Epitaxially grown strained pentacene thin film on graphene membrane.

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Epitaxy

Epitaxially Grown Strained Pentacene Thin Film on Graphene Membrane Kwanpyo Kim, Elton J. G. Santos, Tae Hoon Lee, Yoshio Nishi, and Zhenan Bao*

Organic-graphene system has emerged as a new platform for various applications such as flexible organic photovoltaics and organic light emitting diodes. Due to its important implication in charge transport, the study and reliable control of molecular packing structures at the graphene–molecule interface are of great importance for successful incorporation of graphene in related organic devices. Here, an ideal membrane of suspended graphene as a molecular assembly template is utilized to investigate thin-film epitaxial behaviors. Using transmission electron microscopy, two distinct molecular packing structures of pentacene on graphene are found. One observed packing structure is similar to the well-known bulk-phase, which adapts a face-on molecular orientation on graphene substrate. On the other hand, a rare polymorph of pentacene crystal, which shows significant strain along the c-axis, is identified. In particular, the strained film exhibits a specific molecular orientation and a strong azimuthal correlation with underlying graphene. Through ab initio electronic structure calculations, including van der Waals interactions, the unusual polymorph is attributed to the strong graphene–pentacene interaction. The observed strained organic film growth on graphene demonstrates the possibility to tune molecular packing via graphene–molecule interactions.

1. Introduction Organic electronics has been drawing intense research effort and opened up various applications such as organic light emitting diodes (OLEDs), thin film transistors (OTFTs), and photovoltaic (OPV) devices.[1,2] For these organic electronics Prof. K. Kim, Dr. E. J. G. Santos, Prof. Z. Bao Department of Chemical Engineering Stanford University Stanford, California 94305, USA E-mail: [email protected] Prof. K. Kim Department of Physics Ulsan National Institute of Science and Technology (UNIST) Ulsan 689–798, South Korea T. H. Lee, Prof. Y. Nishi Department of Electrical Engineering Stanford University Stanford, California 94305, USA DOI: 10.1002/smll.201403006 small 2015, DOI: 10.1002/smll.201403006

applications, the molecular packing structures at the interface with electrode and dielectric materials are crucial. For example, the first few layers of molecules near the dielectric/ semiconductor interface mainly determine the charge transport behavior for organic transistors[3] and obtaining optimized molecular packing structures at the interfaces is being actively pursued.[4,5] Similarly, the molecular interface with electrode materials is important for achieving efficient charge injection. Graphene has recently emerged as a promising candidate in applications such as optoelectronics[6] and transparent electrodes.[7,8] For many applications, the well-controlled assembly of inorganic or organic materials on graphene is required to take advantage of graphene's superior properties. The assembly of inorganic materials, such as transition metal dichalcogenide, on graphene has shown promising results for transistors and optoelectronics applications.[9] In case of organic-graphene systems, graphene electrodes have demonstrated their potential for stretchable and flexible organic devices, such as OPVs,[10] OLEDs,[11] and OTFTs.[12] The molecular packing structure on graphene is of great

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importance for charge transport at the graphene–molecule interface.[13] The molecular structure and their orientation at the interface are highly sensitive to the nature of the substrate because the molecular assembly is determined by competition between the molecule-interface and intermolecular interactions. Taking advantage of substrate–molecule interaction to control the morphology and packing structure of organic molecules is highly desirable[14] and graphene can be an ideal substrate for this goal. Previously, highly ordered pyrolytic graphite (HOPG) and epitaxial graphene on various metal substrates have been investigated to study epitaxial growth of various organic molecules.[4,13,15–17] Compared to graphene on metallic substrates, isolated graphene on insulating substrates is more relevant for practical applications.[12,18] On the other hand, suspended graphene can be an ideal sample to study intrinsic molecule–graphene interactions by removing any underlying substrate effect. When graphene is suspended, its single layer thickness also offers unique opportunity for detailed structural investigation using techniques such as transmission electron microscopy (TEM).[19] Pentacene, an organic semiconducting molecule, has shown promising charge transport behaviors and has attracted great interest for transistor applications.[1,12,20] Previously, X-ray scattering has been mainly used to determine the crystal structures of pentacene either in single-crystal bulk-phase or in thin film phase grown on substrates, such as SiO2 or on self-assembled monolayer modified SiO2.[12,21–26] Scanning tunneling microscopy (STM) has been used to study sub-monolayer and thin multi-layer growth modes on various substrates.[27] Notably, pentacene and related molecules at sub-monolayer coverage adapt flat face-on geometry on graphite or graphene-grown metal substrates, due to π–π interaction of molecules with graphene substrates.[16,17] Even though STM techniques have provided important information on the initial nucleation process and have elucidated the generally-face-on pentacene orientation on graphitic substrate, it is not well-suited for determination of thin film crystal structures grown on graphene due to the lack of precise measurements of lattice parameters. Additionally, X-ray diffraction is not usually suitable to study sub-micrometer scale objects, especially if the structure of a specimen is not homogeneous. In this paper we study the molecular packing structures of pentacene thin films grown on graphene via TEM including selected area electron diffraction (SAED). Compared to other characterization tools, TEM imaging and SAED can be applied to nanometer dimensions and very thin samples,[28] which allows us to precisely investigate thin film structures of organic crystals on graphene. Indeed we find two distinct packing structures including a new polymorph phase of pentacene. One observed packing structure is similar to the well-known bulk-phase, which adapts a faceon molecular orientation on graphene. A rarely reported polymorph of pentacene film, which shows significant strain along c-axis, is also observed. In particular, the strained film exhibits a strong azimuthal correlation with underneath graphene lattice together with a specific molecular orientation toward graphene surface. From ab initio calculations with

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taking into account van der Waals interactions, we attribute the unusual polymorph to the strong graphene–pentacene interaction. The observed strained film growth on graphene demonstrates the possibility to control molecular packing via graphene–molecule interactions. The significant elongation of d-spacing in the horizontal layers indicates the reduced vertical interlayer distance of pentacene crystal, which may be useful to obtain enhanced vertical charge transport.

2. Results and Discussion We prepare suspended graphene membrane and utilize it as the pentacene growth template. Suspended graphene membrane has been used as a TEM imaging substrate such as for small molecules adsorbed on graphene membranes.[19,29] Because graphene is an atomic-thin membrane, the background scattering from the graphene is very low and it can serve as an ideal substrate for TEM imaging. Figure 1 shows the specimen schematic and the corresponding TEM images. The suspended graphene cover an area with a diameter of 1.2 µm. The synthesized polycrystalline graphene has the average grain size bigger than 5 µm[7,30] and, therefore, the graphene inside a circular hole is usually a single grain (Figure S1, Supporting Information). After the graphene

Figure 1. TEM images of pentacene film on graphene membrane. a) A schematic of pentacene film grown on graphene membrane. Pentacene is thermally evaporated to a graphene TEM grid. The central circle region has a suspended graphene membrane. Outside the circle, the pentacene is evaporated onto amorphous carbon film. Bottom image shows a side-view of sample geometry. b) TEM image of pentacene crystals on graphene. The pentacene crystals show directional film morphology with one of the crystal axis predominantly aligned in the vertical. c) TEM image of different area where pentacene shows nonparallel crystal axis directions. Inset shows the electron diffraction of underlying graphene membrane. The axes of pentacene crystals have relative misorientation Θ of 17° from the graphene zigzag directions. d) Pentacene growth morphology on amorphous carbon film.


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TEM grid preparation, pentacene molecules were thermally evaporated with the substrate temperature held at 60 °C. For the pentacene deposition, we placed the TEM grid such that the pentacene was also grown on amorphous carbon film outside the suspended graphene area (Figure 1a). This allows us to directly compare the pentacene growth morphology on graphene membrane and amorphous carbon film. The TEM images clearly show that pentacene films on graphene and amorphous carbon have very different growth morphology as displayed in Figure 1b–d. On the graphene membrane, the pentacene crystals exhibit elongated grain morphology with well-defined crystal axis. Figure 1b shows the directional pentacene film morphology with one of the crystal axis predominantly aligned in the vertical direction. TEM images in different regions show that the pentacene crystal axis is aligned with the underlying graphene lattice direction (Figure 1c). Using SAED and high-resolution TEM imaging, we can identify the graphene lattice direction (inset of Figure 1c). Interestingly, all the pentacene crystal axes show a misorientation angle Θ of 17° with respect to the zigzag (ZZ) graphene lattice direction. This is an indirect evidence that pentacene film grows epitaxially on graphene substrate. Additional evidence for epitaxial growth of pentacene on graphene will follow in the article. As shown in Figure 1d, the pentacene film on amorphous carbon exhibits isotropic island-growth mode, which is very distinct from the morphology on graphene. First, we will discuss the pentacene crystal structures on amorphous carbon. Figure 2 summarizes SAED data of a pentacene island grown on amorphous carbon. Figure 2a

Figure 2. Selected area electron diffraction (SAED) of pentacene film grown on amorphous carbon film. a) SAED of a pentacene island grown on amorphous carbon. The circles indicate graphene diffraction spots (0–110) diffraction. b) Electron diffraction simulation of pentacene thin-film phase with zone axis [001]. c) The view of pentacene packing structure observed from zone axis [001]. d) The side-view of pentacene film growth morphology on amorphous carbon. The pentacene molecules show the edge-on packing structure on amorphous carbon surface. small 2015, DOI: 10.1002/smll.201403006

shows one typical SAED measured from pentacene on amorphous carbon. The graphene diffraction spots, which are marked by circles, are also seen since graphene is underneath the carbon film even though pentacene molecules are evaporated onto the carbon film (see Figure 1a). The graphene diffraction can be used as the precise length caliber for SAED; the graphene diffraction (0–110) spots have the distance of 1/2.13 Å−1. We find that the pentacene on carbon adapt the thin-film phase, which is similar to the structure reported for thin films grown on SiO2 substrate.[24,25] The orientation of the pentacene molecules is edge-on on the amorphous carbon film as shown in Figure 2c. The electron diffraction simulation of the reported thin-film phase of pentacene reproduces the same pattern as the observed SAED pattern precisely. We find that d(200) = 2.98 Å, d(020) = 3.79 Å with γ = 90o. This is in a very good agreement with the thin film-phase.[24,25] Therefore, we conclude that the evaporated pentacene on carbon film shows the same crystal phase grown on SiO2 substrate. We now move to the discussion of pentacene film grown on graphene. On graphene, we observe two different structures. These two structures have a very similar packing motif. Figure 3 shows data on one of the observed pentacene crystal structure on graphene, similar to the well-known

Figure 3. Face-on bulk-phase pentacene crystal on graphene. a) TEM image of a pentacene crystal on graphene. The inset shows the zoom-in image at the left edge of the crystal. The lattice fringes of d(001) spacing (≈14 Å) is observed, which is parallel to the crystal edge direction. b) SAED signal from the crystal shown in panel a). The diffraction signals marked by circles are graphene (0–110) spot. The c*-axis is perpendicular to a pentacene crystal axis. The d(001) spacing is observed as 14.2 Å, which is consistent with a bulk-phase pentacene. The c*-axis of the pentacene and graphene (0–110) spot are misoriented by 17°. The streaks are observed and its direction is perpendicular to long axis of pentacene molecules. c) Face-on molecular packing structure on graphene with zone axis [−110]. Bottom image shows a top view of the pentacene film on graphene. d) Electron diffraction simulation with a bulk-phase pentacene with zone axis [−110].


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Figure 4. Strained pentacene crystal on graphene. a) TEM image of a strained pentacene crystal on graphene. b) SAED from the crystal shown in panel a). The diffraction signals marked by circles are graphene (0–110) spots. The c*-axis is aligned with a pentacene crystal axis. The d(001) spacing is observed as 14.9 Å, which is 5% longer than a bulkphase pentacene. The c*-axis of the pentacene and graphene (0–110) spot are misaligned by 17°. The streaks are also observed. c) Molecular packing structure on graphene with zone axis [0–10]. Bottom image shows a top view of the pentacene film on graphene. d) Electron diffraction simulation with a pentacene crystal with zone axis [0–10].

bulk-phase.[21] For a precise length calibration, the graphene diffraction spots are used again. In the SAED pattern, we observe multiple lines of diffraction spots with inter-spot distance of 14.2 Å, which is consistent with the previously known bulk phase.[21] The observation of d(001) spacing from SAED indicates that the pentacene molecules adapt a faceon molecular orientation on graphene. This result is also consistent with the previous observations of pentacene grown on graphene transferred to SiO2 (or graphite).[12,17] Together with the real-space TEM image, the SAED pattern shows that the c*-axis is perpendicular to the one of the observed pentacene crystal axes (Figure 3a). The zoom-in image at the edge of the crystal also shows lattice fringes parallel to the crystal edge, with an approximately 14 Å spacing (inset in Figure 3a). These observations confirm that the well-defined crystal facets correspond to the boundary surface of pentacene molecular layer. By comparing SAED simulations with different zone axes (Figure S2, Supporting Information), we conclude that the observed SAED in Figure 3b is originated from the crystal with zone axis [−110]. The simulation results precisely reproduce the observed SAED pattern (Figure 3d). Interestingly, in the observed SAED, the c*-axis of pentacene crystal has a misorientation angle Θ of 17° with respect to the graphene (0–110) diffraction spot. The reconstructed model of graphene and pentacene from the SAED shows that the underlying graphene lattice direction has a good rotational alignment with the pentacene molecules when imaged from the TEM e-beam incident direction, that is, each pentacene molecule axis is


aligned to the zigzag lattice direction (Figure 3c). We note that, even though there are some preferred Θ, we observed various Θ other than 17° for the observed bulk-phase with d(001) = 14.2 Å. Another previously reported bulk-phase with d(001) = 14.5 Å[22] was not observed in our study. More interestingly, we also observe a new pentacene polymorph on graphene, which is significantly strained along the c*-axis. Figure 4 shows the SAED data from the strainedphase pentacene. The SAED shows the presence of c*-axis in the diffraction plane (Figure 4b), similar to the observed common bulk-phase. Among SAED simulations with different zone axes using bulk-phase pentacene with d(001) = 14.2 Å, zone [0–10] produces a very similar diffraction pattern with the observed SAED pattern (Figure 4d). The intensity distributions among different diffraction spots are also in a good agreement to the experimental observation. However, the simulation (Figure 4d) and the experimentally observed (Figure 4b) SAED patterns show one important discrepancy, i.e., the value of d(001). The experimental observation, 14.9 ± 0.1 Å, is 5% larger compared to the bulk-phase value, 14.2 Å. We emphasize that the presence of graphene diffraction peaks allows us to measure the d-spacing quite accurately and the difference (0.7 Å) is much bigger than the error range (0.1 Å). On the other hand, we experimentally find that d(100) is 6.23 ± 0.03 Å, which is in a good agreement with the bulk-phase value d(100) = 6.23 ± 0.01 Å. Given that the diffraction intensity distribution among various diffraction spots are also quite similar to the bulk-phase, we conclude that the observed polymorph has a crystal structure which is under 5% elongation (or strain) along the c*axis from the known bulk-phase. The strained pentacene film with d(001) = 14.9 Å also exclusively show zone axis [0–10]. Extra TEM images and SAED patterns verify that the epitaxial angle and the zone axis are consistent for the strained phase (Figure S3, Supporting Information). We note that previously a pentacene polymorph with a similar d(001) spacing with 15.0 Å has been reported from pentacene evaporated on polyimide but its detailed crystal structure was not characterized.[23] The diffused lines are also observed from most of SAED in pentacene crystals on graphene. The thermal diffuse scattering lines have been recently observed for 6,13-bistriisopropyl-silylethynyl (TIPS)-pentacene crystals.[31] From the line direction, we can deduce the directional thermal motion of molecules in the crystals at room temperature. We find that the thermal motion in pentacene is similar to that of TIPSpentacene; the diffusion line and the direction of thermal diffusion motion are normal to the long axis of the molecular backbone direction. We observe that the c*-axis of the strained pentacene crystals have exclusively Θ of 17° with respect to the graphene (0–110) diffraction spot. Figure 5a shows the histogram of relative misorientation angles Θ between pentacene c*-axis and graphene armchair (AC) lattice directions. Due to the hexagonal symmetry of graphene lattice, Θ can be defined between 0° and 30°. Please note that c*-axis of pentacene crystal is perpendicular to the usually observed microstructure crystal axis; Θ can be also defined as the misorientation angle between pentacene crystal axis and


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Figure 5. Epitaxial relation between pentacene c*-axis and graphene lattice directions. a) Histogram of relative rotation angles between pentacene c*-axis and graphene armchair lattice direction. Due to the hexagonal symmetry of graphene lattice, the misorientation Θ can be defined between 0° and 30°. Strained phase (filled) exclusively shows a 17° misorientation. The bulk-phase (mesh) exhibits somewhat broad distribution of relative misorientation. b) Rotational angle dependence of adsorption energy of pentacene crystal on graphene surface. Θ ≈ 17° misorientation is the lowest energy configuration. The pentacene molecular axis shows a good alignment with the graphene zigzag direction in the case of 17° misorientation. c) The schematic of different interfaces of pentacene crystal toward graphene substrate. d) Molecular energy calculation of pentacene crystal with difference interfaces. Interface with the zone [0–10] axis shows the lower energy configuration compared to zone [100].

graphene ZZ lattice direction as shown in Figure 1c. The bulk phase shows somewhat scattered distributions as shown in Figure 5a. On the other hand, the strained phase exclusively shows Θ around 17°. The strained pentacene crystals can be explained by the strong interaction between pentacene and graphene during molecular assembly process. The observed zone axis of [0–10] has the most favorable molecular tilting angle toward graphene surface (Figure S2, Supporting Information) with both molecules in a unit cell exhibiting low contact angles. On the other hand, other molecular interfaces, including zone axis [−110], show high molecular tilting orientation toward graphene surface. We can infer that the case with zone axis [010] (or [0–10]) can have the stronger interaction between pentacene crystal and graphene, compared to other pentacene crystal zone axes. Moreover, the longer d(001) for the strained phase means that the pentacene molecules adapt a more face-on structure compared to the bulk-phase. This also supports that the higher degree of interaction between pentacene and graphene is present for the strained phase. To further confirm the above hypothesis, we perform first-principles electronic structure calculations. With density functional theory simulations, including van der Waals interactions between graphene–pentacene interface and pentacene–pentacene intermolecular interactions, we first calculate the misorientation angle Θ dependence of the interfacial energy between pentacene molecules and graphene small 2015, DOI: 10.1002/smll.201403006

(Figure 5b). After we construct a pentacene multilayer thin film with zone axis [010], we calculate the system adsorption energy per molecule as we rotated the pentacene crystal on graphene. Indeed, we find that the system shows the lowest energy at a misorientation angle around 17°. At this configuration, pentacene molecular axis is well aligned with the graphene zigzag lattice direction as previously described in Figures 3c and 4c. This suggests that this stacking relation contributes to the energetically favorable configurations, as the interactions mediated by pz orbitals at the graphene surface and at molecules are enhanced along this graphene crystal axis. We also calculate the interface energy between graphene and different surfaces of pentacene crystal. Figure 5c shows a schematic of different crystal surface of pentacene which can make contact with graphene substrate. From our ab initio calculations, indeed we find that the [0–10] zone axis shows a stronger interaction between the graphene substrate compared to [100] zone axis as the interface energy decreases by about 60 meV per pentacene molecule as shown in Figure 5d. This agreement with experimental observations points that stronger interactions between graphene and this particular zone axis of pentacene may be the reason for inducing selective epitaxial relation as well as the strained phase growth of pentacene. We note that Raman spectra of suspended graphene show no D peak, which indicates the very low concentration of defects in graphene (Figure S4, Supporting



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Information).[32]

Moreover, the surface of suspended graphene after preparation can be quite clean showing contamination-free area bigger than 100 × 100 nm (Figure S1, Supporting Information). The ideal graphene sample allows us to observe the intrinsic epitaxial growth from graphene– pentacene interaction. Scattered residual contamination on graphene surface may be responsible for the growth of nonepitaxial bulk phase of pentacene film.

3. Conclusion We report the rarely-reported strained polymorph of pentacene crystal grown on graphene. The polymorph with d(001) = 14.9 Å shows a specific molecular interfacial orientation and azimuthal alignment to the graphene substrates. The strained phase may be originated from the strong graphene–pentacene interface interaction. Given that various polymorphs show similar values of unit-cell volume within 3%,[21,22,25,26] we find that d(010), layer–layer distance in the vertical direction from graphene surface, should reduce at least 2% from the known bulk-phase. (Please note that the observed d(100) value for the strained phase is the same as the bulk-phase value.) As the vertical interlayer distance is shorter for the strained crystal, this unusual polymorph may exhibit enhanced vertical transport behavior in pentacene– graphene system.[33] The observed strained film growth on graphene demonstrates the possibility to control molecular packing and related vertical charge transport via graphene– molecule interactions with other molecular systems.

4. Experimental Section Graphene Sample Preparation: Graphene was synthesized by chemical vapor deposition on 25 µm thick copper foil.[7] Graphene was transferred to Quantifoil holey carbon TEM grids, following a direct transfer method without using PMMA (poly (methyl methacrylate)) support.[34] With the direct transfer process, the graphene surface is free of polymer residue. The graphene is mostly monolayer and polycrystalline with an average of 5 µm grain size.[30] Pentacene Evaporation: Pentacene was thermally evaporated onto graphene TEM grids. The TEM grids were pre-annealed in air at 200 °C for 30 min to minimize the possible adsorbates on graphene surface prior to the pentacene evaporation. After the pre-annealing process the edge of TEM grids was attached onto substrate with kapton tapes. The substrate temperature was held at 60 °C during deposition and the evaporation rate of 0.2 Å s−1 was used. The vacuum level was around 3 × 10−6 Torr during deposition. Transmission Electron Microscopy: TEM imaging and selective area electron diffraction (SAED) were performed with a FEI Tecnai G2 F20 X-TWIN, operated at 200 kV. The reduced electron beam intensity was used to minimize the e-beam induced damage to pentacene crystals. The diffraction simulations were performed using MacTempas and CrystalKit. The conditions for image acquisition including acceleration voltage of 200 kV were used for diffraction simulations.


Computational Details: Calculations were based on ab initio density-functional-theory using the SIESTA code.[35] The generalized gradient approximation[36] and nonlocal van der Waals density functional[37] were used together with double-zeta plus polarized basis set, norm-conserving Troullier–Martins pseudopotentials[38] and a mesh cutoff of 150 Ry. Atomic coordinates were allowed to relax using a conjugate-gradient algorithm until all forces were smaller in magnitude than 0.01 eV Å−1. Relevant lattice constants (in-plane and out-of-plane) were optimized for each system. To model the system studied in the experiments, we created supercells containing up to 1128 atoms to simulate the interface between pentacene crystals and graphene layers. To avoid interactions between supercell images the distance between periodic images of the pentacene/graphene structures along the direction perpendicular to the graphene-plane was always larger than 20 Å. The resolution of the real-space grid used to calculate the Hartree and exchange-correlation contribution to the total energy was chosen to be equivalent to 150 Ry plane-wave cutoff. The number of k-points was chosen according to the Monkhorst–Pack scheme,[39] and was set to the equivalent of a 45 × 45 × 1 grid in the two-atom primitive unit cell of graphene, which gives well converged values for all the calculated properties. We used a Fermi– Dirac distribution with an electronic temperature of kBT = 21 meV.

Supporting Information Supporting Information is available from the Wiley Online Library or from the author.

Acknowledgements Z.B. acknowledges support from the Stanford Global Climate and Energy Program and the National Science Foundation (Grant No. DMR-1303178). K.K. acknowledges support from the 2014 Research Fund (1.140060.01) of UNIST. E.J.G.S. acknowledges the use of computational resources provided by the Extreme Science and Engineering Discovery Environment (XSEDE), supported by NSF Grant Nos. TG-DMR120049 and TG-PHY120021. T.H.L. acknowledges support from Toshiba Corporation through CIS-FMA program and a fellowship from ILJU foundation in South Korea.

[1] Organic Field-Effect Transistors, (Eds: Z. Bao, J. Lockin), CRC Press, Boca Raton, FL, USA 2007. [2] a) M. Muccini, Nat. Mater. 2006, 5, 605; b) H. Yan, Z. Chen, Y. Zheng, C. Newman, J. R. Quinn, F. Dotz, M. Kastler, A. Facchetti, Nature 2009, 457, 679; c) P. M. Beaujuge, J. M. J. Fréchet, J. Am. Chem. Soc. 2012, 133, 20009; d) C. Wang, H. Dong, W. Hu, Y. Liu, D. Zhu, Chem. Rev. 2012, 112, 2208. [3] a) A. Dodabalapur, L. Torsi, H. E. Katz, Science 1995, 268, 270; b) S. Wang, A. Kiersnowski, W. Pisula, K. Müllen, J. Am. Chem. Soc. 2012, 134, 4015. [4] H. Li, G. Giri, J. B. H. Tok, Z. Bao, MRS Bull. 2013, 38, 34. [5] a) C.-H. Lee, T. Schiros, E. J. G. Santos, B. Kim, K. G. Yager, S. J. Kang, S. Lee, J. Yu, K. Watanabe, T. Taniguchi, J. Hone,


small 2015, DOI: 10.1002/smll.201403006


[6] [7]

[8]

[9]

[10]

[11]

[12] [13]

[14]

[15]

[16]

[17] [18]

[19]

E. Kaxiras, C. Nuckolls, P. Kim, Adv. Mater. 2014, 26, 2812; b) D. He, Y. Zhang, Q. Wu, R. Xu, H. Nan, J. Liu, J. Yao, Z. Wang, S. Yuan, Y. Li, Y. Shi, J. Wang, Z. Ni, L. He, F. Miao, F. Song, H. Xu, K. Watanabe, T. Taniguchi, J.-B. Xu, X. Wang, Nat. Commun. 2014, 5, 5162. M. Liu, X. Yin, E. Ulin-Avila, B. Geng, T. Zentgraf, L. Ju, F. Wang, X. Zhang, Nature 2011, 474, 64. X. S. Li, W. W. Cai, J. H. An, S. Kim, J. Nah, D. X. Yang, R. Piner, A. Velamakanni, I. Jung, E. Tutuc, S. K. Banerjee, L. Colombo, R. S. Ruoff, Science 2009, 324, 1312. a) X. S. Li, Y. W. Zhu, W. W. Cai, M. Borysiak, B. Y. Han, D. Chen, R. D. Piner, L. Colombo, R. S. Ruoff, Nano Lett. 2009, 9, 4359; b) K. S. Kim, Y. Zhao, H. Jang, S. Y. Lee, J. M. Kim, K. S. Kim, J. H. Ahn, P. Kim, J. Y. Choi, B. H. Hong, Nature 2009, 457, 706. a) A. K. Geim, I. V. Grigorieva, Nature 2013, 499, 419; b) L. Britnell, R. V. Gorbachev, R. Jalil, B. D. Belle, F. Schedin, A. Mishchenko, T. Georgiou, M. I. Katsnelson, L. Eaves, S. V. Morozov, N. M. R. Peres, J. Leist, A. K. Geim, K. S. Novoselov, L. A. Ponomarenko, Science 2012, 335, 947; c) L. Britnell, R. M. Ribeiro, A. Eckmann, R. Jalil, B. D. Belle, A. Mishchenko, Y. J. Kim, R. V. Gorbachev, T. Georgiou, S. V. Morozov, A. N. Grigorenko, A. K. Geim, C. Casiraghi, A. H. C. Neto, K. S. Novoselov, Science 2013, 340, 1311; d) Y. Shi, W. Zhou, A.-Y. Lu, W. Fang, Y.-H. Lee, A. L. Hsu, S. M. Kim, K. K. Kim, H. Y. Yang, L.-J. Li, J.-C. Idrobo, J. Kong, Nano Lett. 2012, 12, 2784; e) W. Dang, H. Peng, H. Li, P. Wang, Z. Liu, Nano Lett. 2010, 10, 2870. a) J. Wu, H. A. Becerril, Z. Bao, Z. Liu, Y. Chen, P. Peumans, Appl. Phys. Lett. 2008, 92, 263302; b) L. Gomez De Arco, Y. Zhang, C. W. Schlenker, K. Ryu, M. E. Thompson, C. Zhou, ACS Nano 2010, 4, 2865; c) Y. Wang, S. W. Tong, X. F. Xu, B. Özyilmaz, K. P. Loh, Adv. Mater. 2011, 23, 1514. a) J. Wu, M. Agrawal, H. A. Becerril, Z. Bao, Z. Liu, Y. Chen, P. Peumans, ACS Nano 2009, 4, 43; b) T.-H. Han, Y. Lee, M.-R. Choi, S.-H. Woo, S.-H. Bae, B. H. Hong, J.-H. Ahn, T.-W. Lee, Nat. Photon. 2012, 6, 105. W. H. Lee, J. Park, S. H. Sim, S. Lim, K. S. Kim, B. H. Hong, K. Cho, J. Am. Chem. Soc. 2011, 133, 4447. a) G. Hong, Q.-H. Wu, J. Ren, C. Wang, W. Zhang, S.-T. Lee, Nano Today 2013, 8, 388; b) J. M. MacLeod, F. Rosei, Small 2014, 10, 1038. G. Giri, E. Verploegen, S. C. B. Mannsfeld, S. Atahan-Evrenk, D. H. Kim, S. Y. Lee, H. A. Becerril, A. Aspuru-Guzik, M. F. Toney, Z. Bao, Nature 2011, 480, 504. a) N. Koch, A. Vollmer, I. Salzmann, B. Nickel, H. Weiss, J. P. Rabe, Phys. Rev. Lett. 2006, 96, 156803; b) D. E. Hooks, T. Fritz, M. D. Ward, Adv. Mater. 2001, 13, 227. a) Q. H. Wang, M. C. Hersam, Nat. Chem. 2009, 1, 206; b)W. Chen, H. Huang, A. Thye, S. Wee, Chem. Commun. 2008, 36, 4276; c) H. Zhou, L. Zhang, J. Mao, G. Li, Y. Zhang, Y. Wang, S. Du, W. Hofer, H.-J. Gao, Nano Res. 2013, 6, 131. J. Götzen, D. Käfer, C. Wöll, G. Witte, Phys. Rev. B 2010, 81, 085440. I. Salzmann, A. Moser, M. Oehzelt, T. Breuer, X. Feng, Z.-Y. Juang, D. Nabok, R. G. Della Valle, S. Duhm, G. Heimel, A. Brillante, E. Venuti, I. Bilotti, C. Christodoulou, J. Frisch, P. Puschnig, C. Draxl, G. Witte, K. Müllen, N. Koch, ACS Nano 2012, 6, 10874. a) J. C. Meyer, C. O. Girit, M. F. Crommie, A. Zettl, Nature 2008, 454, 319; b) R. Erni, M. D. Rossell, M.-T. Nguyen, S. Blankenburg, D. Passerone, P. Hartel, N. Alem, K. Erickson, W. Gannett, A. Zettl, Phys. Rev. B 2010, 82, 165443.

small 2015, DOI: 10.1002/smll.201403006

[20] a) C. D. Dimitrakopoulos, A. R. Brown, A. Pomp, J. Appl. Phys. 1996, 80, 2501; b) O. D. Jurchescu, M. Popinciuc, B. J. van Wees, T. T. M. Palstra, Adv. Mater. 2007, 19, 688. [21] a) D. Holmes, S. Kumaraswamy, A. J. Matzger, K. P. C. Vollhardt, Chem. Eur. J. 1999, 5, 3399; b) C. C. Mattheus, A. B. Dros, J. Baas, A. Meetsma, J. L. d. Boer, T. T. M. Palstra, Acta Cryst. 2001, C57, 939; c) T. Siegrist, C. Kloc, J. H. Schön, B. Batlogg, R. C. Haddon, S. Berg, G. A. Thomas, Angew. Chem. Int. Ed. 2001, 40, 1732. [22] a) R. B. Campbell, J. M. Robertson, J. Trotter, Acta Cryst. 1961, 14, 705; b) R. B. Campbell, J. M. Robertson, J. Trotter, Acta Cryst. 1962, 15, 289. [23] C. C. Mattheus, A. B. Dros, J. Baas, G. T. Oostergetel, A. Meetsma, J. L. de Boer, T. T. M. Palstra, Synth. Met. 2003, 138, 475. [24] a) S. E. Fritz, S. M. Martin, C. D. Frisbie, M. D. Ward, M. F. Toney, J. Am. Chem. Soc. 2004, 126, 4084; b) S. C. B. Mannsfeld, A. Virkar, C. Reese, M. F. Toney, Z. Bao, Adv. Mater. 2009, 21, 2294. [25] a) H. Yoshida, K. Inaba, N. Sato, Appl. Phys. Lett. 2007, 90, 181930; b) S. Schiefer, M. Huth, A. Dobrinevski, B. Nickel, J. Am. Chem. Soc. 2007, 129, 10316; c) D. Nabok, P. Puschnig, C. Ambrosch-Draxl, O. Werzer, R. Resel, D.-M. Smilgies, Phys. Rev. B 2007, 76, 235322. [26] a) T. Siegrist, C. Besnard, S. Haas, M. Schiltz, P. Pattison, D. Chernyshov, B. Batlogg, C. Kloc, Adv. Mater. 2007, 19, 2079; b) L. Farina, A. Brillante, R. G. Della Valle, E. Venuti, M. Amboage, K. Syassen, Chem. Phys. Lett. 2003, 375, 490. [27] G. E. Thayer, J. T. Sadowski, F. Meyer zu Heringdorf, T. Sakurai, R. M. Tromp, Phys. Rev. Lett. 2005, 95, 256106. [28] a) J. S. Wu, J. C. H. Spence, J. Appl. Cryst. 2004, 37, 78; b) Y. Murakami, S. Tomiya, N. Koshitani, Y. Kudo, K. Satori, M. Itabashi, N. Kobayashi, K. Nomoto, Phys. Rev. Lett. 2009, 103, 146102; c) L. F. Drummy, D. C. Martin, Adv. Mater. 2005, 17, 903. [29] F. Balzer, H. H. Henrichsen, M. B. Klarskov, T. J. Booth, R. Sun, J. Parisi, M. Schiek, P. Bøggild, Nanotechnology 2014, 25, 035602. [30] K. Kim, Z. Lee, W. Regan, C. Kisielowski, M. F. Crommie, A. Zettl, ACS Nano 2011, 5, 2142. [31] A. S. Eggeman, S. Illig, A. Troisi, H. Sirringhaus, P. A. Midgley, Nat. Mater. 2013, 12, 1045. [32] a) L. M. Malard, M. A. Pimenta, G. Dresselhaus, M. S. Dresselhaus, Phys. Rep. 2009, 473, 51; b) L. G. Cançado, A. Jorio, E. H. Martins Ferreria, F. Stavale, C. A. Achete, R. B. Capaz, M. V. O. Moutinho, A. Lombardo, T. S. Kulmala, A. C. Ferrari, Nano Lett. 2011, 11, 3190. [33] a) C. Ojeda-Aristizabal, W. Bao, M. S. Fuhrer, Phys. Rev. B 2013, 88, 035435; b) K. Berke, S. Tongay, M. A. McCarthy, A. G. Rinzler, B. R. Appleton, F. Hebard, J. Phys.: Condens. Matter 2012, 24, 255802. [34] W. Regan, N. Alem, B. Aleman, B. S. Geng, C. Girit, L. Maserati, F. Wang, M. Crommie, A. Zettl, Appl. Phys. Lett. 2010, 96, 113102. [35] J. M. Soler, E. Artacho, J. D. Gale, A. Garcia, J. Junquera, P. Ordejon, D. Sanchez-Portal, J. Phys.: Condens. Matter 2002, 14, 2745. [36] J. P. Perdew, K. Burke, M. Ernzerhof, Phys. Rev. Lett. 1996, 77, 3865. [37] M. Dion, H. Rydberg, E. Schröder, D. C. Langreth, B. I. Lundqvist, Phys. Rev. Lett. 2004, 92, 246401. [38] N. Troullier, J. L. Martins, Phys. Rev. B 1991, 43, 1993. [39] H. J. Monkhorst, J. D. Pack, Phys. Rev. B 1976, 13, 5188.


Received: October 9, 2014 Revised: November 20, 2014 Published online:


7

Graphene-based electrodes for enhanced organic thin film transistors based on pentacene.

Thin film MnGe grown on Si(111).

Effects of the F₄TCNQ-Doped Pentacene Interlayers on Performance Improvement of Top-Contact Pentacene-Based Organic Thin-Film Transistors.

Epitaxially stabilized thin films of ε-Fe2O3 (001) grown on YSZ (100).

Blending effect of 6,13-bis(triisopropylsilylethynyl) pentacene-graphene composite layers for flexible thin film transistors with a polymer gate dielectric.

Straining graphene using thin film shrinkage methods.

Epitaxially grown BaM hexaferrite films having uniaxial axis in the film plane for self-biased devices.

h-BN Heterostructure.

BiVO4 thin film photoanodes grown by chemical vapor deposition.

Development of an ultra-thin film comprised of a graphene membrane and carbon nanotube vein support.

Optimizing pentacene thin-film transistor performance: Temperature and surface condition induced layer growth modification.

Interaction of oxygen with samarium on Al2O3 thin film grown on Ni3Al(111).

Novel top-contact monolayer pentacene-based thin-film transistor for ammonia gas detection.

In situ preparation, electrical and surface analytical characterization of pentacene thin film transistors.

High-mobility ambipolar ZnO-graphene hybrid thin film transistors.

Suppression of copper thin film loss during graphene synthesis.

Wrinkling of a stiff thin film bonded to a pre-strained, compliant substrate with finite thickness.

Ferromagnetism and Ru-Ru distance in SrRuO3 thin film grown on SrTiO3 (111) substrate.

Ultrasensitive room temperature NH3 sensor based on a graphene-polyaniline hybrid loaded on PET thin film.

Quasi-stabilized hydration layers on muscovite mica under a thin water film grown from humid air.

Magnetotransport in an aluminum thin film on a GaAs substrate grown by molecular beam epitaxy.

Optical properties of epitaxial BiFeO3 thin film grown on SrRuO3-buffered SrTiO3 substrate.

Persistent topological surface state at the interface of Bi2Se3 film grown on patterned graphene.

Nanofiber supported thin-film composite membrane for pressure-retarded osmosis.