barrier-guided growth of micro- and nano-structured graphene
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Nathaniel S. Safron, Myungwoong Kim, Padma Gopalan, and Michael S. Arnold*
Barrier-Guided Growth of Micro- and Nano-Structured Graphene
ION
The prospects of exploiting graphene’s exceptional electronic,[1–3] mechanical,[4,5] and thermal properties[6] in widespread applications has recently been advanced by rapid progress in the chemical vapor deposition of continuous membranes of the material on metal surfaces.[5,7,8] For many applications, however, it is not continuous graphene that is desired but rather graphene that is patterned on the micron and nanometer scales. Here, we report on a new strategy for the rational synthesis of lowdefect density, patterned graphene from the bottomup, called barrierguided chemical vapor deposition (BGCVD). In BGCVD, graphene growth is laterally restricted on planar metal surfaces by selectively passivating the catalytic activity of the metal with patterned barrier templates designed to (i) locally limit the generation of atomic C species, and (ii) confine their migration. We have successfully implemented BGCVD using aluminum oxide barriers on Cu substrates to fabricate highly crystalline, singlelayered structures including channels, nanoribbon arrays, and nanoperforated membranes, over largeareas. We show here that the barriers can restrict the nucleation of graphene to the exposed Cu and then guide its growth, remarkably, with 1 nm lateral precision. We also show that by avoiding damaging topdown etching, the direct synthesis of micro and nanostructured graphene by BGCVD produces superior edges with less disorder.
The patterning of graphene is a powerful approach for tuning its physical and electronic structure and for deviceintegration. Graphene patterned on the micronscale has been employed to create ultrahigh frequency analog amplifiers,[1] electrical interconnects,[9] conduits for heat dissipation,[6] and mechanical resonator membranes.[4] At the nanometerscale, the patterning of graphene opens up a bandgap, making it intriguing for semiconductor electronics and sensing.[10–18] Nanostructured graphene materials are furthermore attractive for energy storage because of their ultrahigh surface area; and, nanoperforated graphene membranes with high pore density have been proposed as ultrafiltration membranes.[19,20]
A critical advantage of graphene over other highperformance carbon materials such as nanotubes, is that its twodimensional form factor lends itself to patterning via scalable and standardized planar processing tools. Patterned graphene has typically been achieved by first exfoliating or growing graphene as a continuous membrane and then patterning it via top-down
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N. S. Safron, M. Kim, Prof. P. Gopalan, Prof. M. S. ArnoldDepartment of Materials Science and Engineering University of Wisconsin-Madison, USA E-mail: [email protected]
DOI: 10.1002/adma.201104195
Adv. Mater. 2012, 24, 1041–1045
subtractive etching. Topdown processing, however, is limited in fidelity by the etching tools that are available, resulting in structural and chemical disorder, at edges and extending several nanometers away from edges, that degrade graphene’s exceptional properties.[10–13,17]
The challenges with topdown processing have motivated the exploration of superior bottomup synthetic methods to avoid etching and to achieve more abrupt edges with fewer defects. Bottomup approaches have been investigated but thus far limited by difficulties in controlling growth orientation or assembly.[16,21] BGCVD overcomes these challenges, relying on selfterminating growth processes rather than harsh chemical etchants to abruptly define edges, while preserving the capability for creating rationallydesigned patterns and compatibility with planar processing. By patterning the catalytic activity of the metal using barriers (rather than the metal itself), BGCVD maintains the planarity of the metal growth substrate and ultimately achieves superior fidelity, especially on the nanoscale, by averting growth on the sidewalls of patterned metal features and avoiding the deterioration and restructuring of patterned metal templates at hightemperature.[22]
The BGCVD process is schematically depicted in Figure 1 for atmospheric pressure methane CVD on the aluminum oxide barrier/Cu system. Prior to growth, 10 nm of aluminum oxide is deposited on the Cu and patterned. We have chosen aluminum oxide due to its high temperature stability, low C solubility, and relative chemical inertness with respect to Cu and C.[23] During growth, methane decomposes selectively on the exposed Cu to produce C (Figure 1(i)), which does not appreciably dissolve into the bulk Cu due to low solubility.[24] The C instead laterally diffuses and accumulates on the exposed Cu until a supersaturation is reached and graphene nucleates (Figure 1(ii)). The preferential deposition and accumulation of C on Cu is driven by the relative inertness of the oxide and the strong CuC bonding energy ∼5 eV.[25] Following nucleation, the graphene crystallites grow up to the Cu/barrier edgeboundary (Figure 1(iii)). At this stage, the supply of C locally depletes, and after the entire Cu surface is passivated by either graphene or the barrier, the catalytic decomposition of methane ceases, terminating growth (Figure 1(iv)). Following BGCVD, the patterned graphene is transferred to arbitrary substrates via a temporary polymer support and etching of the Cu and oxide.
We have successfully implemented BGCVD to create micronscale boxes (Figure 2a–c), arbitrary patterns (Figure 2e–g), nanoribbon arrays (Figure 2h), and nanoperforated graphene (Figure 3b), showing the versatility and scalability of the method. We have employed the (60 μm)2 boxes in particular to better understand the function of the barriers and the crystallinity of the resulting patterned graphene. Figure 2a shows an optical micrograph of a (60 μm)2 region of oxide barrier on a Cu
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Figure 1. Schematic of BG-CVD: (i) methane decomposes into C, which (ii) diffuses and nucle-ates graphene, (iii) growing until (iv) the entire Cu surface is covered.
foil immediately following CVD at 1000 °C. Both the resulting graphene on the exposed Cu and the oxide are nearly transparent. However, the exposed regions of the Cu are smoother after growth indicating a restructuring of the foil. In contrast,
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Figure 2. Demonstration of BG-CVD: a-c) optical micrograph (top) and schematic (bottom) of Bon Cu a) immediately following BG-CVD, b) after Cu etch and transfer to Si/SiO2, and c) after Cspectroscopy of (i) graphene, (ii) the barrier region after growth, and (iii) the barrier region afteC. Escher pattern, f) Maze, g) Bucky the Badger, and h) Nanoribbon array, scale bars = 20, 20, 10corresponds to mono- and bi-layered graphene, respectively, while light grey regions are where nothere, which were subsequently etched for imaging. In (h) white lines are wrinkles.
the Cu underneath the barrier retains its initial coarse morphology, suggesting that the oxide is wellbonded to the Cu, frustrating restructuring. After transferring the graphene and oxide to a Si/SiO2(89 nm) substrate for enhanced optical contrast (Figure 2b), single and fewlayered graphene are observed on the previously unmasked regions contiguous to the barrier. After chemical etching of the barrier, a sharp boundary is observed between the graphene and the previously masked Cu, observed as bare Si/SiO2 in Figure 2c.
The graphene on the unmasked Cu is highly crystalline with low defectdensity, evidenced by a Raman 2Dband approximately twice the intensity of the Gband and the absence of a measurable Dband (Figure 2d). In contrast, nongraphitic carbon is detected in the barrier region following growth. The frequency and width of the Dband and absence of 2Dband in this region is consistent with that of small clusters of C or poly
aromatic hydrocarbons[26] rather than graphene. These small clusters may arise from C species that are unable to appreciably crystallize, kinetically trapped on the oxide upon cooling, or form in pinholes of the oxide. Regardless, this C is removed
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G-CVD with a 60 μm square aluminum oxide barrier u and barrier etch and transfer to Si/SiO2. d) Raman r growth and oxide etching. SEM of BG-CVD: e) M. , 0.1 μm, respectively. In (e-g), medium to dark grey graphene growth occurred due to barriers deposited
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Figure 3. BG-CVD of nanoperforated graphene. a) SEM of aluminum oxide nanodot array on Cu and b) nanoperforated graphene grown from (a), after transfer to SiO2/Si (scale bars = 100 nm). c) Histogram of oxide dot (blue, forward-slash) and perforation (red, back-slash) diameters. Schematics (d, f) and optical micrographs (e, g) of graphene crystallites on pristine Cu (d, e) and Cu with oxide nanodots (f, g), scale bars = 2 μm. h) TEM of BG-CVD nanoperforated graphene with 800 nm selective-area aperture, i) Corresponding electron-diffraction.
concurrently with the etching of the oxide, as verified by Raman (Figure 2c(iii)). Importantly, the latter two observations demonstrate that BGCVD is fundamentally distinct from previously reported liftoff mechanisms.[27] Overall, Figure 2 shows the effectiveness of oxide barriers at terminating the growth of graphene crystallites, thereby preventing their continuation over barriers and at inhibiting the nucleation and growth of graphene under and on top of the barriers.
In addition to the (60 μm)2 box pattern, BGCVD can be implemented using barrier templates that incorporate many discrete and disconnected regions of exposed Cu, in which case the nucleation of graphene in each region will be stochastic (e.g. some regions will nucleate and grow before others). However, one advantage of implementing BGCVD on Cu, in particular, is that it is selfpassivating. Once graphene nucleates and grows to completion in one discrete region, the growth will selfterminate in that region. Nucleation and growth will then proceed in other regions, until all have finished. Figure 2e, for example, demonstrates the effectiveness of BGCVD at reproducing a template with hundreds of discrete regions.
To explore the resolution limits of the BGCVD method, we have fabricated graphene nanoribbons (Figure 2h) and nanoperforated graphene (also semiconducting[11,12]) using electronbeam and block copolymer (BCP) lithography, respectively, to create nanopatterned aluminum oxide barrier templates on Cu (Figure 3b). For the latter, we adopted BCP lithography to create a hexagonal array of 16 nm aluminum oxide dots with a periodicity of 41 nm (Figure 3a) on Cu. Remarkably, these dots guide the growth of graphene through the network of ∼25 nm channels of exposed Cu (Figure 3b), resulting in largearea,
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bottomup graphene perforated by arrays of holes (see also Supporting Figures S1S3 for additional characterization via SEM, optical microscopy and AFM topography). The dotsize and holesize distributions differ only by 1 nm (Figure 3c), revealing that the oxide barriers are stable and terminate the growth of graphene with nearly atomic exactness. These data suggest that the BGCVD strategy could be used to template nanostructured graphene materials with sub5 nm features and <1 nm resolution.
Of key interest is not only resolution but the effect of the barriers on crystallinity and edgestructure. To understand growth mechanisms and crystallinity, we have prematurely terminated growth at ∼30% coverage with and without nanoscale oxide dot arrays (Figure 3d–g). On the bare Cu substrates, hexagonal graphene crystallites are observed with an average nucleation density of 0.02 μm−2 (counted via optical microscopy), characteristic of the employed atmospheric CVD conditions in which zigzag edges are favored[8] (Figure 3e). On the Cu substrates with the oxide dot array, similar hexagonal crystallites are observed, with a smaller nucleation density of 0.005 μm−2 (Figure 3g). This finding demonstrates that, as graphene grows through the network of 25 nm constrictions, the preference for the zigzag growth front is remarkably preserved, resulting in hexagonal edgefaceting. Additionally, it indicates that nucleation is largely unaffected by the oxide barriers and that if the growth proceeds to completion, the different domains of graphene zip together to form a continuous nanoperforated film with a large crystallite size.
To further confirm this finding, we have performed selectedarea electron diffraction on a continuous BGCVD
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Figure 4. Characterization of BG-CVD nanostructures: a) Raman spectra of nanoperforated graphene grown via BG-CVD (top, blue) and top-down etching (bottom, red), b) D-band defect ratio (left-axis, nanoribbons = squares and nanoperforated sheets = diamonds) and D′-band defect ratio (right-axis, nanoribbons = circles and nanoperforated sheets = hexagons) versus edge density for BG-CVD (blue, filled) and top-down etched nanostructures (open, red). Dotted lines show linear fit. c) Square sheet conductance for (top to bottom): w = 30 nm nanoribbon array (black), nanoperforated graphene at T = 295K (blue), 150K (green), 58 K (red). Dotted lines show fits for mobility. Insert: SEM of nanoribbon device (scale bar = 200 nm).
nanoperforated graphene sheet, using 800 nm (Figure 3h) and 4 μm (Supporting Information, Figure S4) selectedarea apertures. The hexagonal diffraction patterns from the 2.13 and 1.23 Å lattice periodicities of graphene (Figures 3i and S5) correspond to one domain of highly crystalline and orientated graphene inside the aperture area. These data indicate a nominal grain size >4 μm, in quantitative agreement with optical microscopy, confirming that nucleation is largely unaffected by the barriers and that the graphene maintains its crystallographic orientation as it grows around the barriers.
We have characterized the edge and defectrelated D and D′ Raman bands[8,26] of the nanostructured graphene to compare the quality of edges produced by topdown etching (via oxygen reactive ion etching) and by BGCVD (Figure 4a). We have specifically quantified the integrated intensity of the D and D′ bands, normalized to the integrated intensity of the Gband, referred as “Raman defect ratio”, for nanoribbon arrays and nanoperforated graphene structures, with various edge densities (Figure 4b). Both the D and D′ Raman defect ratios increase linearly with edge density indicating that the ratios are related to disorder at edges. On average, the topdown etched samples have 1.5x and 2x larger D and D′ Raman defect ratios, respectively, than the more pristine BGCVD samples. Furthermore, the BGCVD materials are also substantially better than topdown samples from literature as well,[12,28,29] which have 510x higher Dband Raman defect ratios (see Supporting Information, Figure S7). We hypothesize that the more substantial defect bands in the topdown samples arise from increased edge roughness (e.g. more edge atoms) and increased structural and chemical disorder and defects near edges that result from limitations inherent to topdown etching. For example, oxygen reactive ion etching through a polymer mask is not perfectly abrupt and therefore is expected to result in a region of disorder and defects in the interior of etched structures, adjacent to the edges.[29] In contrast, the superiority of the edges created by BGCVD results from its selflimited nature in which growth abruptly terminates at the barrier boundary, circumventing etchinduced disorder near edges.
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We have quantified charge transport through BGCVD graphene materials in order to further characterize their structure and continuity (Figure 4c). The materials are highly conductive, for example with a square sheet conductance of 0.4 and 0.1 mS at a charge density of ∼1.5 × 1013 cm−2, for the nanoribbons and nanoperforated structures, respectively–again showing that the BGCVD structures are highly continuous. The roomtemperature ON/OFF conductance modulation of the nanoperforated graphene is 3.9 ± 0.6 (14 devices), improving marginally to 10 at 58 K, as expected for the w = 25 nm dimensions and small band gap ∼ 15 meV produced in this study (see Supporting Figures S9S12). The graphene nanoribbon array (nanoribbon width w = 30 nm) shows a similar room temperature ON/OFF conductance modulation of 6.4. Because of the BGCVD method’s sub1 nm pattern reproducibility, in the future, it should be possible to realize largebandgap semiconducting graphene materials by combining BGCVD with state of the art lithography at w < 10 nm.
The transconductance at negative gate bias indicates hole transport mobility of 40 ± 7 cm2/V·s and 215 cm2/V·s in the nanoperforated and nanoribbon structures (Figure 4c), respectively, comparable to the mobility typically reported for similarly sized, topdown etched nanoribbons, which ranges from 30–300 cm2/V·s (see Supporting Information). In addition to edgedefects and disorder, several other factors are expected to affect mobility including the smoothness of the edges, the transferdependent conformation of the graphene to the underlying gate oxide, charged surface impurities, and the chemical termination of the edges. The particular scattering mechanisms limiting the mobility of our w = 25–30 nm nanostructures, in particular, are still unclear. Nonetheless, edgedefects and roughnessinduced scattering are expected to become the limiting factors as w is further decreased. Thus, because BGCVD reduces edgedefects compared to topdown etching (as shown by Raman scattering), this method is expected to be a promising strategy for realizing highmobility, largeband gap graphene materials in the future.
Overall, the Raman spectroscopy, hexagonal faceting, selected area diffraction, and high electronic conductivity and mobility,
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taken together indicate that BGCVD using the Cu/aluminum oxide system is a robust approach for the bottomup and laterally guided synthesis of graphene. The data show that BGCVD can produce highly crystalline, rationally patterned micro and nanostructured graphene materials, with less edgedisorder than topdown etching and with 1 nm patternreproduction, in a scalable fashion that is compatible with largearea, planar processing tools. Ultimately, as BGCVD avoids chemical etchants and instead relies on a selflimiting growth processes to abruptly define edges, we expect that this method will lead to graphene materials with superior performance.
Experimental SectionBGCVD process: Electron-beam lithography (PMMA, 950k) and block
copolymer lithography (P(S(46k)-b-MMA(21k)) were used to define aluminum oxide barriers directly on Cu foil. Electron-beam evaporation was used to deposit 10 nm of aluminum oxide through the respective polymer mask followed by liftoff (see Supporting Information for more details). The aluminum oxide/copper foil samples were heated to 1000 °C under a 340 sccm flow of forming gas (95% Ar, 5% H2). After annealing for 20 minutes, 66 ppm of CH4 was introduced into the flow and graphene was allowed to grow for 4 hrs. The BG-CVD graphene was transferred from the Cu foil by spin-coating PMMA as structural support for the graphene and backside-etching the Cu by floating in a 0.2 M FeCl3, 0.2 M HCl solution. The PMMA/graphene was collected on a Si/SiO2(89 nm) wafer, followed by dissolution of the PMMA.
Characterization: Raman spectra taken via confocal Raman microscopy (λ = 633 nm, ∼1 micron spot-size); TEM performed at 200 kV; electrical measurements performed on an Si(p++)/SiO2(89 nm) substrate, using the doped silicon as a backgate.
Supporting InformationSupporting Information is available from the Wiley Online Library or from the author.
AcknowledgementsResearch and analysis of BG-CVD synthesis, growth, and mechanisms and of the resulting patterned graphene’s structure and properties supported by the DOE Office of Science Early Career Research Program (Grant number DE- SC0006414) through the Office of Basic Energy Sciences (MSA and NSS). Research and analysis of aluminum oxide nanodot patterning, BCP preassembly and floating transfer, and nanodot adhesion supported by the University of Wisconsin-Madison, Robert Draper Technology Innovation Fund (MK and PG), and the National Science Foundation (Grant number CMMI - 1129802, MK and PG). Prof. T. Kuech is acknowledged for assistance in low-temperature measurement. J. Wood is acknowledged for useful discussions regarding flow conditions for ambient CVD of unpatterned graphene.
Received: November 1, 2011 Revised: December 12, 2011
Published online: January 26, 2012
© 2012 WILEY-VCH Verlag GmbAdv. Mater. 2012, 24, 1041–1045
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