Role of Molecular Weight Distribution on Charge Transport in

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Role of Molecular Weight Distribution on Charge Transport in Semiconducting Polymers Scott Himmelberger,† Koen Vandewal,† Zhuping Fei,‡ Martin Heeney,‡ and Alberto Salleo*,† †

Department of Materials Science and Engineering, Stanford University, Stanford, California 94305, United States Department of Chemistry and Centre for Plastic Electronics, Imperial College London, London SW7 2AZ, U.K.

‡

S Supporting Information *

ABSTRACT: Model semiconducting polymer blends of wellcontrolled molecular weight distributions are fabricated and demonstrated to be a simple method to control intermolecular disorder without affecting intramolecular order or degree of aggregation. Mobility measurements exhibit that even small amounts of low molecular weight material are detrimental to charge transport. Trends in charge carrier mobility can be reproduced by a simple analytical model which indicates that carriers have no preference for high or low molecular weight chains and that charge transport is limited by interchain hopping. These results quantify the role of long polymer tie-chains and demonstrate the need for controlled polydispersity for achieving high carrier mobilities.

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INTRODUCTION Semiconducting polymers are of interest for applications in large area, low cost, and flexible electronics. Significant progress has been made in recent years toward improving charge transport in these materials, necessary for good device performance. High mobilities (>20 cm2/(V s)) have recently been achieved, where a high degree of orientation of the polymer backbone was obtained by taking advantage of a particular, but slow deposition technique.1 However, using more conventional solution casting methods, the same polymer exhibits more modest mobilities (∼1 cm2/(V s)).2 Such a large influence of processing conditions is illustrative of the challenges associated with controlling the complex microstructures of semiconducting polymers, which generally exhibit significant disorder and high levels of entanglement. In order to develop new, high-throughput processing techniques that maintain high carrier mobilities, there is an urgent need for a better understanding of the aspects of polymer microstructure necessary for efficient charge transport. The striking effect of molecular orientation described above highlights the importance of understanding the macromolecular nature of the material, and in particular the role of extended tie-molecules. Indeed, the effects of polydispersity and molecular weight distribution, whose roles are well-established in commodity polymers, are much less studied in the context of charge transport. In this work we clarify the relationship between chain length distribution and carrier mobility and seek to provide guidelines regarding the design and synthesis of polymers for optimal electrical performance. Recently, Noriega et al. have shown that charge transport in semicrystalline conjugated polymers is limited by interchain disorder, with charge carriers able to quickly move along the polymer backbone.3 Hence local scale aggregation is crucial for charges to efficiently jump between chains. In addition to © 2014 American Chemical Society

aggregation, the key requirement for achieving high carrier mobility is the presence of tie-molecules, long polymer chains that bridge the aggregates. This necessity is clearly borne out by the lowering of mobility in low molecular weight (MW) polymer films. Noriega also reported that for all practical carrier densities, charge carriers are confined to the aggregated regions of the film. Together, these observations explain the high performance of weakly ordered polymers which have recently been reported; while not highly crystalline, these materials possess sufficient aggregation as well as chains long enough to provide interaggregate connectivity. The nature of such bridging chains and their optimum density are however still unclear. In this context, engineering MW distributions beyond monodisperse materials is potentially an important avenue to obtain optimal performance. Indeed, low MW polymers possess low levels of disorder but poor interaggregate connectivity, while high MW fractions are well connected but limited by intermolecular disorder.3 Therefore, the question remains if systems which contain low degrees of interchain disorder as well as excellent connectivity between locally aggregated regions can be easily fabricated and are in fact desirable. The impact of MW on microstructure and charge transport in semiconducting polymers has been thoroughly studied. Kline et al. showed that the charge carrier mobility of P3HT could be varied by several orders of magnitude by changing the MW and attributed the drop off in very low MW films to poor connectivity between well-defined, highly crystalline grains.4 Zen et al. observed the same trend in mobilities but attributed the spread to differences in backbone conformation between high and low MW chains.5,6 Numerous studies on P3HT and Received: July 22, 2014 Revised: September 8, 2014 Published: October 16, 2014 7151

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Figure 1. (a) Paracrystalline disorder (g) vs molecular weight for P3HT. Reproduced with permission.12 Copyright 2013 Elsevier. (b) Differential scanning calorimetry of low (8 kDa) and high (42 kDa) molecular weight P3HT as well as a blend of the two materials. (c) Grazing incidence X-ray diffraction scans along the Qxy axis of P3HT blends with varying ratios of high to low MW material. (d) Paracrystalline disorder parameter (g) of P3HT blends along the π-stacking direction as a function of high MW P3HT fraction.

other materials have shown that while low MW materials can form highly ordered structures, efficient charge transport is not achieved until the molecular weight is sufficiently high for these ordered domains to be electrically well connected.7−10 In an effort to obtain the benefits of both high crystallinity and good connectivity, Ma et al. fabricated solar cells from a blend of high and low MW P3HT.11 While they found a higher power conversion efficiency for these blend devices, no further characterization was performed, leaving many unanswered questions as to the reasons behind the improved device performance. Indeed, while the effects of changing the average molecular weight have been thoroughly investigated, the impact of the molecular weight distribution on charge transport remains unexplored. In this work, we study blends of low-disorder, low MW poly3-hexylthiophene (P3HT) with well-connected high MW P3HT, in controlled ratios to investigate systematically the effect of disorder, aggregation, and connectivity on charge transport. Structural features of the films are characterized by grazing incidence X-ray diffraction (GIXD) and absorption spectrum shape analysis while electrical characterization is performed through the use of field-effect transistors (FETs). P3HT’s of modestly low (8 kDa) and high (42 kDa) MW are found to mix completely, forming solid solutions with intermediate levels of disorder and aggregation relative to the low polydispersity samples. The trends in hole mobilities with the blend ratio can be qualitatively reproduced from a simple

transport model, which indicates that charge carriers have no preference for high or low MW chains and supports the assertion that interchain hopping is the rate limiting step in charge transport. These results illustrate the importance of molecular weight and polydispersity control for achieving high carrier mobilities in these materials and identify the basic physics needed to model the transport behavior of polydisperse systems.

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RESULTS/DISCUSSION The effect of molecular weight on disorder in thin films of P3HT has been studied in detail by applying Warren−Averbach peak shape analysis to X-ray diffraction (XRD) patterns.12,13 The paracrystalline disorder parameter (g) is defined as the variance of interplanar spacing (equation S1, Supporting Information). For relatively disordered systems, such as πstacked P3HT chains, it is directly related to the width of the πstacking peak at qz ∼ 1.7 Å−1 (equation S2). For P3HT (Figure 1a), g is relatively low in short chain materials and increases with molecular weight until Mn ∼ 20 kDa; g is independent of chain length in higher molecular weight P3HT, suggesting that π-stacking order is limited by chain folding and entanglements. To investigate the effects of connectivity and disorder and their interplay on charge transport, we chose to study blends of P3HT with MWs above and below this disorder plateau threshold. For the low MW fraction, we chose Mn ∼ 8 kDa, as such a molecular weight 7152

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Figure 2. (a) Absorption coefficient of P3HT blends of low (8 kDa) and high (42 kDa) molecular weight. (b−d) Gaussian line width (σ), excitonic bandwidth (W), and percent aggregation for P3HT blends as a function of high MW P3HT fraction.

and crystallites. This intimate intermixing provides the ideal microstructure to study the effect of aggregate connectivity, as the density of aggregate connecting tie-molecules can be directly varied by adjusting the ratio of high to low MW chains in the film. The electronic structure of the thin-film blends was further characterized by measuring their absorption coefficients by performing transmission and reflection experiments with the aid of an integrating sphere (Figure 2a). It is interesting to note that we see no difference in the absorption onset between the 8 kDa and 42 kDa P3HT, indicating that despite the large difference in molecular weight, conjugation length is not limited by chain length in these samples. Additionally, photoelectron spectroscopy in air (PESA) measurements indicate that the ionization potential is essentially identical for high and low MW P3HT at 4.82 ± 0.05 eV and 4.80 ± 0.05 eV respectively (Figure S3). To complement the GIXD analysis, the absorption peak shape analysis developed by Spano et al. was used in order to quantify the degree of aggregation as well as inter- and intrachain disorder (Figure S4).15,16 This model was developed to describe weakly coupled H-aggregates composed of cofacially packed conjugated chains, as is the case with P3HT, and is implemented by fitting the relative intensities of the absorption peaks in the vibronic progression.17,18 It is immediately apparent that (Figure 2b) the interchain disorder (σ) in locally aggregated regions of the film shows the same trend as the paracrystalline disorder, g, which is calculated from X-ray scattering from crystalline regions of the film large

is low enough to provide a relatively small g (∼7%) but it is still high enough not to be completely excluded during typical synthesis and purification procedures of conventional P3HTs. P3HT was synthesized by the Grignard metathesis route and was fractionated by preparative gel permeation chromatography (GPC).14 The fractions used in this study had Mn ∼ 8, 29, 42, and 61 kDa and each had a low polydispersity (PDI) of 1.2. As phase separation is a concern in polymer blends of vastly different molecular weights, we performed differential scanning calorimetry (DSC) (Figure 1b) to ensure that materials of different molecular weight mix well. The melting points of high (42 kDa) and low (8 kDa) MW P3HT differ by approximately 25 °C while that of a 50−50 blend codissolved in chloroform shows a single melting peak at a temperature between those of the neat materials, indicating molecular mixing. In contrast, the control sample consisting of the two neat materials placed in the same DSC pan showed two distinct melting temperatures corresponding to the pure high and low MW fractions. In order to study the effect of mixing disparate molecular weights on the blend microstructure, thin films of the high and low MW P3HT in varying ratios were spin-coated from chloroform. GIXD was used to quantify g in the π-stacking direction (Figures 1c,d). The width of the π-stacking peak increases monotonically and consequently, the paracrystallinity, as measured by g, is found to vary linearly between the low and high MW samples. This result provides further evidence that the materials are well mixed and suggests that chains of high and low MW are being incorporated into the same aggregates 7153

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Figure 3. (a−d) Mobility vs percentage high molecular weight material for various molecular weight P3HT blends. Dashed lines are fits to the simple model described in the text.

enough to produce discernible diffraction peaks.19 This correspondence indicates that crystallites and aggregates are similarly disordered. Furthermore, the exciton bandwidth W (Figure 2d) has been shown to be directly related to a number of interacting thiophene repeat units along a backbone (N) that constitute an aggregate and as a result gives a measure of the degree of order along the chain backbone.20 W does not vary significantly or systematically with the ratio of high to low MW material and corresponds to an aggregate length of 20−25 interacting monomer units. The percent aggregation of the blends (Figure 2c) increases by a small amount with greater fractions of high MW material, but all blends have a degree of aggregation located in a narrow band between 38 and 45%. In summary, mixing high and low MW P3HT does not affect the intrachain disorder or ability of the chains to aggregate. The primary effect of blending the two materials is to control the number of tie-chains bridging the aggregates −which is related to the concentration of high MW P3HT− and to linearly vary the interchain disorder of these aggregates. It should be noted that while we are able to vary the tie-chain density in these samples, all of the films contain sufficient connectivity between aggregates to be above the percolation threshold. While we are unable to determine the number of interacting chains composing each aggregate in our samples, work by Noriega et al. indicates that extremely local aggregation is sufficient for good charge transport and that once sufficient aggregation is

present, the connectivity of these aggregates becomes more important.3 The impact of the increasing fraction of high MW material on charge transport was investigated through the use of bottom-gate, bottom-contact FETs (Figure 3). Kline et al. demonstrated that the dielectric interface can act as a nucleation site for P3HT crystallites.25 They showed that for films spuncast on octyltrichlorosilane (OTS), the same surface treatment as used in the present work, both high and low MW P3HT were able to nucleate crystals at this interface. Our DSC, XRD, and optical absorption measurements indicate complete mixing between high and low MW chains and a similar degree of aggregation in all samples. Therefore, whatever crystallites do exist at this interface are expected to contain both high and low MW material in the same proportions as in the bulk, with minor effects on the observed trends in mobility. Low MW P3HT (8 kDa), was found to have a mobility approximately a factor of 5 less than neat samples of the three different high MW P3HTs (29, 42, 61 kDa) used in this study. In mixtures, the blends of low MW with each of the three high MW materials showed the same mobility trend: mobility remains constant at the low MW P3HT value with increasing high MW fraction until 40−50% high MW, at which point mobility monotonically increases to the high MW value. It is interesting to note that despite 40% of the film being composed of high MW material, which could provide interaggregate ties, the mobility remains low. The high MW P3HT mobility value 7154

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is not achieved until the film is entirely composed of high MW material. These transport results indicate the blends do not behave according to a simple percolation behavior through the high MW chains. As a control, fractions of two high MW P3HTs (29 and 61 kDa) were mixed and a nearly flat mobility trend was observed (Figure 3d). For all combinations, the mobility trends can be qualitatively reproduced (Figure 3) by a simple model, which makes three basic assumptions: 1 Mobility is limited by interchain charge transport (i.e., movement along the chain backbone is faster than interchain hopping via coupling of π orbitals) 2 The probability of finding a charge on a chain of a given molecular weight is proportional to its concentration (i.e., charges have no energetic preference for high vs low MW chains) 3 Charges can travel farther on high MW chains than on low MW chains This simple model provides an expression for the charge carrier mobility given by eq 1 (see Supporting Information for the full derivation): μ∝

1 ⎛ α + (β − α)wH ⎞ ⎟ ⎜ th ⎝ α + (1 − α)wH ⎠

even with a relatively large torsional angle in the backbone between conjugated segments.23 Thus, even when encountering a kink in the polymer backbone, charge carriers will likely stay on a single chain until that chain either ends or turns away from the electric field.24 Low molecular weight chains are relatively extended due to aggregation and the moderately stiff P3HT backbone. However, as chains become longer they begin to entangle and fold back on themselves. For the mixture of 8 kDa and 29 kDa P3HT we find that the ratio of distance traveled on long chains relative to short chains, β, is equal to α, the molar mass ratio (Table 1), suggesting that the high and low MW chains are about equally extended. However, when the molar mass of the high MW fraction is increased further to 42 kDa and 61 kDa, while still increasing slightly, β begins to level off and diverge from α, consistent with the onset of significant chain folding and entanglement which is expected to occur at the higher MWs. This high MW entanglement effect is especially noticeable in the mixture of only high MW fractions of P3HT (29 + 61 kDa) as a β value close to 1 indicates that charge carriers can travel about equally far on chains of 29 kDa and 61 kDa P3HT Since the electronic energy of the high and low MW fractions is identical, as shown by the PESA measurements, and the polymers form a solid solution, it is unlikely that the intermolecular charge transfer rate differs significantly between high and low MW P3HT chains. It is likely that interchain transport is modestly slower in blends with higher levels of disorder, as modeling shows that paracrystalline disorder introduces shallow trap states into the band gap of the polymer.26 Higher levels of disorder correspond to deeper trap levels, which will have a greater detrimental effect on carrier mobility. However, the disorder in our samples does not span a very large range and we find that the quality of our model’s fit is not highly sensitive to modest changes in the hopping time th, and therefore set it to a constant. Additionally, in low MW P3HT, boundaries between crystalline lamellae are expected to be a significant barrier to charge transport.27,28 While Kline et al. observed significantly different morphologies by atomic force microscopy (AFM) between high (36 kDa) and very low (3 kDa) MW P3HT, as well as the presence of distinct grain boundaries in the very low MW fraction,4 AFM images of our samples show similar morphologies (Figure S6) for all the samples studied, suggesting the grain boundaries are not as prominent in the low MW fractions used here. While these factors likely play a role in transport, the model is able to fit the experimental data quite well without taking them into account. Indeed, while transport in the complex microstructure of a semicrystalline polymer29,30 is certainly more complicated than our model, the fact that the mobility trend can be satisfactorily reproduced with such simple and reasonable assumptions suggests that these factors play a major role in charge transport for many semiconducting polymers. In contrast to our work, Gasperini et al. performed similar mobility measurements on blends of low MW (5.8 kDa) and high MW (89.7 kDa) poly(2,5-bis(3-alkylthiophen-2-yl)thieno[3,2b]thiophene) (PBTTT) and found the mobility reached the high MW value at only a 5% weight fraction of high MW material suggesting that charge carriers in those blends were percolating through only the high MW PBTTT.31 The comparison with our work is not straightforward as the difference in molecular weights used was much larger than the present work, and it is not known whether such different

(1)

where μ is the charge carrier mobility, wH is the weight fraction of high MW P3HT, β is the ratio of the distance a charge carrier can travel on a high MW chain relative to a low MW chain, α is the molar mass ratio between the high and low MW material, and th is the hopping time for intermolecular charge transport. The only fitting parameter is β (Table 1) as α and wH are defined by the MW ratio and weight fractions respectively, and th was chosen to be a constant for all mixtures as discussed in more detail below. Table 1. Parameters Used in Mobility Modela blend

8 + 29 kDa

8 + 42 kDa

8 + 61 kDa

29 + 61 kDa

α β th

3.60 3.60 1.00

5.25 4.33 1.00

7.62 4.38 1.00

2.10 1.09 1.00

α is defined by the molecular weight ratio and β, the ratio of distances traveled on high MW chains relative to low MW chains, was a fitted parameter which had the constraint β ≤ α. a

The first assumption of the model has been demonstrated both experimentally21 and theoretically22,23 for P3HT; charge transport along the polymer backbone is faster than interchain transport owing to the much larger transfer integral in the intrachain direction. The second assumption derives from our observed absence of a difference in the HOMO levels for high and low MW P3HT as measured by PESA. As the energy levels of the materials appear not to differ, there should be no energetic preference for charges to reside on either high or low MW chains. Thus, we assume that the path a charge takes across a device is merely a statistical process related to the ratio of high to low MW polymer chains. Finally, the third assumption comes from the fact that long chains act as tiemolecules, bridging aggregated regions of the film, and allow charges to travel farther before having to hop to another chain. Despite the similar conjugation lengths for high and low MW P3HT, it has been demonstrated that the interchain transfer integral is significantly less than along the intrachain direction, 7155

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Transistor Device Fabrication. Silicon substrates with 200 nm of thermally grown oxide were photolithographically patterned with gold electrodes. The substrates were then sonicated with acetone, methanol, and isopropyl achohol (IPA) before being cleaned by UV−ozone for 20 min. A self-assembled monolayer was applied by immersion of the substrates in (0.2 vol %) octyltrichlorosilane (OTS) in hexane overnight inside a nitrogen-filled glovebox. The substrates were sonicated in toluene for 60 s and rinsed with IPA to remove any multilayers. P3HT was dissolved in chloroform (10 mg/mL) and stirred overnight at room temperature before spincasting. All films were spuncast at 2000 rpm for 30 s. The now complete bottom-gate, bottom-contact field effect transistors were then measured in a vacuum probe station. All devices had 50 μm channel lengths and were 1 mm in width. The mobility was measured in the saturation regime for two substrates and 6−8 devices per composition. GIXD Measurements. Samples were prepared on silicon substrates with a native oxide processed as described above for transistors. X-ray scattering was performed at the Stanford Synchrotron Radiation Lightsource (SSRL) on beamline 11-3 (2D scattering with an area detector, MAR345 image plate, at grazing incidence) with an incident energy of 12.7 keV. Absorption Measurements and Spano Analysis. Samples were prepared on glass substrates as described above for transistors. The absorption coefficient was determined from transmission and reflection measurements made on a Varian Cary 5000 UV−vis-NIR Spectrophotometer equipped with an integrating sphere. Absorption coefficients were fit using equation S3, taking the intermolecular vibrational energy, Ep, to be 0.179 eV.18 Photoemission Spectroscopy in Air Measurements. Samples were prepared as described above for GIXD. Measurements were made on a Riken Keiki AC-2 spectrometer with 10 nW of UV intensity. Atomic Force Microscopy Measurements. Samples were prepared as described above for GIXD. Measurements were made on a Park XE-70 in noncontact tapping mode. Polymer Synthesis/Separation. Polymer synthesis and separation is described in detail in the Supporting Information. The fractions used in the study had molecular weights of Mn ∼ 8, 29, 42, and 61 kDa and all had a low PDI of 1.2.

molecular weight fractions of PBTTT form solid solutions as in the P3HT fractions shown here. In work complementary to that presented here, Wang et al. showed the effect on charge transport of mixing regioregular P3HT (rr-P3HT) with regiorandom P3HT (RRa-P3HT).32 RRa-P3HT is unable to π-stack effectively and contains significant disorder along the chain backbone due to steric hindrance caused by the regio-irregular side chains, resulting in a charge carrier mobility many orders of magnitude lower than rr-P3HT. However, upon mixing the two materials, blends achieve mobilities equal to films of pure rr-P3HT when containing only a few percent of the regioregular material. Wang et al. found that the rr-P3HT preferentially phase separates to the semiconductor/dielectric interface and can form a percolating network when constituting only a few percent of the entire film. Because of the immiscibility of the two materials, the paracrystalline disorder of the rr-P3HT is largely unaffected by the concentration of RRa-P3HT in the film. Additionally, it is energetically more favorable for charge carriers to stay on rr-P3HT due to its longer conjugation length, and thus the RRa-P3HT has essentially no effect on the charge transport properties of the film. The same phase separation and percolation behavior has been observed in mixtures of P3HT and other immiscible insulators.33 These examples illustrate a fundamental and important difference between solid solutions and phase-separated systems in determining charge transport pathways. In conclusion, we have built a model microstructure which allows us to better understand how polydisperse films behave and demonstrated the importance of tie-molecules and low polydispersity in P3HT, a material which should have relevance to many other semicrystalline semiconducting polymers.3 P3HTs of molecular weights spanning nearly an order of magnitude mix at the molecular level to form compounds with levels of paracrystalline disorder intermediate between the neat low and high MW materials. Mobility trends can be reproduced with a simple model, which indicates that charge transport is limited by interchain hopping and that, at least for the MWs studied here, carriers have no preference for low or high MW chains. Our results illustrate the importance of controlling polymer polydispersity to achieve efficient charge transport and indicate that in order to maintain high mobilities, one should take precautions to ensure that the entire molecular weight distribution lies above the threshold at which aggregates are well connected and transport is disorder limited. Even small amounts of low molecular weight material can have deleterious effects on charge transport, due to molecular mixing of high and low molecular weight chains, reducing the density of aggregate connecting molecules. Thus, when designing new polymers, chemists must keep in mind that high MWs are likely a prerequisite to achieving good charge transport, and should focus on ways to improve interchain transport and design molecules which are more tolerant to the high levels of disorder which are likely inherent in macromolecules of this size.

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ASSOCIATED CONTENT

S Supporting Information *

Information concerning the mobility model, polymer synthesis and separation, absorption peak shape analysis, X-ray diffraction, photoelectron spectroscopy, activation energy, and atomic force microscopy. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*(A.S.) 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 S.H. would like to thank the National Science Foundation for support in the form of a Graduate Research Fellowship. This work was supported by the Center for Advanced Molecular Photovoltaics (CAMP) (Award No KUS-C1-015-21) made by the King Abdullah University of Science and Technology (KAUST). M.H. and Z.F. thank the Engineering and Physical Sciences Research Council (Grant Number EP/G060738/1)

EXPERIMENTAL SECTION

Differential Scanning Calorimetry. P3HT was dissolved in chloroform and stirred overnight before being dropcast onto glass slides. The films were allowed to dry before being scraped off with a razor blade, weighed, and placed in a DSC pan. The blend film contained high and low MW P3HT dissolved in the same bottle before dropcasting while the separated sample contained separately dropcast low and high MW P3HT which were placed in the same pan. 7156

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(27) Brinkmann, M.; Rannou, P. Adv. Funct. Mater. 2007, 17, 101− 108. (28) Jimison, L. H.; Toney, M. F.; McCulloch, I.; Heeney, M.; Salleo, A. Adv. Mater. 2009, 21, 1568−1572. (29) Jimison, L. H.; Himmelberger, S.; Duong, D. T.; Rivnay, J.; Toney, M. F.; Salleo, A. J. Polym. Sci., Part B: Polym. Phys. 2013, 51, 611−620. (30) Himmelberger, S.; Dacuña, J.; Rivnay, J.; Jimison, L. H.; McCarthy-Ward, T.; Heeney, M.; McCulloch, I.; Toney, M. F.; Salleo, A. Adv. Funct. Mater. 2013, 23, 2091−2098. (31) Gasperini, A.; Sivula, K. Macromolecules 2013, 46, 9349−9358. (32) Wang, C.; Rivnay, J.; Himmelberger, S.; Vakhshouri, K.; Toney, M. F.; Gomez, E. D.; Salleo, A. ACS Appl. Mater. Interfaces 2013, 5, 2342−2346. (33) Lu, G.; Blakesley, J.; Himmelberger, S.; Pingel, P.; Frisch, J.; Lieberwirth, I.; Salzmann, I.; Oehzelt, M.; Di Pietro, R.; Salleo, A.; Koch, N.; Neher, D. Nat. Commun. 2013, 4, 1588.

for funding. Use of the Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Laboratory, is supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Contract No. DE-AC02-76SF00515.

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REFERENCES

(1) Tseng, H.-R.; Phan, H.; Luo, C.; Wang, M.; Perez, L. A.; Patel, S. N.; Ying, L.; Kramer, E. J.; Nguyen, T.-Q.; Bazan, G. C.; Heeger, A. J. Adv. Mater. 2014, 26, 2993−2998. (2) Ying, L.; Hsu, B. B. Y.; Zhan, H.; Welch, G. C.; Zalar, P.; Perez, L. A.; Kramer, E. J.; Nguyen, T.-Q.; Heeger, A. J.; Wong, W.-Y.; Bazan, G. C. J. Am. Chem. Soc. 2011, 133, 18538−18541. (3) Noriega, R.; Rivnay, J.; Vandewal, K.; Koch, F. P. V.; Stingelin, N.; Smith, P.; Toney, M. F.; Salleo, A. Nat. Mater. 2013, 12, 1038− 1044. (4) Kline, R. J.; McGehee, M. d.; Kadnikova, E. n.; Liu, J.; Fréchet, J. m. j. Adv. Mater. 2003, 15, 1519−1522. (5) Zen, A.; Pflaum, J.; Hirschmann, S.; Zhuang, W.; Jaiser, F.; Asawapirom, U.; Rabe, J. P.; Scherf, U.; Neher, D. Adv. Funct. Mater. 2004, 14, 757−764. (6) Zen, A.; Saphiannikova, M.; Neher, D.; Grenzer, J.; Grigorian, S.; Pietsch, U.; Asawapirom, U.; Janietz, S.; Scherf, U.; Lieberwirth, I.; Wegner, G. Macromolecules 2006, 39, 2162−2171. (7) Schilinsky, P.; Asawapirom, U.; Scherf, U.; Biele, M.; Brabec, C. J. Chem. Mater. 2005, 17, 2175−2180. (8) Müller, C.; Wang, E.; Andersson, L. M.; Tvingstedt, K.; Zhou, Y.; Andersson, M. R.; Inganäs, O. Adv. Funct. Mater. 2010, 20, 2124− 2131. (9) Kline, R. J.; McGehee, M. D.; Kadnikova, E. N.; Liu, J.; Fréchet, J. M. J.; Toney, M. F. Macromolecules 2005, 38, 3312−3319. (10) Goh, C.; Kline, R. J.; McGehee, M. D.; Kadnikova, E. N.; Fréchet, J. M. J. Appl. Phys. Lett. 2005, 86, 122110. (11) Ma, W.; Kim, J. Y.; Lee, K.; Heeger, A. J. Macromol. Rapid Commun. 2007, 28, 1776−1780. (12) Koch, F. P. V.; Rivnay, J.; Foster, S.; Müller, C.; Downing, J. M.; Buchaca-Domingo, E.; Westacott, P.; Yu, L.; Yuan, M.; Baklar, M.; Fei, Z.; Luscombe, C.; McLachlan, M. A.; Heeney, M.; Rumbles, G.; Silva, C.; Salleo, A.; Nelson, J.; Smith, P.; Stingelin, N. Prog. Polym. Sci. 2013, 38, 1978−1989. (13) Rivnay, J.; Noriega, R.; Kline, R. J.; Salleo, A.; Toney, M. F. Phys. Rev. B 2011, 84, 045203. (14) Loewe, R. S.; Khersonsky, S. M.; McCullough, R. D. Adv. Mater. 1999, 11, 250−253. (15) Spano, F. C. J. Chem. Phys. 2005, 122, 234701−234701−15. (16) Spano, F. C. Chem. Phys. 2006, 325, 22−35. (17) Pingel, P.; Zen, A.; Abellón, R. D.; Grozema, F. C.; Siebbeles, L. D. A.; Neher, D. Adv. Funct. Mater. 2010, 20, 2286−2295. (18) Turner, S. T.; Pingel, P.; Steyrleuthner, R.; Crossland, E. J. W.; Ludwigs, S.; Neher, D. Adv. Funct. Mater. 2011, 21, 4640−4652. (19) Duong, D. T.; Toney, M. F.; Salleo, A. Phys. Rev. B 2012, 86, 205205. (20) Gierschner, J.; Huang, Y.-S.; Averbeke, B. V.; Cornil, J.; Friend, R. H.; Beljonne, D. J. Chem. Phys. 2009, 130, 044105. (21) O’Connor, B.; Kline, R. J.; Conrad, B. R.; Richter, L. J.; Gundlach, D.; Toney, M. F.; DeLongchamp, D. M. Adv. Funct. Mater. 2011, 21, 3697−3705. (22) Lan, Y.-K.; Yang, C. H.; Yang, H.-C. Polym. Int. 2010, 59, 16− 21. (23) Lan, Y.-K.; Huang, C.-I. J. Phys. Chem. B 2009, 113, 14555− 14564. (24) Noriega, R.; Salleo, A.; Spakowitz, A. Proc. Natl. Acad. Sci. U.S.A. 2013, 110, 16315−16320. (25) Kline, R. J.; McGehee, M. D.; Toney, M. F. Nat. Mater. 2006, 5, 222−228. (26) Rivnay, J.; Noriega, R.; Northrup, J. E.; Kline, R. J.; Toney, M. F.; Salleo, A. Phys. Rev. B 2011, 83, 121306. 7157

dx.doi.org/10.1021/ma501508j | Macromolecules 2014, 47, 7151−7157