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Macroscopic 3D Nanographene with Dynamically Tunable Bulk Properties Juergen Biener,* Subho Dasgupta, Lihua Shao, Di Wang, Marcus A. Worsley, Arne Wittstock, Jonathan R. I. Lee, Monika M. Biener, Christine A. Orme, Sergei O. Kucheyev, Brandon C. Wood, Trevor M. Willey, Alex V. Hamza, Jörg Weissmüller,Horst Hahn, and Theodore F. Baumann Surface-dominated bulk materials provide the unique opportunity to dynamically control their physical bulk properties by modification of their surfaces through interfacial phenomena.[1,2] Generating a sizable effect, however, requires that surface atoms constitute a large fraction of the total number of atoms. This requirement limits the technical potential of surface-dominated bulk materials, as structures with more than 10% surface atoms are typically not very stable and tend to reduce their surface energy by coarsening.[3] Here, graphene is an exception — it combines a very high surface area of up to 2630 m2 g−1 with the chemical and thermal stability intrinsic to the two-dimensional (2D) structure of sp2-bonded carbon. This makes graphene an interesting building block for realization of stable ultrahigh surface area bulk materials. Graphene also possesses many other remarkable properties,[4,5] including extremely high electrical and thermal conductivities, and exceptional mechanical strength and elasticity, which, in principle, should all be amenable to dynamic control via interfacial phenomena. As such, graphene-based bulk materials hold great technological potential beyond their obvious applications in the fields of energy storage and sensing[6] that so far have driven the development Dr. J. Biener, Dr. M. A. Worsley, Dr. A. Wittstock, Dr. J. R. I. Lee, Dr. M. M. Biener, Dr. C. A. Orme, Dr. S. O. Kucheyev, Dr. B. C. Wood, Dr. T. M. Willey, Dr. A. V. Hamza, Dr. T. F. Baumann Nanoscale Synthesis and Characterization Laboratory Lawrence Livermore National Laboratory 7000 East Avenue, Livermore, CA 94550, USA E-mail: [email protected] Dr. S. Dasgupta, Dr. D. Wang, Prof. H. Hahn Institute for Nanotechnology Karlsruhe Institute of Technology 76021 Karlsruhe, Germany Prof. H. Hahn Joint Research Laboratory Nanomaterials Technische Universität Darmstadt 64287 Darmstadt, Germany Dr. L.-H. Shao, Prof. J. Weissmüller Institut für Werkstoffphysik und -Technologie Technische Universität Hamburg-Harburg Eissendorfer Stasse 42, 21073 Hamburg, Germany Prof. J. Weissmüller Institut für Werkstoffforschung Werkstoffmechanik, Helmholtz-Zentrum Geesthacht Max-Planck-Strasse 1, 21502 Geesthacht, Germany

DOI: 10.1002/adma.201202289

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of various “bottom up” assembly approaches using graphene and graphene oxide towards 3D architectures.[7–10] However, despite recent progress in large-scale production of graphene and graphene oxide (GO) sheets,[11,12] graphene is still prohibitively expensive for large-scale manufacturing of graphenebased bulk materials. In addition, loss of surface area through aggregation remains a challenge.[10] In this Communication, we describe a “top down” strategy to fabricate mass-producible graphene-based bulk materials from low-cost polymer-derived carbon foams through the controlled removal of carbon atoms from a network composed of both amorphous carbon and graphite nanoplatelets (Figure 1a). This approach is inherently inexpensive (a few dollars per kilogram of the material), scalable, and yields mechanically robust, centimeter-sized monolithic samples (Figure 1b) that are composed almost entirely of interconnected networks of singlelayer graphene nanoplatelets. The specific surface area (up to 3000 m2 g−1, see Figure S1 in the Supporting Information) of this 3D nanographene (3D-NG) bulk material is comparable to that of a freestanding graphene layer, yet it has an open macroporosity (Figure 1c) that facilitates rapid mass transport throughout the bulk. Despite its high surface area, 3D-NG has a relatively high density (∼200 kg m−3), which makes the material surprisingly robust. For example, nanomechanical tests revealed a modulus E of 300–1000 MPa and a Meyer hardness of 20–100 MPa (Figure S3 in the Supporting Information). The latter suggests a weight-bearing capacity of 200 kg cm−2, which is surprisingly high for a material that consists of 100% surface atoms. The starting point for fabrication of 3D-NG is a macroporous carbon network prepared by organic sol-gel chemistry. The solgel process involves the catalyzed polymerization of organic precursors to yield a highly cross-linked organic gel that is then dried under ambient conditions and subsequently converted to carbon through pyrolysis in an inert atmosphere.[13] The important advantages of this approach are that the gel can be cast into any desired size or shape, and that architectural features, such as the pore structure and the ligament size, can be controlled through the reaction conditions. In the work reported here, we developed reaction conditions that yield pore sizes that are several orders of magnitude larger than those of traditional organic and carbon aerogels. After the gel has been dried under ambient conditions, pyrolysis then transforms the organic polymer into a porous sp2-bonded carbon network comprising both amorphous regions and multilayer graphene nanoplatelets.

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Figure 1. Processing and architecture of 3D-NG. a) Schematic of the polymer-based top-down approach. b) Illustration of the mechanical robustness of a centimeter-sized 3D-NG sample. c) Open macroporous network architecture composed of micrometer-sized pores and ligaments. d) Internal structure of ligaments composed of curved and intertwined graphene sheets. Selected double-layer regions are marked by ovals, and a typical intensity profile across such a double-layer region is shown on the left. e) X-ray diffraction data confirming the transformation of the initial multilayer graphene component to one that is dominated by single-layer graphene. f) Comparison of the Raman spectra in the D/G band region from 3D-NG and multiplelayer graphene.

3D-NG is then obtained by etching away the more reactive amorphous carbon components and a partial etching of multilayer graphite components (Figure 1a). The challenge is to preserve the macroporous architecture, yielding the final product as large and uniform monoliths. The required homogeneous etching is achieved by controlled burn-off in an oxidizing atmosphere at elevated temperatures. For our purposes, carbon dioxide[14] using the Boudouard equilibrium (C + CO2 2CO) proved to be most effective, as the low reactivity of CO2 provides uniform preferential removal of more reactive components throughout monolithic samples. As would be expected, the bulk density of the material decreases with increasing thermal activation, reaching a typical value of ∼200 kg m−3 for 3D-NG (∼9% of the full density of massive graphite). The conversion of the carbon foam to 3D-NG can be followed by X-ray diffraction (XRD) (Figure 1e). The absence of the stacking-related (002) diffraction peak in the XRD pattern for 3D-NG indicates the transition from a structure containing graphite nanoplatelets to one consisting of single-layer graphene. The in-plane crystallite size (La) of the graphene sheets in 3D-NG is 2–5 nm, as obtained by a Debye– Scherrer analysis of the (100) diffraction peak. These results are consistent with the chemical composition (Figure S2 in the Supporting Information) and were further confirmed by high-resolution transmission electron microscopy (HRTEM) and Raman spectroscopy. Examination of the material by HRTEM (Figure 1d) revealed short, linear features, which,

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owing to their extremely small lateral extension and spacing, could be identified as individual graphene sheets viewed edgeon. Consequently, the skeletal network of 3D-NG consists almost entirely of curved and intertwined monolayer graphene sheets; only in a few regions are stacks of multiple (mostly just twin) graphene layers observed. The Raman features for 3D-NG are similar to those observed for commercial graphene samples grown by chemical vapor deposition (Figure 1f). Analysis of the D/G band ratio (∼3) suggests La values around ∼5 nm (based on the Tuinstra–Koenig relationship).[15] The broadness of the Raman D band indicates considerable bond disorder, such as the presence of five- and seven-membered rings within individual sheets.[15,16] These defects can explain the curved appearance of the graphene sheets in HRTEM images. The electronic structure of 3D-NG was studied to provide the basis to understand how interfacial phenomena affect the bulk properties of the material. Filled and empty electronic states were probed by soft-X-ray emission (SXE)[17] and X-ray absorption spectroscopy (XAS),[18] respectively. The SXE/XAS spectra for 3D-NG (Figure 2a) are very similar to those recorded from highly oriented pyrolytic graphite (HOPG). This confirms the notion that 3D-NG is composed primarily of sp2-bonded carbon, and that the states near the Fermi level have π/π∗ character.[19] Significantly, the spectra collected from HOPG and 3D-NG exhibit a comparable degree of overlap between filled (SXE) and empty (XAS) states at the Fermi level, which suggests that 3D-NG, as graphite, is a semi-metal, although one


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COMMUNICATION Figure 2. Electronic structure of 3D-NG and a schematic of electrochemical gating (EG). a) Valence and conductance band structure of 3D-NG (2800 m2 g−1) and a freshly cleaved highly oriented pyrolytic graphite standard probed by SXE (left) and XAS (right), respectively. The SXE excitation energy was set at 300.5 eV. b) Electrochemically driven ion adsorption causes very strong interfacial electric fields (∼109 V m−1) and charge accumulation/depletion at the electrode/electrolyte interface. c) Experimentally measured charging current I (red line, left ordinate) and area-specific surface charge Q of a 3D-NG electrode (1900 m2 g−1) versus potential E (0.1 M NaClO4 in deionized H2O electrolyte solution, 0.5 mV s−1). The box-shaped charging curve indicates an ideally polarizable electrode.

must exercise caution in inferring ground-state properties from the XAS and SXE spectra owing to core–hole effects. The fact that all atoms in 3D-NG monoliths are surface atoms provides a unique opportunity to dynamically tune the bulk properties of 3D-NG by electrochemical gating (EG). The term EG describes the effect that electrochemically driven ion adsorption causes very strong interfacial electric fields (∼109 V m−1) that can shift the Fermi level through charge carrier accumulation or depletion by a small fraction of an electronic charge per carbon atom (Figure 2b). The effect, analogous to the well-known field effect that forms the foundation of semiconductor device operation, has previously been observed in 2D graphene with EG-induced resistivity changes of several hundred percent.[20,21] The amount of EG-induced charge accumulation/depletion can be obtained from cyclic voltammetry (CV) experiments with 3D-NG as the working electrode. Analysis of CV data (shown in Figure 2c) reveals that the area-specific capacitance of the 3D-NG working electrode in dilute, 0.1 M NaClO4 electrolyte is ∼5.5 μF cm−2 (or 105 F g−1). This translates into accumulation (negative bias) or depletion (positive bias) of approximately 0.01 electrons per carbon surface atom into the π/π∗-bands for a bias of ±1 V, which causes Fermi level shifts of approximately ±0.5 eV (Figure S4 in the Supporting Information). The change in the carrier density that results from EG can thus be many times larger than the intrinsic carrier density and can

therefore change macroscopic physical and chemical properties of the material (for comparison, the intrinsic carrier density of graphite is ∼10−4 free carriers per carbon atom).[22] Indeed, EG induces macroscopic conductance changes in centimeter-sized 3D-NG monoliths that are on the order of several hundred percent and are fully reversible (Figure 3). These conductance changes were measured with a macroscopic four-point probe in a three-electrode electrochemical cell (Figure 3a). The conductance decreases with increasing negative surface charge, and vice versa (Figure 3c). A conductivity minimum that can be attributed to the Dirac point[23] is observed at a potential of –0.6 V (corresponding to accumulation of 0.006 electrons per carbon atom). This suggests that our material is heavily p-doped, presumably related to the ∼0.5 at% oxygen detected by Rutherford backscattering spectrometry (Figure S2 in the Supporting Information). Such changes in conductivity are consistent with those previously measured on graphene[21] and could be exploited to build a low voltage, high power, allcarbon tunable bulk resistor. An even more fascinating prospect would be the realization of transistor-like on/off behavior similar to that observed for bilayer graphene[24] and semiconducting single-wall carbon nanotubes.[25] A promising strategy to open up the required bandgap in 3D-NG is the incorporation of certain defects.[26] Another interfacial charge-related phenomenon is a macroscopic strain response that could be used for actuation.

Figure 3. Interfacial charging induced macroscopic resistance changes. a) Schematic presentation of in situ conductance measurements with a macroscopic four-point probe in an electrochemical cell. b) Applied potential versus time (top) and resulting normalized resistivity changes ρ/ρ0 (bottom) of a 3D-NG working electrode in 0.1 M NaClO4 aqueous electrolyte solution (potential range from −0.6 V to 0.3 V vs. Ag/AgCl). c) Normalized resistance change ρ/ρ0 vs. area-specific elementary charge per surface atom measured with the continuous potential ramping (CV) method.

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Figure 4. Interfacial charging induced macroscopic strain response. a) Schematic presentation of in situ strain measurements performed in a dilatometer equipped with an electrochemical cell. b) Strain Δl/l (red, left ordinate) and applied potential (black, right ordinate) versus time t in a 0.7 M NaF electrolyte solution and potential jumps between –1 V and +1 V. c) strain Δl/lvs. net charge per carbon atom measured during three cyclic potential sweeps at constant scan rate (–0.1 V to +0.3 V vs. Ag/AgCl).

Although this strain effect is known for porous carbon electrodes,[27,28] its utilization largely depends on the availability of a suitable monolithic bulk material that combines high surface area (large strain amplitudes) with high stiffness and strength (high mechanical work density). The length of a 3D-NG electrode changes periodically and reversibly with an amplitude as large as 2.2% (corresponding to a volume strain of 6.6%) as the potential is cycled between –1.0 V and +1.0 V (Figure 4b). The sample expands during negative charging, and vice versa. Both the sign and the maximum value of the strain amplitude are in agreement with graphite intercalation data, where it was found that the C–C in-plane bond length increases if electrons are added into the graphite π bands and vice versa.[29] The strain amplitude is considerably larger than those reported for carbon nanotube (CNT) arrays,[30] nanoporous metals,[31] and piezoelectric materials.[32] Furthermore, and in contrast to CNT arrays, 3D-NG is a bulk material that can be loaded in compression. Thus, given the low cost and the availability of large monolithic samples, 3D-NG is a promising material for effective conversion of electric energy into mechanical work. Details of the performance of carbon nanofoam actuators can be found in the literature.[33] Future applications of 3D-NG, including electrical energy storage based on the electric double-layer technology, will depend on further improvements of the material properties, most importantly on electronic structure engineering aimed at i) increasing the density of states around the Fermi level to increase the quantum capacitance for energy storage applications or ii) opening up a small bandgap for the tunable bulk transistor application. Future investigations will focus on doping and increasing the density of the material while maintaining the high surface area and at least some of the macroporous network structure to avoid any mass transport limitations.

orange polymeric gels were washed with acetone to remove the water from the structure and then dried under ambient conditions. The foams were subsequently carbonized at 1050 °C for 3 h under a N2 atmosphere, yielding porous carbon monoliths with densities of ∼550 kg m−3. The activation of the carbon foam part (2 cm × 3 cm × 4 mm, 1.2 g) to generate 3D nanographene (3D-NG) was carried out under flowing CO2 (10 sccm) at 950 °C for 5 h. The yield of 3D-NG monoliths from the activation process was ca. 25%–30%, based on the weight of the starting carbon foam. Graphene reference samples (Graphene Supermarket) were grown on Ni-coated silicon oxide/silicon substrates (CVDGraphen) and consist of polycrystalline few-layer (3–5) graphene (as confirmed by Raman spectroscopy[34] and atomic force microscopy). The films were transferred after liquid etching of the Ni layer onto silicon oxide/silicon substrates. Details of the characterization of the 3D-NG (surface area, X-ray diffraction, Raman and electron transmission spectroscopy, Rutherford backscattering spectrometry and elastic recoil detection analysis, mechanical properties, soft X-ray emission and soft X-ray absorption near-edge structure spectroscopy, and theoretical calculations) can be found in the Supporting Information.

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

Acknowledgements Work at LLNL was performed under the auspices of the US DOE by LLNL under Contract DE-AC52-07NA27344. Projects 10-LW-045 and 12-ERD035 were funded by the LDRD Program at LLNL. Research at KIT was partially supported by the Center for Functional Nanostructures (CFN). One of us (H.H.) is grateful for financial support by the State of Hesse in the form of a capital equipment grant at TUD. Received: June 7, 2012 Published online: July 31, 2012

Experimental Section The carbon foam material was synthesized through the acid-catalyzed sol-gel polymerization of resorcinol with formaldehyde. Specifically, resorcinol (24.6 g, 0.224 mol) and 37% formaldehyde solution (35.8 g, 0.448 mol) were dissolved in water (40 mL), followed by the addition of glacial acetic acid (0.88 g, 0.014 mol). The reaction mixture was then transferred to glass molds and cured at 70 °C for 72 h. The resultant

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