Optically transparent semiconducting polymernanonetwork for flexible and transparent electronicsKilho Yua,b,c,1, Byoungwook Parka,b,c,1, Geunjin Kimb,c, Chang-Hyun Kima,c, Sungjun Parka, Jehan Kimd, Suhyun Junga,b,c,Soyeong Jeonga,b,c, Sooncheol Kwonb,c, Hongkyu Kangb,c, Junghwan Kimb,c, Myung-Han Yoona,and Kwanghee Leea,b,c,2

aDepartment of Nanobio Materials and Electronics, School of Materials Science and Engineering, Gwangju Institute of Science and Technology, Gwangju61005, Republic of Korea; bHeeger Center for Advanced Materials, Gwangju Institute of Science and Technology, Gwangju 61005, Republic of Korea;cResearch Institute for Solar and Sustainable Energies, Gwangju Institute of Science and Technology, Gwangju 61005, Republic of Korea; and dPohangAccelerator Laboratory, Pohang University of Science and Technology, Pohang 37673, Republic of Korea

Edited by Tobin J. Marks, Northwestern University, Evanston, IL, and approved October 20, 2016 (received for review May 1, 2016)

Simultaneously achieving high optical transparency and excellentcharge mobility in semiconducting polymers has presented achallenge for the application of these materials in future “flexible”and “transparent” electronics (FTEs). Here, by blending only asmall amount (∼15 wt %) of a diketopyrrolopyrrole-based semi-conducting polymer (DPP2T) into an inert polystyrene (PS) matrix,we introduce a polymer blend system that demonstrates bothhigh field-effect transistor (FET) mobility and excellent opticaltransparency that approaches 100%. We discover that in a PS ma-trix, DPP2T forms a web-like, continuously connected nanonet-work that spreads throughout the thin film and provides highlyefficient 2D charge pathways through extended intrachain conju-gation. The remarkable physical properties achieved using our ap-proach enable us to develop prototype high-performance FTEdevices, including colorless all-polymer FET arrays and fully trans-parent FET-integrated polymer light-emitting diodes.

semiconducting polymer | organic electronics | flexible and transparentdevice | polymer blend | charge transport

Optically transparent and mechanically flexible circuitrieshave long been desired for next-generation electronics re-

quiring unprecedented features, such as “see-through” visibility,deformability, and even skin-attachable functionality for healthcare systems (1–3). This new paradigm for electronic applica-tions has motivated researchers to eagerly pursue new innovativesemiconducting materials, and one promising candidate is theclass of materials called semiconducting conjugated polymers(4). Their unique benefits, including mechanical flexibility, lightweight, and processing advantages based on high-throughputfabrication processes using solution-printing technologies, haveaccelerated the development of these materials as key buildingblocks for next-generation ubiquitous systems (2, 5, 6). Never-theless, these materials still cannot fulfill the ultimate require-ments for future “flexible” and “transparent” electronics (FTEs).Together with their inferior charge-carrier mobility because ofconformational and energetic disorder (7), their high light ab-sorption in the visible range, which is inherent to this class ofmaterials (absorption coefficient ∼105 cm−1) (8), makes it diffi-cult to apply these materials in FTEs. Indeed, despite extensiveinvestigations seeking a suitable model system for FTEs byvarying the polymer-structure design and the processing tech-niques used, the simultaneous achievement of optical trans-parency and high mobility in semiconducting polymers remains aformidable challenge (9, 10).Among the various types of semiconducting polymers, low-

bandgap polymers using the donor–acceptor (D-A) copolymeriza-tion scheme are promising candidate materials for FTE applica-tions. These semiconducting copolymers usually exhibit much lessabsorption in the visible range compared with other typical mid-bandgap polymers because of their red-shifted π–π* absorptionspectrum, which exhibits strong absorption in the near-infrared (IR)

region (11). Several D-A copolymers have recently been found toshow exceptionally high mobility (exceeding 1 cm2 V−1·s−1), despitetheir relatively low crystalline order (12–14). However, becauseof their high optical density, even for ultrathin films (thicknesst < 100 nm), it remains difficult to obtain fully transparent andcolorless thin films using copolymers of this type. Moreover,obtaining high mobility typically requires undesirable process-ing techniques, such as high-temperature annealing (12) andmacroscopic alignment processes (15), which are not readilycompatible with flexible electronics. Therefore, the realizationof truly colorless semiconducting layers with high mobility forFTEs remains to be achieved.Multicomponent systems consisting of various polymer blends

have recently attracted particular attention because of the tun-ability of their material properties (16). Recent reports haveshown that blends containing a relatively small amount of semi-conducting polymer in an inert polymer matrix exhibit charge-transport characteristics that are comparable or even superior tothose of the pristine forms (17–22). However, a comprehensiveunderstanding of the underlying mechanism of this intriguingphenomenon has not yet been achieved. We note that this poly-mer-blending approach can provide new opportunities for thedevelopment of innovative polymer systems for FTEs. Here, byintroducing a diketopyrrolopyrrole (DPP)-based semiconducting

Significance

When various electronic appliances used in everyday life be-come deformable and transparent, they will provide tremen-dous versatility in the design and use of see-through, smartmobile applications, exceeding the limitations of the best de-veloped conventional silicon technologies, which are availableonly in rigid, opaque forms. However, even recently discoveredinnovative semiconducting components have failed to simul-taneously achieve such flexibility and transparency. Thus, theexisting options still comprise only hard, planar, or opaquematerials, and obtaining a “key” material for creating trulyflexible and transparent electronics has presented a formidablechallenge. We report an effective means of creating a “trulyflexible, perfectly transparent” and high-mobility semiconductingmaterial and demonstrate several high-end flexible and trans-parent applications based on a polymeric semiconductor system.

Author contributions: K.Y. and K.L. designed research; K.Y. and B.P. performed research;G.K., C.-H.K., S.P., Jehan Kim, S. Jung, S. Jeong, S.K., and M.-H.Y. contributed new re-agents/analytic tools; K.Y., B.P., G.K., C.-H.K., Jehan Kim, H.K., and Junghwan Kim ana-lyzed data; and K.Y. and K.L. wrote the paper.

The authors declare no conflict of interest.

This article is a PNAS Direct Submission.1K.Y. and B.P. contributed equally to this work.2To whom correspondence should be addressed. Email: klee@gist.ac.kr.

This article contains supporting information online at www.pnas.org/lookup/suppl/doi:10.1073/pnas.1606947113/-/DCSupplemental.

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copolymer (DPP2T) into an inert polystyrene (PS) matrix, wecreate a polymer blend system that demonstrates both highmobility and high transparency approaching 100% without de-veloping any color. We discover that a small amount of DPP2Tin an amorphous PS matrix forms a web-like, continuously con-nected nanonetwork that spreads throughout the thin filmformed during solution deposition while remaining confined in athin fibrous structure. Detailed study reveals that this networkstructure of DPP2T provides highly efficient charge pathways withsubstantially reduced structural and energetic disorder through itsextended intrachain conjugation. This approach therefore enablesus to fabricate prototype high-performance FTE devices.

Results and DiscussionThe molecular structures of DPP2T and PS are depicted in Fig. 1A.DPP2T consists of a strong electron-deficient unit, namely, DPP andan electron-rich segment of 2,5-di-2-thienylthieno[3,2-b]thiophene.In general, combining strongly fused moieties causes the resultingpolymer backbone to be planar and rigid (12). We used PS as theinert polymer (Eg ∼4 eV, SI Appendix, Fig. S1) to prepare a fullytransparent 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 ontosubstrates to obtain thin films (thickness t ∼10 nm).DPP2T exhibits strong near-IR absorption (onset at a wave-

length of λ = 920 nm; energy gap of Eg ∼1.35 eV) arising from thealternating structure of repeating units of DPP and 2,5-di-2-thie-nylthieno[3,2-b]thiophene (23) (Fig. 1B). This characteristic ab-sorption spectrum makes DPP2T far more transparent in thevisible range than other midbandgap polymers such as poly(3-hexylthiophene), but it still exhibits considerable absorption in thisspectral region (SI Appendix, Fig. S2) and has a greenish tint (23),as shown in Fig. 1C. By contrast, a blend system of DPP2T/PS thatcontains a small proportion of DPP2T (15 wt % with respect toPS) is almost perfectly transparent throughout the visible range(average transmittance of Ta ∼99% for 380–700 nm; SI Appendix,Fig. S3) and develops no color (Fig. 1C, Inset).Comparing the nanomorphology of the films using trans-

mission electron microscopy (TEM) reveals a striking feature ofthe DPP2T/PS blend system (Fig. 1D). Whereas the TEM imageof pure DPP2T is typical of this class of materials, showingstacked polymer aggregates with random orientations, DPP2T/PS shows a phase-separated heterostructure with a web-likepolymer network that consists of linked fibrils with widths of afew tens of nanometers. The continuously connected 2D nano-network structure consists solely of DPP2T, as confirmedthrough the elemental mapping of sulfur (S) via energy-disper-sive X-ray spectroscopy line-scan analysis in high-angle annulardark-field scanning-transmission electron microscopy mode (SIAppendix, Fig. S4).A more detailed structural analysis using grazing-incidence

wide-angle X-ray scattering (GIWAXS) measurements also re-veals a large difference in the nanomorphology of the polymerchains between the two films (Fig. 1E). The 2DGIWAXS patternsof DPP2T exhibit a dominant edge-on orientation, as inferred bythe existence of two patterns: one consisting of (h00) Bragg peaksin the out-of-plane direction with a lamellar spacing of 23.2 Åand the other consisting of a (010) peak in the in-plane directionwith a π–π-spacing of 3.85 Å. However, DPP2T is estimated topossess rather randomly oriented crystallites and considerablydisordered regions, as indicated by the broad (h00) peaks and thewide arc pattern around q ∼1.32 Å−1, which arises from theamorphous polymer phase (7) (Fig. 1F). By contrast, DPP2T/PSexhibits no characteristic peaks but rather a wide halo attributedto amorphous PS (and probably also to DPP2T) (SI Appendix, Fig.S5), indicating no crystalline-ordered structures of the polymerchains (detailed 1D profile studies are presented in SI Appendix,Fig. S6). The thin, fiber-like DPP2T bundles in the nanonetworkstructure of DPP2T/PS, which seem to be composed of only a fewpolymer chains, as shown in Fig. 1G, produce quite differentGIWAXS patterns (see SI Appendix, Note S1 for additionalstructural studies based on optical absorption spectra).Based on the remarkable contrast in the nanostructures of the

films, we can expect to observe fundamentally different chargetransport between the two systems. Therefore, we analyzed thecharge-transport characteristics of the two films by fabricating top-gate, bottom-contact (TGBC) field-effect transistors (FETs) usingthe DPP2T and DPP2T/PS layers (SI Appendix, Fig. S7). Fig. 2Acompares the devices’ transfer characteristics at room tempera-ture. Both devices exhibit clear p-channel FET characteristics withnegligible hysteresis and low contact resistance (SI Appendix, Fig.S8). However, whereas the saturation- and linear-regime field-effectmobilities (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 μasat =0.67 cm2 V−1·s−1 and μalin = 0.40 cm2 V−1·s−1), the DPP2T/PS FETexhibits much higher maximum saturation- and linear-regimeμ-values of 3.1 cm2 V−1·s−1 and 1.6 cm2 V−1·s−1, respectively (withμasat = 1.6 cm2 V−1·s−1 and μalin = 1.0 cm2 V−1·s−1) (SI Appendix,Figs. S9 and S10). Therefore, we can clearly observe that thecharge-transport characteristics of the DPP2T/PS devices are

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substantially enhanced compared with those of the pure DPP2T de-vices. Considering that the content of semiconducting DPP2T in theseDPP2T/PS devices is only ∼15% (corresponding to an insulating PScontent of ∼85%), which is much less than that in pure DPP2T de-vices, this observation is certainly worth further investigation.We hypothesize that the improved transport properties of the

DPP2T/PS devices are associated with the 2D nanonetworkformation in the DPP2T/PS blend films. Therefore, to more di-rectly observe the correlation between the nanomorphologiesand the charge-transport properties of DPP2T/PS films, we in-vestigated 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 foundto show quite different nanomorphologies for different blendingratios, as shown in Fig. 2C. The morphological phase forma-tion of different concentration of DPP2T in a PS matrix can beroughly divided into three regimes: (i) at low DPP2T concen-trations, the DPP2T forms long, thin fibrils, which becomethicker and more highly interconnected as the DPP2T concen-tration increases; (ii) at a moderate DPP2T concentration, theDPP2T forms completely percolated 2D nanonetworks; and(iii) at high DPP2T concentrations, the fibrillar bundles are ag-gregated and stacked, and the morphology thus approaches that ofthe pure DPP2T phase. These morphological changes are expectedto affect the charge-transport properties of the blend films. Fig. 2Bshows the dependence of the saturation-regime μ on the DPP2T/PSblending ratio (SI Appendix, Table S1, Fig. S11, and Note S2). Withan increasing DPP2T content, the DPP2T/PS devices exhibit rap-idly increasing μ beginning at 3 wt % DPP2T (with respect to thetotal solid) and peak μ at 15 wt %. In this range, the best chargepathways seem to form as the fibrous structures become perco-lated. Further increasing the DPP2T content above 15 wt % re-sults in a gradual decrease in μ as the characteristics of DPP2T/PSapproach those of pure DPP2T. Therefore, we can confirm thatthe trend in μ observed for the devices directly parallels themorphological 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 var-ious low drain-source voltages (VDS), which were assumed to yield auniform electric field strength (F) along the channel (SI Appendix,Figs. S12 and S13). Arrhenius plots of the T-dependent μ-values arepresented in Fig. 3A, from which we extracted the activation en-ergies (EA) using the Arrhenius relation, μ ∝ exp(−EA/kT),where k is the Boltzmann constant. Interestingly, the plots showtwo different activation regimes, with a transition occurring atT ∼190 K (24, 25). In the low-T regime, DPP2T shows an EA of∼19 meV, which is fairly low compared with the values reportedfor other high-mobility polymers (26, 27). By contrast, a surpris-ingly low EA of ∼5 meV is observed for DPP2T/PS (28). Similarly,a lower EA value for DPP2T/PS (compared with that for pureDPP2T) is observed even in the high-T regime. The transitiontemperature observed in Fig. 3A is expected to correspond to thethermal energy necessary to overcome the local hole barrier be-tween aggregated and amorphous DPP2T regions, which causesthe charge transport to be spatially confined within these orderedregions at low T (SI Appendix, Fig. S14). This explains why thesame transition occurs regardless of the level of PS blending. Froma structural perspective, the large, randomly oriented aggregatessurrounded by amorphous regions that are present in the pureDPP2T film cause the paracrystallinity-dominated π–π transportto act as the rate-limiting factor (7), whereas efficient chain-backbone transport becomes the most relevant contribution in theDPP2T/PS system (29, 30) (Fig. 1 F and G and SI Appendix, Fig.S15). Consequently, the carrier drift in the DPP2T/PS film exhibitsexceptionally low thermal coefficients, reflecting single-chainconduction with occasional hopping at interchain connections. Inaddition, the decrease in the EA of DPP2T as F increases in thehigh-T regime implies the occurrence of F-assisted tunneling andhopping 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 F1/2

(Fig. 3C), the linear increase in μ with F1/2 indicates that DPP2Texhibits Poole–Frenkel-like behavior, which is characteristic ofdisorder-limited charge transport (32, 33). By contrast, DPP2T/PS

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Fig. 2. Correlation between the transport characteristics and morphological properties of DPP2T/PS blend films. (A) Representative room-temperaturetransfer characteristics of pure DPP2T and DPP2T/PS FETs under a VDS 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 observedwith the variation of the DPP2T concentration. The vertical lines (whiskers) indicate the range from the 10th to the 90th percentile. The minimum andmaximum values are indicated by asterisks. (Inset) Schematic illustration of the device structure. (C) TEM images of DPP2T/PS films of various concentrationratios. These images show the evolutionary phase change of the DPP2T in the PS matrix as the concentration of DPP2T increases. (Scale bar, 200 nm.)

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shows an almost negligible F dependence of μ, which indicates theformation of nearly undisturbed charge pathways (see SI Appen-dix, Note S3 and Figs. S16 and S17 for further theoretical analysesof the charge transport in DPP2T/PS).It is interesting that the difference in the morphological state

modulated by PS blending can result in such large variations inthe charge-transport characteristics of DPP2T. DPP2T has arigid backbone structure because of its strongly fused moieties,which are responsible for its strong intrachain and interchaincharge transfer and relatively high intermolecular stacking.However, its crystallization is disturbed because of the complexchain entanglement that unavoidably occurs during solutionprocessing because of the limited space available for the move-ment of the crowded DPP2T chains and fibrillar bundles (34) (SIAppendix, Fig. S18). Therefore, a pure DPP2T film is composedof not only well-ordered interchain stacks but also disorderedamorphous 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 con-nected web-like structure, in which large-scale π–π and lamellarstacking 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 entangle-ment of the rigid DPP2T chains by providing a more flexiblesurrounding environment (SI Appendix, Note S4 and Fig. S19).Therefore, the unique structure of DPP2T/PS is assumed to consistof a series of DPP2T units extending along the long axis of eachfibril, in which the structural discontinuity is significantly reducedas the interchain aggregation decreases (see the schematic illus-tration in Fig. 1G). Therefore, DPP2T/PS is expected to supportefficient intramolecular charge transport through continuouslyconnected “clean” pathways without severe structural or ener-getic disorder because the extended DPP2T units promote agreater extent of intramolecular charge delocalization (26, 36),and thus, its near-intrinsic molecular performance manifests atthe length scale of the device channels (see SI Appendix, NotesS5 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 appli-cations. 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

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fabricated from polymeric materials via solution processing.Fig. 4A shows a photograph and an illustration of the devicestructure of our large-area FT-FET array (10 cm × 10 cm, 1,650FETs). To fabricate the transparent source/drain and top-gateelectrodes, poly(3,4-ethylenedioxythiophene):poly(styrenesulfonate)(PEDOT:PSS) was inkjet-printed onto the PEN substrate andthe gate insulating layer (SI Appendix, Fig. S25). Fig. 4B showsthe optical transmittance spectra of the FT-FETs measuredunder different layering conditions (SI Appendix, Table S2).Because all of the layers on the PEN substrate are thin andhighly transparent, the transmittance of the FET device is quitehigh throughout the visible range (Ta ∼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 presentedin Fig. 4C. Although a slight hysteresis and wavy curves are ob-served in the output characteristics (SI Appendix, Fig. S27), be-cause of the relatively low conductivity of the PEDOT:PSSelectrodes (σ < 700 S cm−1), overall, FT-FETs were uniformlyfabricated over the entire device area and showed stable andclear field-effect characteristics, with a high saturation-regime μof 0.80 cm2 V−1·s−1 (μa = 0.64 cm2 V−1·s−1) and a linear-regime μof 0.23 cm2 V−1·s−1 (μa = 0.15 cm2 V−1·s−1). Our FT-FETs arealso 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 perfor-mance degradation (Fig. 4E and SI Appendix, Fig. S28).To take full advantage of the advantageous characteristics of

DPP2T/PS and to demonstrate a potential electronic appli-cation 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 ourFT-FET-PLED devices, which were prepared by fabricatingtransparent 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 thesemitransparent anode of the PLED, and the generated light isemitted through the transparent FET layers on top (SI Appendix,Fig. S30). The molybdenum oxide (MoOx)/PEDOT:PSS hole-injecting layer offers a supporting surface for sequential solutionprocessing for the fabrication of the top FET. The top FEToperates by supplying a gate voltage (VGS)-modulated drain-source current (IDS) to the PLED, which has an aryl-substitutedpoly(p-phenylene vinylene) derivative (PDY-132) emissive layer,resulting in the stable modulation of yellow light with a maximumluminance (L) of ∼252 cd m−2 (Fig. 5C). The corresponding elec-troluminescence (EL) spectrum does not show any significant dif-ferences from that of the PLED without an FET because of the hightransparency of the FET components (Fig. 5D). Additionally, FT-FET-PLED devices with different emission colors were preparedusing various emissive layers consisting of the white-light-emittingpolymer SPW-111, poly[2-methoxy-5-(2-ethylhexyloxy)-1,4-phe-nylenevinylene] (MEH-PPV) and poly(9,9-di-n-octylfluorene-alt-benzothiadiazole) (F8BT), and the resulting devices successfullydemonstrate the modulation of the emitted light by the top FET evenduring bending (Fig. 5E and SI Appendix, Figs. S31 and S32 andMovie S1). Thus, our system represents a major step toward achievingtransparent and deformable all-polymer active-matrix displays.In summary, our results describe a 2D copolymer network with

both ultratransparency and high μ prepared using the DPP2T/PSblend system. The dimensionally confined 2D nanonetworkstructure of the DPP2T in the PS enables much more efficientcharge transport through the extended intrachain conjugation ofdisorder-reduced clean pathways. Simultaneously, the DPP2T/PS blend system also offers exceptional transparency that ap-proaches 100% because of its low visible-light absorption andthin net-like structure. We demonstrated the fabrication of FTEdevices using our completely colorless, high-μ DPP2T/PS system.

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Our results provide scientific insights into the physics of chargetransport in conjugated polymer systems and a pathway towardunprecedented technological applications in organic electronics.

Materials and MethodsDPP2T/PS Blend Solution. DPP2T (Mr ∼73 kg mol−1, 1-Material) and PS (Mr

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

TGBC FET Fabrication. Glass slides (Eagle XGTM, Corning) were cleaned viasequential ultrasonication in water, acetone, and isopropyl alcohol. Ther-mally evaporated Au (15 nm) source/drain electrodes were patterned usingshadow masks. The blend solution was spin cast onto the substrate in aninert nitrogen atmosphere, and the films were subsequently annealed at 80 °Cfor 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 semiconductinglayer, resulting in a thickness of ∼550 nm, and the films were subsequentlyannealed at 80 °C for 30 min. The measured capacitance of the CYTOP layerwas ∼3.5 nF cm−2. The devices were completed with the thermal deposition of50 nm of Al through a shadow mask to form the top-gate electrode.

More experimental details are shown in SI Appendix.

ACKNOWLEDGMENTS.We thank the Research Institute of Solar and Sustain-able Energies at the Gwangju Institute of Science and Technology (GIST) ofKorea for supporting the project by the GIST Research Institute Project 2016.K.L. acknowledges support from Basic Science Research Program through theNational Research Foundation of Korea (NRF) funded by the Ministry of Science,ICT and Future Planning (MSIP) (Grants NRF-2014R1A2A1A09006137 and NRF-2015K1A3A1A16002247). This work was also supported by the R&D pro-gram of MSIP/Commercializations Promotion Agency for R&D Outcomes(Grant 2015K000199).

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