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Stretchable Conjugated Rod−Coil Poly(3-hexylthiophene)-blockpoly(butyl acrylate) Thin Films for Field Effect Transistor Applications Jau-Tzeng Wang,† Shoichi Takshima,‡ Hung-Chin Wu,† Chien-Chung Shih,† Takuya Isono,§ Toyoji Kakuchi,§ Toshifumi Satoh,*,§ and Wen-Chang Chen*,† †

Department of Chemical Engineering, National Taiwan University, Taipei 10617, Taiwan Graduate School of Chemical Sciences and Engineering and §Faculty of Engineering, Hokkaido University, Sapporo 060-8628, Japan

‡

S Supporting Information *

ABSTRACT: We report the synthesis, morphology, and properties of poly(3-hexylthiophene)-block-poly(butyl acrylate) (P3HT-b-PBA) for stretchable electronics applications, which are consisted of semiconducting P3HT and low glass transistion temperature (Tg) PBA blocks. The P3HT-bPBA thin films self-assembled into fibrillar-like nanostructures and maintained the edge-on oreientation even at a low P3HT composition, based on the results from atomic force microscopy (AFM) and grazing incidence X-ray diffraction (GIXD). By varying the P3HT/PBA ratio, the tensile modulus decreased as the block length of PBA increased, from 0.93 GPa for P3HT to 0.19 GPa for P3HT-b-PBA12k. The field effect transistor (FET) using P3HTb-PBA as the active layer exhibited a high p-type mobility over 10−2 cm2 V−1 s−1, indicating its good charge transporting ability. Furthermore, the P3HT-b-PBA6k based FET under 100% strain had a high mobility of 2.5 × 10−2 cm2 V−1 s−1 with an on/off ratio of 7.2 × 106, and it maintained over 10−2 cm2 V−1 s−1 for 1000 cycles, suggesting the promising stability and reproducbility. The result demonstrated that the newly designed conjugated rod−coil block copolymers could have potential applications in stretchable electronic devices.

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for stretchable electrodes or circuits.16,26,27 For example, Jeong and co-workers23 discovered that the blends of polythiophene (P3HT) nanofibers and polystyrene-block-poly(ethylene-cobutylene)-block-polystyrene (SEBS) rubber exhibited a field effect transistor (FET) mobility larger than 2 × 10−3 cm2 V−1 s−1 with the strain up to 50%. Recently, Cho and co-workers28 reported the blend of P3HT nanowires embedded in poly(dimethylsiloxane) (PDMS) matrix had the mobility of around 3 × 10−4 cm2 V−1 s−1 with the strain up 100%. Also, controlling the regioregularity of conjugated polymers could maintain the advantages of semiconducting characteristics as well as mechanical durance. For example, Kim and co-workers have established that the mechanically and electrical properties of P3HT can be systemitically tuned by controlling the RR to meet the purposes of various organic electronic applications.29 Another approach is synthesize block copolymers consisting of conjuagted polymer and rubbery soft polymer blocks In addition, the phase separation of blcok copolymers resulted in the ordered microstructures for controlling the electronic properties.30,31 In our previous work,32 we explored rod−coil donor accptor polyfluorene-block-poly(pendent isoindigo) (PFb-Piso) for stretchable memory applciations and allowed reliable stretchability at strain of 50% but could not sustain a larger strain level. Moreover, the P3HT-based rod−coil block

INTRODUCTION Stretchable electronic devices have attracted extensive research interest becasuse of their potentials applications in biomedical instruments, smart skins, stretchable displays, and photovoltaics.1−11 Apart from the active materials, a stretchable electronic device requires compliant electrodes and substrate, which need to sustain a large mechanical strain and maintain a high conductivity. Coventional semiconductor films could not meet the request for stretchable devices because of their brittle characteristics. Therefore, it is inevitable to develop highly conducting materials under a high stretchability. There are two methods reported in the literature on fabricating stretchable conductors. The first method is to take metal wires and carbon nanotubes into a curved or tortuous shapes.1,12,13 Another approach is to blend fillers, such as signal-wall carbon nanotubes (SWNT),14 metal nanoparticles,15,16 and garphene,17−19 into an elastic polymer to form a conducting elastic gel. However, the above approaches were limited by the adhesion of the semiconductor to the rubber sustrate and its dispersion uniformity in the rubber matrix. Conjugated polymers have been considered for stretchable electrodes due to their wide range of applications including organic field-effect transistors (OFETs),20−22 photovoltaics (OPVs),23 and light-emitting diodes (OLEDs).24,25 However, their high mechanical flexible was limited by the poor elasticity due to their rigid backbone. The improvement on the stretchability of the conjugated films typically through incorporating rubber-like materials into conjugated polymers © XXXX American Chemical Society

Received: December 19, 2016 Revised: February 7, 2017

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DOI: 10.1021/acs.macromol.6b02722 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Scheme 1. Synthesic Scheme of P3HT-b-PBA Rod−Coil Diblock Copolymers

hexylthiophene was purchased from TCI and purified by vacuum distillation over CaH2. 6-Azidohexyl-2-bromo-2-methylpropanoate (AHBMP) initiator was prepared using our reported precedure.48 Synthesis of Alkynyl-Functionalized P3HT. In an argon-filled glovebox, anhydrous LiCl (263 mg, 6.19 mmol), THF (12.3 mL), 2,5dibromo-3-hexylthiophene (2.00 g, 6.13 mmol), and t-BuMgCl (1.0 M in THF; 5.83 mL, 5.83 mmol) were sequentially added to a vial, and the whole mixture was stirred for 4.5 h at 27 °C to prepare the activated Grignard monomer. The monomer solution was then added to a suspension of Ni(dppp)Cl2 (108 mg, 204 μmol). After stirring for 6 min, the mixture was cooled to −10 °C for 3 min. To this reacting mixture, styrene (21.3 g, 204 mmol) and the mixture of TMSA (2.9 M in THF; 3.52 mL, 10.2 mmol) and t-BuMgCl (1.0 M in THF; 10.2 mL, 10.2 mmol) which had been stirred for 4.5 h were added. After stirring at −10 °C for 30 min, a small amount of 5 M HCl was added to the reacting mixture, and the polymer was collected by the precipitation in MeOH. The obtained polymer was further purified by washing with MeOH, acetone, and hexane using a Soxhlet apparatus. Finally, TMS-protected alkynyl-functionalized P3HT (500 mg, 49.1%) was obtained by the chloroform extraction followed by precipitation in MeOH. Mn,NMR = 5660 g mol−1; Mw/Mn = 1.08. 1H NMR (400 MHz, CDCl3, δ (ppm)): 6.97 (s, thiophene), 2.80 (t, J = 5.33 Hz, thiophene−CH2−), 1.76−1.64 (m, thiophene-CH2−CH2−), 1.48− 1.24 (m, CH3−(CH2)3−), 0.95−0.78 (m, −CH3), 0.26 (s, −Si− (CH3)3). The obtained trimethylsilylacetylene (TMS)-protected product (255 mg) was dissolved in dry THF (60 mL). TBAF (6 drops, 1.0 M in THF) and MeOH (6 drops) were then added to the polymer solution at room temperature. After stirring for 15 min, the reaction mixture was passed through a pad of alumina and eluted with THF. Finally, alkynyl-functionalized P3HT (240 mg, 94.1%) was obtained by the precipitation in MeOH. Mn,NMR = 5500 g mol−1; Mw/Mn = 1.07. 1 H NMR (400 MHz, CDCl3, δ (ppm)): 6.97 (s, thiophene), 3.52 (s, −CCH), 2.80 (t, J = 5.33 Hz, thiophene−CH2−), 1.76−1.64 (m, thiophene−CH2−CH2−), 1.48−1.24 (m, CH3−(CH2)3−), 0.95−0.78 (m, −CH3). Synthesis of Azido-Terminated Poly(butyl acrylate) Homopolymer (2). The synthetic scheme for azido-terminated poly(butyl acrylate) (N3-PBA) is shown in Scheme 1. The preparation of N3PBA3k under the typical procedure is as follows: CuBr (0.224 g, 1.56 × 10−3 mol) was evacuated for 30 min in a Schlenk flask and backfilled with argon. Just prior to polymerization, butyl acrylate monomer was passed through basic Al2O3 column in order to remove the inhibitor. A mixture of butyl acrylate (10.00 g, 0.078 mol), PMDETA (0.27 g, 1.56 × 10−3 mol), and AHBMP initiator (0.456 g, 1.56 × 10−3 mol) in anhydrous toluene (13.7 mL) was prepared and then degassed by three freeze−pump−thaw cycles. The liquid mixture was transferred to the other flask including CuBr. The polymerization was performed in a preheated oil bath at 70 °C, and the monomer conversion was monitored by 1H NMR. The polymerization was terminated by

copolymers containing a pendent acceptor group such as C60,33,34 perylene bisimide,35,36 and isoindigo37 or the polyacrylate38−40 were used to control the self-assembled behavior, optoelectronic properties, and flexibility. Also, there were several reports on the synthesis of conjugated segment block with soft polymer block.41−48 For example, Muller et al.43 reported that diblock copolymers of polyethylene (PE) and P3HT display outstanding flexibility and toughness with elongations at break exceeding 600%. However, the above highly extended block copolymers were only discussed in the bulk state and could not be fully released back to their original state after stretchaing. The charge transport characteristics of the polymer thin film under strain have not been fully explored. Therefore, P3HT-based block copolymers with both electrical properties and mechanical properties under a high strain need to be developed. In this study, we report the synthesis of poly(3-hexylthiophene)-block-poly(butyl acrylate) rod−coil diblock copolymers (P3HT-b-PBA, Scheme 1), which were synthesized via the click reaction with alkynyl-functionalized P3HT and azidoterminated PBA homopolymers. The semiconducting P3HT and soft PBA blocks were prepared by the Grignard metathesis polymerization and atom transfer radical polymerization (ATRP), respectively. Note that the Tg of PBA is around −54 °C, which serves as an elastic strand in the matrix. The strain-dependent nanofiber morphologies were characterized by atomic force microscopy (AFM), X-ray diffraction (GIXD), and small-angle X-ray scattering (GISAXS). Top contact FET was used to investigate the charger carrier transport characteristics of the studied polymers under a different strain level. The effect of the P3HT/PBA block ratio on the mechanical property and morphology was evaluated and correlated with the stretchable FET performance.

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EXPERIMENTAL SECTION

Materials. Butyl acrylate (>99%), N,N,N′,N″,N″-pentamethyldiethylenetriamine (PMDETA, 99%), copper(I) bromide (CuBr, 98%), anisole (anhydrous, ≥99.7%), tetrahydrofuran (THF, anhydrous, ≥99.9%), methanol (MeOH, ≥99%), tetra-n-butylammonium fluoride hydrate (TBAF, 98%), acetone (≥99.5%), and chloroform (CHCl3, ≥99.9%, stabilized with amylenes) were purchased from Sigma-Aldrich Co. (St Louis, MO). LiCl was purchased from Kanto Chemical Co. (Tokyo, Japan) and was dried overnight at 160 °C under high vacuum. Trimethylsilylacetylene (TMSA) was received from Shin-Etsu Chemical Co. (Tokyo, Japan) and purified by distillation over CaH2 under an argon atmosphere. 2,5-Dibromo-3B

DOI: 10.1021/acs.macromol.6b02722 Macromolecules XXXX, XXX, XXX−XXX

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calorimetry (DSC) curves were obtained using a TA Instruments Q100 at a heating rate of 5 °C min−1 from −90 to 250 °C. UV−vis absorption spectra were recorded on a Hitachi U-4100 spectrophotometer. For the thin film spectra, polymers were dissolved in o-dichlorobenzene (10 mg mL−1), followed by filtering through a PTFE membrane syringe filter with pore size of 0.22 μm, and then spin-coated onto a quartz substrate at a spinning rate of 1000 rpm for 60 s. Cyclic voltammetry (CV) was performed on a CHI 611B electrochemical analyzer using a three-electrode cell in which ITO (polymer films area were about 0.5 × 0.7 cm2) was used as a working electrode. A platinum wire was used as an auxiliary electrode. All cell potentials were obtained with the use of a homemade Ag/AgCl, KCl(sat.) reference electrode. The electrochemical properties of the polymer films were measured in 0.1 M dry acetonitrile solution containing tetra-n-butylammonium perchlorate as the electrolyte. Atomic force microscopy (AFM) images were performed on a MultiMode AFM system with a Nanoscope 3D controller (Digital Instruments, Santa Barbara, CA) under tapping mode. Commercial silicon cantilevers (Nanosensor PPP-SEIHR) were used with a spring constant of 15 N m−1 and a resonant frequency of 130 kHz. The thickness of polymer film was measured with a Microfigure Measuring Instrument (Surfcorder ET3000, Kosaka Laboratory Ltd.). Grazingincidence X-ray diffraction (GIXD) and small-angle X-ray scattering (GISAXS) experiments were collected on beamline BL13A1 and BL23A1 in National Synchrotron Radiation Research Center (NSRRC, Taiwan) with monochromatic beams of λ = 0.827 (15 keV) and 1.0219 Å (12 keV), respectively. The incident angle set as 0.2° and the scattering intensities are reported as intensity versus q, where q = (4π/λ) sin(θ/2); λ is the wavelength of incident X-rays, and θ is the scattering angle. A rotational polarizer was used to measure the absorption intensity with the incident light polarized parallel (A∥) or perpendicular (A⊥) to the stretching direction and then define the dichroic ratio (R) of the stretched polymer film as R = A∥/A⊥. Note that the preparation of polymer thin film samples was the same as that of the device fabrications for the measurements of GIXD, UV, and AFM. Mechanical Characterization by the Buckling-Based Method. The tensile modulus was measured using the mechanical buckling technique for analyzing the mechanical properties for various thin-film systems including conjugated polymer films.49 The well-known buckling-based method originally developed by Hutchinson, Whitesides, and co-workers50 and further explored by Stafford et al.51 This technique has been used to study the wide range of film types and over a wide range of tensile moduli. The elastomer poly(dimethylsiloxane) (PDMS) was chosen as the relatively soft substrate for all experiments. The studied polymer was spin-coated onto octadecyltrimethoxysilane (OTS) modified SiO2 substrate and then transferred to prestrain 4% PDMS rectangular substrate. After transfer, the PDMS substrate was relaxed to original state; this action created a compressive strain that forced the studied polymer film to adopt sinusoidal buckles. The buckling wavelength, λb, and the thickness of the film, df, can be related to the tensile moduli of the film and the substrate, the substrate, Ef and Es, and the Poisson ratios of the two materials, νf and νs by the following equation:

bubbling air into the solution. The mixture was passed through a neutral Al2O3 column and eluted with THF to remove the catalyst. Finally, the concentrated polymer solution was purified by dialysis (Por dialysis membrane, molecular weight cutoff 1000) in acetone, which was changed on a regular basis for 48 h at room temperature, followed by drying in vacuum to give the a high viscosity liquid (4.60 g, 92%). Monomer conversion: 45%; Mn,NMR = 2930 g mol−1; Mw/Mn = 1.21. 1H NMR (400 MHz, CDCl3, δ (ppm)): 4.12−3.88 (−CH2CH2O−), 3.22−3.14 (−N3CH2CH2), 2.19−1.68, 1.48−1.16, and 1.12−0.87 (br, polymer backbone and butyl side chain). IR: 2096 cm−1 (−NNN stretching). The azido-terminated PBAs with different molecular weight, i.e., N3-PBA6k and N3-PBA12k, were prepared by the same method with different feed ratios ([butyl acrylate]0/[AHBMP]0 of 100 for N3-PBA6k and 200 for N3-PBA12k). N3-PBA6k: monomer conversion: 48%; Mn,NMR = 6100 g mol−1; Mw/ Mn = 1.16. N3-PBA12k: monomer conversion: 46%; Mn,NMR = 11 800 g mol−1; Mw/Mn = 1.13. Synthesis of P3HT-b-PBA Rod−Coil Diblock Comopolymer (3). Click reaction between alkynyl-terminated P3HT and azido-terminated PBA was performed according to Scheme 1. A typical procedure is as follows: The alkynyl-terminated P3HT (250 mg, 0.045 mol) and CuBr (13.0 mg, 0.091 mol) were placed in a Schlenk flask. The flask was evacuated 1 h and then backfilled with argon. A mixture of N3PBA3k (138 mg, 0.047 mol), anhydrous THF (5.0 mL), and PMDETA (15.8 mg, 0.091 mol) were degassed by three freeze−pump−thaw cycles; the liquid mixture was transferred to the other flask including P3HT homopolymer and CuBr. The sealed reaction flask was stirred for 24 h at room temperature. The 1H NMR, IR, and SEC were used to check the click reaction finished and terminated. The polymer solution was purified over a short pad of neutral Al2O3 column to remove the copper catalyst. After the evaporation of solvent, the residue was precipitated in cold methanol several times to remove the excessive PBA homopolymer. The final product of P3HT-b-PBA3k was obtained as a dark purple solid (333.7 mg, 86% yield). Mn,NMR = 8900 g mol−1, Mw/Mn = 1.13. 1H NMR (400 MHz, CDCl3, δ (ppm)): 6.98 (s, Har P3HT), 4.14−3.93 (−CH2CH2O−), 2.80 (t, α-CH2 P3HT), 2.39−2.15, 1.95−1.86, 1.75−1.19, and 0.96−0.82 (br, polymer backbone and alkyl side chain of P3HT and PBA). Anal. Calcd for [(C10H16S1)36 + (C7H12O2)23]: C 69.98%, H 9.54%, S 12.89%. Found: C 70.24%; H 9.27%; S 12.52%. P3HT-b-PBA6k and P3HT-b-PBA12k were synthesized by a similar method. P3HT-b-PBA6k (389.8 mg, 88% yield): Anal. Calcd for [(C10H16S1)36 + (C7H12O2)48]: C 68.81%, H 9.49%, S 9.49%; found: C 69.04%; H 9.57%; S 10.22%. P3HT-bPBA12k (366.8 mg, 92% yield): Anal. Calcd for [(C10H16S1)36 + (C7H12O2)92]: C 67.77%, H 9.45%, S 6.48%. Found: C 68.08%; H 9.88%; S 6.38%. Characterization. The 1H NMR spectra were recorded using a Bruker Avance DRX 400 MHz FT-NMR system. Size exclusion chromatography (SEC) measurements of the obtained polymers were performed at 40 °C using a Lab Alliance RI2000 instrument equipped with Waters Styragel HR2 and HR4 THF 7.8 × 300 mm columns and a refractive index detector. THF was used as the eluent solvent at the flow rate of 1.0 mL min−1. The number-average molecular weight (Mn,(SEC)) and polydispersity (Mw/Mn) of the polymers were calculated on the basis of a polystyrene calibration. The matrixassisted laser desorption/ionization time-of-flight mass spectral (MALDI-TOF MS) measurement of the polymer was performed in reflector mode with an Applied Biosystems Voyager-DE STR-H equipped with 337 nm nitrogen laser (3 ns pulse width). A sample for the measurement was prepared by mixing a polymer solution (15 g L−1 in THF; 2 μL) and a matrix solution (dithranol; 28.3 g L−1 in THF; 50 μL) and then depositing it on a sample plate. Fourier transform infrared spectroscopy (FT-IR) analysis was carried out using a PerkinElmer 100 Model FT-IR spectrophotometer. The elemental analysis was performed using the “Elementar Vario EL cube” elemental analyzer (for NCSH, Germany). Thermogravimetric analysis (TGA) was performed on a TA Instruments Q50. 3−5 mg powder samples were heated under a nitrogen flow (flow rate 100 cm3 min−1) at a heating rate of 10 °C min−1 from 100 to 800 °C. Differential scanning

⎛ 1 − ν 2 ⎞⎛ λ ⎞3 b f ⎟ Ef = 3Es⎜ ⎟ 2 ⎜ ⎝ 1 − νs ⎠⎝ 2πdf ⎠

(1)

We measured the tensile modulus of the substrate, Es (using a commercial pull tester), the buckling wavelength, λb (by optical microscopy), and the film thickness, df (by stylus profilometry). The slope of a plot of λb vs df for four different film thicknesses was inserted into eq 1. The tensile modulus of the PDMS, Es, was obtained around 0.9 MPa. The Poisson’s ratios were taken as 0.5 and 0.35 for PDMS and the conjugated polymer film. Fabrication and Characterization of Field-Effect Transistors. The field-effect transistors (FETs) using the studied copolymers with a top-contact configuration. A 300 nm SiO2 layer (capacitance per unit area = 10 nF cm−2) as a gate dielectric was thermally grown onto the highly n-type doped Si (100) substrates. The Si substrates were C

DOI: 10.1021/acs.macromol.6b02722 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules modified with an octadecyltrimethoxysilane (OTS) self-assembled monolayer according to a reported procedure.52 The polymer solutions (P3HT and P3HT-b-PBA) were prepared in o-dichlorobenzene (10 mg mL−1) and dissolved at 100 °C for at least 2 h. The polymer thin films were spin-coated onto OTS-modified SiO2/Si substrates using a spin rate of 1000 rpm for 60 s. In addition, an annealing process (140 °C under a nitrogen atmosphere for 1 h) was introduced to induce the molecular packing structures and enhance the FET performance. 70 nm thick gold source and drain contacts were subsequently thermal evaporated through a regular shadow mask, and the channel length (L) and width (W) were 50 and 1000 μm, respectively. For measuring the electrical characteristics of the stretched studied polymer thin films (i.e., P3HT, P3HT-b-PBA3k, P3HT-b-PBA6k, and P3HT-b-PBA12k), the polymer films were first spin-coated on a OTSmodified SiO2/Si substrate and annealed at 140 °C for 1 h. Afterward, those films were transferred to an elastomeric PDMS (20:1 base to cross-linker ratio by mass) slab. Various tensile strain levels were applied to the polymer/PDMS matrix, and thus the studied polymer thin films were stretched along with PDMS substrate. Such stretched films were transferred back to a silicon substrate with 300 nm SiO2 dielectric layer. Finally, the top-contact source/drain Au contacts were defined through a regular shadow mask with the channel length (L) and width (W) of 50 and 1000 μm, respectively. Electrical characteristics of FETs were recorded in a N2-filled glovebox using a Keithley 4200 semiconductor parametric analyzer (Keithley Instruments Inc., Cleveland, OH).

improved procedure, a well-defined and narrowly dispersed P3HT was indeed obtained without the high molecular weight byproduct as well as the dicapped byproduct. The Mn,NMR and Mw/Mn values of the obtained TMS-protected alkynylfunctionalized P3HT were determined to be 5660 g mol−1 and 1.08, respectively. More importantly, the MALDI-TOF MS spectrum displayed only one series of peaks corresponding to the P3HT having a bromine α-end and a trimethylsilylethynyl ω-end, which confirmed the quantitative introduction of the ethynyl group without the contamination of the dicapped byproduct. (1H NMR and MALDI-TOF MS spectra are shown in Figure S1, Supporting Information). The TMS protection was quantitatively removed by the treatment with TBAF to give the desired alkynyl-functionalized P3HT with the Mn,NMR of 5500 g mol−1 and Mw/Mn of 1.07 (1H NMR and SEC data are shown in Figure S2). The azido-terminated poly(butyl acrylate)s (N3-PBA) were synthesized by ATRP of butyl acrylate using AHBMP as the initiator, in which the [monomer]0/[AHBMP]0 ratios of 50, 100, and 200 were applied to produce the polymer with different molecular weights. Figure S3 shows the 1H NMR spectra and SEC traces of the obtained poly(butyl acrylate) with the azido end group (N3-PBAn, n = 24, 48, 92). As shown in Figure S3a, the butyl side chain is observed in the ranges of 4.12−3.88 ppm (peak a), which is attributed to the oxygen neighboring methylene group. The relative small proton signal is clearly observed at 3.21 ppm (peak b) due to the azido group derived from the azido-methylene proton. From the relative integration of the proton signals a and b, we calculate the molecular weight of PBA hompolymers. In addition, the signals at 2.19−1.68 ppm (peaks c and d), 1.48−1.16 ppm (peaks e and f), and 1.12−0.87 ppm (peak h) are assigned to the polymer backbone and butyl side chain chemical structure of N3-PBA. The SEC trace of each N3-PBA exhibits a unimodal elution peak with the Mw/Mn value of 1.13−1.21 (Figure S3b). The molecular characteristics of the obtained N3-PBA hmopolymers are listed in Table 1.

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RESULTS AND DISCUSSION Polymer Synthesis and Chemical Structure Characterization. Alkynyl-functionalized P3HT can be synthesized by Grignard metathesis (GRIM) polymerization and subsequent end-capping with an alkynyl group. For example, McCullough et al. established a one-pot protocol to synthesize the alkynylfunctionalized P3HT by the GRIM of the activated Grignard monomer derived from 2,5-dibromo-3-hexylthiopene followed by the addition of ethynylmagnesium bromide.53 The main drawback of this protocol is the contamination of dicapped P3HT along with the desired monocapped one, which eventually results in the contamination of ABA-type triblock copolymer in the final product after the click reaction. The other inherent problem in the synthesis of the alkynylfunctionalized P3HT is the contamination of a high molecular weight byproduct that corresponding to nearly twice that of the main product. The origin of the high molecular weight byproduct seems to be alkyne−alkyne homocoupling reactions, and such reaction gradually proceeds during the storage under the ambient condition. Indeed, Bielawski et al. reported that the protection of the ethynyl end group with a trimethylsilyl (TMS) group can suppress the homocoupling reaction even under ambient conditions.54 Considering these aspects, we have established an improved end-capping procedure to produce the alkynyl-functionalized P3HT, with a high level of monocapped end group fidelity as well as narrow polydispersity index (Mw/ Mn), by modifying McCullough’s procedure. Here, a Grignard reagent of trimethylsilylacetylene (trimethylsilylethynylmagnesium chloride) was employed instead of ethynylmagnesium chloride for the end-capping, affording a TMSprotected alkynyl-functionalized P3HT. In our improved procedure, a large amount of styrene was added to the polymerization mixture before the addition of trimethylsilylethynylmagnesium chloride. The unsaturated group of styrene can coordinate with reactive Ni(0) to form a stable π-complex, which contributes to prevent any further reaction with the bromine to form the dicapped byproduct.55 Using our

Table 1. Molecular Weight, Polydispersity Index (Mw/Mn), and Composition of P3HT of P3HT-b-PBA sample N3-PBA3k N3-PBA6k N3-PBA12k P3HT P3HT-bPBA3k P3HT-bPBA6k P3HT-bPBA12k

Mn,NMRa (g mol−1)

P3HTa (wt %)

Mn,secb (g mol−1)

Mw/Mnb

2930 6100 11800 5500 8900

0 0 0 100 61.8

3500 5500 7800 8700 11500

1.21 1.16 1.13 1.07 1.13

11100

49.5

14100

1.15

18300

30.1

24200

1.12

a

Determined by 1H NMR in CDCl3. bDetermined by SEC with RI detector in THF using polystyrene standard.

The coupling reaction of the alkynyl-functionalized P3HT and N3-PBA6k was carried out by the Cu-catalyzed azido-alkyne click reaction to produce the desired block copolymer, i.e., P3HT-b-PBA6k. The 1H NMR spectrum of the P3HT-b-PBA6k in CDCl3 is shown in Figure 1, in which the proton signals at a (6.98 ppm), b (4.12−3.88 ppm), c (2.80 ppm), d (2.28−2.08 ppm), and e (1.98−1.78 ppm) are attributed to the protons of the thiophene ring, the butyl and hexyl side chain, and the vinyl D

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Figure 1. 1H NMR spectrum of P3HT-b-PBA6k in CDCl3.

group of the butyl acrylate block, respectively. The effective coupling between the azido and alkynyl groups was further confirmed by the fact that the characteristic absorption band at 2097 cm−1 in the FTIR spectrum completely disappeared after the click reaction. Also, the SEC trace of P3HT-b-PBA6k (Mn,NMR = 11 100, Mw/Mn = 1.21) clearly demonstrates the shift in the elution peak to the high-molecular-weight region in comparison with that of its precursor PBA6k, as shown in Figure S4. Furthermore, the elemental analyses of the carbon, hydrogen, and sulfur contents are in a good agreement with the theoretical content. Similarly, 3a and 3c with respective block compositions of P3HT-b-PBA3k and P3HT-b-PBA12k are also confirmed by 1H NMR, FTIR, and SEC as shown in Figure S5. The number-averaged molecular weight and polydispersity of the prepared block copolymers are summarized in Table 1. All the synthesized polymers are easily soluble in common organic solvents, such as o-dichlorobenzene, chloroform, tetrahydrofuran, and toluene. Thermal Properties. The thermal properties of P3HT, N3PBA12k and P3HT-b-PBA were investigated by differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA), as shown in the Figure 2a and Figure S6, respectively. As shown in the DSC curves, the glass transition temperature (Tg) of P3HT-b-PBA observed at −33 to −55 °C is systematically decreased as the PBA block length increased, suggesting the effect of PBA coils (i.e., with a Tg of −57 °C). In addition, the melting temperature (Tm) peak of the P3HT-bPBA is shown in a range of 196−203 °C, which is dominated by the P3HT block. With the increase of the PBA coil length, the P3HT endothermic peak becomes weaker, and the Tm is generally shifted to a lower temperature for the prepared P3HT-b-PBA block polymers. Note that the appearance of individual phase transitions due to each block indicates a possible phase separation in these block copolymers.45 On the other hand, the thermal decomposition temperatures (Td, 5% weight loss) of P3HT and N3-PBA12k are observed at 430 and 302 °C, respectively. As expected, the Tds of the studied P3HTb-PBA block copolymers are between those of the parent P3HT and PBA homopolymers. Optical and Electrochemical Properties. Figure 2b shows the normalized solid state UV−vis absorption spectra of P3HT, N3-PBA12k, and P3HT-b-PBA, and their corresponding maximum absorption wavelengths (λmax) and optical band gaps (Eopt g ) are summarized in Table 2. The P3HT-b-PBA block copolymer thin films exhibit the clear λmax wavelengths at 520 and 545 nm and accompany a vibrational peak at 600 nm. Such features are similar to those of the P3HT homopolymer,

Figure 2. Polymer characterstics, including (a) DSC traces, (b) solid state UV−vis absorption spectra, and (c) cyclic voltammery curves, of the studied block copolymers.

indicating that the soft PBA coils does not interrupt the interor intrachain packing as well as π−π stacking structures of the P3HT block. The optical band gaps (Eopt g ) of P3HT and P3HTb-PBA thin films estimated from the onsets of the absorption spectra are around 1.9 eV. The Eopt values of P3HT-b-PBA g block copolymers show that the P3HT/PBA block ratio has no effect on their optoelectronic properties. The energy levels of the P3HT, P3HT-b-PBA, and N3-PBA were investigated using cyclic voltammetry (CV), as shown in Figure 2c. The highest occupied molecular orbital (HOMO) levels were determined from the onset oxidation potentials ox (Eox onset), based on the equations of HOMO = −e[Eonset − 1/2 Eferrocene + 4.8] V with reference to ferrocene (4.8 eV). On the other hand, the lowest unoccupied molecular orbital (LUMO) levels were calculated by difference between HOMO and optical band gap. The obtained HOMO and LUMO energy levels are calculated as −4.74 to −4.79 eV and −2.89 to −2.87 eV, respectively, as listed in Table 2. The optical and electrochemical properties of the P3HT-b-PBA block copolymers are mainly attributed to the P3HT conjugated backbone, in which the soft PBA segment length does not disturb the energy levels as expected. Polymer Thin Film Morphologies. The surface and crystalline structures of the P3HT-based block copolymers are characterized using atomic force microscopy (AFM) and grazing incidence X-ray diffraction (GIXD) analyses, as shown in Figure 3. The AFM topographies (Figure 3a) present the clear nanofibrillar-like surface structure and interconnected networks. P3HT-b-PBA3k possesses the nanofiber structure similar to that of P3HT due to the low PBA ratio. With increasing the PBA coil ratio, the morphology shows different E

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Macromolecules Table 2. Physical Properties of N3-PBA, P3HT, and P3HT-b-PBA sample P3HT P3HT-b-PBA3k P3HT-b-PBA6k P3HT-b-PBA12k N3-PBA12k

a λfilm max (nm)

520, 518, 520, 519,

546, 544, 543, 546,

603 605 603 608

b Eopt (eV) g

HOMO (eV)

LUMOc (eV)

Td (°C)

Tg (°C)

1.85 1.90 1.91 1.92

−4.74 −4.76 −4.78 −4.79 −5.77

−2.89 −2.86 −2.87 −2.87

430 349 326 316 302

−38 −51 −55 −57

Tm (°C) 209 203 198 196

a Spin-cast film from o-dichlorobenzene solution onto a quartz substrate. bEstimated from onset wavelength of the UV−vis absorption spectrum c opt d using the formula Eopt g = 1240/λonset (nm). The LUMO energy level was estimated by the equation LUMO (eV) = HOMO + Eg . Determined by GISAXS measurement using thermal-annealed thin films.

structures. The similar lamellar spacing and π−π stacking distance of the studied P3HT-b-PBA block copolymers indicates that the packing structure is dominate by P3HT conjugated backbone.46 Tensile Modulus and Polarized UV/Vis Properties. The ductility of studied polymer thin films was measured by polarized UV−vis spectroscopy and tensile modulus. The dichroic ratio (R) is calculated from the polarized UV−vis spctroscopy, as shown in Figure S8. Figure 4a shows the R

Figure 3. (a) AFM topographies and (b) 2D GIXD patterns of the studied polymer thin films. Note that all polymer thin films were annealed at 140 °C under vacuum for 1 h. Figure 4. (a) Dichroic ratio of the studied polymer films under a strain ranging from 0 to 100%. (b) Mechanical properties of the studied polymer thin films. The thin film tensile modulus was determined from the relationship between buckling wavelength and film thickness, in which the modulus is summarized in the inset.

types, such as fiber bundle, fiber network, and soft fiber nanostructures. It indicates the PBA segment highly influences the thin film nanoscale morphology and reduces the P3HT domain aggregation. To further study the microphase separation between the P3HT and PBA domains, the grazingincidence small-angle X-ray scattering (GISAXS) is employed, as shown in Figure S7. The 1D-GISAXS profiles of P3HT-bPBA3k, P3HT-b-PBA6k, and P3HT-b-PBA12k thin films show the first-order scattering peaks (q*) at 0.0424, 0.0348, and 0.0274 Å−1, respectively, which correlate to the center-to-center average distances of 14.82, 18.06, and 22.93 nm for the P3HT domain. The molecular packing nanostructures between polymer chains are further studied by the 2D GIXD patterns. As shown in Figure 3b, the studied polymer thin films of P3HT and P3HT-b-PBA possessed diffraction signals in the qz direction (i.e., (n00)) with the similar lamellar spacing of around 15.72− 15.8 Å. The π−π stacking diffraction peak is observed in the inplane direction with the spacing distances in a range of 3.8− 3.86 Å, indicating the well-ordered edge-on molecular packing

values of P3HT and P3HT-b-PBA3k have a similar trend: progressively increase to 1.1−1.2 as the strain is applied from 0 to 50% and then become steady with increasing the tensile strain. However, the R values of P3HT-b-PBA6k and P3HT-bPBA12k rapidly increase to 1.4−1.6 with the strain applied from 0 to 100%. It indicates that the conjugated chain alignment by the strain could be enhanced. To further measure the thin film ductility, the tensile modulus was estimated using the buckling method and measured the buckled films, as shown in Figure S9. The variation of buckling wavelengths (λb) with different film thickness (df) is shown in Figure 4b. The tensile modulus of pure P3HT is evaluated approximately 0.93 GPa, similar to that of the previously reported values.49 As expected, the modulus of the P3HT-b-PBA thin films (inset of Figure 4a) decreases from F

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Macromolecules 0.63 GPa (P3HT-b-PBA3k) to 0.19 GPa (P3HT-b-PBA12k) with increasing the length of PBA segment. The thin film ductility is enhanced by incorporating the low-Tg PBA phase into the rigid P3HT phase. Strain-Dependent Polymer Thin Film Morphologies. Under the mechanical strain, the surface structure of polymer thin film using the second transfer method56 was first characterized by the optical microscopy, as summarized in Figure 5a and Figure S10. In the case of unstrained state, a

Figure 6. (a) Lamellar packing spacing and (b) π−π stacking distance of the studied polymer thin films under different strain levels. Note that the X-ray incident light is controlled to be perpendicular (perp.) or parallel (para.) to the strain direction for the lamellar spacing and π−π stacking analyses.

P3HT-b-PBA6k film under the strain of 100% is only increased to 4−5 nm. By further increasing the PBA ratio, P3HT-bPBA12k thin film exhibits a softer fiber-like structure but changes to the noncontinuous fiber domains under the high strain state (Figure 5b and Figure S11). Although the high PBA composition could enhance the ductility of thin film, the nanofiber structure is disrupted at a high strain. Thus, the optimized block ratio (i.e., P3HT-b-PBA6k) could maintain both good mechanical properties and well-organized nanofiber structure of the prepared thin film. In order to analyze the molecular packing of the prepared polymers, the grazing-incidence wide-angle X-ray diffraction (GIWAXS) was used, as shown in Figures S12 and S13. The results of the lamellar and π−π spacing variation from the 2D GIWAXS patterns are summarized in Figures 6a and 6b. The incident X-ray beam was controlled to be on perpendicular or parallel to the stretching direction. The lamellar spacing distances of the P3HT and P3HT-b-PBA3k thin films have a similar trend. When the applied strain to 50%, the spacing in the direction of the X-ray beam perpendicular to the strain direction is decreased; however, the lamellar spacing does not have an obvious change in the other direction (Figure 6a). It suggests that the layer-by-layer lamellar stacking is significantly compressed in the perpendicular direction. As the tensile strain increased to 50−100%, the lamellar spacing is recovered back to the original state, which results in the crack formation to release the tensile strain on the polymer thin film.57 For the case of P3HT-b-PBA6k and P3HT-b-PBA12k thin film, the lamellar spacing in the X-ray beam perpendicular to the strain direction becomes smaller than that in the parallel direction under the applied strain. It suggests that the lamellar packing is pressed in the perpendicular direction but does not to the original value since no crack is occurred in the film. Nevertheless, the π−π spacing of the P3HT and P3HT-bPBA3k thin films is slightly elongated in the X-ray beam parallel to the strain direction compared to the other direction (Figure 6b), indicating that the crystalline regions are elongated and

Figure 5. (a) OM images of the studied polymer thin films at 100% strain (the scale bar is 20 μm). (b) AFM topographies of the selected areas on a stretched polymer films.

smooth surface is observed for all of the studied samples (Figure S10a). As the strain applied, the surface structure significantly changes with a different block ratio. The P3HT film exhibits the microscale crack under the 25% strain (Figure S10b), and the cracks are continuously enlarged with increasing the strain force. As the PBA block increased, the P3HT-bPBA3k thin film improves film ductility, but nanoscale cracks are observed at the strain of 50%−100% (Figure S10c and Figure 5a). On the contrary, the thin films of P3HT-b-PBA6k and P3HT-b-PBA12k possess the smooth surface without cracking even at the strain of 100%, as shown in Figure 5a. It suggests that the thin film ductility is significantly enhanced by incorporating the low-Tg PBA in the rod−coil block polymers. The thin film structure of the prepared thin films were further characterized using AFM and GIXD, as shown in Figures 5b and 6a,b and Figures S11−S14. The AFM topographies for the higher P3HT content (i.e., P3HT and P3HT-b-PBA3k) gradually show the aggregated and rough surfaces at different strain levels (25%, 50%, and 100), in which the roughness is increased from 2−3 to 15−20 nm (Figure 5b and Figure S11). On the contrast, the P3HT-b-PBA6k with a clear nanofibrillar structure is maintained even under the strain of 100% (Figure 5b). Note that the surface roughness of the G

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Macromolecules aligned in the stretching direction. Similarly, the π−π spacing of P3HT-b-PBA6k and P3HT-b-PBA 12k thin films is also progressively increased as enhancing the strain force (Figure 6b). But only the lamellar signals ((200), (300)) and π−π interaction of P3HT-b-PBA12k thin film under the strain of 100% become weaker. It explains that the stretchable P3HT-bPBA12k FET could not keep a high carrier mobility under a high strain condition. The grazing-incidence small-angle X-ray scattering (GISAXS) reveals the microphase separation structural change between the P3HT and PBA domains under a different strain level, as shown in Figure S14. The domain spacings (dspacings) of the P3HT-b-PBA thin films with the incident X-ray perpendicular or parallel to strain direction are summarized in Figure S14d. As P3HT-b-PBA3k thin film under the strain, the d-spacing values change from 16.93 nm (25%) to 18.11 nm (50%) and 19.21 nm (100%). It indicates that the higher P3HT ratio with a strong π−π interaction limits the elongated behavior of the P3HT and PBA blocks, and then the nanoscale cracks are occurred to release the strain force. In the case of the P3HT-b-PBA6k and P3HT-b-PBA12k thin films, the d-spacing value linearly increases in the X-ray beam parallel and perpendicular stretching directions. As applied strain 25%− 100%, the d-spacing values of P3HT-b-PBA6k and P3HT-bPBA12k thin film change from 22.12 to 30.05 nm and 28.43 to 45.2 nm, respectively. The d-spacing value of the P3HT-bPBA6k thin film becomes steady under the strain 50%−100% region. However, the d-spacing value of P3HT-b-PBA12k thin films significantly increases as the strain enhanced, suggesting the fiber-to-fiber distance is enlarged and thus the nanofiber structure could be disrupted. Note that the AFM height images under an extreme strained level (Figure 5b) agree with that from the GISAXS result; the P3HT-b-PBA6k thin film retains the tight and network-like nanofiber structure, but P3HT-bPBA12k exhibits the noncontinuous or fractured fiber-like surface. The above result indicates that the P3HT-b-PBA6k thin film could be potentially used in stretchable electronic devices as the active layer due to its morphology stability under a high strain. Strain-Dependent Field-Effect Transistor Characteristics. The charge carrier mobility of the prepared polymers was probed using the top-contact field effect transistor (FET) device, as shown in Figure 7a. We measured the transfer curves at tensile strain values of 0−100% with the charge transport direction parallel or perpendicular to the stretching direction as shown in Figure 7b and Figure S15 and summarized in Table 3. The FET hole mobility of the P3HT device rapidly decreases form 7.9 × 10−2 to 1.1 × 10−4 cm2 V−1 s−1 as the strain level increased to 100%. Also, the FET mobility of P3HT-b-PBA3k and P3HT-b-PBA12k devices decrease from 7.3 × 10−2 and 1.4 × 10−2 m2 V−1 s−1 to 1.6 × 10−3 and 1.9 × 10−3 cm2 V−1 s−1, respectively, as the strain increased to 100%. The FET mobility in the stretching direction parallel or perpendicular to electrode shows no obviously change. The variation on the FET mobility of P3HT-b-PBA3k-based device shown in Figure 7c is separated into two regions: 0−50% and 50%−100%. As shown in the OM and AFM images, the tiny nanoscale cracks are observed in the P3HT-b-PBA3k thin film under the strain of 50%, but the crack area and density increase significantly as strain increased to 100%. Therefore, the FET mobility decreases faster in the strain region of 50−100% than that in the lower strain of 0−50% region. In the case of P3HT-b-PBA12k-based FET device under the applied strain, the mobility decreases slower than that of the

Figure 7. (a) Schematic illustration of the FET device with a P3HT-bPBA active layer. (b) FET transfer characteristics of the P3HT-b-PBA polymer films under unstrain (left) and strain of 100% with the charge transport direction parallel (median) and perpendicular (right) to the stretching direction. (c) The charge carrier mobility of P3HT and P3HT-b-PBA thin films under different strain level.

P3HT-b-PBA3k device but the initial mobility is lower (1.4 × 10−2 cm2 V−1 s−1). On the contrary, the FET P3HT-b-PBA6k based device shows a stable FET mobility around 0.02 cm2 V−1 s−1 even under the strain of 100%. Figure 7b shows the transfer curves as the charge transporting direction parallel or perpendicular to the stretching direction, and the FET mobility maintains a similar value under a different strain level (Figure 7c). The related output characteristics are shown in the Figures S16 and S17. Also, the on/off current ratio and threshold voltage of the P3HT-b-PBA6k thin film show an insignificant change under the strain condition. Moreover, the devices based on the P3HT-b-PBA6k thin film stretched for 1000 stretching/ release cycles show stable values at the strains of 25%, 50%, and 100%, as shown in Figure 8 and Figure S18. At the strain levels of 25% and 50%, the mobility and threshold voltage maintain the stable performance after 1000 stretching cycles (Figure 8a,b). When strain up to 100% (Figure 8c), the mobility and threshold voltage decreases to 1.5 × 10−2 cm2 V−1 s−1 and −31.7 V after 200 stretching cycles, but no obvious decay even after 1000 cycles of strain (1.1 × 10−2 cm2 V−1 s−1 and −34.3 V). It indicates that the optimized block ratio (P3HT-b-PBA6k) can provide a stable and high charge transport characteristics (larger than 10−2 cm2 V−1 s−1) under strain up to 100%.

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CONCLUSION In this study, we have successfully synthesized P3HT-b-PBA rod−coil block copolymers for stretching FET device applications, in which the P3HT and PBA blocks are responsible for the semiconducting properties and mechanical endurance, respectively. Thin film morphology with the nanofibrillar-like structures after thermal annealing were confirmed by AFM and X-ray diffraction analyses. The polymer thin films ductility were enhanced by incorporating the low-Tg PBA block, based on the results from the polarized optical H

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Macromolecules Table 3. FET Performance of the P3HT and P3HT-b-PBA6k Stretched up to 100% P3HT strain levelb (%) 0 25 25 50 50 100 100

strain direction parallel perpendicular parallel perpendicular parallel perpendicular

μ (cm2 V−1 s−1) 7.9 1.2 1.1 7.4 4.3 2.2 1.1

× × × × × × ×

10−2 10−2 10−2 10−4 10−4 10−4 10−4

P3HT-b-PBA6k

Ion/Ioff

Vth (V)

∼104 ∼104 ∼104 ∼103 ∼103 ∼103 ∼103

3 −6 −3 −9 −5 −12 −8

μ (cm2 V−1 s−1) 6.1 5.1 5.1 3.9 3.4 2.5 2.5

× × × × × × ×

10−2 10−2 10−2 10−2 10−2 10−2 10−2

Ion/Ioff

Vth (V)

∼107 ∼107 ∼107 ∼107 ∼107 ∼107 ∼107

−10 −13 −13 −15 −18 −21 −20

a All the electrical properties are averaged from at least 10 devices with three different batches. bThe FET devices were fabricated using a transferred polymer thin film with various strain levels.

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protected alkynyl-functionalized P3HT; 1H NMR spectrum and SEC traces of alkynyl-functionalized P3HT36 azido-terminated, N3-PBA homopolymers; P3HT, N3-PBA12k, P3HT-b-PBA3k, P3HT-b-PBA6k, and P3HT-b-PBA12k TGA curves, GISAXS, optical micrographs, AFM images, GIWAXS of the studied polymers; FET output characteristics of the studied polymer films under different strains (PDF)

AUTHOR INFORMATION

Corresponding Authors

*(W.-C.C.) E-mail [email protected]. *(T.S.) E-mail [email protected]. ORCID

Toshifumi Satoh: 0000-0001-5449-9642 Wen-Chang Chen: 0000-0003-3170-7220 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors highly appreciate the financial support from the Ministry of Science and Technology of Taiwan. The authors also acknowledge NSRRC, Taiwan, for facilitating the X-ray experiments.

Figure 8. Endurance of the P3HT-b-PBA6k thin film-based FET device under the applied strain of (a) 25%, (b) 50%, and (c) 100% for 1000 stretching/releasing cycles. The charge transport direction is controlled to be parallel or perpendicular to the strain direction.

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absorption and optical microscopy. The field effect transistors (FET) using P3HT-b-PBA as the active layer showed a high ptype mobility over 10−2 cm2 V−1 s−1, indicating its good charge transporting ability. Furthermore, the P3HT-b-PBA6k based FET under 100% strain had a high mobility of 2.5 × 10−2 cm2 V−1 s−1 with an on/off ratio of 7.2 × 106 and it maintained over 10−2 cm2 V−1 s−1 for 1000 cycles, suggesting the promising stability and reproducbility. Our results indicate the newly designed conjugated rod−coil block copolymers could maintain the excellent charge transporting mobility even under a high strain, showing their potentials for high performance stretchable transistors.

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b02722. Synthesis of azido-terminated poly(butyl acrylate) and P3HT-b-PBA; FET performance of the P3HT-b-PBA3k and P3HT-b-PBA12k stretched up to 100%; 1H NMR spectrum and MALID-TOF MS spectrum of the TMSI

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DOI: 10.1021/acs.macromol.6b02722 Macromolecules XXXX, XXX, XXX−XXX