTOP layer was ∼3.5 nF cm⁻². The devices were completed with the thermal deposition of 50 nm of Al through a shadow mask to form the top-gate electrode.

More experimental details are shown in SI Appendix.