Naphthodithiophenediimide–Benzobisthiadiazole-Based Polymers

Jan 25, 2017 - New π-conjugated polymers with strong electron affinity, PNDTI-BBTs, consisting of naphtho[2,3-b:6,7-b′]dithiophenediimide (NDTI) an...
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Naphthodithiophenediimide−Benzobisthiadiazole-Based Polymers: Versatile n‑Type Materials for Field-Effect Transistors and Thermoelectric Devices Yang Wang,†,‡ Masahiro Nakano,*,† Tsuyoshi Michinobu,‡ Yasuhiro Kiyota,‡ Takehiko Mori,‡ and Kazuo Takimiya*,† †

Emergent Molecular Function Research Group, RIKEN Center for Emergent Matter Science (CEMS), 2-1 Hirosawa, Wako, Saitama 351-0198, Japan ‡ Department of Organic and Polymeric Materials, Tokyo Institute of Technology, 2-12-1 Ookayama, Meguro-ku, Tokyo 152-8552, Japan S Supporting Information *

ABSTRACT: New π-conjugated polymers with strong electron affinity, PNDTI-BBTs, consisting of naphtho[2,3b:6,7-b′]dithiophenediimide (NDTI) and benzo[1,2-c:4,5-c′]bis[1,2,5]thiadiazole (BBT) units, were synthesized. PNDTIBBTs have low-lying LUMO energy levels (∼−4.4 eV), which is sufficiently low for air-stable electron transport in organic field-effect transistors and for being readily doped by a wellknown n-dopant, N,N-dimethyl-2-phenyl-2,3-dihydro-1H-benzoimidazole (N-DMBI), affording doped polymer films with relatively high conductivities and Seebeck coefficients. Depending on the solubilizing alkyl groups (2-decyltetradecyl, PNDTI-BBT-DT, or 3-decylpentadecyl groups, PNDTI-BBT-DP), not only the electron mobility in the transistor devices with the pristine polymer thin films (PNDTI-BBT-DT: ∼0.096 cm2 V−1 s−1; PNDTI-BBT-DP: ∼0.31 cm2 V−1 s−1) but also the conductivity and power factor of the doped thins films (PNDTI-BBT-DT: ∼0.18 S cm−1 and ∼0.6 μW m−1 K−2; PNDTI-BBTDP: ∼5.0 S cm−1 and ∼14 μW m−1 K−2) were drastically changed. The differences in the electric properties were primarily ascribed to the better crystalline nature of the PNDTI-BBT-DP than those of PNDTI-BBT-DT in the thin-film state. Furthermore, UV−vis and ESR spectra demonstrated that doping effectiveness was largely affected by the alkyl groups: the PNDTI-BBT-DP films with better crystalline order prevented overdoping, resulting in ca. 20 times higher conductivity and power factors. From these results, it can be concluded that tuning the intermolecular interaction and consequently obtaining the thin-film with well-ordered polymers by the alkyl side chains is a promising strategy for developing superior thermoelectric materials.

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K−1), TE properties can also be evaluated by the power factor (PF, W m−1 K−2, eq 2)

INTRODUCTION

π-Conjugated polymers are promising materials for the manufacture of lightweight, mechanically flexible, and largearea printed electronic devices, such as organic field-effect transistors (OFETs) and polymer solar cells (PSCs).1−6 Recently, π-conjugated polymers have also captured the interest of thermoelectric (TE) researchers owing to the unique advantages that they have over their inorganic counterparts in realizing the low cost and high-throughput production of flexible TE modules.7−10 In general, TE devices require both hole and electron conducting materials. Whereas TE materials based on p-type polymers have been intensively studied and shown rapid advances,11−15 there are few examples of efficient n-type polymer TE materials.16,17 TE properties are quantified by the unitless figure of merit, ZT (eq 1), and because of the extremely low thermal conductivities (κ) of polymer materials (κ = 0.01−1 W m−1 © XXXX American Chemical Society

ZT = TS2σ /κ

(1)

PF = S2σ

(2)

where S (V K−1, positive for p-type and negative for n-type) is the Seebeck coefficient or thermopower and σ (S cm−1) is electrical conductivity, which is expressed as σ = nqμ (3) where n (cm−3) is carrier concentration, q is carrier charge, and μ (cm2 V−1 s−1) is carrier mobility. Consequently, a high electron mobility (μe) and the capability to be n-type doped, Received: October 25, 2016 Revised: January 3, 2017

A

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RESULTS AND DISCUSSION In order to validate our design strategy for highly electrondeficient polymers, we first carried out theoretical calculations of a triad-type model compound for PNDTI-BBTs, where two BBT units sandwich the NDTI core (namely, BBT-NDTIBBT), by using DFT methods at the B3LYP/6-31G(d) level.32 In the calculations, methyl groups were substituted for the Nalkyl side chains to reduce computational cost. The ELUMO of BBT-NDTI-BBT was calculated to be −3.99 eV, which is much lower than those of the model triads of previously reported NDTI-based polymers containing benzo[c][1,2,5]thiadiazole and naphtho[1,2-c:5,6-c′]bis[1,2,5]thiadiazole (−3.49 and −3.58 eV, respectively),22 indicating that the targeted PNDTI-BBTs are quite attractive as n-type polymers (see Supporting Information for details and Figure S1). The syntheses of PNDTI-BBTs were carried out by the copper iodide (CuI)-assisted Stille-coupling polycondensation of stannylated NDTI22 and brominated BBT derivatives33 (Scheme S1). The polymers were purified by consecutive Soxhlet extraction with methanol for 12 h, hexane for 12 h, and dichloromethane for 12 h to remove low molecular weight fractions, and then fractions with reasonably high molecular weight were finally extracted with chloroform. The numberaveraged molecular weights (Mn) were 16.8 and 20.5 kDa for PNDTI-BBT-DT and PNDTI-BBT-DP, respectively, as evaluated by gel permeation chromatography (GPC) using polystyrene as the standard and o-dichlorobenzene as the eluent at 140 °C (Table 1 and Figure S2). As the resulting polymers were sufficiently soluble in hot chloroform, chlorobenzene, and o-dichlorobenzene, they were used to fabricate organic electronic devices via the solution process. Figure 2 shows the cyclic voltammograms, the UV−vis−NIR absorption spectra in solution, and those of the polymer thin

i.e., electron-doped, are the prerequisites for realizing highperformance n-type polymer TE materials. Although polymer semiconductors with high μes have been developed,18,19 polymers capable of being efficiently electron-doped and consequently showing high electrical conductivities have been scarcely reported.4 Thus, the relationship between molecular structures of n-type polymers and TE properties remains vague. Polymers with energy levels of the lowest unoccupied molecular orbital (ELUMO) lower than −4.0 eV are potential candidates for the development of versatile n-type semiconducting polymers.20 To realize such low-lying ELUMO, employing or combining π-building block(s) with high electron deficiency in the polymer backbone is a promising strategy. 21−23 In this context, naphtho[2,3-b:6,7-b′]dithiophenediimide (NDTI)24−29 and benzo[1,2-c:4,5-c′]bis[1,2,5]thiadiazole (BBT)30,31 are among the most electrondeficient building blocks capable of being incorporated into a conjugated polymer backbone. By combining these two building blocks, new polymers PNDTI-BBT-DT and PNDTIBBT-DP, which have different solubilizing alkyl chains, were designed and synthesized (Figure 1). As expected from the high

Figure 1. Molecular structures of NDTI, BBT, PNDTI-BBTs, and NDMBI.

electron deficiencies of the units, the ELUMOs of the polymers are approximately 4.4 eV below the vacuum level, which are among the lowest for conjugated polymers reported to date.18 We have also investigated their OFET properties, solution doping behaviors, electrical conductivities, and TE properties. Through these investigations, we have found that the solubilizing side chains play a critical role in dictating the electronic properties of the polymer thin films, particularly the doping effectiveness, conductivity, and TE properties. Together, our findings offer insightful information on the design strategy for developing TE materials based on conjugated polymers.

Figure 2. (a) Cyclic voltammograms of the polymer thin films deposited on Pt plate working electrode. (b) UV−vis−NIR absorption spectra in chloroform solution and thin-film of PNDTI-BBTs.

films. From the reduction (Ered) onsets in the cyclic voltammograms, the ELUMOs of the polymers were estimated to be ca. −4.4 eV (Table 1). Note that these polymers are among the most electron-deficient polymers ever reported.18,22

Table 1. Molecular Weights and Electrochemical and Optical Properties of Polymers polymer

Mn (kDa)/PDIa

Eonset (V)b ox

b Eonset red (V)

EHOMO (eV)c

ELUMO (eV)c

λmax (nm)

λonset (nm)d

Egopt (eV)e

PNDTI-BBT-DT PNDTI-BBT-DP

16.8/2.52 20.5/2.53

+1.23 +1.26

+0.09 +0.11

−5.5 −5.5

−4.4 −4.4

884 914

1240 1300

1.00 0.96

a onset Determined by GPC using polystyrene as standard and o-dichlorobenzene as eluent at 140 °C. bOxidation (Eonset ox ) and reduction (Ered ) onset potentials observed in cyclic voltammograms. cElectrochemical EHOMO and ELUMO were estimated from E (eV) = −4.31 − Eonset. dEstimated from the absorption edge. eOptical band gap (Eopt g ) estimated from the onset wavelength of the thin films.

B

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Figure 3. Transfer characteristics of OFETs based on PNDTI-BBTs (-DT: a and -DP: b). Out-of-plane and in-plane XRD patterns of drop-cast thinfilms of PNDTI-BBT-DT (c, e) and -DP (d, f). The thin films of PNDTI-BBTs were annealed at 270 °C (-DT) or 300 °C (-DP).

lying ELUMOs, the OFETs fabricated from the polymer thin films afforded unipolar n-type charge transport properties. The spin-coated thin film of PNDTI-BBT-DT without thermal annealing exhibited μe of 3.3 × 10−3 cm2 V−1 s−1 and an on−off ratio (Ion/Ioff) of 1.3 × 102 (Table S1). Annealing the polymer thin film led to a drastic increase of μe, and the maximum μe of 0.096 cm2 V−1 s−1 with Ion/Ioff of 3.8 × 103 was achieved at the annealing temperature of 270 °C (Figure 3a). Similarly, the PNDTI-BBT-DP based devices also showed improved performances upon thermal annealing: μe of 1.1 × 10−2 cm2 V−1 s−1 was obtained on the pristine film (Table S1), whereas thermal annealing at 300 °C enhanced μe to as high as 0.31 cm2 V−1 s−1 with Ion/Ioff larger than 105 (Figure 3b). It is interesting to note that the μes of PNDTI-BBT-DP based OFETs are approximately 3 times higher than those of PNDTI-BBT-DT based devices, even though the polymer backbones are identical. This can be explained by the difference of molecular ordering in the thin-film state: as demonstrated in the thin-film X-ray diffraction (XRD) patterns (Figure 3c−f), both PNDTI-

In contrast, the oxidation peaks were irreversible and had rather small current response compared to the reduction peaks, indicative of the fairly n-type nature of these polymers. The absorption spectra of the two polymer thin films exhibited intense absorption bands in the near-infrared region (λmax ∼ 900 nm) corresponding to the optical energy gap of ca. 1.0 eV estimated from the absorption onsets (Table 1 and Figure 2b, solid line). In addition, the absorption spectrum of PNDTIBBT-DP with the alkyl groups with the far branching position showed a slight bathochromic shift relative to that of PNDTIBBT-DT, indicating that the steric effect caused by the branched side chains could affect the efficiency of π-conjugation along the polymer backbone (Figure 2b, dashed line). OFETs with the bottom-gate/top-contact configuration were fabricated with spin-coated thin films of the polymers in chlorobenzene solution (5 g L−1) on octadecyltrichlorosilane (ODTS)-treated Si/SiO 2 substrates. The devices were evaluated under ambient conditions, and the mobilities were extracted from the saturation regime. As expected from the lowC

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Figure 4. Thin-film absorption spectra of polymers under various doping conditions (a: PNDTI-BBT-DT; b: PNDTI-BBT-DP); ESR spectra (c: -DT; d: -DP), Seebeck coefficient (S), conductivity (σ), and power factors (PF) (e: -DT; f: -DP) of doped polymer films; out-of-plane XRDs of pristine and doped thin films (g) of PNDTI-BBT-DT and PNDTI-BBT-DP and S−σ plots for two polymers (h).

stacking, i.e., (010) diffraction, was observed in the out-of-plane XRD pattern of PNDTI-BBT-DT thin film (Figure 3c), implying that small face-on orientation fraction may exist in the PNDTI-BBT-DT thin film. As a result, the formation of the better crystalline order in the thin film and the edge-on structure with the lamellar motif of PNDTI-BBT-DP account for the improved performance of the thin-film OFETs. Similar improvements of mobility of OFETs by altering the side-chain structure have been reported for small molecules34−36 and polymers.37,38 As the ELUMOs of PNDTI-BBTs are as low as −4.4 eV, chemical n-type doping is expected to be operative. In fact, the doping with N,N-dimethyl-2-phenyl-2,3-dihydro-1H-benzoimidazole (N-DMBI),39,40 which is an efficient dopant for such ntype semiconductors as [6,6]-phenyl-C61-butyric acid methyl

BBT-DT and PNDTI-BBT-DP polymers are expected to have a lamellar structure with d-spacings of 2.30 and 2.65 nm, respectively, estimated from the (100) peaks in the out-ofplane XRDs. The longer d-spacing for the latter can be understood by the longer alkyl chains. On the other hand, judging from the intense (100) peak and the appearance of the (200) peak in the XRD pattern of the PNDTI-BBT-DP thin film (Figure 3d), it is likely that the crystalline nature of PNDTI-BBT-DP is better than that of PNDTI-BBT-DT in the thin-film state. Furthermore, although quite broad, a peak assignable to the π−π stacking (d = ca. 0.4 nm) designated as (010) in the in-plane XRD is observed for PNDTI-BBT-DP (Figure 3f), whereas no such peak can be detected for PNDTIBBT-DT (Figure 3e). Furthermore, although it is not very clear, a broad peak around 2θ = ca. 20° assignable to the π−π D

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thin film showed lower conductivity than the PNDTI-BBT-DP thin film, implying that the mobility in the doped PNDTI-BBTDT thin film could be significantly decreased (vide infra). We also evaluated TE properties of the doped films by imposing a temperature difference across the sample (Figure 4e,f). The Seebeck coefficient (S, blue traces) of the doped polymer thin films gave negative values, confirming that the electron transport was dominant. The absolute values of S for the PNDTI-BBT-DT thin films were lower than those for the PNDTI-BBT-DP thin films at any nominal dopant ratios. For example, S for PNDTI-BBT-DT was −56 μV K−1 at 33 mol %, whereas it was −169 μV K−1 for PNDTI-BBT-DP. By combining the electrical conductivity (σ) and S, PF was estimated according to eq 2 (Figure 4e,f, green traces). The best PF was 14 μW m−1 K−2 for PNDTI-BBT-DP at room temperature, which was significantly higher than that of PNDTI-BBT-DT, 0.6 μW m−1 K−2, and relatively high among the n-type TE materials based on semiconducting polymers recently reported.42 From the results of n-type doping behavior, electrical conductivity, and TE property of the two polymers with the same π-conjugated backbone and different alkyl side chains, the following hypothesis can be formulated: in order to design better semiconducting polymers for TE application, it is crucial to control the ordering of the polymer by tuning the alkyl side chain. In the case of the present two polymers, the higher OFET mobility of PNDTI-BBT-DP than that of PNDTI-BBTDT is primarily explained by better crystalline nature of formers in the thin films, which is in line with the widely accepted concepts that the crystalline polymers can afford high mobilities and “the side chain engineering” is a powerful tool to achieve the better crystallinity.37,38 Furthermore, higher conductivity by more than 20 times of the doped film of PNDTI-BBT-DP than that of PNDTI-BBTDT can also be explained by the ordering nature in the thinfilms; judging from the absorption and ESR spectra of the doped polymer films, PNDTI-BBT-DT is easily dopable by NDMBI, affording a high carrier concentration in the thin-film state (Figure 4a,c), whereas PNDTI-BBT-DP would resist doping, resulting in the limited doped fraction in the thin film (Figure 4b,d). The better crystalline nature of PNDTI-BBT-DP in thin-film state suppresses the penetration of N-DMBI molecules into the crystalline domains, whereas the poorly ordered PNDTI-BBT-DT in the thin-films allows the dopant molecules to be in close proximity to the polymer backbone, thereby stabilizing the charge-transfer state and resulting in a high actual doping level. This is supported by the XRD patterns of the doped thin films (Figure 4g). When the dopant ratio was increased, the lamellar d-spacing of PNDTI-BBT-DT became larger, from 2.30 nm in the pristine state to 2.73 nm after the addition of 47 mol % dopant. Furthermore, the lamellar peak became quite broad after doping, indicative of a significant decrease of the crystalline order in the doped thin films. Such a largely disordered, amorphous-like structure is likely not suitable for facile carrier transport, resulting in low mobility in the doped film and consequently low conductivity even with the high carrier concentration. In sharp contrast, the lamellar dspacing of PNDTI-BBT-DP was maintained at 2.65 nm without any obvious broadening of the lamellar peak when the nominal dopant ratio was increased from 0 to 47 mol %. Even with the relatively low carrier concentration, the highly ordered structure in the doped PNDTI-BBT-DP thin film can contribute to the high electrical conductivity.

ester (PCBM),39 C60,40 poly{N,N′-bis(2-octyl-dodecyl)-1,4,5,8naphthalene-1,4,5,8-bis(dicarboximide)-2,6-diyl]-alt-5,5′-(2,2′bithiophene)} (P(NDIOD-2T)),16 and benzodifurandionebased polyphenylenevinylene (BDPPV) derivatives,17 is quite effective (see Experimental Section). Figure 4a,b shows the vis−NIR absorption spectra of undoped and doped polymer thin films with different dopant ratios. Upon the addition of 13 mol % dopant, a new absorption band from 1400 to 1800 nm, which can be assigned as polaron and/or other doped species, such as bipolaron,41 emerged together with a neutral polymer band at 500−1200 nm for both polymer films. A further increase of the dopant ratio, however, afforded different spectra for the two polymer thin films. For PNDTI-BBT-DT, further growth of the polaron-like band was observed (Figure 4a). In sharp contrast, even at the highest dopant ratio of 47 mol %, the relative intensity of the polaron-like band to the neutral band in PNDT-BBT-DP is quite low compared to that in PNDTI-BBT-DT, indicating that the actual doping effectiveness of the two polymer thin films is markedly different (Figure 4b). These observations were qualitatively consistent with the results from the ESR measurements (Figure 4c,d), where the peak intensity is continuously increasing for PNDTI-BBT-DT (Figure 4c), whereas the intensity is saturated at the dopant ratio of 33−47% for PNDTI-BBT-DP (Figure 4d). On the other hand, it should be also noted that the absolute spin densities of two doped polymers are not significantly different from each other at lower dopant ratio (13 or 33 mol %), which is not quite consistent with the absorption spectra (Figure 4a,b), where at each nominal dopant ratio, the PNDTI-BBTDT film showed more intensive polaron-like band than the PNDTI-BBT-DP film did (Figure 4a,b). This strongly indicates the existence of spin-less species, such as bipolaron, in the doped polymer film, making difficult to directly estimate carrier concentration by the ESR measurement. Although the quantitative discussion on the carrier concentration or actual doping level is rather difficult, the doping level and thus the carrier concentration in the PNDTI-BBT-DT thin film should be always higher than that in the PNDTI-BBT-DP film, which in turn is indicative of higher electrical conductivity according to eq 3. The electrical conductivities of the polymer thin films with various dopant ratios were evaluated by the four-probe method, and the results are summarized in Figures 4e and 4f (red traces). The conductivities of neutral polymer thin films were below the measurement limit, in accordance with the fact that they were inherently closed-shell molecules. Upon doping, the conductivity rose, indicative of the generation of carrier species and consistent with the polaron-like bands observed in the absorption spectra (Figure 4a,b). However, the conductivity dependence on the nominal doping ratio (amount of dopant used relative to polymer) for the two polymer thin films was quite different. For PNDTI-BBT-DT, the conductivity almost saturated at the nominal doping ratio of 23 mol %, and the maximum conductivity of 0.18 S cm−1 was obtained at the dopant ratio of 33 mol %. In the case of PNDTI-BBT-DP, the conductivity gradually increased to 5 S cm−1 at the 33 mol % dopant ratio and then slowly decreased at much higher dopant ratios. The maximum conductivity of PNDTI-BBT-DP is quite high for n-type polymers. It should be noted that the conductivities of the doped polymer thin films were not consistent with the carrier concentrations estimated from the absorption and ESR spectra; even with the higher actual doping ratio, i.e., higher carrier concentration, the PNDTI-BBT-DT E

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identical, the difference in crystalline order in the thin films can be caused by the side chains, namely, the one carbon atom difference in branching position from the polymer backbone. It is quite surprising that such a subtle change in the side chain can cause significant differences in the electronic properties, namely, 3 times higher μe and 20 times higher conductivity and power factor. Thus, it can be concluded that tuning the intermolecular interaction and consequently obtaining the thin film with well-ordered polymers by the alkyl side chain is a promising strategy for developing superior TE materials, as it can simultaneously enhance carrier mobility and tune the doping effectiveness of conjugated polymers in the thin-film state, both of which are key parameters for determining TE properties.

Finally, higher Seebeck coefficient, S, for the PNDTI-BBTDP film than that of PNDTI-BBT-DT can also be related to the crystalline order of polymers in the thin film state. It has been reported that S can be expressed as the function of carrier number (n), namely S = C(ln Nc − ln n + A)

(4)

where n is the carrier concentration, NC is the effective DOS, and C and A represent the fitting parameter and a scattering factor of the semiconductor, respectively.43 From this equation, it is clear that the low carrier concentration is beneficial to achieving a high S. In the S−σ plots (Figure 4h), conductivity (σ) is employed instead of n because, as we already discussed, quantitative estimation of n is virtually impossible and σ should be proportional to n. In both polymers, S is proportional to ln σ in the entire conductivity regions. This is qualitatively consistent with eq 4. More importantly, the effective DOS, NC in eq 4, can be directly correlated to the density of semiconducting polymers, i.e., the number of sites in the thin film.44 As obviously observed in the XRD data of the doped polymer films (Figure 4g), the PNDTI-BBT-DP thin film with better crystalline order should have higher density of molecules in the thin film than those of PNDTI-BBT-DT with the fairly amorphous nature. The higher density of polymer molecules in the thin film can imply larger number of sites, which is consistent with higher S for PNDTI-BBT-DP than that of PNDTI-BBT-DT in all the conductivity range in Figure 4h. From these analyses of two doped polymers, it is safe to say that the better crystalline nature of the PNDTI-BBT-DP relative to that of the PNDTI-BBT-DT thin film affords all these positive effects as TE material.45 It should be pointed out that the difference in crystalline nature of the two polymers in the thin-film state is caused by a one carbon atom difference in the branched alkyl group or, more precisely, a one carbon atom difference in the branching position from the NDTI core. Although the mobility enhancement by tuning the alkyl side chain has been well-documented,34−38 the significant difference in doping effectiveness caused by the one carbon atom difference in the branched alkyl chain was initially unexpected, but it turned out to be inherently important for the enhancement of conductivity and TE properties. Note that in the present study the mobility of OFETs using neutral polymer thin films increased by ca. 3 times, whereas the conductivity and power factor of the doped polymer thin films increased by ca. 20 times. The present results for two very similar polymers suggest that the tuning of the ordering in the thin-film state of conjugated polymers by side chains is one of the promising strategies for the development of superior polymer TE materials. In summary, two new n-type semiconducting polymers composed of NDTI and BBT units with different alkyl side chains were synthesized, characterized, and evaluated as electronic materials. As both polymers had very low-lying LUMOs of ca. − 4.4 eV, they acted as versatile n-type semiconductors. PNDTI-BBT-DP with the 3-decylpentadecyl group demonstrated a high μe of 0.31 cm2 V−1s−1 in OFETs, σ of 5 S cm−1, and PF of 14.2 μW m−1 K−2 in the doped state, whereas PNDTI-BBT-DT with the 2-decyltetradecyl group showed μe of 0.096 cm2 V−1s−1, σ of 0.18 S cm−1, and PF of 0.6 μW m−1 K−2. Interestingly, the large differences in transport and TE properties can be explained by the difference in crystalline nature of PNDTI-BBTs in the thin-film state. As the π-conjugated polymer backbones of the two polymers are

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EXPERIMENTAL SECTION

Synthesis of Polymers. To a 0.5−2.0 mL microwave-pressurized vial equipped with a stirring bar, 2,7-bis(tributylstannyl)-NDTI-DT22 or -NDTI-DP (0.05 mmol), 4,8-dibromo-BBT33 (0.05 mmol), tris(dibenzylideneacetone)dipalladium(0) (0.8 mg, 0.9 μmol), copper iodide (0.57 mg, 3.0 μmol), triphenylarsine (1.1 mg, 3.5 μmol), and toluene (1.8 mL) were added. Then, the vial was sealed, refilled with argon, placed inside a microwave reactor, and heated at 180 °C for 60 min. After cooling to room temperature, the reaction mixture was poured into methanol (50 mL) containing hydrochloric acid (1 M, 1 mL) and stirred for 6 h. The resulting precipitate was collected by filtration and subjected to sequential Soxhlet extraction with methanol, hexane, and dichloromethane. The residue was extracted with chloroform, and the concentrated fraction was precipitated in 50 mL of methanol to yield PNDTI-BBTs; PNDTI-BBT-DT: 66% yield. Anal. Calcd for (C72H100N6O4S4)n: C, 69.63; H, 8.12; N, 6.77. Found: C, 69.05; H, 8.15; N, 6.42. 1H NMR (400 MHz, 1,1,2,2tetrachloroethane-d2): δ 1.9−0.8 (broad multiplet), owing to strong aggregation, only the aliphatic protons assignable to branched alkyl groups were observed. PNDTI-BBT-DP: 68% yield. Anal. Calcd for (C74H104N6O4S4)n: C, 69.99; H, 8.25; N, 6.62. Found: C, 69.65 H, 8.38; N, 6.14. 1H NMR (400 MHz, 1,1,2,2-tetrachloroethane-d2): δ 1.9−0.8 (broad multiplet), owing to strong aggregation, only the aliphatic protons assignable to branched alkyl groups were observed. Fabrication and Characterization of OFETs. Top-contact/ bottom-gate OFET devices were fabricated on heavily doped n+-Si (100) wafers with 200 nm thick thermally grown SiO2 (Ci = 17.3 nF cm−2). The cleaned substrates were modified with octadecyltrichlorosilane (ODTS) to form a self-assembled monolayer (SAM). Thin films of the polymers were deposited on the SAM-modified substrate by spin-coating chlorobenzene solution (5 g L−1) in a glovebox, optionally followed by thermal annealing at 100, 150, 250, 270, or 300 °C in the glovebox. On top of the polymer thin film, ca. 80 nm thick gold was deposited as source and drain contacts using a shadow mask. The OFET devices had a channel length (L) of 40 μm and a channel width (W) of 1.5 mm. Current−voltage characteristics of the OFET devices were measured at room temperature in air with a Keithley 4200-SCS semiconductor characterization system. Field-effect mobilities were calculated in the saturation regime (Vd = Vg = 60 V) using the equation

Id = CiμFET

W (Vg − Vth)2 2L

(5)

where Id is the drain current in the saturated regime, W and L are the semiconductor channel width and length, respectively, Ci is the capacitance per unit area of the gate dielectric layer, and Vg and Vth are the gate voltage and the threshold voltage, respectively. Vg − Vth of the devices was determined from the square root values of Id in the saturation regime. Current on/off ratios (Ion/Ioff) were determined from the minimum current at approximately Vg = −20 to 0 V (Ioff) and the current at Vg = 60 V (Ion). Doping of polymers and evaluation of conductivity and thermoelectric properties: Chloroform solutions of polymers and (4-(1,3F

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dimethyl-2,3-dihydro-1H-benzoimidazol-2-yl)phenyl)dimethylamine (N-DMBI, 3 g L−1) were prepared separately.34 Aliquots of N-DMBI and polymer solutions were mixed at room temperature, and then the thin films were fabricated by spin-coating or drop-casting on glass substrates and annealed at 120 °C for 8 h. Four-probe conductivity measurements were carried out in air with an MCP-T610 (Mitsubishi Chemical Analytech). For Seebeck coefficient measurements, all devices were fabricated on PEN substrates. The polymer thin films were deposited on the PEN substrate and then pasted on the sample stage. Finally, by imposing a temperature difference across the sample around room temperature (for PNDTI-BBT-DT, ΔT ≈ 1 K; for PNDTI-BBT-DP, ΔT ≈ 0.5 K; the temperatures were measured with an error less than 0.01 K46) and measuring the thermovoltages, the Seebeck coefficients were calculated with the equation S = ΔV /ΔT

(6)

(7)

where a is the peak intensity and ΔHpp is the line width.

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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.6b02313. Figures S1−S7, Scheme S1, and Table S1 (PDF)

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REFERENCES

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where ΔV is the thermal voltage obtained between the two ends of the device subjected to temperature gradient ΔT. Physicochemical Studies. UV−vis absorption spectra were measured using a Shimadzu UV-3600 spectrometer. Cyclic voltammograms (CVs) were recorded on an ALS Electrochemical Analyzer Model 612D in benzonitrile containing tetrabutylammonium hexafluorophosphate (Bu4NPF6, 0.1 M) as supporting electrolyte at a scan rate of 100 mV/s. Counter and working electrodes were made of Pt, and the reference electrode was Ag/AgCl. All the potentials were calibrated with the standard ferrocene/ferrocenium redox couple (Fc/ Fc+: E1/2 = +0.49 V measured under identical conditions). AFM images were obtained on a Nanotechnology, Inc., scanning probe microscope Nanocute system. X-ray diffractions of thin films deposited on the Si/SiO2 substrate after annealing (PNDTI-BBT-DT: 270 °C; PNDTI-BBT-DP: 300 °C) were obtained with a Cu Kα source (λ = 1.541 Å) in the air using Ultima IV (Rigaku Corporation). ESR measurements were carried out for thin films on PEN substitutes using an X-band spectrometer (JEOL JES-TE100). The intensity was standardized by the polymer weight. The magnetic susceptibility is calculated with the following equation:

χ = aΔHpp2

Article

AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected] (M.N.). *E-mail [email protected] (K.T.). ORCID

Masahiro Nakano: 0000-0002-9231-4124 Tsuyoshi Michinobu: 0000-0001-6948-1189 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by JSPS KAKENHI Grants 15H02196 and 16K05900 and Iketani Science and Technology Foundation. HRMS measurements were carried out at the Molecular Structure Characterization Unit, RIKEN Center for Sustainable Resource Science (CSRS). DFT calculations using Gaussian 09 were performed by using the RIKEN Integrated Cluster of Clusters (RICC). G

DOI: 10.1021/acs.macromol.6b02313 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

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DOI: 10.1021/acs.macromol.6b02313 Macromolecules XXXX, XXX, XXX−XXX