Intrinsic and Extrinsic Parameters for Controlling the Growth of

Jian-Yao Zheng , Hongjun Xu , Jing Jing Wang , Sinéad Winters , Carlo Motta .... A. Fischer , Dean M. DeLongchamp , Todd Emrick , and Alejandro L. Br...
0 downloads 0 Views 8MB Size
Download PDF
Letter pubs.acs.org/NanoLett

Intrinsic and Extrinsic Parameters for Controlling the Growth of Organic Single-Crystalline Nanopillars in Photovoltaics Yue Zhang,† Ying Diao,‡,§ Hyunbok Lee,† Timothy J. Mirabito,† Richard W. Johnson,† Egle Puodziukynaite,† Jacob John,† Kenneth R. Carter,† Todd Emrick,† Stefan C. B. Mannsfeld,∥,§ and Alejandro L. Briseno*,† †

Department of Polymer Science and Engineering, University of Massachusetts, Amherst, Massachusetts 01003, United States Department of Chemical Engineering, Stanford University, Stanford, California 94305, United States § Stanford Synchrotron Radiation Light Source, SLAC National Accelerator Laboratory, Menlo Park, California 94025, United States ∥ Dresden University of Technology, Dresden, 01062, Germany ‡

S Supporting Information *

ABSTRACT: The most efficient architecture for achieving high donor/acceptor interfacial area in organic photovoltaics (OPVs) would employ arrays of vertically interdigitated p- and n- type semiconductor nanopillars (NPs). Such morphology could have an advantage in bulk heterojunction systems; however, precise control of the dimension morphology in a crystalline, interpenetrating architecture has not yet been realized. Here we present a simple, yet facile, crystallization technique for the growth of vertically oriented NPs utilizing a modified thermal evaporation technique that hinges on a fast deposition rate, short substrate−source distance, and ballistic mass transport. A broad range of organic semiconductor materials is beneficial from the technique to generate NP geometries. Moreover, this technique can also be generalized to various substrates, namely, graphene, PEDOT−PSS, ZnO, CuI, MoO3, and MoS2. The advantage of the NP architecture over the conventional thin film counterpart is demonstrated with an increase of power conversion efficiency of 32% in photovoltaics. This technique will advance the knowledge of organic semiconductor crystallization and create opportunities for the fabrication and processing of NPs for applications that include solar cells, charge storage devices, sensors, and vertical transistors. KEYWORDS: 3D-nanopillar, 2D-3D thin film transition, graphene, single-crystalline, organic photovoltaic (OPV), crystallization mechanism

M

spaced single-crystalline NPs with diameters equal to or less than twice the exciton diffusion length (∼10−25 nm) of the material under study.11 However, to date, this ideal architecture has not yet been realized due to (i) the difficulty of controlling molecular orientation, packing, and crystallization at the substrate interface, and (ii) the inability to adapt inorganic deposition techniques to organic materials. For semiconductor inorganic materials,14−18 vertical architectures have been demonstrated with materials such as ZnO,19 Si,20 and GaAs.21,22 In regards to organic materials, NP-like geometries have been demonstrated with small-molecule organic and polymer semiconductor materials.11,13,23−25 However, for such systems, the length scales of such architectures are not ideal, and the materials used are either amorphous or semicrystalline which limits the exciton diffusion length. In addition, the adapted processing techniques such as the template-induced or

olecular orientation and ordering at interfaces in organic semiconductors significantly affects the electronic structure and charge transport properties.1 Therefore, the anisotropy of the transport properties in organic semiconductor thin films in particular must be taken into account. However, it is not only the orientation of the film as a whole that is important, but the molecular orientation of the first few monolayers near the interfaces which may significantly affect the electronic interfacial properties,2,3 such as the electronic trap states,4 contact resistances,5 or interface dipoles.6 In recent years, tremendous efforts have been made to produce the ideal architecture for harvesting energy from the sunlight.7−9 Organic photovoltaics (OPVs) have, to date, demonstrated a power conversion efficiency exceeding 10%.10 OPV designs have evolved through different architectures in efforts to optimize energy harvesting and charge transport. In principle, the most efficient geometry would employ nanopillarlike vertical arrays of p- and n- type organic semiconductor materials.7,11−13 The ideal architecture should have a ∼50 nm thick electron and donor layer, structured as an array of equally © XXXX American Chemical Society

Received: May 24, 2014 Revised: September 2, 2014

A

dx.doi.org/10.1021/nl501933q | Nano Lett. XXXX, XXX, XXX−XXX

Nano Letters

Letter

Figure 1. Technique for growing organic 3-dimensional NPs and growth strategy. (a) Illustration of the crystallization apparatus and the corresponding temperature gradient along an evacuated tube furnace. (b) Procedure for crystallizing vertically oriented single crystal NPs onto graphene-coated substrates.

Figure 2. Representative SEM images of vertically oriented nanopillars grown on graphene substrates from p-type materials: (a) CuPc, (b) ZnPc, and (c) pentacene. Nanopillars grown from n-type materials: (d) F16CuPc (e) PTCDA, and (f) C60. Nanopillars selectively crystallized on patterned graphene from (g, h) CuPc, (i) PTCDA, (j) patterned “UMASS” from CuPc, and (k) cross-section SEM image of CuPc NPs. B

dx.doi.org/10.1021/nl501933q | Nano Lett. XXXX, XXX, XXX−XXX

Nano Letters

Letter

Figure 3. SEM micrographs of the morphology evolution and crystallization kinetics of growth behavior. (a) Morphology evolution of ttb-CuPc films deposited on graphene with different growth times of 0−16 min (scale bars = 1 μm). (b) Film thickness and temperature profile as a function of time for the crystallization of ttb-CuPc on Si/graphene substrates. The film growth rate dramatically increases at the time mark of ∼5 min. (c) Nucleation density and aspect ratio of ttb-CuPc nanostructures as a function of time. A clear distinction of morphology transition occurs at 4.5 min of crystallization. The data from (c) were acquired from images in (a).

seeded growth method are complicated.26 Moreover, there remains a lack of understanding of the critical factors and governing mechanisms in creating vertically oriented, threedimensional, and highly crystalline NPs from organic semiconductors. Our approach to form arrays of single crystalline organic semiconductor NPs employs a modified thermal evaporation technique that enables rapid and reproducible crystallization of NPs on graphene from a wide variety of organic semiconductors. Initially, we obtain a thin, crystalline “wetting layer” grown at slower deposition rates and low substrate temperatures, followed by a strain-induced morphology evolution that yields an array of densely packed oriented NPs grown at considerably faster deposition rates and higher substrate temperatures on graphene substrates. In this process, a morphology transition from a two-dimensional (2D) thin film to a three-dimensional (3D) NP morphology is observed. The technique affords dense crystalline arrays of vertical NPs from a broad range of p- and n- type organic semiconductors.

Moreover, we discovered that by using different parameters, the technique can be generalized to most substrates. We describe the crystallization mechanism and support our model with Stranski−Krastanov (SK) growth kinetics and X-ray diffraction experiments. To demonstrate the NP application in devices, OPVs with active layers fabricated by this method exhibit a 32% increase in power conversion efficiency over thin films of the same materials, with efficiencies reaching 3%. The outcome from this study will open new routes for improving performance of organic electronics devices with large surface area single-crystalline vertically oriented NPs and be applicable to vertical transistors, sensors, photodetectors, and energy storage devices. We first demonstrate our approach of oriented growth on graphene-coated substrates as a model crystallization study. Graphene, a two-dimensional sheet of carbon atoms, has been shown to induce epitaxial growth of organic adsorbate molecules and is used to characterize self-assembled organic monolayers with molecular resolution.27−30 It has also been C

dx.doi.org/10.1021/nl501933q | Nano Lett. XXXX, XXX, XXX−XXX

Nano Letters

Letter

Figure 4. (a) GIXD scattering patterns of ttb-CuPc on graphene at 1−3, 3.5, and 6 min deposition, and (b) the schematic of the two different growth regimes: 2D growth within 1−2 min (left) and 3D growth >3 min (right).

growth of crystal nuclei. At ∼4.5 min a distinct morphology transition was observed as the nucleation density became saturated. We monitored crystallization from the moment the evacuated tube was inserted into the furnace. Approximately 4 min is required to reach a steady state temperature at graphene substrate. It should be noted that after 4 min, the source/ substrate temperature and the film deposition achieved local thermal equilibrium and a constant rate of deposition, respectively. To understand the crystallization kinetics, we determined the deposition rate and film thickness by depositing organic materials onto a graphene-coated quartz crystal (Figure 3b). From the slope of film thickness versus time data, we determined that the deposition rate at steady state was ∼4.3 Å/ s after about 4 min of crystallization. Such a deposition rate is nearly 1 order of magnitude faster than a typical deposition rate for small molecule semiconductor thin films.40 Source temperatures varied and were dependent on the organic material used. A plot of temperature and film thickness as a function of time is shown in Figure 3b. Figure 3c shows the nucleation density and aspect ratio of the crystals as a function of crystallization time. The initial formation of crystal nuclei is obvious from the shoulder at 2−4 min from the nucleation density profile. The profile also shows how the nucleation density stabilizes at about 4−5 min. Beyond this time frame, the discrete nuclei evolve into 3D NPs as shown from the SEM micrographs in Figure 3a. It is also clear that, beyond the 4−5 min mark, the aspect ratio continuously increases, roughly increasing exponentially with time. To review these observations, we identified three morphology evolution events. First, a wetting layer consisting of three to six molecular layers with a critical thickness of 1−2 nm is observed within the first few minutes. The formation of the 2D wetting layer is likely driven by the strong graphene− adsorbates interactions. Subsequent nucleation events follow, which initiate the 2D to 3D growth transition, and ultimately the nucleation density saturates at about 4−5 min. Finally, the most obvious morphology evolution is observed beyond 5 min, where the growth of individual single crystalline NPs now becomes anisotropic along the growth axis. To better interpret the morphology evolution events, and track molecular packing and orientation, we performed grazingincidence X-ray diffraction (GIXD) experiments, as shown in

used as a template for growing highly crystalline oriented thin films from several organic semiconductors such as P3HT,31 perfluoropentacene,32 CuPc,33,34 PTCDA,28,35 and tetraazaterrylene (TAT).36,37 In spite of these efforts, no reports have documented the growth of high density arrays of vertically oriented single crystalline NPs on graphene from a broad range of organic compounds. While several groups have demonstrated vertical growth of organic NPs on nongraphene substrates,7,24,38,39 there is no mention of a detailed growth mechanism or applicability to a library of organic semiconductors. Figure 1 shows a schematic of the crystallization apparatus and temperature profile along the furnace zone from source to substrate and the process by which crystalline organic NPs are formed on graphene substrates. As a model growth system, we crystallized tetra-tert-butyl-copper phthalocyanine (ttb-CuPc) on graphene to understand the growth kinetics and crystallization behavior. As shown in Figure 1a, graphene substrates were placed about 5 cm above the source material in an evacuated tube at a reduced pressure, varied from 10−2 mbar to 10−5 mbar. The experiment was conducted by inserting a tube containing the source compound and graphene substrate into a temperature gradient furnace at a set elevated source temperature. We found that an ultrathin film was first formed, followed by a strain-induced morphology transition that produced 3D single-crystalline NPs (Figure 1b). As shown in Figure 2, the crystallization process is applicable to a variety of p-type semiconductors including metal phthalocyanine derivatives (ttb-CuPc, ZnPc), tetraphenyldibenzoperiflanthene (DBP), and pentacene (SI). In addition, n-type semiconductors such as fluorinated CuPc (F16CuPc), perylene tetracarboxylic− dianhydride (PTCDA), and C60/C70 were also employed in our studies. We also demonstrate selective crystallization by crystallizing NPs onto patterned graphene using nanoimprint lithography to define features (Figure 2g−j). To understand the nucleation and growth behavior on graphene substrates, a systematic study on thin film morphology evolution was conducted using ttb-CuPc as a model small molecule semiconductor. Figure 3a shows SEM images of ttb-CuPc films grown at different stages of growth. We observed a wetting layer within 1 min, followed by the D

dx.doi.org/10.1021/nl501933q | Nano Lett. XXXX, XXX, XXX−XXX

Nano Letters

Letter

Figure 5. Characterization of 2D film and 3D NPs. (a) TEM image of a ttb-CuPc crystal grown at 3.5 min on graphene with the corresponding reciprocal electron diffraction and the indicated [010] growth direction. (b) A high-resolution image of the 2D and 3D (i.e., nanopillar) layer with indicated d-spacings obtained from the fast Fourier transform (FFT) analysis.

whether there is actually epitaxial growth of the wetting layer in the present case, not only the unit cell of the ttb-CuPc initial layers but also its azimuthal angle relative to the graphene lattice would need to be known. Therefore, we can neither rule out nor confirm that a combination of initial epitaxial growth and subsequent strain release (SK model) is responsible for the observed 2D−3D transition. Even if it were in the case of ttbCuPc/graphene, it would still be difficult to explain the relative robustness with which we detect the 2D−3D transitions for a variety of different materials since the exact mode of epitaxy (coincident, commensurate, etc.) of organic materials on crystalline substrates depends on the specific material/substrate combination.44,45 However, the formation of the initial wetting layer can also be discussed in the context of a model for thin film growth that is more general and does not require any epitaxial growth on crystalline substrates. According to Verlaak et al.,46 2D growth is energetically favorable for any supersaturation (and even undersaturation) of the molecular vapor so long as the intermolecular interactions between a given molecule and all surrounding molecules in the same molecular layer (ψc) are smaller than the interaction strength between this molecule and the substrate (ψmol−sub), i.e. ψc−ψmol−sub < 0. Previous DFT calculations on the CuPc system indicate that, indeed, the CuPc/graphene binding energy ψc (3.37 eV33) is higher than the CuPc interlayer binding energy ψmol−sub (2.57 eV47), and a 2D wetting layer formation at the initial stage of deposition is therefore energetically favorable for this and similar phthalocyanine derivatives on graphene. This has also been observed experimentally for many large aromatic molecules on graphite by thermal desorption spectroscopy that shows that the first monolayer is bound significantly stronger to a graphite substrate than the second and subsequent layers.48−50 The Verlaak model predicts that 2D nucleation is favorable in the initial growth stage and even at low supersaturation levels. After a few layers, however, ψmol−sub is no longer significantly larger than ψc since with increasing distance from the more strongly binding, semimetallic graphene surface, the top layer of ttbCuPc effectively becomes the substrate for subsequent deposition. Beyond this point and especially for nonequilibrium conditions such as low substrate temperatures and high molecular flux, the question of whether the growth mode remains 2D or changes to 3D cannot be sufficiently described by the thermodynamic equilibrium energetics. In the following, we want to discuss some of the dynamic and kinetic aspects involved in this growth process. Vertical NPs

Figure 4. After only 1 min deposition, GIXD patterns show the signature of a 2D crystalline layer with an in-plane d-spacing of 13.6 Å; after 2 min, this spacing expanded to 16.7 Å. The absence of out-of-plane peaks or any out-of-plane modulation of the in-plane Bragg rod intensity suggests that at this stage, the film is only a few molecular layers thick (i.e., 2D). Both inplane d-spacing values are compatible with a face-on configuration of the ttb-CuPc molecules, and it is plausible that an initially strained, possibly epitaxial phase (13.6 Å) that forms directly on the graphene surface is replaced by an equilibrium phase (16.7 Å) as material is accumulated. The latter value is also close to the dimensions of a theoretically calculated 2D unit cell of flat-lying ttb-CuPc molecules (optimized potentials for liquid simulations molecular force field calculations yield 17.6 Å). At the 2D to 3D transition after 3.5 min of deposition, a faceon (π−π) stacking peak appears along the Qz vector, with a dspacing of 3.3 Å. This d-spacing corresponds to the well-known π−π stacking distance found in CuPc crystals.41 We verified the intermolecular distances by carrying out TEM diffraction experiments on discrete NPs, as shown in Figure 5a. Figure 5b shows a TEM image of a 2D film with crystal nuclei that was grown on graphene film. High-resolution TEM (Figure 5c) indicates that molecules in the 2D wetting layer crystallize onto graphene with a face-on packing motif. Fast Fourier transform (FFT) analysis of these images yielded an estimated in-plane spacing of 16.8 Å which is comparable with the value obtained from the GIXD images. The mechanism responsible for the formation of vertical NPs atop the initial wetting layer appears largely materialindependent. In the following we will discuss the formation of a wetting layer and the transition to 3D island growth from both energetic (thermodynamic) and kinetic points of view, starting with the energy-based arguments. One model discussed in literature which describes 2D−3D morphology transitions in inorganic thin film growth is the SK growth model.42,43 According to the SK model, the first few layers exhibit commensurate epitaxy with respect to the substrate surface lattice thereby minimizing the interface energy and promoting 2D growth. However, the associated energetic penalty within the growing film (strain) accumulates during the growth of subsequently deposited layers, providing an energetic driving force for the eventual 2D−3D transition. As already noted above, there is indeed some evidence for an initially compressed (strained) ttb-CuPc lattice in the GIXD data from the initial layer in Figure 4a. However, in order to assess E

dx.doi.org/10.1021/nl501933q | Nano Lett. XXXX, XXX, XXX−XXX

Nano Letters

Letter

Figure 6. Characterization of electronic devices. Parts a, b, and c show schematic configuration of ttb-CuPc hole-only, C60 electron-only devices, and inverted OPV cells. Parts d and e show J−V characteristics of carrier-only devices. (f) J−V characteristics of OPVs under illumination with AM 1.5 conditions. The performance of OPVs with different configurations are compared, specifically, with and without PSBMA-Py interlayer (IL) and between C70 TF and C70 NP, respectively. The OPV device of C70 NP with PSBMA-Py IL (red filled circle) shows the highest PCE of 2.9%.

growth requires that the rate of crystal growth in the vertical (out-of-plane) direction be much higher than in the horizontal (in-plane) directions. A large difference in the crystal growth rates can be caused by either intrinsic or extrinsic parameters, or a combination of both. Intrinsic parameters refer to material specific factors that dictate crystal growth rate anisotropy, such as the intermolecular coupling strength and its spatial anisotropy.51 Extrinsic parameters are those that strongly influence the system kinetics and mass transport such as the substrate temperature and the rates of evaporation and desorption. Owing to the intrinsic shape anisotropy of the ttb-CuPc molecule and the corresponding anisotropy in the intermolecular coupling, one would expect the fastest growth axis to be along the π−π stacking direction for ttb-CuPc, which would indeed favor vertical growth. The role of intrinsic parameters manifests in the observation that the graphene-directed NP growth is not entirely system-independent. For instance, at similar conditions as those used for ttb-CuPc, C60 and pentacene form dome-like features instead. Neither of the latter two molecules exhibits a coupling anisotropy as large as that of ttb-CuPc or other phthalocyanines which exhibit a dominant π−π stacking direction (which leads to their “quasi3D” electro-optical properties). Therefore, under close-toequilibrium crystal growth conditions (e.g., at high substrate temperatures near the source temperature), it is expected that C60 and pentacene, unlike ttb-CuPc and similar derivatives, form lower aspect ratio crystallites. However, the extrinsic parameters play a dominant role in the formation of vertical NP from the materials used in this study. This is evident from the fact that the NP formation critically depends on a short source−substrate distance, fast deposition rate (0.4−0.5 nm/s), and relatively low substrate temperatures (80 °C). These conditions increase the chemical potential difference between the source and the substrate and thus the supersaturation at the growth front and, therefore, favor crystal growth toward the source, i.e., in the vertical direction. The low substrate temperature promotes the vertical

growth in two ways: (i) the low temperature suppresses an efficient thermally activated (diffusion) mass transport of monomers across the surface of the growing wires that could otherwise offset the vertical mass transport imposed by the molecular beam; (ii) a low substrate temperature also limits the rate of desorption, essentially keeping the (isotropic) molecular vapor pressure low near the substrate and the vertical chemical potential gradient (and supersaturation) high. A related view was given in earlier works discussing the role of the ballistic regime of the gas flow on the film growth.52,53 When the substrate is placed at a distance from the source less than the mean free path, the transport of organic molecules from the source to the substrate is ballistic with a molecular flux that is predominantly vertical to the substrate. In the ballistic regime out-of-substrate plane growth that increases the roughness is favored over in-plane growth that reduces film roughness.47 In our control study with a low chamber pressure of 5 × 10−6 mbar, the mean free path of the gas molecules is ∼2 m. Summarizing this discussion, we assert that the initial wetting layer is the result of the fairly strong molecule−substrate interactions with the graphene substrate. The transition to 3D growth which results in the formation of the vertical nanowires is governed by both intrinsic factors such as crystal growth rate (fastest along the π−π stacking direction) but more so extrinsic factors such as a large chemical potential gradient and a low substrate temperature which were found to be the key factors in the vertical nanopillar growth across a wide range of material systems. To demonstrate the advantage of vertically oriented crystalline nanopillars (NP) over thin films (TF) on the charge transport, we investigated the hole- and electron-only devices. The one-carrier devices are comprised of a 200-nm-thick layer of ttb-CuPc (hole-only) or C60 (electron-only) deposited on top of ITO/graphene glass substrate with the top electrode, as shown in Figure 6a and b. The single-carrier devices of NP and TF were controlled by both intrinsic (with vs without graphene) and extrinsic (fast vs slow deposition rate) factors. F

dx.doi.org/10.1021/nl501933q | Nano Lett. XXXX, XXX, XXX−XXX

Nano Letters

Letter

This efficiency increase is largely attributed to the enhanced interfacial area based on the interconnected vertical NPs, creating more available sites for exciton dissociation. This study will provide a route toward realizing efficient OPVs and may have implications in various other organic electronic applications.

A comparison of the current density−voltage (J−V) characteristics of such devices is shown in Figure 6d and e. For ttb-CuPc hole-only device, the charge transport of NP devices appears two orders of magnitude higher than that of TF devices. This improvement can be attributed to well-ordered and face-on stacking of ttb-CuPc molecules on a graphene substrate, which enables the efficient out-of-plane charge transport along surface normal direction.34,54 Similarly, in the case of C60, a greatly improved charge transport was observed in C60 NP electrononly device as well. Note that a thin electron injection interlayer (IL) of the zwitterionic poly(sulfobetaine methyl methacrylate)−pyrene (PSBMA-Py) was inserted between the C60 layer and the ITO/graphene substrate to reduce the electron injection barrier between C60 and ITO (SI). Therefore, both hole- and electron-only devices with NP configuration perform more efficiently with out-of-plane charge transport than TF morphology counterparts. Furthermore, the vertically oriented crystalline NPs can offer an ideal device architecture for OPVs of interdigitated heterojunction. In this regard, we fabricated inverted OPVs using C70 as an acceptor layer and tetraphenyldibenzoperiflanthene (DBP) as a donor layer. The schematic device configuration and J−V characteristics of C70/DBP inverted OPVs are shown in Figure 6c and f. The C70 layer was controlled by varying deposition rates for comparison between the TF and the NP structure, while a graphene layer insertion, an intrinsic factor, is not changed. The device parameters of OPVs are summarized in Table S1. The power conversion efficiency (PCE) increases from 2.2% with the C70 TF from a slow deposition rate (0.1 nm/s) to 2.9% with the vertically oriented C70 NPs from a fast deposition rate (0.4 nm/s). This improvement is mainly attributed to the increased short circuit current density (JSC) from 5.37 mA/cm2 of thin film OPVs to 6.24 mA/cm2 of NP OPVs, which originates from the enlarged interfacial area at donor/acceptor interface from its interdigitated geometry. We show that, with the IL, OPVs show higher PCEs than the counterpart thin film samples. We note that our vertically oriented crystalline NP OPVs do not deteriorate the fill factor (FF), rather increased, although it clearly increases the JSC. Generally, it is known that OPVs with bulk heterojunction (BHJ) architecture show higher JSC due to the larger donor/ acceptor interface area to dissociate the photogenerated excitons but deteriorates its FF due to the disordered charge transport to an anode and a cathode. However, our device design can realize enhancement of both JSC and FF with vertically oriented C70 crystals by simple morphology control. In conclusion, our general and easy route to generate vertically oriented NPs will provide a new efficient method for ideal OPV architectures which have both a higher charge transport and larger interfacial area. In summary, we have demonstrated a facile and general approach to obtain vertically oriented NPs of single-crystalline small molecule organic semiconductors. The method presented here represents a powerful way to grow vertical NPs from different p- and n-type semiconductors onto various common substrates used in electronic devices. A clear transition from 2D to 3D growth mode is illustrated and explained where both the intrinsic (competition between intermolecular interactions and substrate−molecule interactions) and extrinsic factors (fast deposition and ballistic mass transport) played crucial roles. Furthermore, the structure superiority of vertically oriented NPs over a conventional bilayer thin film system is evidenced by the achieved PCE of 2.9%, which is a 32% improvement.

■

ASSOCIATED CONTENT

S Supporting Information *

Graphene preparation, crystallization method, morphology and crystallographic characterization, fabrication of patterned graphene substrates and one-carrier only device and OPV fabrication methods, absorbance of devices with C70 TF and C70 NP, and cross-section SEM images of DBP/C70. This material is available free of charge via the Internet at http:// pubs.acs.org.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS

■

REFERENCES

Y.Z., H.L., T.J.M., and A.L.B. acknowledge support by the Offi ce of Naval Research (N000141110636 and N0001471410053). T.E. is thankful for the support by NSFCHE 1152360. K.R.C and J.J. thank the NSF Nanoscale Science and Engineering Center for Hierarchical Manufacturing (CHM) at the University of Massachusetts, Amherst, MA, USA (Grant No. CMMI-1025020). Portions of this research were carried out at the Stanford Synchrotron Radiation Lightsource, a Directorate of SLAC National Accelerator Laboratory and an Office of Science user Facility operated for the U.S. Department of Energy Office of Science by Stanford University. We also acknowledge William Dichtel for teaching us growing graphene. We thank Yue Wang for the generous help on HRTEM characterization and Cheng Li for SEM images for Figure S6.

(1) Oosterhout, S. D.; Wienk, M. M.; van Bavel, S. S.; Thiedmann, R.; Koster, L. J. A.; Gilot, J.; Loos, J.; Schmidt, V.; Janssen, R. A. J. Nat. Mater. 2009, 8, 818−824. (2) Schunemann, C.; Wynands, D.; Eichhorn, K. J.; Stamm, M.; Leo, K.; Riede, M. J. Phys. Chem. C 2013, 117, 11600−11609. (3) Rand, B. P.; Cheyns, D.; Vasseur, K.; Giebink, N. C.; Mothy, S.; Yi, Y.; Coropceanu, V.; Beljonne, D.; Cornil, J.; Brédas, J.-L.; Genoe, J. Adv. Funct. Mater. 2012, 22, 2987−2995. (4) Cattin, L.; Bernede, J. C.; Lare, Y.; Dabos-Seignon, S.; Stephant, N.; Morsli, M.; Zamora, P. P.; Diaz, F. R.; del Valle, M. A. Phys. Status Solidi A 2013, 210, 802−808. (5) Makha, M.; Cattin, L.; Lare, Y.; Barkat, L.; Morsli, M.; Addou, M.; Khelil, A.; Bernede, J. C. Appl. Phys. Lett. 2012, 101, 233307. (6) Vasseur, K.; Rand, B. P.; Cheyns, D.; Temst, K.; Froyen, L.; Heremans, P. J. Phys. Chem. Lett. 2012, 3, 2395−2400. (7) Van Dijken, J. G.; Fleischauer, M. D.; Brett, M. J. J. Mater. Chem. 2011, 21, 1013−1019. (8) Heremans, P.; Cheyns, D.; Rand, B. P. Acc. Chem. Res. 2009, 42, 1740−1747. (9) Farinhas, J.; Ferreira, Q.; Di Paolo, R. E.; Alcacer, L.; Morgado, J.; Charas, A. J. Mater. Chem. 2011, 21, 12511−12519.

G

dx.doi.org/10.1021/nl501933q | Nano Lett. XXXX, XXX, XXX−XXX

Nano Letters

Letter

(10) You, J.; Dou, L.; Yoshimura, K.; Kato, T.; Ohya, K.; Moriarty, T.; Emery, K.; Chen, C. C.; Gao, J.; Li, G.; Yang, Y. Nat. Commun. 2013, 4, 1446−1455. (11) Matsuo, Y.; Sato, Y.; Niinomi, T.; Soga, I.; Tanaka, H.; Nakamura, E. J. Am. Chem. Soc. 2009, 131, 16048−16050. (12) Hirade, M.; Nakanotani, H.; Yahiro, M.; Adachi, C. ACS Appl. Mater. Interfaces 2010, 3, 80−83. (13) Chang, C. Y.; Wu, C. E.; Chen, S. Y.; Cui, C.; Cheng, Y. J.; Hsu, C. S.; Wang, Y. L.; Li, Y. Angew. Chem., Int. Ed. Engl. 2011, 50, 9386− 90. (14) Hochbaum, A. I.; Yang, P. D. Chem. Rev. 2010, 110, 527−546. (15) Garnett, E. C.; Brongersma, M. L.; Cui, Y.; McGehee, M. D. Nanowire Solar Cells. In Annual Review of Materials Research, Vol. 41; Clarke, D. R., Fratzl, P., Eds.; Annual Reviews: Palo Alto, 2011; Vol. 41, pp 269−295. (16) Yu, M.; Long, Y. Z.; Sun, B.; Fan, Z. Y. Nanoscale 2012, 4, 2783−2796. (17) Wei, W.; Bao, X. Y.; Soci, C.; Ding, Y.; Wang, Z. L.; Wang, D. Nano Lett. 2009, 9, 2926−2934. (18) Yu, K. H.; Chen, J. H. Nanoscale Res. Lett. 2009, 4, 1−10. (19) Gonzalez-Valls, I.; Lira-Cantu, M. Energy Environ. Sci. 2009, 2, 19−34. (20) Tsakalakos, L.; Balch, J.; Fronheiser, J.; Korevaar, B. A.; Sulima, O.; Rand, J. Appl. Phys. Lett. 2007, 91, 233117. (21) Czaban, J. A.; Thompson, D. A.; LaPierre, R. R. Nano Lett. 2008, 9, 148−154. (22) Mariani, G.; Zhou, Z.; Scofield, A.; Huffaker, D. L. Nano Lett. 2013, 13, 1632−1637. (23) Garcia-Frutos, E. M. J. Mater. Chem. C 2013, 1, 3633−3645. (24) Zhao, Y. S.; Zhan, P.; Kim, J.; Sun, C.; Huang, J. ACS Nano 2010, 4, 1630−1636. (25) Yoon, S. M.; Lou, S. J.; Loser, S.; Smith, J.; Chen, L. X.; Facchetti, A.; Marks, T. Nano Lett. 2012, 12, 6315−6321. (26) Haberkorn, N.; Gutmann, J. S.; Theato, P. ACS Nano 2009, 3, 1415−1422. (27) Liu, Z.; Liu, Q.; Huang, Y.; Ma, Y.; Yin, S.; Zhang, X.; Sun, W.; Chen, Y. Adv. Mater. 2008, 20, 3924−3930. (28) Wang, Q. H.; Hersam, M. C. Nat. Chem. 2009, 1, 206−11. (29) Xiao, K.; Yoon, M.; Rondinone, A. J.; Payzant, E. A.; Geohegan, D. B. J. Am. Chem. Soc. 2012, 134, 14353−14361. (30) Zhong, S.; Zhong, J. Q.; Mao, H. Y.; Wang, R.; Wang, Y.; Qi, D. C.; Loh, K. P.; Wee, A. T.; Chen, Z. K.; Chen, W. ACS Appl. Mater. Interfaces 2012, 4, 3134−3140. (31) Kim, J. S.; Park, Y.; Lee, D. Y.; Lee, J. H.; Park, J. H.; Kim, J. K.; Cho, K. Adv. Funct. Mater. 2010, 20, 540−545. (32) Salzmann, I.; Moser, A.; Oehzelt, M.; Breuer, T.; Feng, X.; Juang, Z.-Y.; Nabok, D.; Della Valle, R. G.; Duhm, S.; Heimel, G.; Brillante, A.; Venuti, E.; Bilotti, I.; Christodoulou, C.; Frisch, J.; Puschnig, P.; Draxl, C.; Witte, G.; Müllen, K.; Koch, N. ACS Nano 2012, 6, 10874−10883. (33) Xiao, K.; Deng, W.; Keum, J. K.; Yoon, M.; Vlassiouk, I. V.; Clark, K. W.; Li, A. P.; Kravchenko, I. I.; Gu, G.; Payzant, E. A.; Sumpter, B. G.; Smith, S. C.; Browning, J. F.; Geohegan, D. B. J. Am. Chem. Soc. 2013, 135, 3680−7. (34) Mativetsky, J. M.; Wang, H.; Lee, S. S.; Whittaker-Brooks, L.; Loo, Y.-L. Chem. Commun. 2014, 50, 5319−5321. (35) Huang, H.; Chen, S.; Gao, X.; Chen, W.; Wee, A. T. S. ACS Nano 2009, 3, 3431−3436. (36) Fan, J.; Zhang, L.; Briseno, A. L.; Wudl, F. Org. Lett. 2012, 14, 1024−1026. (37) Wise, A. J.; Zhang, Y.; Fan, J.; Wudl, F.; Briseno, A. L.; Barnes, M. D. Phys. Chem. Chem. Phys. 2014, 16, 15825−15830. (38) Lu, Z.; Zhan, C.; Yu, X.; He, W.; Jia, H.; Chen, L.; Tang, A.; Huang, J.; Yao, J. J. Mater. Chem. 2012, 22, 23492−23496. (39) Yang, F.; Shtein, M.; Forrest, S. R. Nat. Mater. 2005, 4, 37−41. (40) Ling, M. M.; Bao, Z. Chem. Mater. 2004, 16, 4824−4840. (41) Zhang, Y.; Dong, H.; Tang, Q.; Ferdous, S.; Liu, F.; Mannsfeld, S. C. B.; Hu, W.; Briseno, A. L. J. Am. Chem. Soc. 2010, 132, 11580− 11584.

(42) Barrena, E.; de Oteyza, D.; Sellner, S.; Dosch, H.; Ossó, J.; Struth, B. Phys. Rev. Lett. 2006, 97, 076102. (43) Fenter, P.; Schreiber, F.; Zhou, L.; Eisenberger, P.; Forrest, S. R. Phys. Rev. B 1997, 56, 3046−3053. (44) Forrest, S. R. Chem. Rev. 1997, 97, 1793−1896. (45) Hooks, D. E.; Fritz, T.; Ward, M. D. Adv. Mater. 2001, 13, 227− 241. (46) Verlaak, S.; Steudel, S.; Heremans, P.; Janssen, D.; Deleuze, M. Phys. Rev. B 2003, 68, 195409. (47) Yim, S.; Heutz, S.; Jones, T. Phys. Rev. B 2003, 67, 165308. (48) Zacharia, R.; Ulbricht, H.; Hertel, T. Phys. Rev. B 2004, 69, 155406. (49) Günther, C.; Karl, N.; Pflaum, J.; Strohmaier, R.; Gompf, B.; Eisenmenger, W.; Müller, M.; Müllen, K. Langmuir 2004, 21, 656− 665. (50) Shimada, T.; Ohuchi, F. S.; Koma, A. Jpn. J. Appl. Phys., Part 1 1993, 32, 1182−1185. (51) Carter, P. W.; Hillier, A. C.; Ward, M. D. J. Am. Chem. Soc. 1994, 116, 944−953. (52) Yang, F.; Shtein, M.; Forrest, S. R. Nat. Mater. 2004, 4, 37−41. (53) Shtein, M.; Peumans, P.; Benziger, J. B.; Forrest, S. R. J. Appl. Phys. 2003, 93, 4005−4016. (54) Singha Roy, S.; Bindl, D. J.; Arnold, M. S. J. Phys. Chem. Lett. 2012, 3, 873−878.

H

dx.doi.org/10.1021/nl501933q | Nano Lett. XXXX, XXX, XXX−XXX