Effect of Solution Shearing Method on Packing and Disorder of

Article pubs.acs.org/cm

Effect of Solution Shearing Method on Packing and Disorder of Organic Semiconductor Polymers Gaurav Giri,† Dean M. DeLongchamp,‡ Julia Reinspach,†,§ Daniel A. Fischer,‡ Lee J. Richter,‡ Jie Xu,∥ Stephanie Benight,† Alex Ayzner,§ Mingqian He,⊥ Lei Fang,† Gi Xue,∥ Michael F. Toney,§ and Zhenan Bao*,† †

Department of Chemical Engineering, Stanford University, Stanford, California 94305, United States National Institute of Standards and Technology, 100 Bureau Drive, Gaithersburg, Maryland 20899, United States § Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Laboratory, Menlo Park, California 94025, United States ∥ Department of Polymer Science and Engineering, School of Chemistry and Chemical Engineering, Nanjing University, Nanjing, 210093, People’s Republic of China ⊥ Corning Incorporated, Corning, New York 14831, United States ‡

S Supporting Information *

ABSTRACT: The solution shearing method has previously been used to tune the molecular packing and crystal thin film morphology of small molecular organic semiconductors (OSCs). Here, we study how the solution shearing method impacts the thin film morphology and causes structural rearrangements of two polymeric OSCs with interdigitated side chain packing, namely P2TDC17FT4 and PBTTT-C16. The conjugated backbone tilt angle and the thin film morphology of the P2TDC17FT4 polymer were changed by the solution shearing conditions, and an accompanying change in the charge carrier mobility was observed. For PBTTT-C16, the out-of-plane lamellar spacing was increased by solution shearing, due to increased disorder of side chains. The ability to induce structural rearrangement of polymers through solution shearing allows for an easy and alternative method to modify OSC charge transport properties.

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43 cm2/(V s), both from inkjet printed single crystals and solution coated thin films.17,36,37 Recently, new polymers have been reported with comparably high charge carrier mobilities.23,33,38,39 For example, polymers based on diketopyrrolopyrrole (DPP) units have exhibited hole mobility higher than 10 cm2/(V s).22,30 More recently, with an alignment layer, a polymer semiconductor was reported to give TFT mobility as high as 36.3 cm2/(V s).40 For polymeric OSCs, a promising design has been to incorporate monomers with fused ring aromatic structures to give large intermolecular π−π stacking overlapping areas. Additionally, to increase the electronic coupling, it is reported that the donor−acceptor copolymer design enhances the intermolecular interaction and shortens the distance between the polymer backbones with proper side chains.23,25 In general, the TFT́ s charge carrier transport can be tuned by changing the molecular packing of the OSC. Both the π−π stacking distance (as measured by distance between the polymer backbone) and polymer backbone orientation are critical parameters that

INTRODUCTION Organic electronic devices are actively being pursued for flexible, transparent, large-area, and low-cost applications.1−6 The basic device element of organic electronics is the thin film transistor (TFT), and current research effort is focused on improving the charge transport properties of the organic semiconductors (OSCs) in TFTs. TFTs can be fabricated over a large area by using solution deposition methods such as rollto-roll coating, spin coating, inkjet printing, and so on.7−17 Solution processable conjugated polymers are promising candidates for organic electronics due to their excellent mechanical properties and low temperature deposition from solution.18−20 However, the charge transport properties of polymeric OSCs, specifically the charge carrier mobility, must be improved to create high-performance TFTs.21−23 Recently, several high-mobility polymers have been reported, showing great promise of these materials.23−33 Charge transport in polymers is generally limited by a low structural order and poor interchain and intergrain connectivity in the thin films.34,35 Rubrene, a small-molecule OSC, has shown mobility as high as 40 cm2/(V s) as a vapor-grown single crystal, and 2,7-dioctyl[1]benzothieno[3, 2-b][1]benzothiophene (C8−BTBT) has shown mobilities as high as © 2015 American Chemical Society

Received: October 14, 2014 Revised: March 16, 2015 Published: March 17, 2015 2350

DOI: 10.1021/cm503780u Chem. Mater. 2015, 27, 2350−2359

Chemistry of Materials

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impact charge transport mobility.23 Polymers with interdigitated lamellar packing have been shown to give highly crystalline thin films.38,39,41,42 In addition to molecular design, processing conditions can also have a great impact on polymer morphology, packing and disorder. Thus, it is important to gain understanding on ways to modify polymer packing structure and control disorder using solution processing methods. Using solution processing for polymer structural rearrangements would allow for expeditious roll-to-roll coating of high performance organic electronics. We have previously demonstrated that the solution shearing method can be used to tune the molecular packing and crystalline texture of OSC small molecules.11 Figure 1a shows a

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EXPERIMENTAL SECTION

Materials. Poly(2,5-bis(thiophene-2-yl)-(3,7-dihepta-decanyl tetrathienoacene) (P2TDC17FT4) (MW = 14000 g/mol), poly(3hexylthiophen-2,5-diyl) (P3HT) (MW = 12000 g/mol) (purchased from BASF) and poly(2,5-bis(3-hexadecylthiophene-2-yl)thieno[3,2b]thiophene)(PBTTT-C16) (MW = 40000−80000 g/mol, purchased from Ossila Co.) were used without further purification. Phenyltrichlorosilane (PTS) and octadecyltrichlorosilane (OTS) was purchased from Sigma-Aldrich and were used as received (storage under an argon atmosphere to prevent hydrolysis). Highly doped ntype silicon wafers (resistivity − 1)⎥ ⎢ ⎦ ⎣ 3 2

(1)

in which I is the intensity, θ is the incident X-ray angle and measured cos2α is the azimuthal average of the square of the angle between the substrate and the conjugated backbone. The fit used to calculate the tilt angle of the π-system (also assumed to be the conjugated backbone tilt) is shown in Figure 3b,d, and the tilt angles are listed in Table S1 (Supporting Information). The π-system average tilt for the thin film solution-sheared at 0.05 mm/s was 82.5 ± 2.1°. Assuming the polymer chain axis lies in-plane, consistent with the GIXD, this tilt angle indicates that the conjugated backbones are highly “edge-on” with respect to the substrate. The highly edge-on α of the film rules out almost all orientation distributions except for a uniform, almost perfectly edge-on population. The backbone α became less edge-on as the shearing speed was increased. For the thin film solution-sheared at 0.5 mm/s, the α further decreased to 72.7 ± 1.0°. As the casting speed of the thin films was increased to 2.5 mm/s, the α decreased to 67.6 ± 1.5°. Because α is the average of an orientation distribution, its decrease can be interpreted many ways. It could represent the change of a monomodal orientation distribution, with the assumption that the whole film has the same molecular orientation. Or it could be interpreted as an increase in disorder, where an edge-on fraction is accompanied by an increasingly large disordered fraction. Other distributions are possible. The technique is less ambiguous for the thin film solution-sheared at 0.05 mm/s. Although the NEXAFS of the P2TDC17FT4 thin film is somewhat ambiguous with respect to tilt relative to the substrate, it is consistent with the GIXD data as the conjugated backbones tilt can impact the lamellar spacing of the unit cell. A conjugated backbone that has a lower tilt angle when sheared at lower speeds is one possible interpretation of the NEXAFS results. The optical properties of the polymer thin films were characterized by UV−vis spectroscopy. Figure 3e shows the UV−vis spectra of the P2TDC17FT4 thin films sheared at different speeds. The P2TDC17FT4 thin film sheared at 0.05 mm/s had a broader absorption spectrum compared to those sheared at faster speeds (2.5 mm/s), especially the presence of a large shoulder at λ ≈500 nm. The onset of absorption also showed a shift to shorter wavelengths (i.e., higher energies) with increasing shearing speed. The wavelength of maximum absorption (λ1,max) changed with increasing shearing speed, but a simple trend is lacking (Table S1, Supporting Information). Next, we measured the impact of the changing conjugated backbone orientation on the electronic performance of the P2TDC17FT4 TFTs. The in-plane mobility was measured via a bottom-gate, top-contact field-effect transistor structure. Figure 4a,b(c,d) shows the transfer and output characteristics for a TFT made from the P2TDC17FT4 thin film, as cast by solution shearing at a speed of 0.05 mm/s (2.5 mm/s). The performance of the P2TDC17FT4 polymer cast at a faster speed is higher than that of the P2TDC17FT4 polymer cast at lower speeds, and the hysteresis between the forward and backward scan is low for both shearing speeds (Figure S4, Supporting Information). For both conditions, there is a wide plateau where the mobility does not vary significantly with the

Figure 3. (a) Carbon K-edge NEXAFS spectra for P2TDC17FT4 thin films solution sheared at 0.05 mm/s, with varying incident angles, with the substrate aligned parallel to the beam. (b) Analysis of C 1s → π* intensity data, showing extrapolation to 0 and 90°, to calculate the average tilt of the π-system normal. (c) Carbon K-edge NEXAFS spectra for P2TDC17FT4 thin films solution sheared at 2.5 mm/s, with varying incident angles, with the substrate aligned parallel to the beam. (d) Analysis of C 1s → π* intensity data, showing extrapolation to 0 and 90°, to calculate the average tilt of the π-system normal. (e) UV−vis spectroscopy of P2TDC17FT4 thin film solution sheared at various speeds.

spectra for a P2TDC17FT4 thin film sheared at 0.05 and 2.5 mm/s, respectively. The NEXAFS spectra were identical for the orthogonal directions parallel and perpendicular to the shearing direction (Figure S2, Supporting Information), indicating that the polymer backbones in the thin film do not have a preferred alignment in the plane of the substrate, also confirmed via polarized UV−vis spectroscopy (Figure S3, Supporting 2353


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5c), which we hypothesized enhanced charge transport by having smooth grains. Structural Change and Disorder in PBTTT-C16 Polymer. The existence of various structures in PBTTT thin films cast from solution shearing was also investigated. PBTTT has been extensively studied for organic electronic applications due to its high charge carrier mobility, where its high performance was attributed to greater structural order.38,43,48,54,56 Previous literature has shown that PBTTT side chains are packed in an interdigitated structure, and that the conjugated planes of PBTTT tilt significantly within the crystalline lamellae.42,48 The existence of a second polymorph has been identified through differential scanning calorimetry (DSC) for PBTTT.38 The PBTTT-C16 used here was solution sheared from a heated decalin solution at 150 °C, composed of 5 mg/mL of PBTTT-C16. Figure 6a is a representative GIXD image of a solutionsheared PBTTT-C16 film with the (h00) Bragg peaks indicating lamellar spacing, while the Bragg peak at Qxy ≈ 1.71 Å−1 indicates the π−π stacking distance.58−60 Figure 6b,c focuses on the (300) Bragg peak for PBTTT-C16 thin film

Figure 4. (a) Transfer and (b) output characteristics of the P2TDC17FT4 film solution sheared at 0.05 mm/s. (c) Transfer and (d) output characteristics of the P2TDC17FT4 film solution sheared at 2.5 mm/s. (e) Hole mobilities of the solution sheared P2TDC17FT4 TFTs as a function of shearing speeds. The error bars show the standard deviation.

gate voltage (Figure S5, Supporting Information). Figure 4e shows the mobility of TFTs as a function of increasing shearing speeds. Charge mobilities increased significantly from 0.020 cm2/(V s), for the metastable P2TDC17FT4 films solution sheared at 0.05 mm/s, to 0.072 cm 2/(V s), for the P2TDC17FT4 films sheared at 2.5 mm/s. The electrical characteristics are listed in Table S2 (Supporting Information). As both molecular packing and thin film morphology affect charge transport behavior of TFTs, AFM was used to characterize the surface morphology of the thin films. Figure 5 shows the AFM topographical image of a thin film cast at 0.05

Figure 6. (a) Grazing incidence X-ray diffraction image of the PBTTT-C16 polymer sheared at 0.4 mm/s. (b) The (300) Bragg peak position of a PBTTT-C16 thin film solution sheared at 0.4 mm/s. The white dashed line indicates the (300) Bragg peak position of a dropcast sample. The red line marks the position of the metastable (300) Bragg peak. (c) The (300) Bragg peak position of a PBTTT-C16 thin film solution sheared at 10 mm/s. (d) Lamellar spacing change in PBTTTC16 thin films as a function of shearing speed. The error bars show the standard deviation. The (300) Bragg peak position of a PBTTT-C16 thin film solution sheared at 10 mm/s, (e) before and (f) after thermal annealing. The white dashed line indicates (300) Bragg peak position of a thin film solution sheared at 0.4 mm/s. The red dashed line is present to guide the eye.

Figure 5. AFM images of P2TDC17FT4 thin films solution sheared at (a and b) 0.05 mm/s at different regions and (c) 2.5 mm/s. The image is 5 × 5 μm and the AFM height scale bar is 20 nm.

and 2.5 mm/s, respectively. The thin film cast at 0.05 mm/s exhibits well-defined terraces, but did not seem to form a wellconnected film (Figure 5a,b). The poor connection between regions may be responsible for the lower charge transport. On the other hand, films cast at 2.5 mm/s formed smooth, connected films and exhibited terrace-like morphology (Figure 2354


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function of sin2θ (Figure 7b) was used to calculate the α of the conjugated backbone of PBTTT-C16. The α for PBTTT-C16 cast at 0.4 mm/s was observed to be 66.1 ± 0.2°. Assuming the chain axis is comprehensively in plane, this gives a tilt of the backbone of 24° from the surface normal, consistent with earlier reports on PBTTT-C14.48 NEXAFS data collected at the shearing speed of 10 mm/s are shown in Figure 7c,d, and the α of the PBTTT-C16 backbone was similar at 66.3 ± 0.4°. Thus, in the case of PBTTT-C16, the increase in lamellar spacing is not due to the changing tilt of the polymer backbone. Previously, DeLongchamp et al. have used Fourier transform infrared (FTIR) spectroscopy to identify the mesophase of PBTTT-C14 thin films, that exhibited a 2 Å lamellae expansion, a liquid crystal phase defined by melted and presumed noninterdigitated alkane side chains.63 To investigate if the change in lamellar distance during solution shearing was also correlated to side chain order, we performed additional characterization of the solution-sheared PBTTT-C16 films using FTIR spectroscopy. Figure 8a shows the C−H

solution sheared at 0.4 and 10 mm/s, respectively. At faster speeds (i.e., 10 mm/s), the (300) Bragg peak position decreased to a lower qz value, indicating that the lamellar spacing had increased. Figure 6d shows the change in lamellar spacing as a function of shearing speed. Up to 1 Å expansion of the lamellar spacing was seen at the shearing speed of 8 mm/s compared to the shearing speed of 0.4 mm/s. Figure 6e,f show the (300) Bragg peak positions of a thin film sheared at 10 mm/s, before and after thermal annealing, respectively. After thermally annealing the film at 150 °C for 10−60 min and slowly cooling to room temperature, the (300) Bragg peak position shifted to a higher value (qz ≈ 0.84 Å−1), and the lamellar spacing decreased to 22.5 Å, similar to values obtained in previous literature.61,62 Our thermal annealing results showed that, unlike for P2TDC17FT4 polymer, a slower shearing speed for PBTTTC16 promoted the formation of a molecular packing structure that was stable at room temperature, while the faster shearing speeds promoted the formation of a metastable, slightly rearranged molecular packing structure. However, for both polymers, the metastable packing structure had a larger lamellar spacing than the equilibrium packing structure. We utilized UV−vis spectroscopy, NEXAFS and IR spectroscopy to understand the molecular packing change in PBTTT-C16 that generated larger lamellar spacing. Figure 7a

Figure 8. (a) IR spectra of PBTTT-C16 solution sheared at various speeds. (b) UV−vis spectra of PBTTT-C16 thin film solution sheared at various speeds. Figure 7. (a) Carbon K-edge NEXAFS spectra for PBTTT-C16 thin films solution sheared at 0.4 mm/s, with varying incident angles with the substrate aligned perpendicular to the beam. (b) Analysis of C 1s → π* intensity data, showing extrapolation to 0 and 90°, to calculate the average tilt of the π-system normal. (c) Carbon K-edge NEXAFS spectra for PBTTT-C16 thin films solution sheared at 10 mm/s, with varying incident angles with the substrate aligned parallel to the beam. (d) Analysis of C 1s → π* intensity data, showing extrapolation to 0 and 90°, to calculate the average tilt of the π-system normal.

asymmetric stretch frequency (CH2 vasy) of PBTTT-C16 thin films solution-sheared at 0.4 and 10 mm/s. The CH2 vasy peak of the PBTTT-C16 thin film solution-sheared at 0.4 mm/s was 2919.7 cm−1, representative of a moderately well-ordered alkyl side chain without the presence of many gauche defects.63 Conversely, the CH2 vasy peak of the PBTTT-C16 thin film solution sheared at 10 mm/s was higher, 2924.1 cm−1, indicating that the alkyl side chains became increasingly disordered as the shearing speed increased. We hypothesized that the increased disorder in the alkyl side chain led to a larger average lamellar spacing when the shearing speed was increased. The UV−vis spectroscopy for PBTTT-C16 thin films sheared at different speeds is shown in Figure 8b. The wavelength where maximum absorption (λ1,max) occurred for π−π* transition was blue-shifted to shorter wavelengths upon increasing the shearing speed, that is, from 554 nm for solution sheared thin films at 0.4 mm/s to 537 nm sheared at 10 mm/s

shows the surface sensitive NEXAFS spectra recorded at incident X-ray angles varied at 20, 33, 44, 55, 70, and 90° for a PBTTT-C16 thin film sheared at 0.4 mm/s. The NEXAFS spectra were identical for the orthogonal directions parallel and perpendicular to the shearing direction (Figure S6, Supporting Information), indicating that the thin film was isotropic in the plane of the substrate, also corroborated by polarized UV−vis spectroscopy (Figure S7, Supporting Information). In this case, eq 1 is again used to calculate the α of the C 1s → π* with respect to the substrate plane. The trend in the π* intensity as a 2355


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indicate any observable trend upon increasing the shearing speed (Figure S8, Supporting Information). Solution Shearing of P3HT. Last, a noninterdigitated polymer, regioregular P3HT was solution sheared in this study. We observed that there was no change in its molecular packing under similar shearing conditions (Table S5 and Figure S11, Supporting Information ). The films were isotropic in the plane of the substrate. The UV−vis spectroscopy for P3HT sheared at different speeds show the presence of increased amorphous fractions at the slower shearing speeds (Figure S12, Supporting Information).35

(Table S3, Supporting Information). There was also a reduction in the intensity of a shoulder peak (λ ≈ 600 nm). The blueshift in λ1,max and the decrease in the left shoulder with increased shearing speed could be due to increased disorder in the thin film.45 The electronic performance of the TFTs, as fabricated from solution sheared PBTTT-C16 thin films, were also measured. Figure 9a,b(c,d) shows the transfer and output characteristics

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DISCUSSION Semicrystalline polymers, such as polyethylene and polypropylene, are ubiquitous in our everyday lives. It is well established that the crystalline structures of these polymers can greatly impact their mechanical properties. For example, isotactic polypropylene has up to three different polymorphs, each with different fracture toughness and ductility.66 For polymers, the processing conditions can also have a significant impact on their packing structure.66,67 Poly(vinylidene fluoride), a piezoelectrically active polymer, has been shown to form different packing structures depending on whether it was electrospun or spincoated.68,69 Because some structures are piezoelectrically inactive, precise control in producing the right packing structure is highly important. Change in the crystalline structure of polymeric OSCs has also been demonstrated by several groups. Rivnay et. al, reported that two polymer packing structures may exist for poly([N,N9-bis(2-octyldodecyl)-naphthalene-1,4,5,8-bis(dicarboximide)-2,6-diyl]-alt-5,59-(2,29-bithiophene)) P(NDI2OD-T2) through extensive GIXD experiments.70 Rogers et al. reported the formation of different structures of poly diyl[4,4-bis(2-ethylhexyl)-4H-cyclopenta[2,1-b:3,4-b′]dithiophene-2,6-diyl] (PCPDTBT) during bulk heterojunction solar cell formation.71 Although morphology and packing have been established to be critical factors in impacting polymeric OSC charge transport properties, a thorough study to correlate how solution processing changes polymer packing structure remains lacking. We speculate that the increase in viscosity that occurs during the final stages of solvent evaporation (due to increase of relative polymer concentration) and fast drying conditions during solution shearing can both cause molecular confinement and trapping of metastable structures. Metastable structure formation has also been observed previously in other polymeric systems, such as polyethylene when supercooled from melt, or by crystallizing under high hydrostatic pressure.72,73 Polymers with a rigid backbone like P2TDC17FT4 may also form liquid crystalline states which can then be isolated as metastable structures. More work is necessary to deconvolute the different effects that give rise to polymeric disorder and change in packing structure. In this study, we investigated thin film formation of three polymers, P2TDC17FT4, PBTTT-C16, and regioregular P3HT, fabricated using the solution shearing method. For P2TDC17FT4, we observed the formation of a metastable structure at slow shearing speeds (0.05 mm/s). This structure had a larger lamellar stacking distance as compared to the equilibrium structure (Figure 2). Previous work attributes the increased lamellar spacing to a larger amount of defects on the alkyl side chain for PBTTT, and we expect the same effect to occur in the solution sheared P2TDC17FT4.63 NEXAFS

Figure 9. (a) Transfer and (b) output characteristics of the PBTTTC16 thin films solution sheared at 2 mm/s. (c) Transfer and (d) output characteristics of the PBTTT-C16 thin films solution sheared at 10 mm/s. (e) Hole mobilities of the solution sheared PBTTT-C16 TFTs as a function of shearing speeds. The error bars show the standard deviation.

for a TFT made from PBTTT-C16 thin film, as cast by solution shearing at a speed of 2 mm/s (10 mm/s). The mobilities obtained from PBTTT-C16 TFTs at different shearing speeds are shown in Figure 9e. At the slow shearing speed of 0.4 mm/ s, nonideal TFT characteristics were obtained for the PBTTTC16 thin films (Figure S9, Supporting Information).64,65 The GIXD image of the thin film sheared at 2 mm/s still shows a lamellar spacing similar to the thin films solution sheared at 0.4 mm/s. However, the transfer curve shows the square root of the drain current follows a linear dependence on the gate voltage, and the mobility does not vary significantly over a large range of gate voltage (Figure S10, Supporting Information). There was no significant dependence of charge carrier mobility on its shearing speed (Table S4, Supporting Information). This result followed the reasoning that both the backbone tilt and the π−π stacking distance did not change, and these interactions are important factors in determining the charge transport in TFTs. The coherence length, as calculated by Scherrer analysis of the π−π stacking Bragg peak, also did not 2356


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These morphological changes also influenced both the optical and charge transport properties of PBTTT-C16 and P2TDC17FT4. Hence, solution shearing can be used to change the electronic and potentially mechanical properties of polymeric OSCs.

analysis indicated that the metastable structure was accompanied by the conjugated backbone being highly edge-on with respect to the substrate. (Figures 3 and 4). UV−vis spectroscopy showed a red shift of the thin films exhibiting the larger lamellar spacing, which becomes blue-shifted as a function of increasing shearing speed (Figure 3e). This observation can be attributed to a combination of disorder in the thin film, the formation of domains with smaller conjugation lengths or to a change in the electronic overlap in the structure. Unfortunately, direct correlations between the polymer microstructure and the transport are difficult to draw due to the poor film morphology at low speeds. For PBTTT-C16, various structures with increasing lamellar spacing were formed with increasing shearing speed (Figure 6). Our obtained NEXAFS results indicated that there was no change in the conjugated backbone tilt. IR spectroscopy showed that the polymer’s alkyl side chains became more disordered with increasing shearing speed. This disorder in alkyl spacing was also previously observed and was attributed to the increase in gauche defects in the alkyl side chain.63 TFTs formed from these thin films did not show a charge carrier mobility dependence on the PBTTT structural rearrangement. It appears that the change in alkyl packing did not affect charge transport. Previous studies showed that for PBTTT, even though the interdigitated packing helped in forming larger crystalline domains, the paracrystallinity along the π−π stacking direction was ≈7.3%, indicating disordered behavior.42,43,48,74 If such a large disorder existed in the charge transport direction for our solution sheared films, charge transport may hence be limited by this disorder. On the other hand, the noninterdigitated polymer P3HT did not show any changes in its molecular packing using similar shearing conditions. Our results suggest that side chain interactions may be a parameter that impacts the formation of metastable structures and determines whether they can be stabilized at room temperature. For the case of the interdigitated polymers such as P2TDC17FT4 and PBTTTC16, solution shearing can influence the disorder present in the alkyl side chain packing. Increased disorder then enables the formation of metastable structures with a larger average lamellar spacing. The P2TDC17FT4 polymer also shows a change in the conjugated backbone tilt, as studied by NEXAFS. We hypothesized that balanced and strong intermolecular interactions between the interdigitated alkyl side chains and the π−π stacking plane enabled stabilization and hence, easier formation of various structures during solution shearing. Controlling disorder in polymers is an attractive avenue for controlling charge transport for organic semiconductors. In the future, solution shearing can be used to exert a greater control of the π−π stacking orientation and distance in polymers, as seen in small molecular OSCs.11 Also, the molecular weight and the polydispersity of the polymer could be additional important factors that can be explored in conjunction with solution shearing. In conclusion, we have demonstrated tuning of polymer packing structures in interdigitated polymeric OSCs using the solution shearing method. The change in packing showed an increased lamellar spacing for both PBTTT-C16 and P2TDC17FT4, indicating that the metastable structures had a disordered alkyl side chain packing compared to the equilibrium structure. For P2TDC17FT4, this metastable, disordered alkyl packing is accompanied by a change in the conjugated backbone tilt angle, as studied through NEXAFS.

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ASSOCIATED CONTENT

S Supporting Information *

GIXD, NEXAFS, UV−vis spectra and analysis for polymers, P3HT lamellar spacing and shearing conditions, and TFT results for P2TDC17FT4 and PBTTT-C16. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Portions of this research were carried out at the Stanford Synchrotron Lightsource, a national user facility operated by Stanford University on behalf of the U.S. Department of Energy. We thank I. McCulloch and M. J. Heeney for providing high purity PBTTT-C16. Z.B. acknowledges partial financial support from the National Science Foundation Solid-State Chemistry Program (DMR-1303178).

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REFERENCES

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