Letter pubs.acs.org/NanoLett
High Mobility Field Effect Transistors Based on Macroscopically Oriented Regioregular Copolymers Hsin-Rong Tseng,† Lei Ying,† Ben B. Y. Hsu,† Louis A. Perez,‡ Christopher J. Takacs,§ Guillermo C. Bazan,*,† and Alan J. Heeger*,† †
Center for Polymers and Organic Solids, ‡Department of Materials, and §Department of Physics, University of CaliforniaSanta Barbara, Santa Barbara, California 93106, United States S Supporting Information *
ABSTRACT: Field-effect transistors fabricated from semiconducting conjugated polymers are candidates for flexible and low-cost electronic applications. Here, we demonstrate that the mobility of high molecular weight (300 kDa) regioregular, poly[4-(4,4dihexadecyl-4H-cyclopenta[1,2-b:5,4-b′]dithiophen-2-yl)-alt[1,2,5]thiadiazolo[3,4-c]pyridine] can be significantly improved by introducing long-range orientation of the polymer chains. By annealing for short periods, hole mobilities of 6.7 cm2/(V s) have been demonstrated. The transport is anisotropic, with a higher mobility (approximately 6:1) parallel to the polymer backbone than that perpendicular to the polymer backbone. KEYWORDS: Organic field-effect transistor, conjugated polymer, nanostructure, long-range orientation, macroscopic alignment, regioregular polymer rganic field-effect transistors (OFETs) continue to be of interest because of continuing improvements of the charge carrier mobility to values that are promising for use in “plastic electronics”.1,2 Significant progress has been made in the performance of solution-processed OFETs based on smallmolecule or small-molecules mixed with polymer semiconductors (4−5 cm2/(V s)).3,4 For small-molecule OFETs, however, the device performance varies in different regions due to their crystallinity (the grain size is comparable to the channel width). Solution-processed conjugated polymers offer better film forming and mechanical properties, resulting in improved properties such as, for example, flexibility and large-area uniformity.5 However, the hole mobility typically is in the range as 0.1−3 cm2/(V s),6−8 in part due to difficulties in achieving macroscopic order/orientation. To solve the problem of polymer organization, the development of alternating donor−acceptor (D−A) structures provides an effective way to achieve closer intermolecular π−π stacking.9 A significant increase in mobility to 8.2 cm2/(V s) has been reported based for (E)-2-(2-(thiophene-2-yl)vinyl)thiophene donor and the coplanar diketopyrrolopyrole acceptor, as a result of better intermolecular packing.10 Recent studies have demonstrated that, with elevated molecular weight, the copolymer poly[2,6-(4,4-bis-alkyl-4H-cyclopenta-[2,1-b;3,4-b′]dithiophene)-alt-4,7-(2,1,3-benzothiadiazole)] exhibited increased macroscopic order, and that the hole mobility can be improved to values as high as 3.3 cm2/(V s).9 It has also been reported that regioregular versions of the D−A copolymer poly[4-(4,4-dihexadecyl-4H-cyclopenta[1,2-b:5,4-b′]dithiophen-2-yl)-alt-[1,2,5]thiadiazolo[3,4-c]pyridine]
O
© 2012 American Chemical Society
(PCDTPT) exhibit close to 2 orders of magnitude larger hole mobilities compared to similar polymers with less wellorganized structural units.11,12 In addition to molecular design, anisotropic order can be induced during the fabrication process via methods such as mechanical rubbing,13,14 directional solidification,15 and surface treatment techniques, such as aligning polymers on rubbed or grooved substrates.16,17 Directional film deposition methods, such as zone casting,18−21 flow coating,22 or dip-coating process23−25 have also been utilized. In this contribution, we show how anisotropic order and macroscopic alignment of polymer fibers can be achieved by introducing directional solvent evaporation of high molecular weight PCDTPT solutions on nanoscale grooved substrates, ultimately leading to excellent hole mobilities within a transistor configuration. Based on the molecular weight dependence of the mobility as described in the previous study,9 we used preparative gel permeation chromatography to isolate fractions of PCDTPT with number average molecular weights (Mn) of 300 kDa and 100 kDa and polydispersities of approximately 1.5 (Figure S1 in the Supporting Information). The molecular structure of PCDTPT and the regiochemistry of the [1,2,5]thiadiazolo[3,4-c]pyridine structural units relative to the backbone vector are shown in Figure 1a. Bottom-contact, bottom-gate OFET devices with PCDTPT deposited by drop casting from chlorobenzene showed that hole mobilities increased from 0.8 Received: September 27, 2012 Revised: November 1, 2012 Published: November 21, 2012 6353
dx.doi.org/10.1021/nl303612z | Nano Lett. 2012, 12, 6353−6357
Nano Letters
Letter
under a nitrogen atmosphere. All of the annealing processes were carried out for 6 min. The thicknesses of SiO2 before/after scratching were measured by ellipsometry (XLS-100, J. A. Woollam Co., Inc.). The thicknesses are 209.5, 209.1, 208.1, and 207 nm for nonstructured, 100-nm-, 250-nm-, and 500-nmstructured substrates, respectively. The grooves are at most 1−5 nm in depth; thus the thickness is reduced by less than 1% after scratching. Field effect mobilities were obtained in the saturation region of transistor operation by using the equation, IDS = (W/2L) Ciμ(VGS − Vth)2, where W/L is the channel width/length, Ci is the gate dielectric layer capacitance per unit area, and VGS and Vth are the gate voltage and threshold voltages. Tapping mode AFM images were obtained using an Asylum MFP-3D Standard System at room temperature to study the topography of the nanostructured surface and polymer thin film. Polarized UV− vis absorption spectra were recorded at room temperature on a PerkinElmer (Lambda 750) spectrophotometer, with the incident light polarized along a certain direction by a broadband thin film polarizer (Melles Griot). The samples for the polarized absorption were prepared on nanostructured glass substrates with same process used for the device fabrication. Grazing incidence wide-angle X-ray scattering (GIWAXS) measurements were performed at beamline 11-3 at the Stanford Synchrotron Radiation Lightsouce (SSRL) with an X-ray wavelength of 0.9752 Å, at a 400 mm sample to detector distance. Samples were scanned for 90 s in a He environment at an incident angle of 0.1°. The measurements were calibrated using a LaB6 standard. Figure 2 shows surface topography images of the dielectric surfaces and the resulting PCDTPT layers as obtained by atomic force microscopy (AFM) measurements. Figure 2a−d shows the effect of scratching with the diamond nanoparticles on the gate dielectric. For the untreated surface one observes a root-mean-square (RMS) surface roughness of 0.2 nm. After scratching, the RMS roughness increases to 0.48, 0.49, and 0.86 nm for 100, 250, and 500 nm nanostructured surfaces, respectively. With the increase of the diamond particle size one observes deeper and wider grooves on the surface. The surface uniformity also decreases. Figure 2e−h demonstrates the topography of the PCDTPT thin films. Though the polymer fibers show random orientation on nonstructured substrates (Figure 2e), long-range orientation and alignment are observed when deposited on the nanostructured surfaces (Figure 2f−h). The fiber height determined by AFM is ∼6 nm on all of the substrates. The fiber width varies between 30−40 nm on nonstructured substrate and to 50−100 nm for nanostructured substrates. The 100-nm- and 250-nm-structured substrates give similar aligned fiber networks, but the 500-nmstructured substrate results in less desirable curved fiber formation. This difference may due to the fact that the grooves on 500-nm-structured substrate are wider (see Figure 2d), offering less confinement for polymer alignment. In addition, there are small 10−20 nm branches that provide connectivity between the main polymer fiber bundles. Figure S3 in the Supporting Information provides a magnified view of the fiber bundles; several aligned polymer chains can be observed inside each bundle. From the AFM images, we speculate that the polymer chains assemble into fibers during the slow directional drying process and that the fibers follow the surface grooves and ultimately form larger bundles (or fibers). Grooves that are either too small or too wide are not as effective for achieving long-range fiber orientation.
Figure 1. (a) Molecular structure and regiochemistry of PCDTPT used in these studies. (b) Cartoon illustration of the tunnel-like configuration setup and (c) Cross-sectional illustration for directional solvent drying. The pictures are not to scale.
cm2/(V s) (Mn = 100 kDa) to 2.5 cm2/(V s) (Mn = 300 kDa); see Figure S2 in the Supporting Information. From here onward, all of the presented data were obtained by using 300 kDa material. Subsequent efforts focused on managing the dielectric substrate and the mode of film formation. First, surface grooves were introduced on n-doped silicon wafers with 200 nm thick SiO2 gate dielectric insulator layer by scratching the surface with diamond lapping films (from Allied High Tech Products Inc.) with nanoparticle diameters of 100 nm (100-nmstructured), 250 nm (250-nm-structured), and 500 nm (500nm-structured). The scratching distance was 1.5 m at a pressure of 0.1 kg/cm2. Source/drain contacts (50 nm thick gold) were deposited on top of the resulting nanostructured SiO2. A schematic illustration of detailed directional drying process is shown in Figure 1b and c. As shown in Figure 1b, substrates were set face-to-face with two glass spacers, forming a tunnellike configuration that confines the direction of solvent evaporation. A solution of PCDTPT in 1,2-dichlorobenzene with concentration of 0.25 mg/mL was then injected into the tunnel. The solution was dried in a covered Petri dish in a nitrogen environment. This method allows one to slowly grow the organic films on the aligned nanostructures.16,17 The tunnel configuration provides control of the solvent evaporation direction.26 Bottom-gate, bottom-contact field effect transistors with the architecture “Si (500 μm)/SiO2 (200 nm)/Au (50 nm)/ decyltricholosilane/copolymer” were fabricated by slow drying in the tunnel-like configuration with 0.25 mg/mL solutions. The channel length was varied from 10 to 80 μm with channel width of 1 mm. The mobility data in this work mostly came from the devices with 20 μm channel lengths. Average data were calculated from analysis of 10 devices. Prepatterned substrates (Si/SiO2/Au) were first cleaned by ultrasonication in acetone for 3 min and isopropanol for 3 min, then dried in an oven (in air) at 120 °C for 10 min. The samples were then surface-activated with acid hydrolysis and dried again in the oven under the same conditions. After treatment by UVO3 for 15 min, the substrates were passivated by decyltricholosilane from 1 vol % toluene solution at 80 °C for 25 min. After rinsing the substrate with toluene and drying under nitrogen flow, the polymer semiconductors were cast onto the substrate. All film deposition, annealing, and I−V characterization were done 6354
dx.doi.org/10.1021/nl303612z | Nano Lett. 2012, 12, 6353−6357
Nano Letters
Letter
than that fabricated in nitrogen. The mobility increases to 0.8, 1.0, and 2.1 cm2/(V s) after annealing in nitrogen at 100, 150, and 200 °C, respectively. The mobility remains similar (0.6 cm2/(V s)) after annealing in air at 100 °C, but decreases to 0.15 and 0.08 cm2/(V s) after annealing at 150 and 200 °C, respectively. Although the performance decreases due to the fact that organic semiconductors are sensitive to oxygen and humidity,27 PCDTPT shows moderate mobility even after 100 °C annealing in air. Degradation of the performance when operated in air is a general feature of “Plastic Electronics” devices. Thus, at least for the present time, encapsulation will be required for real application. Devices with the carrier transport perpendicular to polymer fibers were made to study anisotropic transport. The devices were fabricated by depositing PCDTPT polymer on 100 nmstructured substrates with the source/drain contacts parallel to the scratch direction. The anisotropic carrier mobility and the polarized absorption data are shown in Figure 4a and b, respectively, as obtained from devices after 200 °C annealing. The mobility shows a dependence on the fiber orientation. Mobility along the fiber (μ//) is always higher than that perpendicular to fiber (μ⊥), that is, the average mobility μ//: μ⊥ = 6:1. Macroscopic optical anisotropy (approximately 2:1) is also observed. The polarized absorption shows that the polymer backbones are aligned along the fibers, implying that carrier transport is maximum along the conjugated backbone direction.9,16,28,29 Grazing incidence wide-angle X-ray diffraction (GIWAXS) was performed to probe the effects of annealing on the molecular organization of solution drop cast samples from 1,2dichlorobenzene on 100 nm-structured substrates.30 The data are collected as an image (see Figure S5 in Supporting Information) by an area detector where azimuthal integrations are performed to produce line cuts. The line cuts for 300 kDa PCDTPT at RT, 100, 150, 200, and 250 °C are plotted as counts (intensity) versus the scattering vector (q (Å−1)) in Figure 5. The two length scale regimes commonly investigated in conjugated polymer thin films are from q ∼ 0.1−1.0 Å−1, which corresponds to the lamellar side chain packing and backbone, or π−π stacking between ordered polymer segments in the q range of ∼1.2−2.0 Å−1.31 Several notable changes occur in the polymer microstructure as the temperature increases. The alkyl stacking distance decreases as the peak shifts from q ∼ 0.18 Å−1 at RT to 0.24 Å−1 at 250 °C which suggests that the side chain interdigitation increases or that they form a different more dense conformation. Another effect of annealing is an increase in long-range order of the alkyl stacking peaks with the appearance of a third order peak at q ∼ 0.72 Å−1. The decrease in the interchain distance and the increase of long-range order is advantageous because it could lower the barrier for interchain hopping and increase charge mobility.29 Thermal annealing effects are also visible in the π−π stacking region, as shown in the inset in Figure 5. In this regime, several peaks are present that suggest that the polymer is well-ordered with strong intermolecular interactions. The hump that the peaks emerge from is most reasonably attributed to amorphous scattering from disordered polymer regions. The change of intensity and position of these peaks as the film is annealed suggests that the film contains polymorphs or different polymer stacking structures. At 200 °C, the peak at q ∼1.3 Å−1 has decreased, while the peaks at q ∼ 1.4 and 1.8 Å−1 have increased. This rearrangement suggests that annealing allows the polymer chains to assemble or form a more favorable
Figure 2. AFM images of dielectric substrate surfaces and related polymer fiber morphology without and with nanostructures. (a) Surface without structures, (b) 100-nm-structured surface, (c) 250nm-structured surface, (d) 500-nm-structured surface, and (e) polymer fibers on nonstructured surface, (f) polymer fibers on 100nm-structured surface, (g) polymer fibers on 250-nm-structured surface, (h) polymer fibers on 500-nm-structured surface.
Figure 3a summarizes the results of devices on 100-nmstructured substrates as a function of annealing temperature. The average mobility increases from 1.9 to 2.6, 3.4, and 6.0 cm2/(V s) after annealing at 100, 150, and 200 °C, respectively. An average mobility of 4.9 cm2/(V s) is observed after annealing at 250 °C. After optimizing the annealing temperature, we further studied the device performance relative to surface scratching by different nanoparticle sizes. As shown in Figure 3b, the average mobility is 2.0, 4.4, and 3.1 cm2/(V s) for nonscratched, 250, and 500 nm nanostructured devices. A maximum mobility equal to 6.7 cm2/(V s) was obtained by using the 100-nm-structured substrate after annealing at 200 °C (see Figure 3a). The transfer and output characteristics are shown in Figure 3c and d. The performance of the devices on 100 nm-structured substrates fabricated in air has been studied. The results of the devices annealed in nitrogen and in air are shown in Figure S4 in the Supporting Information. The average mobility for the ascast device is 0.5 cm2/(V s), which is about factor of 4 lower 6355
dx.doi.org/10.1021/nl303612z | Nano Lett. 2012, 12, 6353−6357
Nano Letters
Letter
Figure 3. Mobility of PCDTPT devices (the quoted value is an average obtained from 10 independent devices). (a) Devices on 100-nm-structured substrates with various annealing temperatures and (b) devices on substrates without and with nanostructures after annealing at 200 °C. The horizontal lines in the box denote the 25th, 50th, and 75th percentile values. The error bars denote the fifth and 95 percentile values. The open square inside the box denotes the mean value. FET characteristics (I−V curves) of PCDTPT with a mobility of 6.7 cm2/(V s). (L = 20 μm, W = 1 mm) are shown in parts c and d: (c) transfer curves taken at VDS = −40 V and (d) output curves taken at various VG. The mobility values were calculated from the dashed lines in panel a. The contact resistance of this device is about 4000 Ω.
Figure 5. Grazing wide-angle X-ray scattering line profiles of PCDTPT films formed by drop casting the solution on 100-nm-structured substrates. The data show that the annealing step produces significant changes in structure and/or assembly. As the sample is heated, the interchain packing becomes closer and the π−π stacking arrangement changes as evidenced by the shift of the first order peak to higher q (0.24 Å−1) and the change of peak intensities in the higher q regime (1.5−2.0 Å−1) shown in the inset.
Figure 4. Anisotropic characteristics of PCDTPT thin films on 100nm-structured substrates. (a) Mobility of devices with carrier transport parallel or perpendicular to polymer fiber orientation (quoted values are again an average obtained from 10 independent devices). (b) Polarized absorption of PCDTPT thin film with the polarization parallel (black) or perpendicular (red) to polymer fiber orientation.
crystallite or polymorph that can be advantageous for interchain charge hopping and therefore hole mobility. To summarize, it has been shown that macroscopic alignment, by a combination of directional solvent evaporation through a tunnel-like configuration and slow drying, of high molecular weight regioregular PCDTPT copolymer on a nanostructured insulator leads to large changes in charge
carrier mobility. Note that the procedure is operationally simple and does not require sophisticated equipment or instrumentation. Long-range alignment of polymer fibers is obtained by creating nanostructured features on the substrate. After thermal annealing, a hole mobility for transport along the conjugated chain of 6.7 cm2/(V s) has been achieved. The modest 6356
dx.doi.org/10.1021/nl303612z | Nano Lett. 2012, 12, 6353−6357
Nano Letters
Letter
(11) 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. (12) Osaka, I.; McCullough, R. D. Acc. Chem. Res. 2008, 41, 1202− 1214. (13) Heil, H.; Finnberg, T.; von Malm, N.; Schmechel, R.; von Seggern, H. J. Appl. Phys. 2003, 93, 1636−1641. (14) Yang, C. Y.; Soci, C.; Moses, D.; Heeger, A. J. Synth. Met. 2005, 155, 639−642. (15) Brinkmann, M.; Rannou, P. Adv. Funct. Mater. 2007, 17, 101− 108. (16) Sirringhaus, H.; Wilson, R. J.; Friend, R. H.; Inbasekaran, M.; Wu, W.; Woo, E. P.; Grell, M.; Bradley, D. D. C. Appl. Phys. Lett. 2000, 77, 406−408. (17) van de Craats, A. M.; Stutzmann, N.; Bunk, O.; Nielsen, M. M.; Watson, M.; Müllen, K.; Chanzy, H. D.; Sirringhaus, H.; Friend, R. H. Adv. Mater. 2003, 15, 495−499. (18) Pisula, W.; Menon, A.; Stepputat, M.; Lieberwirth, I.; Kolb, U.; Tracz, A.; Sirringhaus, H.; Pakula, T.; Müllen, K. Adv. Mater. 2005, 17, 684−689. (19) Miskiewicz, P.; Mas-Torrent, M.; Jung, J.; Kotarba, S.; Glowacki, I.; Gomar-Nadal, E.; Amabilino, D. B.; Veciana, J.; Krause, B.; Carbone, D.; Rovira, C.; Ulanski, J. Chem. Mater. 2006, 18, 4724− 4729. (20) Duffy, C. M.; Andreasen, J. W.; Breiby, D. W.; Nielsen, M. M.; Ando, M.; Minakata, T.; Sirringhaus, H. Chem. Mater. 2008, 20, 7252− 7259. (21) Tracz, A.; Makowski, T.; Masirek, S.; Pisula, W.; Geerts, Y. H. Nanotechnology 2007, 18, 485303. (22) DeLongchamp, D. M.; Kline, R. J.; Jung, Y.; Germack, D. S.; Lin, E. K.; Moad, A. J.; Richter, L. J.; Toney, M. F.; Heeney, M.; McCulloch, I. ACS Nano 2009, 3, 780−787. (23) Guangming, W.; Swensen, J.; Moses, D.; Heeger, A. J. J. Appl. Phys. 2003, 93, 6137−6141. (24) Liu, N.; Zhou, Y.; Wang, L.; Peng, J.; Wang, J.; Pei, J.; Cao, Y. Langmuir 2008, 25, 665−671. (25) Li, L.; Gao, P.; Schuermann, K. C.; Ostendorp, S.; Wang, W.; Du, C.; Lei, Y.; Fuchs, H.; Cola, L. D.; Müllen, K.; Chi, L. J. Am. Chem. Soc. 2010, 132, 8807−8809. (26) Liu, J.; Sun, Y.; Gao, X.; Xing, R.; Zheng, L.; Wu, S.; Geng, Y.; Han, Y. Langmuir 2011, 27, 4212−4219. (27) Sirringhaus, H. Adv. Mater. 2009, 21, 3859−3873. (28) Tsao, H. N.; Cho, D.; Andreasen, J. W.; Rouhanipour, A.; Breiby, D. W.; Pisula, W.; Müllen, K. Adv. Mater. 2009, 21, 209−212. (29) Wang, S.; Kappl, M.; Liebewirth, I.; Müller, M.; Kirchhoff, K.; Pisula, W.; Müllen, K. Adv. Mater. 2012, 24, 417−420. (30) DeLongchamp, D. M.; Kline, R. J.; Fischer, D. A.; Richter, L. J.; Toney, M. F. Adv. Mater. 2011, 23, 319−337. (31) Chabinyc, M. L. Polym. Rev. 2008, 48, 463−492.
orientation/alignment (optical anisotropy of only 2:1) implies that one can expect higher mobility with a higher degree of orientation/alignment by using higher molecular weight PCDTPT.
■
ASSOCIATED CONTENT
S Supporting Information *
Additional Figures S1−S5. This material is available free of charge via the Internet at http://pubs.acs.org.
■
AUTHOR INFORMATION
Corresponding Author
*G.C.B.: Phone: (805) 893-5538. E-mail:
[email protected]. edu; A.J.H.: Phone: (805) 893-3184. E-mail: ajhe@physics. ucsb.edu. Notes
The authors declare no competing financial interest.
■
ACKNOWLEDGMENTS This research was supported by the MC-CAM Program at UCSB sponsored by Mitsubishi Chemical Corporation (Japan). LAP acknowledges support from the ConvEne IGERT Program (NSF-DGE 0801627) and a Graduate Research Fellowship from the National Science Foundation (GRFP). Portions of this research were carried out at the Stanford Synchrotron Radiation Lightsource user facility operated by Stanford University on behalf of the U.S. Department of Energy, Office of Basic Energy Sciences. H.R.T. designed the study and performed the device fabrication, collected the data of FET device performance, and AFM analysis; L.Y. synthesized the polymers, PCDTPT, and conducted polarized absorption measurement; L.A.P. conducted the GIWAX measurement and analysis; H.R.T., L.Y., and L.A.P. prepared the manuscript; B.Y.H. and C.J.T. gave technical support and conceptual advice; G.C.B. and A.J.H. supervised the work, pointed out the importance of the anisotropic optical absorption, and participated in the writing of the manuscript.
■
REFERENCES
(1) Allard, S.; Forster, M.; Souharce, B.; Thiem, H.; Scherf, U. Angew. Chem., Int. Ed. 2008, 47, 4070−4098. (2) Malliaras, G.; Friend, R. Phys. Today 2005, 58, 53−58. (3) Giri, G.; Verploegen, E.; Mannsfeld, S. C. B.; Atahan-Evrenk, S.; Kim, D. H.; Lee, S. Y.; Becerril, H. A.; Aspuru-Guzik, A.; Toney, M. F.; Bao, Z. Nature 2011, 480, 504−508. (4) Smith, J.; Zhang, W.; Sougrat, R.; Zhao, K.; Li, R.; Cha, D.; Amassian, A.; Heeney, M.; McCulloch, I.; Anthopoulos, T. D. Adv. Mater. 2012, 24, 2441−2446. (5) Forrest, S. R. Nature 2004, 428, 911−918. (6) McCulloch, I.; Heeney, M.; Bailey, C.; Genevicius, K.; MacDonald, I.; Shkunov, M.; Sparrowe, D.; Tierney, S.; Wagner, R.; Zhang, W.; Chabinyc, M. L.; Kline, R. J.; McGehee, M. D.; Toney, M. F. Nat. Mater. 2006, 5, 328−333. (7) Kim, J.; Lim, B.; Baeg, K.-J.; Noh, Y.-Y.; Khim, D.; Jeong, H.-G.; Yun, J.-M.; Kim, D.-Y. Chem. Mater. 2011, 23, 4663−4665. (8) Fan, J.; Yuen, J. D.; Cui, W.; Seifter, J.; Mohebbi, A. R.; Wang, M.; Zhou, H.; Heeger, A.; Wudl, F. Adv. Mater. 2012, DOI: 10.1002/ adma.201202195. (9) Tsao, H. N.; Cho, D. M.; Park, I.; Hansen, M. R.; Mavrinskiy, A.; Yoon, D. Y.; Graf, R.; Pisula, W.; Spiess, H. W.; Müllen, K. J. Am. Chem. Soc. 2011, 133, 2605−2612. (10) Chen, H.; Guo, Y.; Yu, G.; Zhao, Y.; Zhang, J.; Gao, D.; Liu, H.; Liu, Y. Adv. Mater. 2012, DOI: 10.1002/adma.201201318. 6357
dx.doi.org/10.1021/nl303612z | Nano Lett. 2012, 12, 6353−6357
