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Cite This: Chem. Mater. 2018, 30, 6353−6360

Stretchable Polymer Gate Dielectric with Segmented Elastomeric Network for Organic Soft Electronics Boseok Kang,†,⊥ Eunjoo Song,†,⊥ Seon Baek Lee,† Byeong-Gyu Chae,‡ Hyungju Ahn,§ and Kilwon Cho*,† †

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Department of Chemical Engineering and Center for Advanced Soft Electronics, Pohang University of Science and Technology, Pohang 37673, Korea ‡ Department of Material Science and Engineering, Pohang University of Science and Technology, Pohang 37673, Korea § Pohang Accelerator Laboratory, Pohang University of Science and Technology, Pohang 37673, Korea S Supporting Information *

ABSTRACT: Stretchable electronics has emerged as a key technology for human-friendly soft electronic applications. However, the difficulty of preparing stretchable gate dielectrics is a major impediment to the realization of stretchable electronic devices. Here, we present a stretchable polymer gate dielectric for use in stretchable organic thin-film transistors (OTFTs). Our strategy is to form a segmented elastomeric network inside an end-functionalized reactive liquid rubber with a high dielectric constant by using an organosilane cross-linking agent. The developed dielectric layer consists of hard segments of self-organized nanocrystals and soft rubber segments. The nanocrystals distributed in the rubber matrix resist the growth of electrical treeing inside the film and thus significantly increase the dielectric strength and reduce the leakage current density (∼10−8 A cm−2 at 2.0 MV cm−1). The superior insulating properties are maintained well during tensile stretching with moderate strains ε ≤ 30%. Furthermore, fully stretchable OTFTs based on the developed gate dielectric were demonstrated to stably operate without electrical breakdown under stretching of ε ≤ 34%. These results prove that our method is a promising approach to the development of stretchable polymer gate dielectrics.

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ε0k/d, where ε0 = 8.854 × 10−12 F m−1 is the vacuum permittivity and k is the dielectric constant of the insulating material. In this case, Ci may become quite low for the operation of TFT devices. In addition, commercially available silicone rubbers such as polydimethylsiloxane (PDMS) and Ecoflex have typically small k values (e.g., 2 ∼ 3),12−14 which also contributes to reduce Ci. To fabricate a stretchable GD layer with superior insulating properties and sufficient Ci to operate an OTFT device, we here used an end-functionalized liquid reactive rubber, dicarboxy-terminated poly(acrylonitrile-co-butadiene) (CTBN), and an organosilane cross-linking agent, 1,6bis(trichlorosilyl)hexane (C 6 ) (Scheme 1a). Poly(acrylonitrile-co-butadiene) rubber contains an abundance of polar acrylonitrile groups and thus has a dielectric constant (3 ≤ k ≤ 10) higher than that of most of nonpolar silicon rubbers; k has been known to vary with the content of nitrile groups.15,16 This relatively high k means that it is not necessary to significantly lower the dielectric thickness to secure Ci, and

INTRODUCTION Interest in flexible and stretchable organic electronics has rapidly grown in recent years, and there have been tremendous efforts to develop stretchable conductors and stretchable organic semiconductors with the aim of achieving stretchable organic thin-film transistors (OTFTs).1−6 Several research groups have tried to fabricate stretchable gate dielectric (GD) layers for use in OTFTs, but there seems to still be a lack of successful studies on such GDs.7 The development of a stretchable GD is essential to the full realization of stretchable OTFTs, and the ideal stretchable dielectric film requires following properties: (i) sufficient elasticity, (ii) superior insulating properties, (iii) capacitance high enough to drive TFTs, and (iv) resistance to organic solvents for postdeposition processes. Common elastomers such as silicon rubber and polyurethane have been tested for use as stretchable GDs,8,9 but the elastomers have large free volume that can permit the easy migration of charges into their inside when subjected to a moderate electric field, resulting in a high leakage current density J.10,11 If the thickness d of the dielectrics is increased to suppress J, the unit capacitance Ci, which is defined as the dielectric capacitance per unit area, decreases as given by Ci = © 2018 American Chemical Society

Received: June 6, 2018 Revised: August 23, 2018 Published: September 7, 2018 6353

DOI: 10.1021/acs.chemmater.8b02388 Chem. Mater. 2018, 30, 6353−6360

Article

Chemistry of Materials

of 12 h is sufficient for the mixed compounds to form a crosslinked elastomeric network. Defect-free and smooth surfaces are necessary conditions for the use of GDs in OTFTs.24 To optimize the formation of pinhole-free and smooth dielectric films with respect to the C6 content ([C6]), CTBN was cross-linked in the presence of various [C6] from 0 to 10 wt %, and then the surface morphologies of the films were investigated (Figure 2a). The surface of the un-cross-linked CTBN film had an uneven color distribution in the optical microscopy image, which indicates that the film was rough on the macroscale. As [C6] increased, the surfaces of the polymer films gradually became smooth, and at [C6] ≥ 7 wt %, uniform films were obtained after crosslinking. All of the cross-linked CTBN films had flat and smooth surfaces on the microscale (the root mean squared roughnesses RRMS ≤ 0.8 nm) according to the atomic force microscopy (AFM) height images (Figure 2a and Figure S1) and were free of pinholes. The film thicknesses of the cross-linked CTBN films were in the range 635 ≤ d ≤ 570 nm for 2 ≤ [C6] ≤ 10 wt %, which is 15% less than that of the un-cross-linked film (d = 747 nm; see Figure S2). The reduction in d suggests that the free volume in the films is largely reduced and that they are densified as the degree of cross-linking increases.25 The effects of varying the C6 content on the insulating properties of the dielectric films were systematically investigated by fabricating metal−insulator−metal (MIM) devices on rigid substrates and obtaining their current density vs electric field (J−E) curves (Figures 2b and c). The pure CTBN film was broken down electrically even at a relatively small electric field, E = 0.08 MV cm−1. By contrast, the chemically cross-linked CTBN rubber showed noticeable increase in the dielectric strength and reduction in the leakage current of the dielectric films. As [C6] increased from 2 to 10 wt %, J decreased, and the J−E curves became stabilized features. The average J at E = 0.5 MV cm−1 in Figure 2b is significantly reduced to 10−9 < J < 10−8 A cm−2 after cross-linking with [C6] ≥ 3 wt % (Figure 2c, filled circles). In particular, the films cross-linked with [C6] ≥ 7 wt % well maintained a low J < 3.0 × 10−8 A cm−2 without electrical breakdown as E was increased to 2.5 MV cm−1 (Figure 2b). These insulating properties of the cross-linked CTBN films are significantly superior to those of PDMS films, a reference elastomer material (Figure 2b, black solid line). The unit capacitance of the cross-linked CTBN was 4.75 ≤ Ci ≤ 5.8 nF cm−2 for 2 ≤ [C6] ≤ 10 wt % (Figure 2c, open circles; see Table S1), which is higher than that of the uncross-linked CTBN film (3.8 nF cm−2) and satisfactory for device operation. The initial increase in the capacitance with increasing [C6] can be attributed to the reduction in the film thickness (Figure S2) but could also be due to an increase in the k value of the dielectric material. In practice, estimated k values of the dielectric films increased slightly from 3.2 to 3.7 as [C6] increases from 0 to 10 wt % (Table S1); this improvement would be due to the fact that the density of polar acrylonitrile groups in cross-linked and densified CTBN films is higher than that in un-cross-linked films.26 Stretchable MIM devices employing the cross-linked CTBN film ([C6] = 7 wt %) as a dielectric were fabricated on a PDMS substrate with embedded carbon nanotubes (CNTs) that is a bottom electrode, and their J−E characteristics were quantified under various stretching conditions (0 ≤ ε ≤ 50%). The elongation at break of the CTBN films on a PDMS substrate was approximately ε = ∼60%. The morphology and electrical

Scheme 1. (a) Chemical Structures of CTBN and C6 and a Scheme to Form Segmented Elastomeric Network Film by a Crosslinking Reactiona and (b) Photograph of a CrossLinked CTBN Rubber Sampleb

The film contains C6-assembled crystals in elastomeric CTBN matrix. Green lines: CTBN polymer chains. b[C6] = 10 wt %.

a

therefore the required insulation would be provided. CTBN is chemically reactive at elevated temperatures, and cross-linking of CTBN can be achieved by a condensation reaction with C6 which contains two silanetriols (−Si(OH)3) at the chain ends after exposure to moisture.17,18 The six reactive hydroxyl substituents of C6 are expected to self-polymerize to form hard segments consisting of C6 aggregates in the soft rubber matrix. The self-organized C6 aggregates may act as nanofillers in the dielectric film and would enhance the dielectric strength by inhibiting electrical treeing growth.19,20 As a result, the crosslinking of the CTBN rubber with C6 was found to form an elastomeric network in which hard segments consisting of selforganized C6 crystals are present (Scheme 1b). This rubbery GD had superior insulating properties with a low J ∼ 10−8 A cm−2 and a moderate Ci ∼ 5 nF cm−2 that was sufficient for the operation of OTFTs.

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RESULTS AND DISCUSSION To verify the cross-linking reaction in the CTBN blend films containing C6 (10 wt %) at an elevated temperature of 90 °C, the changes in the absorbances of chemical bonds related to the major functional groups (silanol, siloxane, esters, and carboxyl) during the curing process (0−12 h) were monitored by using Fourier transform infrared spectroscopy (FTIR) (Figures 1a−c). The intensities of the peaks assigned to Si− OH stretching at 968 cm−1 and CO stretching of carboxyl group at 1714 cm−1 decreased as the curing time increased to 12 h.17,21,22 Absorption bands appeared at 1056 and 1731 cm−1 and increased in intensity during the cross-linking process; these bands are attributed to Si−O−Si and CO in the ester group, respectively.17,22,23 These results indicate that while the silanol and the carboxyl groups in C6 and CTBN are consumed by the condensation reaction, ester groups are newly formed by an esterification reaction, and some siloxane bonds are also formed by the self-condensation of Si−OH in C6 (Figure 1d). The depicted changes in absorbance almost completely stopped after 12 h for all 5 bonds; thus, a cure time 6354


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Figure 1. (a and b) Room-temperature FTIR spectra of the CTBN:C6 films ([C6] = 10 wt %) as a function of cure times (the reaction temperature of 90 °C) and (c) absorbance change of the five major bonds in the films. (d) Two chemical reactions feasible in the CTBN:C6 cross-linking system: esterification reaction between CTBN and C6 molecules and self-condensation reaction of silanol groups in C6 molecules.

are involved in a self-condensation reaction that produces crystalline aggregates in the rubber network. The shape and distribution of the C6 crystalline aggregates in cross-linked CTBN with [C6] = 7 wt % were visualized by using dark-field scanning transmission electron microscopy (DF-STEM) (Figure 3b). The STEM-energy dispersive X-ray spectroscopy (EDS) line profile along the dotted line in Figure 3b contains an intense signal of Si atoms, indicating that the light regions in the STEM image are C6 aggregates (inset, Figure 3b). The C6 nanocrystallites were spherical with a diameter of approximately 8 nm. To determine the size variation with [C6], small-angle X-ray scattering (SAXS) analysis was employed (Figure 3c). The scattering profiles were analyzed by using the unified fitting model in the Irena package of the software IGOR Pro to determine the radius of gyration Rg of the C6 aggregates in the cross-linked CTBN. As [C6] in the cross-linked CTBN film was increased from 5 to 7 to 10 wt %, Rg increased from 7.3 to 7.7 to 9.0 nm. This result is probably due to the increase in the number of C6 molecules available to participate in the self-organization reaction. We suggest that these spherical crystalline aggregates might have a pseudodiamond lattice because the silanetriols at the ends of the C6 chains react to form a tetrahedral structure with a center consisting of the siloxane or the remaining silanol groups

stability of the embedded CNT electrodes under stretching are provided in detail in Figure S3. The CTBN dielectric film on the substrate exhibited superior insulating stability during stretching with strains up to ε = 30%. The MIM devices had a low average J ∼ 10−8 A cm−2 at E = 2.4 MV cm−1 in the strain range 0 ≤ ε ≤ 30% (Figure 2d, filled circles). They maintained this low J even when repeatedly stretched and released over 60 cycles at ε = 10% and were not completely electrically broken down even after 100 cycles of stretching, exhibiting J ∼ 10−6 A cm−2. The nanostructure of the cross-linked CTBN-based GD films was investigated by using X-ray scattering and electron microscopy experiments. The two-dimensional grazing incidence X-ray diffraction (2D GIXD) pattern for a cross-linked CTBN film with [C6] = 10 wt % contained a crystalline peak at qz = 0.56 Å−1 in the out-of-plane direction (Figure 3a). This peak corresponds to a lattice spacing of d = 11.2 Å, which is approximately the theoretical length of the alkyl chain in a C6 molecule (12.2 Å, see Figure S4). The un-cross-linked CTBN thin film did not produce any crystalline peak near qz = 0.56 Å−1; this peak appeared for [C6] above 5 wt %, and its intensity increased as [C6] increased from 5 to 100 wt % (Figures S4). These results strongly support the fact that some of the C6 molecules containing reactive silanetriols at both chain ends 6355


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Figure 2. (a) Surface morphologies of the CTBN:C6 films with various [C6] observed by optical microscopy (left column) and AFM (right column). The numbers on the right upper corners of each AFM images indicate RRMS of the films. (b) J−E curves of the CTBN:C6 films with various [C6]. Black solid line: result from a PDMS film for the comparison experiment. (c) Change of average current density J at 0.5 MV cm−1 in (b) and unit capacitance of the dielectric films as a function of [C6]. (d) Change of J of the CTBN:C6 ([C6] = 7 wt %) films at 2.4 MV cm−1 during stretching (filled circles) and after repeated stretching cycles with ε = 10% (open circles).

determined to be 3.7 MPa by the AFM indentation measurement (Figure S7). Based on the above experimental observations, we propose an explanation of the superior insulating stability of the crosslinked CTBN film under stretching conditions. Electrical treeing is a representative mechanism for explaining the electrical prebreakdown of an polymeric insulator when it is subjected to a high electrical field stress.30 The initiation and growth of electrical trees are accompanied by partial discharges in high stress regions such as gaseous voids and defects.19 Lowdensity regions in polymer insulators are vulnerable to the development of electrical treeing.31 Pure CTBN elastomer has innately sufficient free volume and loosely packed chain regions, so current filaments can grow easily in the dielectric film through the electrical treeing mechanism (Figure 3e, left). However, in the case of the cross-linked CTBN film, the polymer matrix is largely densified, and thus its resistance to electrical treeing growth would be improved.31 In addition, the hard segments of C6 crystallites that form uniformly in the

(Figure 3d). The dimensions of the empty space in the C6 crystal frame are much larger than the cross-sectional diameter of a CTBN chain (see the simulation results in Figure S5), and thus the self-organized C6 crystals are thought to be filled with CTBN polymer chains and could form hard segments with density higher than that of the surrounding CTBN matrix. The effects of cross-linking on the glass transition temperature (Tg) of CTBN were further examined by performing differential scanning calorimetry (DSC) on CTBN cross-linked with [C6] = 0, 2, and 7 wt % (Figure S6). Un-cross-linked neat CTBN ([C6] = 0 wt %) had a Tg of −68.9 °C.27 After CTBN had been cross-linked with C6, the Tg values rose slightly to −67.6 and −67.4 °C for [C6] = 2 and 7 wt %, respectively, which is attributed to chain-end cross-linking and physical entanglements of polymer chains. The Tg value of cross-linked CTBN with [C6] = 7 wt % was still below room temperature, indicating that cross-linked CTBN is still rubbery at room temperature. 28,29 In practice, its elastic modulus was 6356


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Figure 3. (a) 2D GIXD pattern of the cross-linked CTBN ([C6] = 10 wt %) films. (b) DF-STEM image for the cross-linked CTBN ([C6] = 7 wt %) film. Bottom inset: STEM-EDS line profile along the white dotted line. (c) SAXS profiles obtained in transmission mode for bulk cross-linked rubber samples ([C6] = 5, 7, and 10 wt %). Table inset: radius of gyration of the C6 aggregates, calculated using the unified fitting model. (d) Pseudodiamond lattice consisting of tetrahedral structures of C6-self-assembled crystallites. (e) Schematic drawing for the proposed mechanism for the electrical breakdown of the homo-CTBN film (left) and the superior stretching stability of the cross-linked CTBN rubber GD (right).

were stretched with strains up to 34%, they still exhibited common transfer characteristics, with IG ∼ 10−9 A (Figure 4c, dotted lines). The normalized average mobility μ/μ0 is plotted as a function of the strain in Figure 4, where μ0 and μ are the field-effect transistor (FET) mobilities before and after stretching, respectively (Table S2); we think the decrease in the normalized average mobility at early stage of stretching would come from the imperfect stretching stability of the other components of the OTFTs such as the electrodes and the semiconducting layer (Figure 4d, open circles).34 In particular, the FET mobility of the previously reported poly(3-hexyl thiophene) nanowires (P3HT NWs)/PDMS (1:10 w/w) blend film that had been used as a stretchable active layer here had been found to decrease by nearly 1 order of magnitude when stretched to 50%.5,35 Even though the μ/μ0 values of the devices decreased as the strain increased, typical transfer characteristics were observed with a low leakage current and no dielectric breakdown for strains up to ε = 34% (Figure 4d, filled circles), which demonstrates the promise of our developed dielectric as a stretchable gate insulator for use in future electronic devices; we provide the summarized gate leakage current levels of different stretchable dielectric materials reported in the literature (Table S3).

rubber matrix may act as nanofillers and would hinder the propagation of filamentary electrical treeing toward the bottom electrode (Figure 3e, right).19,20 Therefore, the cross-linked rubber GD should have enhanced electrical breakdown strength and maintain low leakage current characteristics during stretching. The huge increase in J at ε = 40% despite the presence of the C6 aggregates in the cross-linked network film occurred because the thickness of the GD film already reached a critical value after stretching, such that a filamentary conducting path might be formed between the top and bottom electrodes. Dielectric films cross-linked with [C6] ≥ 7 wt % have excellent chemical resistance to a wide range of organic solvents such as chloroform and toluene and are therefore quite compatible with continuous solution processes as well as vacuum deposition process.32 A highly ordered semiconducting film of a donor−acceptor-type polymer, poly{2,2′-[(2,5bis(2-octyldodecyl)-3,6-dioxo-2,3,5,6-tetrahydropyrrolo[3,4-c]pyrrole-1,4-diyl)]dithiophene-5,5′-diyl-alt-thieno[3,2-b]thiophen-2,5-diyl} (PDBT-co-TT),33 was obtained by spincoating it from a chloroform solution onto the cross-linked dielectric ([C6] = 7 wt %) (Figures S8a−f). Moreover, the usefulness of our developed material as a gate dielectric for OTFT devices was first verified by fabricating OTFT devices on rigid substrates. The solution-deposited PDBT-co-TT thin layer and vacuum-evaporated pentacene layer were used as an active layer. The films showed high crystallinity on the CTBN GDs, and the devices exhibited ideal transfer and output characteristics, which prove the usefulness and versatility of the developed CTBN GD. Next, all-stretchable OTFTs were demonstrated by using the stretchable cross-linked CTBN GD ([C6] = 7 wt %) (Figure 4a). The other components of the stretchable OTFTs are described in detail in the Experimental Section. Stretchable OTFTs employing the cross-linked CTBN GD were found to exhibit moderate OTFT transfer characteristics without a large gate leakage current, IG (Figures 4b and c). When the devices

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CONCLUSION We presented a cross-linking strategy and nanostructural control to develop an elastomeric GD material with insulating properties that tolerate stretching. The cross-linking of high-k liquid reactive CTBN rubber by using the organosilane compound C6 was found to result in the formation of a segmented elastomeric network with two parts: an acrylonitrile-butadiene rubber network and hard segments of C6 aggregates. The cross-linked rubber films were found to exhibit superior insulating properties under stretching conditions (0 ≤ ε ≤ 30%) with a low J of approximately 10−8 A cm−2 at E = 2.4 MV cm−1. This outcome is rationalized by the increased 6357


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Device Fabrication and Characterizations. To fabricate the rigid MIM devices, heavily doped n-type Si wafers with native oxide and Au patterns (a rate of 3 Å s−1, 50 nm-thick) were used as bottom and top electrodes, respectively. For comparison with CTBN dielectrics, the reference PDMS dielectric film (Dow Corning Sylgard 184, A:B = 1:10) was prepared by spin-coating the PDMS solution diluted with toluene (1000 rpm, 1 min) and curing at 60 °C overnight. The thickness of the PDMS film was about 1 μm. The PDBT-co-TT and pentacene OTFTs on rigid substrates were also built on the 300 nm-thick SiO2/Si wafers on which the gate electrodes were prepatterned with Ti/Au (10/40 nm-thick), and the cross-linked CTBN dielectric films with [C6] = 7 wt % were formed. The PDBTco-TT films were spin-coated from a solution (8 mg mL−1 in chloroform). A 60 nm-thick pentacene layer was thermally evaporated at a rate of 0.2 Å s−1 under vacuum (∼10−7 Torr). The source and drain electrodes (50 nm-thick) were patterned by thermal evaporation with channel length L = 100 μm and channel width W = 1000 μm. Stretchable MIM and OTFT devices were fabricated on PDMS sheets on the surface of which a CNT network had been embedded as a gate electrode. The CNT solution was spray-coated on the octadecyltrichlorosilane-treated SiO2/Si wafers, and PDMS (Dow Corning Sylgard 184, A:B = 1:10) was poured onto it, then the assembly was cured at 60 °C overnight. The cured films were detached to yield CNT-embedded PDMS substrates with a thickness about 1 mm. Then, cross-linked CTBN dielectric films were fabricated on the PDMS substrates by the same method above. A small droplet of liquid metal eutectic gallium−indium (EGaIn) was employed as the top electrode on the CTBN samples. The area of EGain electrodes was calibrated by using optical microscopy images. To fabricate the allstretchable OTFTs, a previously reported poly(3-hexylthiophene) nanowires (P3HT NWs)/PDMS (1:10 w/w) blend material was used as an organic semiconducting layer.5 Au/CNT hybrid contacts were used as substitutes for stretchable electrodes which were prepared by thermal evaporation of Au (a rate of 3 Å s−1, 50 nm-thick) and then spraying a CNT solution on samples covered with shadow masks (L = 50 μm, W = 1000 μm) (see Figure S9). The CNT solution was prepared using multiwall CNT powders (70 mg) dissolved in Nmethyl-2-pyrrolidone solvent (30 mL) and then dispersed by ultrasonication for 30 min. During the Au deposition and subsequent spraying of the CNT solution, the shadow masks were fixed by an attractive magnetic force between the mask (Invar 36) and the below magnet plate. For measuring characteristics of stretchable MIM and OTFT devices, homemade tensile loader was employed. The capacitance of the dielectric films was measured using an Agilent E4980A precision LCR meter at 1 kHz and 1.0 Vrms. The electrical properties of the devices were measured using Keithley 4200SCS units under ambient conditions. FET mobilities of the OTFTs were calculated in the saturation region by using the following equation,

Figure 4. (a) Scheme and photograph showing the structure of a fully stretchable OTFT achieved using the cross-linked CTBN GD with [C6] = 7 wt %. (b) The output characteristic of the fabricated OTFT device at ε = 0%. (c) Transfer characteristics at the saturation region (VDS = −60 V) of the devices with increasing strain from ε = 0 to 34%. Dotted lines: gate current during operation at each strain. (d) Change of the normalized FET mobility (μ/μ0, open circles) and gate leakage current at VGS = −60 V (filled circles) as a function of strain.

density of the film structure and the C6 nanocrystals that develop in the rubber matrix, which prevent electrical treeing growth in the dielectric film. We used the stretchable GD to fabricate fully stretchable OTFTs, which were found to operate stably without dielectric breakdown for strains up to ε = 34%. We believe that our results would provide a guideline for the development of stretchable GD materials to facilitate the realization of future stretchable electronics.

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ID =

W μCi(VGS − Vth)2 2L

where ID is drain current, VGS is gate voltage, and Vth is threshold voltage. For the stretched devices, revised W and L ratio values were used in the equation as considering the Poisson’s ratio (υPDMS = 0.5) of PDMS.

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EXPERIMENTAL SECTION

Preparation of Dielectric Films and Characterizations. CTBN (acrylonitrile 8−12 wt %, M̅ n ∼ 3800, Sigma-Aldrich) was dissolved in toluene (80 mg mL−1), and then various contents of C6 (Sigma-Aldrich) were added to the solution. The solutions were stirred at 70 °C overnight in N2 atmosphere in a glovebox. The solutions were moved into air, and the CTBN-C6 thin films were formed by spin-coating the solution at 1500 rpm for 30 s using polytetrafluoroethylene filters with 0.45 μm pore size. The films were cured at 90 °C overnight. The FTIR spectrum was recorded in transmittance mode (Bruker vertex 70v). Surface morphologies of the films were characterized using a Digital Instruments Multimode AFM in tapping mode. The structural and chemical information were obtained using dark-field imaging and EDS in STEM mode (JEOL2100F).

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.8b02388. Details about characterization of cross-linked CTBN films (AFM images, cross-sectional SEM images, 2D GIXD patterns, DFT calculations, DSC curves, and AFM indentation test results), electrical characterization of OTFTs that were prepared with CTBN:C6 GDs, and electro-stretching stability results of embedded CNT and Au/CNT electrodes on PDMS substrates (PDF) 6358


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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Kilwon Cho: 0000-0003-0321-3629 Author Contributions ⊥

B.K. and E.S. contributed equally to this work.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by a grant from the Center for Advanced Soft-Electronics funded by the Ministry of Science and ICT as Global Frontier Project (Code 2011-0031628). The authors thank J. Kim, G. Y. Bae, and H. Lee for discussions and their help with AFM and SEM imaging. Portions of this research were carried out at 3C, 5A, and 9A beamlines of the Pohang Accelerator Laboratory, Korea.

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