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Nov 30, 2015 - The classical nucleation theory predicts that the nucleation rate (Jss) under steady-state conditions is indicated by Jss = N0Zβ* exp(...
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Effects of Regioregularity and Molecular Weight on the Growth of Polythiophene Nanofibrils and Mixes of Short and Long Nanofibrils To Enhance the Hole Transport Yujeong Lee,†,⊥ Jin Young Oh,‡,⊥,# Sung Yun Son,§ Taiho Park,*,§ and Unyong Jeong*,∥ †

Department of Materials Science and Engineering, Yonsei University, 50 Yonsei-Ro, Seodaemun-Gu, Seoul, Korea 120-749 Research Institute of Iron and Steel Technology, Yonsei University, Seoul 120-749, Korea § Department of Chemical Engineering and ∥Department of Materials Science and Engineering, Pohang University of Science and Technology (POSTECH), 77 Cheongam-Ro, Nam-Gu, Pohang, Gyeongbuk 790-784, Korea ‡

S Supporting Information *

ABSTRACT: Morphological control over polythiophenes has been widely studied; however the impacts of regioregularity (RR) and molecular weight (MW) on their structural development have not been investigated systematically. This study examined a representative polythiophene, poly(3-hexylthiophene) (P3HT), to reveal that small differences in the RR can produce a large difference in the growth of nanofibrils. Low-RR P3HTs generated neat long nanofibrils (LNFs), whereas high-RR P3HTs formed short nanofibrils (SNFs). This study identified a critical RR (96−98%) depending on their MW, below which P3HT grew into LNFs and above which P3HT grew into SNFs. This study also found that the mixing ratio between high-RR P3HT and a low-RR P3HT in the solution phase is strongly correlated with the relative populations of SNF and LNF in the coated film. This study suggested that mixing high-RR and low-RR polymers may be a good strategy to optimize the electrical properties of polythiophenes for target applications. As an example, a mixture of high-RR (75%) P3HT and low-RR P3HT (25%) improved considerably the power conversion efficiency of bulk heterojunction polymer solar cells compared with the values obtained from the pure high-RR P3HT and the pure low-RR P3HT. KEYWORDS: poly(3-hexylthiophene), regioregularity, nanofibril, polythiophenes, organic solar cells

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INTRODUCTION The structures and orientations of conjugated polymer assembly have profound effects on device performances due to the anisotropy in the long-distance charge transport, which results mainly from the directionality of the π−π intermolecular stacking interactions.1−4 Among various conjugated polymers, regioregular polythiophenes have been intensively studied for their applications to organic transistors5,6 and solar cells.6−15 Tremendous efforts have been devoted to control their microstructures,16−20 including their in-plane crystal orientations21−25 and their grain size in the thin film state.26−29 Nanofibrils with controlled alignment are advantageous for use in field-effect transistors. Recently, polythiophene nanofibrils aligned alongside the source and drain electrodes have attracted significant interest because their one-dimensional molecular stacking facilitated by π−π interactions accelerates charge transport by reducing the number of intergrain energy barriers.30−32 Thin films consisting of polythiophene nanofibrils are highly flexible without large changes in their electrical preformances;15,33 however, such films have low packing densities because the nanofibril networks naturally contain empty spaces among the nanofibrils. The low density of the nanofibril films can lead to air trap formation and the electrical © 2015 American Chemical Society

properties are difficult to control, which presents critical weaknesses for use in electronic devices. Nanofibrils lose their advantageous properties when used in organic solar cells because charge transport occurs predominantly along the film thickness direction in solar cells. Optimal power conversion efficiencies (PCEs) have been attained by generating a complicated network of nanometer-scale domains with p−n junctions, in which the polythiophene phase is formed by aggregates of small grains.34 Therefore, the structures of the polythiophene grains must be controlled to prepare optimal structures for target applications. Several approaches to the production of high-quality polythiophene nanofibrils have been suggested, including thermal annealing,35 aging in marginal solvents,8,36,37 selfseeded growth via a cooling-and-heating cycle,38,39 guided solvent evaporation,40,41 and ultrasonically assisted selfassembly.42 Although nanofibril growth can be achieved through many different processes, systematic studies have not Received: September 8, 2015 Accepted: November 30, 2015 Published: November 30, 2015 27694

DOI: 10.1021/acsami.5b08432 ACS Appl. Mater. Interfaces 2015, 7, 27694−27702

Research Article

ACS Applied Materials & Interfaces

kDa P3HT formed relatively thick SNFs (41.8 ± 5.7 nm) that were shorter than 1 μm (Figure 1c). Practical applications of conjugated polymers rely on the use of P3HT films coated onto dry substrates. The solutions were cooled to −15 °C, reheated to 25 °C, and spin-coated onto Si wafers at room temperature. The cooling and heating rates were −8 °C/min and 1.8 °C/ min, respectively. The 87 kDa P3HT generated straight LNFs (Figure 1b), whereas the 50 kDa P3HT produced SNFs (Figure 1d). A MW of 50 kDa was known to be sufficiently high to produce high-quality LNFs.38,39 Although the 87 and 50 kDa polymers have different MWs, which is not ideal for disentangling the influence of MW from RR effect, 50 kDa P3HT can represent a P3HT with a high RR and high MW because the MW effect is not significant when MW is large (>20 kDa) (it will be discussed later). The high-MW P3HT samples shown in Figure 1 were characterized by distinct RR values (93.5% for 87 kDa and 98.5% for 50 kDa). The effect of the RR on the morphology requires a more systematic investigation. Because the MW influences the structural development, we synthesized a 19 kDa P3HT with a 96% RR, and increased the RR by selectively collecting high-RR chains. The polymer solutions were slowly cooled from 40 to −20 °C at a cooling rate of −8 °C/min, and the P3HT solution was centrifuged at −20 °C. The growth of nanofibrils during the cooling process can be monitored with UV−vis spectroscopy. The peaks corresponding to π−π stacking at 560 and 605 nm appeared at 10 °C and became prominent as the solution temperature decreased further.48−51 The UV−vis spectra are shown in Supporting Information Figure S1. The 1H NMR spectra revealed that the RR of the 19 kDa sample increased from 96% to 98.1% (Supporting Information Figure S2). The regioregularity of P3HT has been determined by the area ratio of the peaks at 2.80 ppm (head-to-tail, HT) and 2.56 ppm (head-to-head, HH) in the 1H NMR spectra. The ratio of the peak area [HT/(HT + HH)] represents the RR which is the fraction of HT coupling.51 The MWs of the P3HT sample having 98.1% RR were indistinguishable from the MW of the initial polymer; hence, the morphological differences were attributed to the difference in the RR. It is notable that we obtained repeatedly the same RR values from the NMR analysis for each polymer. The NMR analysis was reliable in distinguishing the changes of RR in this study. The NMR spectra used to obtain the RR values are shown in the Supporting Information (Figure S2). Figure 2 shows the AFM images of the P3HT crystals obtained from the 19 kDa sample with a 96% RR (Figure 2a,b) and a 98.1% RR (Figure 2c,d). The 98.1% RR P3HT sample was dried completely in vacuum and re-dissolved in m-xylene at 60 °C. The polymer concentrations of the solutions were identical (0.2 wt %). The vacuum-filtered structures (Figure 2a) were LNFs with a uniform thickness (21.6 ± 1.5 nm); however, the structures obtained from the 98.1% RR P3HT were SNFs (39.4 ± 6.8 nm) (Figure 2c). The spin-coated films revealed similar results, with LNFs from the 96% RR P3HT (Figure 2b) and short SNFs from the 98.1% RR P3HT (Figure 2d). The same method was applied to other P3HTs with various MWs. P3HTs used in this study are summarized in Table 1. More AFM images of the spin-coated high-RR P3HTs are shown in Supporting Information Figure S3. The structures and their RRs are summarized in Figure 3. The red circles indicate SNF, and the dark squares indicate LNF. These results clearly revealed that even a 2% difference in the RR (96−98%) can considerably affect the P3HT structures.

yet examined the effects of the molecular weight (MW) and regioregularity (RR) on their growth. This study investigated the effects of RR and MW on the structure of P3HT nanofibrils. We identified a critical RR (96− 98%), below which P3HT grew into uniform long nanofibrils (LNFs) and above which P3HT grew into jagged short nanofibrils (SNFs). We studied the electrical properties of the materials prepared with distinct crystal structures. A mixture of SNF and LNF provided the best performances when used in organic solar cells by enhancing charge transport along both the vertical (thickness) and the lateral directions. This study revealed that the mixing ratio of a high-RR P3HT and a low-RR P3HT could tune the population of SNF and LNF in the resulting thin film. The structure was characterized using atomic force microscopy (AFM), X-ray diffraction (XRD), and scanning electron microscopy (SEM). The P3HT film density as a function of the mixing ratio was determined based on the refractive index of the sample.

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RESULTS AND DISCUSSION Effects of the RR and MW on the Structural Development. The morphological development of P3HT depends on the polymer’s MW;43−47 however, it cannot be explained solely by MW, as shown in Figure 1. The

Figure 1. AFM images (tapping mode) of the P3HT crystals prepared by vacuum filtration (a, c) or spin-coating (b, d) from m-xylene solutions (0.2 wt %) of 87 kDa P3HT and 50 kDa P3HT. regioregularities (RRs) of the 87 kDa P3HT and 50 kDa P3HT were 93.5% and 98.5%, respectively.

morphologies obtained from 87 kDa (93.5% RR) and 50 kDa (98.5% RR) P3HT were investigated using AFM in the tapping mode. The m-xylene solutions of the polymers (0.2 wt %) were cooled to 0 °C and vacuum-filtered through a Cu-coated TEM grid placed on a filter paper (ADVANTEC filter paper 131). The vacuum-filtered nanofibrils corresponded to the P3HT structure formed in the solution phase. The P3HT nanofibrils that remained on the TEM grid were transferred onto Si wafers for AFM analysis (Figure 1a,c). The 87 kDa P3HT produced straight LNFs with a uniform thickness of 21.3 ± 2.8 nm and average length larger than 5 μm (Figure 1a), whereas the 50 27695

DOI: 10.1021/acsami.5b08432 ACS Appl. Mater. Interfaces 2015, 7, 27694−27702

Research Article

ACS Applied Materials & Interfaces

fresh m-xylene at 60 °C and spin-coated on a Si wafer. We observed the formation of LNFs, which indicates that the formation of SNFs obtained from the high-RR P3HT is not the effect of impurities in the polymer. The LNFs obtained after the washing process are shown in the Supporting Information (Figure S4). The strong dependence of the nanofibril structure on the P3HT RR is not yet understood. The classical nucleation theory predicts that the nucleation rate (Jss) under steady-state conditions is indicated by Jss = N0Zβ* exp(−(ΔGc*/kT)), where N0 is the total number of nucleation sites, k is the Boltzmann constant, T is the temperature, β* is the atomic attachment rate (which accounts for long-range diffusive transport processes), and Z is the so-called Zeldovich factor that corrects for thermal vibrations that tend to destabilize the nuclei. The Zeldovich factor is the probability that nuclei of the critical size at ΔGc* go forward to grow instead of being dissolved in solution. It was reasonable to assume that No was similar for the nucleation of both the high-RR and low-RR P3HT because the polymer concentrations were equal. Thermodynamically, the volumetric free energy (ΔGv) of the crystal decreased as the molecular attraction among the π−π stacked molecules increased; hence, the critical free energy (ΔGc*) for the formation of stable crystal nuclei decreased in the high-RR P3HTs. This lower critical energy increased the number of molecular concentration fluctuations that reach the critical size of a stable nucleus, N0 exp(−(ΔGc*/kT)). Dynamically, the crystal nuclei formed by the high-RR P3HT are expected to have a comparable or greater thermal stability than the crystal nuclei formed by the low-RR P3HT; therefore, the Zeldovich factor of the high-RR P3HT is expected to be equal to or slightly larger than that of the low-RR P3HT. The molecular deposition rate, β* ∼ exp(−(Qd/kT)), of the highRR P3HT is expected to be high due to the strong van der Waals attraction among the molecules. The preceding discussion suggests that the overall nucleation rates may be compared using a simpler expression, Jss ∼ exp(−(Qd/kT)) exp (−(ΔGc*/kT)). We postulated that the enhanced self-assembly among alkyl groups in the high-RR P3HTs accelerated nucleation and growth of the crystals relative to the nucleation and growth in P3HT samples with considerable regiodefects. The nucleation expression predicts that the nucleation temperature in high-RR P3HT samples may exceed the temperature measured in low-RR P3HT. This trend was confirmed by the UV−vis spectra. We could observe the crystallization kinetics by inspecting the appearance of the shoulder peaks in UV−vis spectra of 19 kDa P3HTs (Supporting Information Figure S5). As the temperature decreased from 40 to −20 °C at a cooling rate of −8 °C/ min, the UV−vis peaks corresponding to the π−π stacking interactions appeared. Figure 4 shows the intensity change of the shoulder peak (λ = 604 nm). The appearance of the shoulder peak was considerably different, at 25 °C in the 98.1% RR 19 kDa P3HT and at 10 °C in the 96% RR 19 kDa P3HT. The results indicate that the high nucleation rate in the high-RR P3HT generated large numbers of nuclei at the higher solution temperature, which resulted in the production of SNF instead of LNF. The high-RR molecules formed π−π stacking, even when the molecules did not perfectly overlap in the radial direction, thereby producing jagged SNF. Mixing P3HT SNFs and LNFs. One of the goals of this study was to investigate the effects of RR on the electrical properties of the thin films. The electrical properties of the thin

Figure 2. AFM images of P3HT crystals obtained by vacuum filtration (a, c) or spin-coating (b, d) from m-xylene solutions (0.2 wt %) of 19 kDa P3HT. The RR increased from 96% to 98.1% by selectively extracting P3HT molecules with a high RR.

Table 1. P3HT Samples with Different MWs and RRs, Used in This Study Mna (kDa)

RRb (%)

9.8 12.7 15.3 19 19.7 50 87

95.4 95.3 95.5 96 96.7 98.5 93.4

RR-Cc (%) 96.9 96.9 98.1 98 94.1

a

Number-average MW (Mn). bRR before centrifugation. cRR after centrifugation.

Figure 3. Formation of LNF and SNF, depending on the RR and MW of the P3HT molecules. Higher RRs were obtained from their corresponding P3HTs with lower RRs. Small differences in the RR (especially 96−98%) determined whether LNF or SNF were formed.

The collection of high-RR polymers may be considered as a process of removing impurities in the polymers. To check the possible effect of impurities, we removed them by forming nanofibrils from relatively dense P3HT solutions (2 wt %) at −20 °C for 6 h. At this condition, the polymers formed a bulky gel consisting of nanofibrils. The polymers were retrieved by centrifuge, and the transparent solution was decanted. Most amounts of the polymers were retrieved by this process (yield was more than 95%); therefore the RR values were preserved the same. The retrieved polymers were dissolved completely in 27696

DOI: 10.1021/acsami.5b08432 ACS Appl. Mater. Interfaces 2015, 7, 27694−27702

Research Article

ACS Applied Materials & Interfaces

The results shown in Figure 5 revealed that mixes of the LNF and SNF may offer a good route to obtaining P3HT thin films with good hole transport in both the lateral and thickness directions. Figure 6 shows the P3HT nanofibrils obtained from

Figure 4. Intensity change of the shoulder peak (λ = 604 nm) from the UV−vis spectra of the 98.1% RR 19 kDa P3HT and the 96% RR 19 kDa P3HT. The temperature of the solutions was cooled from 40 to −20 °C.

films prepared from LNF (93.5% RR 87 kDa P3HT) and SNF (98.5% RR 50 kDa P3HT) were compared in Figure 5. Panels a and b of Figure 5 plot the resistance to hole transport along the lateral direction in the P3HT thin films as a function of the electrode distance. Au electrodes were evaporated onto Si wafers through a line-and-space mask, and the P3HT solutions (0.5 wt % in m-xylene) were spin-coated onto the Si wafers. The channel width was 1.0 mm, and the channel length was varied. The electrode setup permitted hole transport only and blocked the flow of electrons. The Au electrode thickness and the P3HT film thickness were 50 and 100 nm, respectively. The LNF film displayed a lower resistance than the SNF film. As the distance between the electrodes increased, the resistance gap between the two films increased (Figure 5c). Panels d and e of Figures 5 show the hole resistance of the P3HT thin films in the thickness direction as a function of the film thickness. The dimensions of the evaporated Au electrodes were 0.5 cm × 0.5 cm. The SNF film displayed a smaller resistance than the LNF film (Figure 5f), in contrast to the resistance in the lateral direction. The low resistance of the LNF film in the lateral direction was attributed to the reduced number of intergrain barriers. The holes flowed along the LNFs through π−π overlap. At the same time, hole transport along the LNFs in the thickness direction was unfavorable because holes tend to flow through longer paths.

Figure 6. (a−c) AFM images of nanofibrils from dilute mixture solutions (0.08 wt %) at different mixing ratios (low-RR P3HT:highRR P3HT, %:%) of the low-RR P3HT (87 kDa) and the high-RR P3HT (50 kDa). The mixing ratios were (a) 75:25, (b) 50:50, and (c) 25:75. (d) SEM image at the mixing ratio of 50:50 from a higher concentration (0.25 wt %).

dilute mixtures (0.08 wt % in m-xylene) of the 87 kDa P3HT and the 50 kDa P3HT samples at different mixing ratios (87 kDa:50 kDa = 75:25 (a), 50:50 (b), and 25:75 (c)). The polymer solutions were subjected to a cycle of cooling and heating, and spin-coated onto the substrate at room temperature. The figures reveal that the SNFs and LNFs were copresent, indicating that the 87 kDa P3HT produced LNFs and

Figure 5. (a, b) J−V curves obtained in the lateral direction of the P3HT LNF film (a) or in the SNF film (b). (c) Comparison of the resistance values in the P3HT films along the lateral direction, as a function of the electrode distance. (d, e) J−V curves obtained in the thickness direction in the P3HT LNF film (d) or in the SNF film (e). (f) Comparison of the resistances of the P3HT films in the thickness direction as a function of the film thickness. The hole conductance was measured using the Au/P3HT/Au structure. 27697

DOI: 10.1021/acsami.5b08432 ACS Appl. Mater. Interfaces 2015, 7, 27694−27702

Research Article

ACS Applied Materials & Interfaces the 50 kDa P3HT generated SNFs. It is not yet clear how come the growth of LNFs and SNFs took place separately. More details of the separate growth mechanism should be pursued in future studies. We consider that the slightly higher energy required for molecular packing among the low-RR P3HT molecules prevented their deposition onto the high-RR P3HT nanofibril seeds, and vice versa. In a thin film formed from a more concentrated solution (0.25 wt %, Figure 6d), the SNFs were not discernible from the LNFs at the mixing ratio of 50:50, but the mixture film appeared denser than the film prepared entirely from LNF. The density of a material can be monitored by measuring the refractive index. A high refractive index indicates a high density for a given material. The refractive index was measured by ellipsometry (SE MG-Vis 1000, Nanoview). Figure 7 shows the

Figure 7. Changes in the refractive index of the P3HT films as a function of the mixing ratio of low-RR P3HT (87 kDa) and high-RR P3HT (50 kDa) in the m-xylene solutions. Figure 8. Length distribution of P3HT SNF and LNF obtained from mixture solutions containing low-RR P3HT (87 kDa) and high-RR P3HT (50 kDa). (a) High-RR P3HT only, (b) low-RR P3HT only, (c, d) 25 wt % high-RR P3HT, (e, f) 50 wt % high-RR P3HT, and (g, h) 75 wt % high-RR P3HT.

refractive index values as a function of the mixing ratio (87 kDa:50 kDa). The refractive index increased considerably even in samples containing a small fraction (25%) of 50 kDa P3HT and reached a maximum at a 75% weight fraction of 50 kDa, which is a higher value than that obtained from a film consisting of pure 50 kDa P3HT. The large number of interstitial spaces among the SNFs in the pure 50 kDa P3HT is considered to reduce the film density. The fraction of SNF versus LNF in a thin film may differ from the weight fraction of the high-RR P3HT in the solution phase. Because the fraction of SNF in the spin-coated mixture film depends on the coating conditions, their population in the mixture film should be compared with that in the SNF-only film, and vice versa for the LNF population. Figure 8 shows the distributions of SNFs and LNFs in mixture thin films with different mixing ratios. The lengths and numbers of SNF and LNF were measured from AFM images. Four images were taken to be partially overlapped at each mixing ratio. The images were combined to measure the length of the nanofibrils. Some images used to analyze the length distribution are shown in the Supporting Information (Figure S6). Figure 8a plots the distribution of SNF obtained from a pure 50 kDa P3HT (highRR) solution. The number of SNF displayed a Gaussian-like distribution with a maximum population at 500 nm in length. No LNF were found in the high-RR P3HT solution. Figure 8b shows the distribution of LNF obtained from the pure low-RR P3HT (87 kDa). The Gaussian-like distribution was centered at 10 μm. When the two polymers were mixed in a solution and spin-coated to form a film (the total polymer concentration was fixed at 0.25 wt %), the population of SNF and LNF in the resulting film was consistent with the polymer mixing ratio in the solution. A 25% high-RR P3HT mixture (Figure 8c,d) resulted in a film containing a population of SNF that was

considerably smaller than the population present in the pure high-RR film (Figure 8a). The population of SNF increased as the mixing ratio of the high-RR P3HT increased (Figure 8c,e,g). By contrast, the population of LNF decreased as the mixing ratio of high-RR P3HT increased (Figure 8d,f,h). The results presented in Figure 8 were used to plot the relationship between the mixing ratio in the solution and the population in the spin-coated film, as shown in Figure 9. The accumulated total number (NSNF) of the SNFs in mixture films were normalized by the number (NSNF,o) in the pure SNF film (50 kDa); CSNF ≡ NSNF/NSNF,o. The same definition was

Figure 9. Changes of the normalized total numbers (CSNF and CLNF) in the mixture films of SNF and LNF. The accumulated total number (NSNF or NLNF)) in the mixture films were divided by the accumulated total number in the pure SNF film or the pure LNF film; CSNF ≡ NSNF/NSNF,o, and CLNF ≡ NLNF/NLNF,o. 27698

DOI: 10.1021/acsami.5b08432 ACS Appl. Mater. Interfaces 2015, 7, 27694−27702

Research Article

ACS Applied Materials & Interfaces applied to LNFs, CLNF ≡ NLNF/NLNF,o. CNSF increased linearly and CLNF decreased monotonically as the mixing ratio of highRR P3HT increased in the solution. These results indicate that the fraction of SNF and LNF in the mixture films could be controlled by adjusting the mixing ratio of high-RR P3HT and low-RR P3HT in a solution. Enhanced Hole Transport by Mixing SNFs and LNFs. Rapid hole transport in the lateral direction is useful in fieldeffect transistors, and rapid hole transport along the thickness direction is desirable in solar cells. A mixture film consisting of LNF and SNF may display synergistic effects in the electrical properties. Figure 10 shows the hole mobilities of the P3HT

Figure 10. Charge carrier mobilities in the P3HT films as a function of the mixing ratio of low-RR P3HT (87 kDa) and high-RR P3HT (50 kDa) in the lateral direction (left, black) and in the thickness direction (right, blue).

Figure 11. (a) Schematic illustration of the polymer solar cells tested in this study. (b) Representative current density−voltage (J−V) curves obtained from the polymer solar cells as a function of the mixing ratio of low-RR P3HT (87 kDa) and high-RR P3HT (50 kDa). The average photovoltaic parameters as a function of the mixing ratio: (c) opencircuit voltage (Voc), (d) short-circuit current (Jsc), (e) fill factor (FF), and (f) power conversion efficiency (PCE).

films in the lateral and thickness directions as a function of the mixing ratio of low-RR (87 kDa) and high-RR (50 kDa) P3HTs. The hole mobility was measured using the Mott− Gurney law, as defined by the space charge limited current (SCLC) theory.52−54 Overall, the mobility in the lateral direction was 1 order of magnitude higher than the mobility in the thickness direction. The maximum hole mobility in the lateral direction (black) was obtained at 25% mixing ratio of high-RR P3HT, whereas the maximum mobility in the thickness direction (blue) was found at 75% mixing of highRR P3HT. The maximum mobility in the thickness direction was nearly twice the value obtained from the LNF-only film. The synergistic effects by mixing LNF and SNF were obtained in bulk heterojunction (BHJ) polymer solar cells. In the BHJ solar cells, rapid vertical hole transport is critical because hole transport is much slower than electron transport. In addition, rapid hole transport in the lateral direction can enhance the short-circuit current (Jsc). The results shown in Figure 10 indicate that vertical transport increased but lateral transport decreased as the fraction of SNF increased, so an optimal mixing ratio between the high-RR and low-RR P3HTs was obtained. Figure 11a shows a schematic diagram of a P3HT/ICBA BHJ solar cell consisting of a mixture of P3HT LNF (87 kDa) and SNF (50 kDa). The photovoltaic parameters according to the mixing ratios are summarized in Table 2. The average characteristic values were obtained from five solar cells at each mixing ratio. Figure 11b shows representative current density−voltage (J−V) curves obtained from the solar cells as a function of the polymer mixing ratio. The BHJ layer was formed by spin-coating a blend solution of P3HTs and ICBA in a 1:1 weight fraction. Panels c−f of Figure 11 show the average photovoltaic parameters of the solar cells. The open-circuit voltage (Voc) remained constant, regardless of the mixing ratio of high-RR P3HT (Figure 11c). Jsc, the fill factor (FF), and the power conversion efficiency (PCE)

reached maximum values at a 75% mixing ratio of high-RR P3HT. The PCE values were 4.6% and 4.5% from the low-RR P3HT only and the high-RR P3HT only, respectively. The PCE of the BHJ film prepared at a 75% mixing ratio was 5.8%, which is a 25% increase compared to the LNF-only or the SNF-only BHJ solar cells. The enhancement in the PCE resulted from enhanced charge transport obtained by mixing LNFs and SNFs without decreasing the film density.

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CONCLUSIONS The structure of the P3HT nanofibrils was sensitive to the regioregularity (RR), critically at in the range of 96−98%. LowRR P3HT generated uniform long nanofibrils (LNFs), whereas high-RR P3HT formed short nanofibrils (SNFs). The critical RR varied according to the molecular weight (MW) of P3HT, low values (∼96%) for low-MW P3HTs and high values (∼98%) for high-MW P3HTs. The mixing ratio between highRR P3HT and low-RR P3HTs in the solution was reasonably correlated with the population of SNF and LNF present in the coated film, indicating that the relative population of SNF in the film may be controlled by adjusting the polymer mixing ratio in the solution. The films formed from mixture solutions of SNF and LNF had larger densities than the LNF-only films. Hole transport in the P3HT films could be enhanced in both the lateral and thickness directions by mixing high-RR P3HT with the low-RR P3HT. By using a mixture of 75% high-RR P3HT and 25% low-RR P3HT in bulk heterojunction solar cells, power conversion efficiency was improved considerably from the LNF-only or SNF-only solar cells. This study revealed that RR profoundly affects the structural development of polythiophene nanofibrils. Optimal fabrication conditions for 27699

DOI: 10.1021/acsami.5b08432 ACS Appl. Mater. Interfaces 2015, 7, 27694−27702

Research Article

ACS Applied Materials & Interfaces Table 2. Photovoltaic Parameters of the BHJ Solar Cells low RR:high RR 100:0 75:25 50:50 25:75 0:100

Voc (V) 0.77 0.77 0.78 0.77 0.77

± ± ± ± ±

0.005 0.005 0.010 0.005 0.005

Jsc (mA/cm2) 11.18 11.44 12.16 12.37 10.82

± ± ± ± ±

FF

0.18 0.17 0.15 0.22 0.26

0.53 0.54 0.57 0.61 0.54

0.100 0.015 0.010 0.011 0.010

PCE (%) 4.605 4.856 5.269 5.911 4.526

± ± ± ± ±

0.12 0.18 0.09 0.16 0.17

Rs (Ω·cm) 11.01 8.94 7.92 5.73 9.63

± ± ± ± ±

0.35 0.65 0.78 0.60 0.83

spectrometer equipped with an ultrashielded superconducting magnet system (Bruker Biospin, Avance II). All of the P3HTs were dissolved in CDCl3 (5 mg/mL). The 1H NMR spectra were analyzed by using MestRec in order to determine the HT-RR. Refractive index was measured by a spectroscopic ellipsometer (SE MG-Vis 1000, Nanoview).

use in transistors and solar cells should be carefully investigated by adjusting the relative populations of SNF and LNF.

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± ± ± ± ±

EXPERIMENTAL SECTION

Materials. m-xylene was purchased from Yakuri Pure Chemicals (>90.0%) and used as received. P3HT (RR ≥ 90.0%; MW = 87 kDa) was purchased from Sigma-Aldrich. All of the other P3HTs used in this study were synthesized. P3HTs used in this study are summarized in Table 1. Synthesis of P3HT. In a glovebox under a nitrogen atmosphere, isopropyl magnesium chloride in THF (0.5 mL, 1 mmol) was added to a solution of 2,5-dibromo-3-hexylthiophene (326 mg, 1 mmol) in anhydrous THF (20 mL) and the mixture was stirred at room temperature. After 1 h, the polymerization was initiated with Ni(dppp)Cl2 dispersed in THF (5 mL) and the mixture was stirred for 3 h. The polymerization was quenched by adding 5 N HCl. After precipitation in methanol, polymer was purified by Soxhlet extraction in methanol, hexane, and chloroform. The chloroform fraction was concentrated under reduced pressure. The concentrate was precipitated in methanol and filtered. The resulting polymer was dried in a vacuum oven for 24 h at ambient temperature, and the desired purple solid was obtained (120 mg, 72%). GPC: Mn = 19,700. PDI. 1H-NMR (400 MHz, CDCl3, ppm): δ 6.98 (s, 1H); 2.81 (t, J = 7.5 Hz, 2H); 1.71 (br, 2H); 1.36 (br, 6H); 0.91 (t, J = 6.6 Hz, 3H). The MW was controlled by the amount of the initiator Ni(dppp)Cl2: 19.7 kDa (4.5 mg, 0.0083 mmol), 19.0 kDa (5 mg, 0.0087 mmol), and 9.8 kDa (10 mg, 0.0166 mmol). Centrifugation of P3HT Nanofibrils. A 1.0 wt % P3HT solution in m-xylene was cooled to −15 °C and centrifuged for 30 min at 11,000 pm. The temperature was measured by thermocouple. After centrifugation, the supernatant was eliminated and the same quantity of m-xylene was added to the solution. P3HT was dissolved by heating to 60 °C, cooled again to −15 °C, and then centrifuged at the same condition. The supernatant was eliminated and the P3HT crystal precipitate was dried in vacuum for 5 h at 25 °C. When the dried P3HT powder was used in solution, it was dissolved at 80 °C in mxylene. Fabrication of the Bulk Heterojunction Polymer Solar Cells. Poly(3,4-ethylenedioxythiophene): polystyrenesulfonate (PEDOT:PSS) layer (40 nm) was coated on a ITO-coated glass substrate and annealed at 140 °C for 10 min in a heating chamber. The photoactive solution (P3HT:ICBA = 1:1 (w/w), 2 wt % in total) in mxylene was spin-coated on the PEDOT:PSS-coated substrate. Thickness of the BHJ layer was 150 nm. The sample was annealed at 150 °C for 10 min in a glovebox. Finally, a top electrode (LiF/Al) was thermally evaporated onto the BHJ layer under vacuum (below 5 × 10−6 Torr). The device area was 5.4 mm2. The photovoltaic properties were measured using a Keithley 2400 source meter under AM 1.5 solar illumination and 100 mW/cm2 irradiation generated using a 300 W Oriel solar simulator. Characterization. For the UV−visible spectra of the P3HT solutions, a small quantity (1 mL) of the solution was taken from the solution. The structure of P3HT crystal was investigated by AFM (Dimension 3100, Digital Instrument Co.). To investigate the crystal structure in the solution phase, the solution was vacuum-filtered through a Cu-coated TEM grid placed on a filter paper (ADVANTEC filter paper 131). The P3HT crystals on the TEM grid were transferred onto a Si wafer for AFM analysis. The dry films were prepared by spin-coating the P3HT solutions on Si wafer at 3000 rpm for 30 s. 1H NMR spectra were measured by a 400 MHz FT-NMR

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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/acsami.5b08432. UV−vis absorbance spectra, NMR spectra, and AFM images of the nanofibrils (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*(U.J.) E-mail: [email protected]. *(T.P.) E-mail: [email protected]. Present Address #

Department of Chemical Engineering, Stanford University, Stanford, California 94305, USA. Author Contributions ⊥

Y.L. and J.Y.O. contributed equally to this work.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS U.J. thanks the MSIP (Ministry of Science, ICT and Future Planning) for support under the IT Consilience Creative Program (Grant NIPA-2013-H0203-13-1001). T.P. acknowledges the support from the Center for Advanced Soft Electronics under the Global Frontier Research Program (Grant No. NRF-2012M3A6A5055225) through the NRF funded by MSIP (Korea).

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REFERENCES

(1) Schenning, A. P. H. J.; Meijer, E. W. Supramolecular Electronics; Nanowires from Self-Assembled π-Conjugated Systems. Chem. Commun. 2005, 3245−3258. (2) Torsi, L.; Magliulo, M.; Manoli, K.; Palazzo, G. Organic FieldEffect Transistor Sensors: A Tutorial Review. Chem. Soc. Rev. 2013, 42, 8612−8628. (3) Coropceanu, V.; Cornil, J.; da Silva Filho; da Silva Filho, D. A.; Olivier, Y.; Silbey, R.; Bredas, J.-L. Charge Transport in Organic Semiconductors. Chem. Rev. 2007, 107, 926−952. (4) Arias, A. C.; MacKenzie, J. D.; McCulloch, I.; Rivnay, J.; Salleo, A. Materials and Applications for Large Area Electronics: Solution-Based Approaches. Chem. Rev. 2010, 110, 3−24. (5) Kim, D. H.; Han, J. T.; Park, Y. D.; Jang, Y.; Cho, J. H.; Hwang, M.; Cho, K. Single-Crystal Polythiophene Microwires Grown by SelfAssembly. Adv. Mater. 2006, 18, 719−723. (6) Lee, W.; Kim, G. H.; Ko, S. J.; Yum, S.; Hwang, S.; Cho, S.; Shin, Y. H.; Kim, J. Y.; Woo, H. Y. Semicrystalline D-A Copolymers with Different Chain Curvature for Applications in Polymer Optoelectronic Devices. Macromolecules 2014, 47, 1604−1612. 27700

DOI: 10.1021/acsami.5b08432 ACS Appl. Mater. Interfaces 2015, 7, 27694−27702

Research Article

ACS Applied Materials & Interfaces (7) Ma, W.; Kim, J. Y.; Lee, K.; Heeger, A. J. Effect of the Molecular Weight of Poly(3-hexylthiophene) on the Morphology and Performance of Polymer Bulk Heterojunction Solar Cells. Macromol. Rapid Commun. 2007, 28, 1776−1780. (8) Berson, S.; De Bettignies, R.; Bailly, S.; Guillerez, S. Poly(3hexylthiophene) Fibers for Photovoltaic Applications. Adv. Funct. Mater. 2007, 17, 1377−1384. (9) Xin, H.; Kim, F. S.; Jenekhe, S. A. Highly Efficient Solar Cells Based on Poly(3-hexylthiophene) Nanowires. J. Am. Chem. Soc. 2008, 130, 5424−5425. (10) Wu, P.- T.; Xin, H.; Kim, F. S.; Ren, G.; Jenekhe, S. A. Regioregular Poly(3-hexylthiophene): Synthesis, Self-Assembly of Nanowires, High-Mobility Field-Effect Transistors, and Efficient Photovoltaic Cells. Macromolecules 2009, 42, 8817−8826. (11) Xin, H.; Reid, O. F.; Ren, G.; Kim, F. S.; Ginger, D. S.; Jenekhe, S. A. Polymer Nanowire/Fullerene Bulk Heterojunction Solar Cells: How Nanostructure Determines Photovoltaic Properties. ACS Nano 2010, 4, 1861−1872. (12) Kim, F. S.; Ren, G.; Jenekhe, S. A. One-Dimensional Nanostructures of π-Conjugated Molecular Systems: Assembly, Properties, and Applications from Photovoltaics, Sensors, and Nanophotonics to Nanoelectronics. Chem. Mater. 2011, 23, 682−732. (13) Kim, J. H.; Kim, M.; Jinnai, H.; Shin, T. J.; Kim, H.; Park, J. H.; Jo, S. B.; Cho, K. Organic Solar Cells Based on Three-Dimensionally Percolated Polythiophene Nanowires with Enhanced Charge Transport. ACS Appl. Mater. Interfaces 2014, 6, 5640−5650. (14) Kim, Y. J.; Cho, C. H.; Paek, K.; Jo, M.; Park, M. K.; Lee, N. E.; Kim, Y. J.; Kim, B. J.; Lee, E. Precise Control of Quantum Dot Location within the P3HT-b-P2VP/QD Nanowires Formed by Crystallization-Driven 1D Growth of Hybrid Dimeric Seeds. J. Am. Chem. Soc. 2014, 136, 2767−2774. (15) Kim, H. J.; Kim, J. − H.; Ryu, J. − H.; Kim, Y.; Kang, H.; Lee, W. B.; Kim, T. − S.; Kim, B. J. Architectural Engineering of Rod-Coil Compatibilizers for Producing Mechanically and Thermally Stable Polymer Solar Cells. ACS Nano 2014, 8, 10461−10470. (16) Brinkmann, M. Structure and Morphology Control in Thin Films of Regioregular Poly(3-hexylthiophene). J. Polym. Sci., Part B: Polym. Phys. 2011, 49, 1218−1233. (17) Roehling, J. D.; Arslan, I.; Moule, A. J. Controlling Microstructure in Poly(3-hexylthiophene) Nanofibers. J. Mater. Chem. 2012, 22, 2498−2506. (18) Hartmann, L.; Tremel, K.; Uttiya, S.; Crossland, E.; Ludwigs, S.; Kayunkid, N.; Vergnat, C.; Brinkmann, M. 2D Versus 3D Crystalline Order in Thin Films of Regioregular Poly(3-hexylthiophene) Oriented by Mechanical Rubbing and Epitaxy. Adv. Funct. Mater. 2011, 21, 4047−4057. (19) Oh, J. Y.; Lee, T. I.; Jang, W. S.; Chae, S. S.; Park, J. H.; Lee, H. W.; Myoung, J. M.; Baik, H. K. Mass Production of a 3D Non-Woven Nanofabric with Crystalline P3HT Nanofibrils for Organic Solar Cells. Energy Environ. Sci. 2013, 6, 910−917. (20) Aryal, M.; Trivedi, K.; Hu, W. Nano-Confinement Induced Chain Alignment in Ordered P3HT Nanostructures Defined by Nanoimprint Lithography. ACS Nano 2009, 3, 3085−3090. (21) Kline, R. J.; McGehee, M. D.; Toney, M. F. Highly Oriented Crystals at the Buried Interface in Polythiophene Thin Film Transistors. Nat. Mater. 2006, 5, 222−228. (22) O'Connor, B.; Kline, R. J.; Conrad, B. R.; Richter, L. J.; Gundlach, D.; Toney, M. F.; DeLongchamp, D. M. Anisotropic Structure and Charge Transport in Highly Strain-Aligned Regioregular Poly(3-hexylthiophene). Adv. Funct. Mater. 2011, 21, 3697−3705. (23) Aiyar, A. R.; Hong, J. I.; Reichmanis, E. Regioregularity and Intrachain Ordering: Impact on the Nanostructures and Charge Transport in Two-Dimensional Assemblies of Poly(3-hexylthiophene). Chem. Mater. 2012, 24, 2845−2853. (24) Guo, Y.; Jiang, L.; Ma, X.; Hu, W.; Su, Z. Poly(3hexylthiophene) Monolayer Nanowhiskers. Polym. Chem. 2013, 4, 4308−4311. (25) Kim, D. H.; Mei, J.; Ayzner, A. L.; Schmidt, K.; Giri, G.; Appleton, A. L.; Toney, M. F.; Bao, Z. Sequentially Solution-

Processed, Nanostructured Polymer Photovoltaics Using Selective Solvents. Energy Environ. Sci. 2014, 7, 1103−1109. (26) Xue, L.; Gao, X.; Zhao, K.; Liu, J.; Yu, X.; Han, Y. The Formation of Different Structures of Poly(3-hexylthiophene) Film on a Patterned Substrate by Dip Coating from Aged Solution. Nanotechnology 2010, 21, 145303. (27) Wang, S.; Kiersnowski, A.; Pisula, W.; Mullen, K. Microstructure Evolution and Device Performance in Solution-Processed Polymeric Field-Effect Transistors: The Key Role of the First Monolayer. J. Am. Chem. Soc. 2012, 134, 4015−4018. (28) Crossland, E. J.W.; Tremel, K.; Fischer, F.; Rahimi, K.; Reiter, G.; Steiner, U.; Ludwigs, S. Anisotropic Charge Transport in Spherulitic Poly(3-hexylthiophene) Films. Adv. Mater. 2012, 24, 839−844. (29) 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. High-Mobility Field-Effect Transistors Fabricated with Macroscopic Aligned Semiconducting Polymers. Adv. Mater. 2014, 26, 2993−2998. (30) Aleshin, A. N. Polymer Nanofibers and Nanotubes: Charge Transport and Device Applications. Adv. Mater. 2006, 18, 17−27. (31) Briseno, A. L.; Mannsfeld, S. C. B.; Jenekhe, S. A.; Bao, Z.; Xia, Y. Introducing Organic Nanowire Transistors. Mater. Today 2008, 11 (4), 38−47. (32) Jo, S. B.; Lee, W. H.; Qiu, L.; Cho, K. Polymer Blends with Semiconducting Nanowires for Organic Electronics. J. Mater. Chem. 2012, 22, 4244−4260. (33) Savagatrup, S.; Printz, A. D.; O'Connor, T. F.; Zaretski, A. V.; Rodriquez, D.; Sawyer, E. J.; Rajan, K. M.; Acosta, R. I.; Root, S. E.; Lipomi, D. J. Mechanical Degradation and Stability of Organic Solar Cells: Molecular and Microstructural Determinants. Energy Environ. Sci. 2015, 8, 55−80. (34) Nguyen, T. L.; Choi, H.; Ko, S. J.; Uddin, M. A.; Walker, B.; Yum, S.; Jeong, J. E.; Yun, M. H.; Shin, T. J.; Hwang, S.; Kim, J. Y.; Woo, H. Y. Semi-Crystalline Photovoltaic Polymers with Efficiency Exceeding 9% in a ∼ 300 nm Thick Conventional Single-Cell Device. Energy Environ. Sci. 2014, 7, 3040−3051. (35) Ma, W.; Yang, C.; Gong, X.; Lee, K.; Heeger, J. Thermally Stable, Efficient Polymer Solar Cells with Nanoscale Control of the Interpenetrating Network Morphology. Adv. Funct. Mater. 2005, 15, 1617−1622. (36) Li, L.; Lu, G.; Yang, X. Improving Performance of Polymer Photovoltaic Devices Using an Annealing-Free Approach via Construction of Ordered Aggregates in Solution. J. Mater. Chem. 2008, 18, 1984−1990. (37) Kim, J. S.; Lee, J. H.; Park, J. H.; Shim, C.; Sim, M.; Cho, K. High-Efficiency Organic Solar Cells Based on Performed Poly(3hexylthiophene) Nanowires. Adv. Funct. Mater. 2011, 21, 480−486. (38) Oh, J. Y.; Shin, M.; Lee, T. I.; Jang, W. S.; Min, Y. H.; Myoung, J. M.; Baik, H. K.; Jeong, U. Self-Seeded Growth of Poly(3hexylthiophene) (P3HT) Nanofibrils by a Cycle of Cooling and Heating in Solutions. Macromolecules 2012, 45, 7504−7513. (39) Oh, J. Y.; Shin, M.; Lee, T. I.; Jang, W. S.; Lee, Y. J.; Kim, C. S.; Kang, J. W.; Myoung, J. M.; Baik, H. K.; Jeong, U. Highly Bendable Large-Area Printed Bulk Heterojunction Film Prepared by the SelfSeeded Growth of Poly(3-hexylthiophene) Nanofibrils. Macromolecules 2013, 46, 3534−3543. (40) Park, M. S.; Aiyar, A.; Park, J. O.; Reichmanis, E.; Srinivasarao, M. Solvent Evaporation Induced Liquid Crystalline Phase in Poly(3hexylthiophene). J. Am. Chem. Soc. 2011, 133, 7244−7247. (41) Lee, C.-K.; Pao, C.-W. Nanomorphology Evolution of P3HT/ PCBM Blends during Solution-Processing from Coarse-Grained Molecular Simulations. J. Phys. Chem. C 2014, 118, 11224−11233. (42) Kim, B.-G.; Kim, M. − S.; Kim, J. Ultrasonic-Assisted Nanodimensional Self-Assembly of Poly-3-hexylthiophene for Organic Photovoltaic Cells. ACS Nano 2010, 4, 2160−2166. (43) Kline, R. J.; McGehee, M. D.; Kadnikova, E. N.; Liu, J.; Frechet, J. M. J.; Toney, M. F. Dependence of Regioregular Poly(327701

DOI: 10.1021/acsami.5b08432 ACS Appl. Mater. Interfaces 2015, 7, 27694−27702

Research Article

ACS Applied Materials & Interfaces hexylthiophene) Film Morphology and Field-Effect Mobility on Molecular Weight. Macromolecules 2005, 38, 3312−3319. (44) Brinkmann, M.; Rannou, P. Molecular Weight Dependence of Chain Packing and Semicrystalline Structure in Oriented Films of Regioregular Poly(3-hexylthiophene) Revealed by High-Resolution Transmission Electron Microscopy. Macromolecules 2009, 42, 1125− 1130. (45) Kim, D. H.; Lee, J.; Park, J. − I.; Chung, J. W.; Lee, W. H.; Giri, G.; Yoo, B.; Koo, B.; Kim, J. Y.; Jin, Y. W.; Cho, K.; Lee, B. − L.; Lee, S. Molecular Weight-Induced Structural Transition of LiquidCrystalline Polymer Semiconductor for High-Stability Organic Transistor. Adv. Funct. Mater. 2011, 21, 4442−4447. (46) Yang, H.; Shin, T. J.; Bao, Z.; Ryu, C. Y. Structural Transitions of Nanocrystalline Domains in Regioregular Poly(3-hexylthiophene) Thin Films. J. Polym. Sci., Part B: Polym. Phys. 2007, 45, 1303−1312. (47) Liu, J.; Arif, M.; Zou, J.; Khondaker, S. I.; Zhai, L. Controlling Poly(3-hexylthiophene) Crystal Dimensions: Nanowhiskers and Nanoribbons. Macromolecules 2009, 42, 9390−9393. (48) Park, Y. D.; Lee, H. S.; Choi, Y. J.; Kwak, D.; Cho, J. H.; Lee, S.; Cho, K. Solubility-Induced Ordered Polythiophene Precursors for High-Performance Organic Thin-Film Transistors. Adv. Funct. Mater. 2009, 19, 1200−1206. (49) Samitsu, S.; Shimomura, T.; Heike, S.; Hashizume, T.; Ito, K. Effective Production of Poly(3-alkylthiophene) Nanofibers by means of Whisker Method using Anisole Solvent: Structural, Optical, and Electrical Properties. Macromolecules 2008, 41, 8000−8010. (50) Mena-Osteritz, E.; Meyer, A.; Langeveld-Voss, B. M. W.; Janssen, R. A. J.; Meijer, E. W.; Bauerle, P. Two-Dimensional Crystals of Poly(3-Allkyl-thiophene)s: Direct Visualization of Polymer Folds in Submolecular Resolution. Angew. Chem., Int. Ed. 2000, 39, 2679−2684. (51) Mao, H. Y.; Xu, B.; Holdcroft, S. Synthesis and StructureProperty Relationships of Regioregular Poly(3-hexylthiophenes). Macromolecules 1993, 26, 1163−1169. (52) Reid, O. G.; Munechika, K.; Ginger, D. S. Space Charge Limited Current Measurements on Conjugated Polymer Films using Conductive Atomic Force Microscopy. Nano Lett. 2008, 8, 1602− 1609. (53) Nikitenko, V. R.; Heil, H.; von Seggern, H. Space-Charge Limited Current in Regioregular Poly-3-hexyl-thiophene. J. Appl. Phys. 2003, 94, 2480−2485. (54) Chung, D. S.; Lee, D. H.; Yang, C.; Hong, K.; Park, C. E.; Park, J. W.; Kwon, S.-K. Origin of High Mobility within an Amorphous Polymeric Semiconductor: Space-Charge-Limited Current and Trap Distribution. Appl. Phys. Lett. 2008, 93, 033303.

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