Enhancing Long-Range Ordering of P3HT by ... - ACS Publications

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Enhancing Long-Range Ordering of P3HT by Incorporating Thermotropic Biphenyl Mesogens via ATRP Taniya M. S. K. Pathiranage, Minkyung Kim, Hien Q. Nguyen, Katherine E. Washington, Michael C. Biewer, and Mihaela C. Stefan* Department of Chemistry and Biochemistry, University of Texas at Dallas, Richardson, Texas 75080, United States S Supporting Information *

ABSTRACT: The methoxybiphenyl liquid crystalline mesogen 6-((4′-methoxy-[1,1′-biphenyl]-4-yl)oxy)hexyl methacrylate (MeOBP) was incorporated into block copolymers with regioregular poly(3-hexylthiophene) (P3HT) by a combination of Grignard metathesis (GRIM) and atom transfer radical polymerization (ATRP) techniques. Side chain engineering with a thermotropic liquid crystalline segment produced an improved morphology resulting in field-effect mobilities of ∼10−2 cm2/(V s) measured in organic thin film transistors (OTFTs). It was observed that improved π−π stacking distances between edge-on oriented polythiophene segments upon annealing facilitated charge transport which was accomplished with π−π stacked biphenyl mesogens with P3HT block. Variation of surface morphology of polymer thin films upon annealing was investigated by tapping mode atomic force microscopy (TMAFM). A combined study of differential scanning calorimetry (DSC) and polarized optical microscopy (POM) was also performed to identify liquid crystalline mesophase transitions of the synthesized block copolymers in response to the temperature.

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INTRODUCTION Organic polymeric semiconductors can be used as flexible, lightweight, and cost-effective materials in applications such as organic field effect transistors (OFET),1−3 organic photovoltaics (OPV),4−6 sensors,7 and organic light-emitting diodes8 (OLED). In fabricating flexible large area optoelectronics at low cost, organic semiconductors have a high potential due to their good solution processability and relatively good mechanical properties. Furthermore, these properties can become key factors in device printing for certain applications. The structure and morphology of the active layer are crucial for charge separation and charge transport in semiconducting polymers and influence the device performance. Because of the interconnected crystalline nanodomains consisting of wellaligned nanofibrils and π stacked conjugated polythiophene backbones, regioregular poly(3-hexylthiophene) (P3HT) displays relatively high hole mobilities in OFETS9−14 with good environmental and thermal stability. Grignard metathesis polymerization (GRIM) reported by McCullough’s group in 1999 was employed for the synthesis of semiconducting regioregular P3HT block copolymers.15 The same group reported the synthesis and morphological characterization of block copolymers containing regioregular P3HT.1,3,16−19 End functionalization20,21 and a series of end group modifications have been performed to obtain precursors for the synthesis of P3HT block copolymers.22−27 In addition, the livingness of the GRIM method allowed the synthesis of regioregular P3HT with well-defined molecular weights and end groups.20,21,28−30 © XXXX American Chemical Society

ATRP is a living polymerization technique that allows the synthesis of well-defined block copolymers from various vinyl monomers. ATRP can be employed to synthesize P3HT block copolymers with well-defined composition and molar weights in a one-pot synthesis between the P3HT macroinitiator and the liquid crystalline (LC) monomers providing the advantage of tuning the properties of the resulting polymers by varying the molar ratio between P3HT and the liquid crystalline block. Liquid crystalline materials can undergo molecular reorientation with various external stimuli such as electric or magnetic field, temperature, and shear stress. In the side chain liquid crystalline polymers (SCLCP) with thermotropic liquid crystalline mesogens, the connection between the mesogen and the polymer backbone is made by a side chain spacer group. When the semiconducting P3HT is covalently attached to a second block of thermotropic liquid crystalline mesogens, highly improved morphologies can be obtained by cocrystallization upon annealing above the thermotropic mesophase transition. During this process, P3HT can rearrange into welldefined microphase-separated domains with long-range ordering due to the supramolecular assembly within the matrix of highly organized mesogens, resulting in improved optical and electronic properties. The temperature triggered control of the orientation of the polymer nanostructures allows applications for thermotropic SCLCP of P3HT in flexible thin film Received: June 28, 2016 Revised: September 5, 2016

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DOI: 10.1021/acs.macromol.6b01378 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Scheme 1. Synthesis of Block Copolymers of P3HT by ATRP

Table 1. Molecular Weights, UV−Vis Absorbance, and HOMO/LUMO Energy Levels

a

polymer

Mna (g mol−1)

PDI

λmaxb (nm)

P3HT-Br P1 Br-P3HT-Br P2 P3

13000 25600 8500 12700 6000

1.6 1.7 1.7 1.8 2.9

450 282, 442 451 282, 446 285

λmaxc (nm) 530, 331, 526, 318, 290

optical band gap (eV)

HOMOd (eV)

LUMOe (eV)

band gap (eV)

1.90 1.78 1.90 1.81

−5.12 −5.19 −4.97 −5.47

−3.15 −3.46 −3.16 −3.70

1.88 1.73 1.81 1.77

604 560, 587, 642 603 539, 575, 625

Determined by SEC with polystyrene calibration (THF eluent). bAbsorption of chloroform solution. cAbsorption of annealed thin films at 150 °C. Estimated from the onset of oxidation wave of cyclic voltammogram. eEstimated from the onset of reduction wave of cyclic voltammogram.

d

properties by tuning the molecular weight and composition. The field-effect mobilities are correlated with the variation of surface morphology upon application of heat to trigger the mesophase changes. Liquid crystalline properties are investigated in detail in a combined study of X-ray diffraction analysis (XRD) and polarized optical microscopy (POM).

transistors and sensors. Flexible side chains attached to the conjugated backbone support facile deposition, improve flexibility, and increase the mobilities. Methoxy biphenyl (MeOBP) was selected as the mesogen for this study due to its thermotropic behavior with high thermal stability and lamellae forming behavior.31−34 In previous studies of homopolymers and copolymers of MeOBP, lamellar stacking has been observed in both crystalline and smectic phases. The ability of the biphenyl moiety to undergo π stacking in the crystalline and smectic phases enables the formation of welldefined microdomains.35 It has been shown that lamellar and π stacking in both P3HT and LC block improves the compatibility and improved the microdomain morphology of P3HT with the formation of distinct nanowires. Among the recent research on P3HT-based thermotropic liquid crystalline block copolymers,1,3,36−39 Chen’s group has reported a block copolymer containing discotic mesogens which was used as an additive in bulk heterojunction solar cells.40 Our group recently reported a block copolymer of P3HT with methacrylate attached azobenzene that underwent a nematic phase transition. Switching between the cis−trans configurations of azobenzene moiety with response to heat and UV−vis radiation together with thermotropic liquid crystalline phase behavior triggered the alignment of the P3HT block in this material which has potential applications in thermal actuators and OFETs.3 In this study, we report the synthesis and characterization of two block copolymers of P3HT with MeOBP mesogen where we demonstrate both liquid crystalline and semiconducting

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RESULTS AND DISCUSSION

Synthesis of Block Copolymers. The synthesis of P3HT block copolymers with biphenyl attached methacrylate (MeOBP) was accomplished by ATRP as shown in Scheme 1. In situ end-capping of nickel-terminated P3HT with allylmagnesium bromide and hydroboration−oxidation of the allyl-terminated P3HT generated the hydroxypropyl-terminated P3HT. This was converted into monofunctionalized ATRP macroinitiator (P3HT-Br) [1]. In the same way, difunctionalized ATRP macroinitiator (Br-P3HT-Br) [2] was synthesized from difunctionalized hydroxypropyl-terminated P3HT precursor. Dialdehyde CHO/CHO-P3HT was obtained from H/ H-P3HT in a Vilsmeier reaction. Further reduction with LiAlH4 generated the difunctionalized hydroxypropyl-terminated P3HT precursor. The conversion of end-capped P3HT from hydroxypropyl to bromoester was monitored by 1H NMR analysis (Figures S2 and S3 in the Supporting Information). The triplet corresponding to the methylene protons (3.78 ppm) adjacent to the hydroxyl group completely disappears concomitant with the appearance of a singlet at 1.96 ppm that corresponds to the six methyl protons of the bromoester P3HT. Both P1 and P2 B


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Figure 1. UV−vis absorption spectra of copolymers P1 and P2 in solution and films (nonannealed and annealed at 150 °C).

Table 2. Field-Effect Mobilities of the Synthesized Copolymers Measured in OFETs polymer P1

P2

a

device untreated OTS treated FTS treated untreated OTS treated FTS treated

av VT (V) 13.5 −12.2 −17.7 18.9 15.0 11.0

± ± ± ± ± ±

Ion/Ioff

3.6 0.7 1.0 2.4 3.6 2.2

highest mobility (cm2/V s)

3

3.81 7.95 6.31 8.73 3.46 6.58

10 103 103 101 101 101

× × × × × ×

−3

10 10−3 10−2 10−4 10−3 10−3

average mobilitya (cm2/V s) 1.24 × 10−3 7.89 × 10−3 2.35 × 10−2 7.39 × 10−4 1.25 × 10−3 4.68 × 10−3

Measured for 10 devices with the same channel length of 20 μm.

displayed broad peaks at 6.8 and 7.4 ppm due to the presence of the aromatic protons of the biphenyl units (Figures S5 and S6). Protons that are ortho to the methoxy group and oxygenterminated hexyl group gave the upfield peak at 6.8 ppm while the remaining meta protons gave the more downfield peak at 7.4 ppm. Copolymer P1 contained ∼20 mol % of PMeOBP and P2 contained ∼30 mol % as determined by 1H NMR by integrating the peak corresponding to the methylene protons adjacent to the thiophene ring of the P3HT block versus the peak corresponding to the aromatic protons of phenyl rings (1H NMR spectra given in the Supporting Information under Figures S5 and S6). Incorporation of the second block was further confirmed by size exclusion chromatography (SEC) measurements, which showed an increase in molecular weight as compared to the P3HT macroinitiators (Table 1 and Figure S20). A homopolymer of MeOBP (P3) was also synthesized as shown in Scheme 1 to observe and compare the mesophases with P1 and P2. The observed PDI was 2.9 for P3, which is due to slow polymerization caused by bulky nature of the monomer. Characterization. UV−Vis Analysis. The UV−vis analysis was performed in solution and in thin films for both P3HT precursors (Figure 1; Figures S8 and S9) and block copolymers. Absorption maxima in chloroform for P1 (442 nm) and P2 (446 nm) correspond to the P3HT segment while absorption maximum at 282 nm corresponds to the methoxy biphenyl chromophore. Both in solution (285 nm) and film (290 nm) P3 gives a peak for π stacked methoxy biphenyl chromophores. In the thin film UV−vis spectra of P1 and P2, peaks at 331 and 318 nm, respectively, correspond to MeOBP while peaks at 560 nm (P1) and 539 nm (P2) correspond to P3HT π−π* transition (Figure 1). A red-shift of ∼80 nm was observed between the spectra for solution and film which can be attributed to the increased effective conjugation with 2D packing of polythiophene chains within lamellae and close π stacking between them upon crystallization. Vibronic peak

appearance as shoulder peaks around 600 nm in P1 and P2 films was observed upon annealing41 (Figure 1). This indicates better ordering of polymer chains with extended planarity leading to longer effective conjugation with less torsion. The HOMO and LUMO energy levels were determined from the onset potential values of the oxidation and reduction peaks that correspond to P3HT (which are assumed to correspond to the longest π conjugated system) by cyclic voltammetry (Table 1). The multiple small oxidations that were observed prior to the oxidation of P3HT are due to the biphenyl mesogens attached to the side chains of the second block.42 Both copolymers have comparable electrochemical band gaps with the P3HT macroinitiators indicating good optoelectronic properties. The higher PDI (1.6−1.8) observed for block copolymers is due to partial end-capping of P3HT chains during GRIM resulting with end-capped (∼80%) and non-end-capped P3HT chains (∼20%) (Figure S21). Since only the end-capped fraction can undergo conversion into the block copolymer, we assume the different interactions of non-end-capped P3HT and P3HT bromoesters with the SEC column make the peak broader and the PDI higher. The PDI values remaining constant during block copolymer synthesis via ATRP indicated well-controlled polymerization. Field-Effect Mobility. The field-effect mobilities of the synthesized block copolymers were measured in bottom-gate bottom-contact device configuration using the equation μ=

IDS 2L WCi (VGS − VT)2

where IDS is the source−drain current, W is the channel width, L is the channel length, Ci is the capacitance of the dielectric, VGS is the gate voltage, and VT is the threshold voltage. Hole mobility was obtained for untreated devices and also for devices treated with octadecyltrimethoxysilane (OTS) and (heptaC


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Figure 2. Current−voltage characteristics of P1 on FTS treated and annealed OTFT: (a) transfer curve at VDS = −100 V (W = 475 μm, L = 20 μm); (b) output curves at different gate voltages.

Figure 3. Current−voltage characteristics of P1 on OTS treated and annealed OTFT: (a) transfer curve at VDS = −100 V (W = 475 μm, L = 20 μm); (b) output curves at different gate voltages.

separated with insulating methacrylates on periphery interrupts long-range charge transport within the channel region of OFETs in P2. In contrast, P1 with high molecular weight P3HT precursor with extended π−π stacking and rigid planar backbone can efficiently transport charges, through the backbone and hopping between lamellae. Annealing the polymer films above the mesophase transition temperature results in microphase separation, and it has a positive effect on charge transport by hopping between the P3HT lamellae due to reduced π−π staking distance. Formation of distinct nanofibrils after microphase separation observed in TMAFM and TEM images (Figures 4 and 6) implies the formation of edge on lamella of the block copolymers within the nanofibrils. This similar morphological behavior of both block copolymers to pristine P3HT explains their comparable mobility values. Charge transport appears to be efficient in P1 due to the possibility of connecting the adjacent nanodomains with long P3HT tie chains within the matrix of mesogen bound methacrylates.45 This allows extended carrier transport pathways despite the presence of insulating liquid crystalline block. 3.4. Surface Morphology. Morphology of the polymer films was investigated by tapping mode atomic force microscopy (TMAFM). Thin films were formed by drop-casting polymer solutions in chlorobenzene on mica substrate and OFET devices followed by slow evaporation which resulted in distinct

decafluoro-1,1,2,2-tetrahydrodecyl)trimethoxysilane (FTS). The average field-effect mobilities measured for P1 to P2 on untreated devices were in the range of 10−3 and 10−4 cm2 V−1 s−1, respectively, with the highest value of 3.81 × 10−3 cm2 V−1 s−1 for P1 (Table 2). P1 displayed average field effect mobility of same order on OTS-modified devices but displayed 1 order higher hole mobility on FTS treated devices with highest measured mobility of 6.31 × 10−2 cm2 V−1 s−1. Reported fieldeffect mobilities for various P3HT block copolymers vary in the range from 10−2 to 10−5 cm2 V−1 s−1. The relatively high mobilities obtained with OFETs with modified substrates can be attributed to the self-assembly of P3HT chains with high degree of edge-on orientation upon annealing due to the favorable interactions with the hydrophobic substrate. Field-effect mobility of the diblock copolymer P1 is higher than the triblock copolymer P2 in all the conditions which can be related to composition of the polymer and morphology. In P3HT the π stacked crystalline nanodomains are interconnected by polymer chains which ensure the charge transport between the domains minimizing the charge trap at the grain boundaries.43,44 In P2, despite the high crystallinity, shorter nanofibrils (Figure 4) formed with low molecular weight P3HT precursor and the presence of the insulating liquid crystalline segments on both ends may adversely affect the charge transport compared to P1. Isolated P3HT nanodomains D


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Figure 4. TMAFM phase images (3 μm × 3 μm) of P1: (a) without surface treatment, (b) with OTS treatment, and (c) with FTS treatment on OFET channel regions. TMAFM phase images (3 μm × 3 μm) of P2: (d) without surface treatment, (e) with OTS treatment, and (f) with FTS treatment on OFET channel regions.

Figure 5. 3-D TMAFM images (3 μm × 3 μm) of P1 obtained on mica (a) before annealing and (c) after annealing at 150 °C and 3-D TMAFM images of P2 on mica (b) before annealing and (d) after annealing at 150 °C.

helps the supramolecular organization of the highly ordered, crystalline long-range assemblies of the P3HT chains. On OFET device surfaces, better surface morphology was observed on treated devices compared to nontreated (Figure 4). Nanofibrils of P3HT can be clearly observed in treated devices which can be correlated with their hole mobility values. In devices of P1 (both OTS and FTS treated) the long nanofibrils indicate the ordered packing of the polymer that extends longer within the channel region of the device, which

nanostructured morphology of P1 and P2. Chlorobenzene was used as the solvent, since slow evaporation provides sufficient time to allow for an effective packing of polymer chains before evaporation which results in densely packed P3HT nanofibrils in solid films with edge-on orientation with long-range crystallinity.41,46 The devices were annealed above 150 °C for 20 min to allow removal of chlorobenzene and to trigger edgeon orientation of polymer chains which is thermodynamically favored. Annealing the devices above the mesophase transition E


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Figure 6. TEM images of self-assembled nanowires of (a) P1 and (b) P2 obtained after drop-casting 1 mg/mL polymer solutions in chlorobenzene on the copper grid.

effectively connects the source and drain resulting in better charge transport. This further explains why the short chain nanofibrils in the devices of P2 correlate with relatively lower mobility than P1. van der Waals interactions between the alkyl chains of both conducting and nonconducting blocks and the silane monolayer allow better anchoring of the ordered polymer chains on the treated hydrophobic surface creating better morphological features. By contrast, polymer films deposited on mica surface displays a ribbon like morphology in nanoscale with uniformly distributed microphase separation on a wide area (Figure 5 and Figure S15). Thus, P2 from low molecular weight P3HT precursor displays narrow and distinct nanofibrils compared to the wider nanofibrils obtained for P1 with high molecular weight P3HT. The increase of the molecular weight introduces granular surface features, diminishing the nanowire morphology due to folding of polymer chain rather than extending to its full length.47,48 This change in morphology also attributed to difference in surface interactions. These highly oriented crystalline nanostructures become more evident upon annealing above the mesophase transition indicating the positive effect of biphenyl mesogens on P3HT self-assembly in the long-range ordered crystalline nanodomains (Figures 4 and 5, Figure S15). The self-assembly of both P1 and P2 was further confirmed by transmission electron microscopy (TEM) analysis. The formation of distinct nanowires in both P1 and P2 (Figure 6) upon drop-casting the polymers in chlorobenzene indicates the effective packing of the polymer within the film. X-ray Diffraction Studies. Thin film XRD analysis was performed for P1, P2, P3HT precursors, and P3 to further investigate crystallinity of polymer films (Figures 7 and Table 3). XRD data obtained for both P1 and P2 display similar pattern to pristine P3HT indicating the lamellar packing and πstacking of semiconducting P3HT block in the block copolymers was not adversely affected by the addition of the LC block. The sharp intense peak at 5.32° (d = 16.60 Å) for P1 and P2 corresponds to lamellar stacking of P3HT in (100) plane in the direction of hexyl chains. In addition, small broad peaks at 10.79° (d = 8.19 Å) and 16.23° (d = 5.46 Å) in P1 and 10.85° (d = 8.15 Å) and 16.55° (d = 5.35 Å) in P2 corresponds to (200) and (300) crystal planes, respectively, which are higher order reflections of lamellae packing indicating the extended long-range ordering within the crystalline thin films. When correlating the polymer composition to the ordered packing in the condensed matte of both P1 and P2, the consistency of these results with pristine P3HT indicates the

Figure 7. Thin film XRD spectra of P1, P2, and P3 on SiO2 substrate deposited from chlorobenzene by drop-casting.

Table 3. XRD Data Obtained for P1, P2, and P3 on SiO2 Surface polymer

2θ (deg)

d spacing (Å)

P1

5.32 10.79 16.23 20.44 23.98 5.33 10.85 16.55 20.37 23.69 19.67

16.60 8.19 5.46 4.34 3.71 16.57 8.15 5.35 4.36 3.75 4.51

P2

P3

packing (100) lamellar stacking of (200) lamellar stacking of (300) lamellar stacking of π−π stacking of biphenyl π−π stacking of P3HT (100) lamellar stacking of (200) lamellar stacking of (300) lamellar stacking of π−π stacking of biphenyl π−π stacking of P3HT π−π stacking of biphenyl

P3HT P3HT P3HT

P3HT P3HT P3HT

unaltered assemblies of the P3HT block. The π−π stacking perpendicular to the rigid conjugated thiophene backbone in (010) plane can be identified with the small broad peaks at 23.98° (d = 3.71 Å) for P1 and 23.69° (d = 3.75 Å) for P2. This d spacing is smaller than the value for pristine P3HT (d = 3.8 Å) indicating the dense packing of the polythiophene segment in copolymers in edge-on orientation which can correlate with the morphological properties observed on TMAFM and TEM (Figures 4 and 6). The appearance of peak at 20.44° (d = 4.34 Å) for P1 and 20.37° (d = 4.36 Å) for P2 was identified as the F


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Figure 8. DSC thermograms for second cycle of heating (left) and second cycle of cooling (right) for polymers P1−P3.

peak from π−π stacking of biphenyl mesogens, corresponding to the XRD data obtained for P3. The corresponding peak for P3 appears at 19.67° (d = 4.51 Å). According to the data, it appears the π−π stacking distances of thiophene subunits and biphenyl mesogens is lower in the block copolymers compared to the homopolymers, indicating the copolymers are more closely self-assembled upon cocrystallization while preserving the inherent ordering and properties of the two individual blocks. The π−π stacking of both blocks are closer in P1 than in P2 due to the high molecular weight of P3HT macroinitiator in P1 which creates close extended packing resulting with higher hole mobility. The observed edge-on orientation enables fast charge transport along the rigid planar backbone and in the π−π stacking direction through hopping which is highly desirable for OFETs to ensure efficient in-plane charge transport within the channel region between the source and drain.49,50 Liquid Crystallinity. The thermotropic phase behavior and the liquid crystalline mesophase changes for the synthesized polymers were investigated in a combined study of differential scanning calorimetry (DSC) and polarized optical microscopy (POM). All the measurements for melting, glass transition, and liquid crystalline temperatures were obtained from the second heating and the second cooling curves in DSC. The smectic MeOBP is expected to induce increased ordering than the incorporation of a nematic mesogen due to the ability to π stack based on our previous work.3 MeOBP is a thermotropic, smectic liquid-crystalline mesogen that exhibits the lamellae forming ability in the crystalline state and π stacking observed in both crystalline and smectic mesophase which make it a better choice for block copolymer synthesis with P3HT.35,50 In the heating cycles of P1 and P2 (Figure 8), clear firstorder endothermic transitions can be observed for the melting of P3HT in P1 (207.5 °C) and P2 (212.5 °C). The appearance of the shoulder peak in the melting of P1 and the appearance of two distinct adjacent peaks for melting in P2 shows a eutectic behavior of the polymers. This could be due to the presence P3HT chains with no addition of the second block (Figures S20 and S21). Because of the large content of P3HT in the

copolymers, they retain the crystalline nature with the ability to form lamellae, which can be clearly observed in the cooling curves with prominent first-order endothermic crystallization peaks. Glass transition of P3HT in P1 (52.7 °C) and P2 (56.2 °C) appear as less prominent second-order transitions. Liquid crystalline mesophase changes require very small exchange of heat which results in weak transitions both on heating and cooling traces. Liquid crystalline mesophase transitions between 80 and 160 °C can be observed in heating both P1 (101.3 and 133.8 °C) and P2 (128.6 °C). Upon cooling mesophase transitions are visible between 100 and 140 °C. Slow relaxation dynamics and high crystallinity below the glass transition temperature make polymers a more rigid material. The polymer chains become flexible above the glass transition temperature, which make it easier to rearrange in an orderly manner with an external force which could be shear stress, temperature, magnetic field, etc. Since the mesophases lie in between the glass transition and melting of P3HT, the flexibility of the P3HT segment within the mesophase results in organized supramolecular assembly with biphenyl mesogens, which function as an organizing grid for the mesophase. Because of negligible relaxation dynamics of the polymer chains, this packing is maintained in solid films upon cooling back to the crystalline state as demonstrated in XRD and UV−vis. LC block copolymers exist in either microdomain structure or liquid crystalline phase structure. When it comes to the ordering, the mesogen orientation during a phase transition occurs within the microdomains. Since both P1 and P2 contain P3HT as the dominant component, the ordering of MeOBP is expected to take place within the domains of P3HT. With XRD and UV−vis studies, it was found that MeOBP does not perturb the packing arrangement of P3HT but rather enhance it. During the mesophase transition the side-chain mesogens form the smectic layer with uniaxial orientation perpendicular to the interface which supports the edge-on orientation of the semiconducting segment. According to the previous work, MeOBP forms a layered structure with edge-on lamellae G


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Figure 9. Polarized optical micrographs of P1 during smectic phase (110−115 °C) (a) under mid power and (b) under high power. Polarized optical micrographs of P2 during smectic phase (110−115 °C) (c) under mid power and (d) under high power.

Figure 10. Polarized optical micrographs of MeOBP monomer (a) before heating up to mesophase transition and (b) after gaining smectic phase (60 °C). Polarized optical micrographs of polyMeOBP (P3) (c) before heating up to mesophase transition and (d) after gaining smectic phase (63.8 °C).

formation in both crystal and a smectic phase35 which proved to be beneficial in this study. The thin films of polymers were heated to 150 °C (slightly above isotropic temperature) on the heating stage followed by cooling to 110−115 °C until the smectic phase become visible. The polarized optical micrographs of P1 and P2 were obtained at 400× magnification under crossed polarizers (Figure 9). The

temperature ranges correspond to mesophase changes for both P1 and P2 with the heating stage were consistent with the results obtained from DSC thermograms. Upon further cooling of both P1 and P2, formation of smectic features was observed in the temperature range of 110−115 °C, thus indicating the MeOBP block was able to introduce liquid crystalline properties to both P1 and P2. Thermotropic poly(MeOBP) H


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Macromolecules

temperature ramp, 10 °C min−1; final temperature, 280 °C. A Hewlett-Packard fused silica capillary column cross-linked with 5% phenylmethylsiloxane was used with helium as the carrier gas (1 mL min−1). Number-average molecular weight (Mn) and polydispersity index (PDI) were determined with size exclusion chromatography (SEC) on a Viscotek VE 3580 separation module equipped with ViscoGEL columns (GMHHR-M), connected to a refractive index (RI) detector. SEC analysis was done under following conditions: flow rate = 1.0 mL min−1, injector volume = 100 mL, detector temperature = 30 °C, column temperature = 35 °C. HPLC grade THF was used as the eluent with calibration based on polystyrene standards. DSC thermograms were obtained with a Q100 differential scanning calorimeter (TA Instruments) using hermetically sealed aluminum sample holders. Temperature programming was done for three cycles from 20 to 250 °C at the rate of 10 °C/min. Cyclic voltammetry measurements were obtained using a BAS CV50W voltametric analyzer (Bioanalytical Systems, Inc.). The electrochemical cell was composed of a platinum electrode, a platinum wire auxiliary electrode, and an Ag/AgCl reference electrode with acetonitrile (0.1 M TBAP) purged with argon as the electrolyte. All electrochemical shifts were standardized to the ferrocene redox couple at 0.471 V. Evaluation of HOMO and LUMO energy levels was obtained by using the following equations: HOMO (eV) = −e(Eox + 4.71) and LUMO (eV) = −e(Ered + 4.71), where Eox and Ered are the measured potentials relative to Ag/AgCl+. Thin film X-ray diffraction study was performed on a RIGAKU Ultima III diffractometer with Cu Kα (λ = 1.54 Å) as the radiation source from 1° to 40° (2θ) at 0.04° intervals at a rate of 2°/min. Polymer thin films on clean SiO2 substrates were obtained by drop-casting polymer solutions (5.0 mg mL−1) in chlorobenzene. Both solution (in chloroform) and solid state UV−vis spectra of polymers were recorded using an Agilent 8453 UV−vis spectroscopy system. Annealing was done at 120 °C in this study. Transmission electron microscopy (TEM) imaging of thin films of P1 and P2 was performed on a Tecnai G2 Spirit Biotwin microscope by FEI, and images were analyzed using ImageJ software. Samples were prepared by drop-casting 1 mg/mL polymer solutions in chlorobenzene on copper mesh grid. Polymer samples for matrixassisted laser desorption ionization−time of flight mass spectrometry (MALDI-TOF) analysis were prepared by combining the matrix and the analyte in a 2:1 volume ratio to spot the sample plate. The block copolymer samples were dissolved in chloroform to yield solutions in 1 mg/mL concentration. Terthiophene was used as the matrix in a 20 mg/mL concentration solution in chloroform. Field-Effect Transistor Fabrication and Measurements. Fieldeffect mobility measurements were performed on thin-film transistors with bottom-gate bottom-contact configuration. Highly doped, n-type silicon wafers (0.001−0.003 Ω cm) with a 200 nm thick thermal oxide (SiO2) were used as the substrate. Chromium (5 nm) followed by gold (100 nm) was deposited by E-beam evaporation as metal contacts. The source−drain pads were patterned by photolithography. To generate the common bottom-gate backside of the wafer was etched with buffered oxide etchant (7:1 BOE from JT Baker). The resulting transistors had a channel width of 475 μm and channel length ranging from 5 to 80 μm. The polymer films were deposited by drop-casting of polymer solutions in chlorobenzene (1 mg mL−1). A Keithley 4200SCS semiconductor characterization system was used to probe the devices. The probe station used for electrical characterization was a Cascade Microtech Model Summit Microchamber. All the measurements were performed at room temperature in air. For the surface treatments with OTS and FTS, the devices were sequentially rinsed with water, acetone, hexanes, and chloroform and placed in a sealed container in a solution of OTS/FTS of 8 × 10−3 M in distilled toluene for 48 h. TMAFM studies of the thin film morphology were carried out using a Nanoscope IV-Multimode Veeco, equipped with an E-type vertical engage scanner. The AFM measurements were performed both on the channel region of OFET devices that were measured for field-effect mobilities and thin films on a mica substrate. Thin films were obtained by drop-casting 1,2,4-trichlorobenzene (TCB) solution of polymers (1 mg mL−1) on mica. The AFM images were recorded at room

block helps tuning alignment of nanofibrils of P3HT into highly ordered crystalline domains upon cocrystallization after annealing above crystalline to smectic transition. According to the reported XRD data for P3, the smectic layer thickness is 26 Å,35,50 which is higher than the lamellae thickness of P3HT segment (∼16.5 Å). The appearance of the small peak at 3.3° (d = 26 Å) in XRD spectra of P1 indicate the lamellae packing of MeOBP and explains the higher mobility of P1 over P2 obtained with higher organization of both blocks. With the reported XRD data of P350 during the transition from crystalline state to smectic phase, the peak for the π stacking remains while the peak for lamellae packing disappears explaining the distortion of the layered structure. Broadening of interlamellar spacing of MeOBP upon heating cause perturbation of the lamellae in LC block. The micrographs of the P3 (Figure 10) together with DSC data (Figure 8) indicate the formation of characteristic smectic batonnets upon annealing to the mesophase temperature. Before annealing, the film appears less organized with scattered spherical domains.

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CONCLUSIONS

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EXPERIMENTAL SECTION

We have successfully synthesized two new thermotropic side chain liquid crystalline block copolymers of P3HT and PolyMeOBP by a combination of GRIM and ATRP living polymerization methods. Thermotropic liquid crystalline smectic mesophase transitions of the polymers were investigated and confirmed in a combined study of differential scanning calorimetry (DSC) and polarized optical microscopy (POM). The results obtained from UV−vis, XRD, and TMAFM suggest that the annealing of the copolymers above the smectic mesophase transition induces higher crystallinity and long-range ordering to P3HT with improve supramolecular self-assembly. Field-effect mobilities measured were comparable with the P3HT precursors indicating the attached liquid crystalline block had no adverse effect on optoelectronic properties. Field effect transistors fabricated from P3HT-bPMeOBP and PMeOBP-b-P3HT-b-PMeOBP showed average hole mobility of 2.35 × 10−2 and 4.68 × 10−3 cm2 V−1 s−1, respectively, on treated devices with an ON/OFF current ratio of around 103. The results obtained provided new strategies for the design of semiconducting block copolymers with improved morphology and physical properties which can be employed in sensors and memory device applications.

Materials. All commercial chemicals were purchased from Aldrich Chemical Co. Inc. and were used without further purification unless otherwise noted. All reactions were conducted in oven-dried glassware under nitrogen environment. The polymerization glassware was dried at 120 °C for at least 24 h before use and cooled under a nitrogen atmosphere. Copper(I) bromide (99+%) was purified by stirring in glacial acetic acid under nitrogen for 12 h, followed by washing with ethanol and diethyl ether and drying under vacuum for 24 h. 1,1,4,7,10,10-Hexamethylenetriethylenetetramine (97%) (HMTETA) was dried over molecular sieves, and tetrahydrofuran (THF) and toluene were dried over sodium benzophenone ketyl and freshly distilled prior to synthesis of monomer and polymerizations. Characterization. NMR studies of the synthesized monomers and polymers were performed in CDCl3 on a Bruker Avance III 500 NMR spectrometer at 30 °C with tetramethylsilane (TMS) as the internal reference. All GC/MS analyses were performed on an Agilent 68905973 GC/MS workstation under the following conditions: injector and detector temperature, 250 °C; initial temperature, 70 °C; I


Article

Macromolecules

[CuBr]0:[HMTETA]0 = 400:1:2:2. A dry Schlenk flask was charged with P3HT macroinitiator (0.2 g, 0.015 mmol), MeOBP (2.2 g, 6.0 mmol), chlorobenzene (5 mL), and CuBr (∼4.28 mg, 0.030 mmol). After three freeze−pump−thaw cycles, the reaction mixture was immersed in a thermostat oil bath at 85 °C. Then, HMTETA (∼6.28 μL, 0.030 mmol) was added to the reaction mixture via a deoxygenated syringe. The reaction mixture was stirred for 12 h at 85 °C followed by precipition of the reaction mixture in methanol. Block copolymer was isolated by washing with excess cold methanol followed by drying under vacuum. Polymer was characterized by SEC (Mn = 12 700 g mol−1, PDI = 1.8) and 1H NMR (Figure S6). 1H NMR (500 MHz, CDCl3): δ H 0.9 (s, 3H), 1.35−1.43 (m, 6H), 1.69 (t, 2H), 2.80 (t, 2H), 3.85 (s, 3H), 3.94 (s, 4H), 6.86 (s, 4H), 6.98 (s, 1H), 7.47 (s, 4H). SEC: Mn = 12 700 g mol−1; PDI = 1.8. Synthesis of Poly(MeOBP) (P3). Synthesized and recrystallized MeOBP was subjected to ATRP following the same experimental procedure with the molar ratio [MeOBP]0:[ethyl 2-bromoisobutyrate]0:[CuBr]0:[HMTETA]0 = 200:1:2:2. A dry Schlenk flask was charged with the initiator ethyl 2-bromoisobutyrate (0.015 mmol), MeOBP (1.1 g, 3.0 mmol), chlorobenzene (3 mL), and CuBr (∼4.28 mg, 0.030 mmol). After three freeze−pump−thaw cycles, the reaction mixture was immersed in a thermostat oil bath at 85 °C. Then, HMTETA (∼6.28 μL, 0.030 mmol) was added to the reaction mixture via a deoxygenated syringe. The reaction mixture was stirred for 6 h at 85 °C followed by precipition of the reaction mixture in methanol. Polymer was isolated by washing with excess cold methanol followed by drying under vacuum. Polymer was characterized by SEC (Mn = 6000 g mol−1, PDI = 2.9) and 1H NMR (Figure S7). 1H NMR (CDCl3, 500 MHz): δ 7.43, 6.81 (s, 8H), 4.20 (s, 2H), 3.95 (s, 3H), 3.81 (s, 2H), 1.91 (s, 3H), 1.74 (s, 2H), 1.62 (s, 2H), 1.42 (s, 2H), 1.13 (s, 2H), 0.86 (s, 2H). SEC: Mn = 6000 g mol−1, PDI = 2.9.

temperature in air using silicon cantilevers with a nominal spring constant of 42 N m−1 and a nominal resonance frequency of 320 kHz. A typical value of the AFM detector signal corresponding to a rootmean-square (rms) cantilever oscillation amplitude was equal to 1−2 V, and the images were collected at 1 Hz scan frequency in 3 μm × 3 μm scan size. A Meiji polarizing optical microscope (POM) with a Parker Daedal heating stage was used to study the thermotropic liquid crystalline properties of polymer samples, and the images were captured by an Olympus digital camera and were recorded using QCapture software. Experiments were performed for the polymer samples in solid state by sandwiching in between two glass slides and subjecting to heat at the rate of 5 °C/min from room temperature to 190 °C and were allowed to cool down back to room temperature. Monomer and Copolymers Synthesis. Synthesis of Biphenyl Monomer (MeOBP). The biphenyl-containing methacrylate monomer 6-((4′-methoxy-[1,1′-biphenyl]-4-yl)oxy)hexyl methacrylate (MeOBP) was prepared with the procedure described by Zhang and co-workers.51 1H NMR (CDCl3, 500 MHz): δ 7.86, 6.97 (m, 8H), 6.15, 5.58 (s, 2H), 4.20 (t, 2H), 4.03 (t, 2H), 3.88 (s, 3H), 1.98 (s, 3H), 1.86 (m, 2H), 1.76 (m, 2H), 1.55 (m, 2H), 1.48 (m, 2H). Synthesis of Monobromoester-Terminated P3HT (P3HT-Br) [1]. Syntheses of allyl end-capped P3HT and hydroxypropyl end-capped P3HT precursors are given in Scheme S1 of the Supporting Information. Hydroxypropyl-terminated P3HT (2 g, 0.15 mmol) was dissolved in anhydrous THF (100 mL) under inert gas, followed by stirring for 15 min at 40 °C. Then triethylamine (3 mL, 22 mmol) and 2-bromoisobutyryl bromide (2.5 mL, 20 mmol) were added dropwise to the reaction mixture and was stirred for 8 h at 40 °C. The resulting monobromoester-terminated P3HT macroinitiator (P3HTBr) was precipitated in methanol and washed with cold methanol (300 mL), followed by drying under vacuum for 24 h. Polymer was characterized by SEC and 1H NMR (Figure S3). 1H NMR (500 MHz, CDCl3): δ H 0.9 (t, 3H), 1.35−1.43 (m, 6H), 1.69 (t, 2H), 1.95 (s, 6H), 2.80 (t, 2H), 4.28 (m, 2H), 6.98 (s, 1H), SEC: Mn = 13 000 g mol−1; PDI = 1.6. Synthesis of Dibromoester-Terminated P3HT (Br-P3HT-Br) [2]. Synthesis of the precursor dihydroxyl end-capped P3HT is given in Scheme S3 of the Supporting Information. Dihydroxyl-terminated P3HT (1 g, 0.15 mmol) was reacted with excess triethylamine (3 mL, 22 mmol) and 2-bromoisobutyryl bromide (2.5 mL, 20 mmol) following the same procedure for the synthesis of monobromoesterterminated poly(3-hexylthiophene), and the resulted dibromoesterterminated P3HT macroinitiator (Br-P3HT-Br) was characterized by SEC (Mn = 8500 g mol−1, PDI = 1.7) and 1H NMR (Figure S4). 1H NMR (500 MHz, CDCl3): δ H 0.9 (t, 3H), 1.35−1.43 (m, 6H), 1.69 (t, 2H), 1.95 (s, 6H), 2.80 (t, 2H), 6.98 (s, 1H). SEC: Mn = 8500 g mol−1; PDI = 1.7. Synthesis of Poly(3-hexylthiophene)-b-poly(MeOBP) (P1). MeOBP was subjected to atom transfer radical polymerization with P3HT macroinitiator, using CuBr/HMTETA in chlorobenzene as solvent at 85 °C. The molar ratio was [M]0:[P3HT-Br]0:[CuBr]0: [HMTETA]0 = 200:1:2:2. A dry Schlenk flask was charged with P3HT macroinitiator (0.2 g, 0.015 mmol), MeOBP (1.1 g, 3.0 mmol), chlorobenzene (3 mL), and CuBr (∼4.28 mg, 0.030 mmol). After three freeze−pump−thaw cycles, the reaction mixture was immersed in a thermostat oil bath at 85 °C. Then, HMTETA (∼6.28 μL, 0.030 mmol) was added to the reaction mixture via a deoxygenated syringe. The reaction mixture was stirred for 12 h at 85 °C followed by precipition of the reaction mixture in methanol. Block copolymer was isolated by washing with excess cold methanol followed by drying under vacuum. Polymer was characterized by SEC (Mn = 25 600 g mol−1, PDI = 1.7) and 1H NMR (Figure S5). 1H NMR (500 MHz, CDCl3): δ H 0.9 (s, 3H), 1.35−1.43 (m, 6H), 1.69 (t, 2H), 2.80 (t, 2H), 3.62 (s, 3H), 3.94 (s, 4H), 6.76 (s, 4H), 6.98 (s, 1H), 7.43 (s, 4H). SEC: Mn = 25 600 g mol−1; PDI = 1.7. Synthesis of Poly(MeOBP)-b-poly(3-hexylthiophene)-b-poly(MeOBP) (P2). The synthesized dibromoester-terminated P3HT and MeOBP were subjected to ATRP following the same experimental procedure including the molar ratio [MeOBP]0:[Br-P3HT-Br]0:

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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.6b01378. Synthesis of allyl-terminated P3HT, hydroxypropylterminated P3HT, and dihydroxypropyl-terminated P3HT; 1H NMR spectra of allyl-terminated P3HT, hydroxypropyl-terminated P3HT, monobromoester-terminated P3HT, dibromoester-terminated P3HT, P3HTb-poly(MeOBP), poly(MeOBP)-b-P3HT-b-poly(MeOBP), and poly(MeOBP); UV−vis absorption spectra of P3HT-Br and Br-P3HT-Br; current−voltage characteristics and of P1 on untreated devices and current−voltage characteristics and of P2 on untreated, OTS treated, and FTS treated devices; TMAFM height and phase images of P1 and P2 on OFET devices and on mica; cyclic voltammograms of P1 and P2 (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (M.C.S.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Funding from Welch Foundation (AT-1740), NIH (1R21EB019175-01A1), and NSF (DMR-1505950) is gratefully acknowledged.

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REFERENCES

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