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Self-Folded Three-Dimensional Graphene with Tunable Shape and Conductivity Tetsuhiko Fujiwara Teshima, Calum Studart Henderson, Makoto Takamura, Yui Ogawa, Shengnan Wang, Yoshiaki Kashimura, Satoshi Sasaki, Toichiro Goto, Hiroshi Nakashima, and Yuko Ueno Nano Lett., Just Accepted Manuscript • Publication Date (Web): 07 Dec 2018 Downloaded from http://pubs.acs.org on December 7, 2018

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Self-Folded Three-Dimensional Graphene with Tunable Shape and Conductivity Tetsuhiko F. Teshima1,*, Calum S. Henderson1,2,†, Makoto Takamura1, Yui Ogawa1, Shengnan Wang1, Yoshiaki Kashimura1, Satoshi Sasaki1, Toichiro Goto1, Hiroshi Nakashima1, and Yuko Ueno1 1. NTT Basic Research Laboratories, NTT Corporation. 3-1 Morinosato Wakamiya, Atsugi, Kanagawa 243-0198, Japan *E-mail address: [email protected] 2. School of Chemistry, The University of Edinburgh. David Brewster Road, Edinburgh EH9 3FJ, Scotland, United Kingdom †

Present address: ISIS Neutron and Muon Source, Science and Technology Facilities Council. Rutherford Appleton Laboratory, Didcot, OX11 0QX, United Kingdom

KEYWORDS: self-folding, graphene, parylene, non-linear conductivity

ABSTRACT: Three-dimensional (3D) graphene architectures are of great interest as regards applications in flexible electronics and bio-interfaces. In this study, we demonstrate the facile

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formation of predetermined 3D polymeric microstructures simply by transferring monolayer graphene. The graphene adheres to the surface of polymeric films via noncovalent - stacking bonding and induces a sloped internal strain, leading to the self-rolling of 3D microscale architectures. Micro-patterns and varied thicknesses of the 2D films prior to the self-rolling allows for control over the resulting 3D geometries. The strain then present on the hexagonal unit cell of the graphene produces a non-linear electrical conductivity across the device. The driving force behind the self-folding process arises from the reconfiguration of the molecules within the crystalline materials. We believe that this effective and versatile way of realizing a 3D graphene structure is potentially applicable to alternative 2D layered materials as well as other flexible polymeric templates.

Introduction Graphene has emerged as the most promising two-dimensional (2D) materials for applications in flexible electronics and energy storage, given its electrical/thermal conductivity,1,2 mechanical strength,3 and chemical stability.4 Specifically, its exceptionally large open-surface with free edges and non-cytotoxicity allows the anchoring and desorption of ions, molecules, and cells, showing its relevance to supercapacitors,5 catalysis,6 chemical sensing,7 and bioelectrodes.8 Indeed, to extend these properties to practical applications, it is preferable that the inherently planar geometries are assembled on a macroscopic scale with three dimensionalities (3D).5,6 In particular, curved, folded, and wrinkled architectures are extensively featured as regards flexible deformability.9,10 Although these 3D geometries of graphene have hitherto been fabricated by applying external mechanical forces to flexible templates,11 this manual process


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causes technical difficulties when constructing well-controlled micro- or nano-scale 3D structures. Origami-inspired self-rolling polymer bilayers have been used to assemble curved graphene into microscale 3D architectures.12-14 The polymeric bilayer films with heterogeneous mechanical properties are spontaneously rolled into tubular and spherical forms from 2D micropatterns.15,16 Conventionally, graphene has been loaded on these films to induce its wrinkling or folding. Nevertheless, the self-assembly principle in which the graphene itself behaves as the driving forces behind the transformation into 3D polymeric shapes has yet to be demonstrated.17 This difficulty is attributable to the relatively poor adhesion and delamination of graphene with foldable substrates. Provided a graphene monolayer can trigger spontaneous self-folding, the simple transfer of graphene would be attractive for the simple fabrication of microscale 3D curved or folded sensors and actuators.17 This spontaneous transformation may also be tailored to retain or modulate the unique features of pristine graphene in the folded state. Herein, we demonstrate the rapid and easy formation of microscopic self-folded homogeneous polymeric thin films by employing transferred monolayer graphene, known as a micro-roll. We use poly(chloro-p-xylylene) (parylene-C) as a flexible template that adheres to graphene due to π-π stacking sp2 hybridization.18 This interaction confers tight adhesion between graphene and parylene; hence, this is utilized practically to develop implantable bioelectrodes.8 Consequently, the simply transferred graphene endows parylene sheets with interfacial mechanical heterogeneity and electrical conductivity. As a result, the differential strain gradients within the graphene/parylene bilayer initiate the autonomous formation of rolled-up shapes. To highlight the applicability as a curved electronic component, the ordered 3D shapes and their electrical properties are found to be readily controllable by controlling the initial 2D micro-


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patterns and the number of graphene layers. We anticipate that this effective and versatile way of realizing 3D carbon structure will be applicable to alternative 2D layered materials as well as other flexible polymeric templates.

Results and discussion Fabrication of 3D graphene-polymer architectures The process for fabricating the self-foldable multi-layered films consists of three simple steps (Figure 1a, S1). First, Ca-alginate was gelated on a SiO2 substrate in the first step, and then the substrates were wholly laminated with parylene-C.19,20 Monolayer graphene was synthesized on a Cu substrate via a chemical vapor deposition (CVD) process. Following the dissolution of the Cu layers, the graphene was transferred onto the parylene-C surface with a wet transfer technique. A photo-resist was then spin-coated and photo-lithographically micro-patterned to protect the graphene layers from the subsequent etching steps. Thereafter, the multi-layered films were micro-patterned by employing reactive ion etching with oxygen plasma through a photoresist mask (Figure S5). The films have a clearly identifiable 110 nm-thick parylene-C on the 80 nm-thick Ca-alginate layer with highly defined geometries. The dissolution of the sacrificial Ca-alginate layer with ethylene-diaminetetraacetic acid (EDTA) occurs from the edge of the graphene/parylene bilayers. Figure 1b shows the boundary between the pristine parylene and graphene-laden parylene 100×100 m2 wide and 71 nm thick before and after adding EDTA with a final concentration of 5 mM. While the pristine parylene remains flat, the exposure to EDTA drives a batch self-folding of the graphene-laden parylene


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films into cylindrical micro-rolls (Supplementary movie SM1). Time-lapse images show that the complete dissolution of Ca-alginate layer does not directly trigger simultaneous initiation of selffolding rectangular 200×400 m2 bilayers (Figure 1c). Then, the bilayers get wrinkled and immediately curve perpendicular to the longer sides. The micro-roll curvature (1/) reaches an equilibrium value in the typical time scales of less than 2 sec. This deformation is delayed approximately 1 min after the completion of the Ca-alginate dissolution (Figure S2). The time delay in self-folding stems from the time-consuming formation of the tight - adhesion between graphene and parylene. As the self-folding possibilities are illustrated in Figure 1d, the self-folding orientation is random and contributes to the formation of various 3D structures. The overall bending rigidity makes it more favorable for rectangular films to curve along the long or short axis. The other feasible configuration is a corkscrew-shaped micro-roll that diagonally curves toward the intermediate folding angles. Alternatively, both sides of the bilayers curved inward to the points at which they met, thus developing a double-roll configuration. The films instead assume dogears configurations whose corners curve inward. While predicted theoretically, these random bending directions have not been yet experimentally observed in self-foldable polymeric bilayers.21 The polycrystalline graphene employed in this study consists of a patchwork of relatively smaller single-crystalline grains separated by grain boundaries. Since one might expect a single domain of graphene to have the same elastic energy, which induces an isotropic tensile force for folding, the uncontrollable folding orientation originates from domains inside the CVDgrown graphene.


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Controllable 3D architectures of micro-rolls To achieve multi-angle observation, the transition from the 2D bilayers to a curved format micro-roll was viewed horizontally by inclining the microscope lenses. This inclined microscopy system revealed that the bilayer was bent in the direction of the substrates, which is called a “mountain-fold”, while maintaining the bottom parylene on the inward side (Figure S3, Supplementary movie SM2). The graphene on the micro-rolls appears to be exposed to tensile force thus causing it to cover the inner parylene from the outside. The tight - adhesion between graphene and parylene was revealed by a control-experiment in which the graphene was transferred onto silk fibroin films.22 The graphene-laden film remained flat after release in a similar manner to that without graphene (Figure S4). Concurrently, Raman spectroscopy confirmed that the graphene evanesced from silk fibroin. Silk fibroin derived from Bombyx mori silkworm consists of few aromatic rings; therefore, it exhibits fewer π-π stacking interactions with graphene. This result is consistent with the other polymeric films such as polymethyl methacrylate (PMMA), polydimethylsiloxane (PDMS), and negative photoresist (SU-8). These polymeric films with transferred monolayer graphene remained flat, partly owing to graphene’s slippage from the polymer surface with fewer aromatic rings than parylene.11,23 The polymer crystallinity of chemically vapor-deposited parylene would also facilitate well-aligned orientation of aromatic rings and overlap of the  system to the graphene.24-26 These results imply that a rich π-π stacking force is required for self-folding since it confers tight but noncovalent bonding between graphene and parylene. SEM images in Figure 2a show the films after the dissolution of the Ca-alginate and the final 3D structure of the self-folded micro-rolls. After self-folding, the surface of the micro-rolls


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contains few wrinkles or deflections, providing hollow internal spaces. The relatively uniform surface suggests that the strain in encompassing graphene is evenly induced in-plane. Although both ends seem to have collapsed onto the substrate as a result of capillary forces in the lyophilization process, the laminate layers are stuck tightly in position and structurally preserved with multiple rolling. Note that the multi-layers of graphene are not in contact with each other because of the insulating parylene layer. To investigate their durability, the micro-rolls were exposed to external environmental stimuli, namely changes in temperature and pH (Figure S6). When the micro-rolls were immersed in aqueous media on thermo-stages, a temperature range from room temperature to 95°C does not cause either the parylene or graphene to expand. When the pH of the surrounding media is arbitrarily conditioned from 2 to 14, the micro-rolls neither unfold nor become tighter. Hence, no change in strain is imposed on the micro-rolls by changes in temperature and pH. Consequently, the π-π adhesion between parylene and graphene is so physically and chemically inert that it confers sufficiently tight and irreversible mutual adhesion that is capable of enduring a harsh external environment. Subsequently, we investigated the final 3D architectures of the micro-rolls with transferred monolayer graphene. Figure 2b plots the experimentally obtained average curvature radii of the resulting tubular structures () as a function of parylene thickness (tp) ranging from 27 to 322 nm. Given the transferred monolayer graphene, the thinner tp leads to a smaller . When tp is constant,  is unaffected by 2D micropattern design in any of the configurations shown in Figure 1d. Reducing the thickness of the parylene layer decreases the bending rigidity, thereby improving machining controllability of a fine 3D structure. In addition to tp, the ratios of strain () and elastic modulus (E) of each layer have a noticeable effect on . The bending stiffness of the graphene is negligible due to its atomic thickness, and the linear gradient of the in-plane


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strain along the film thickness drives the self-folding. We applied quantitative predictions of  by modelling using bimorph-beam theory.27,28 This theory predicts  based on E, t, and  as follows:

𝜌=

𝐸p(𝑡p + 𝑡g)2 1 + 4𝛼𝛽 + 6𝛼𝛽2 + 4𝛼𝛽3 + 𝛼2𝛽4 6𝐸g𝜀p𝑡g

(1 + 𝛽)3

(1)

where  is the relative elastic modulus (Ep/Eg), and  is the ratio of the thickness (tp/tg). The subscripts “p” and “g” denote parylene and graphene, respectively. For simplicity, tg, Ep, and Eg are approximately 0.3 nm, 3.2 GPa,15 and 1 TPa to estimate  as a function of variable tp and g. We assume g to vary between 1.2% and 2.2% as a fitting parameter (Figure 2d). The experimental values of  largely follow the trend predicted by equation (1) when g lies within the 1.6–1.8% range. The theory assumes that the graphene on the parylene is under a constant uniaxial strain, irrespective of its curvature.

Characterization of self-folding monolayer graphene Using Raman spectroscopy, we characterized the molecular structure of monolayer graphene on a flat parylene surface. Figures 2c and S7a plot representative and entire Raman spectra of the graphene on different flat substrates: parylene and silk fibroin. Specifically, the Raman spectra contain two prominent peaks typifying monolayer graphene, namely G and 2D (G’) peaks. The G band frequency (G) at ~1,584 cm−1 originates from the doubly degenerate in-plane vibrational zone-center E2g phonons. The 2D band (2D) at ~2,675 cm−1 is assigned to overtone of the transverse optical phonon around K points in the Brillouin zone. While the spectra with


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independent peaks at 1,610 cm−1 correspond to the deformation of the C-C bonds in parylene-C, they are negligibly weaker when tp is less than 322 nm. The multipoint raster-scanned Raman spectra over an area of the flat 220×420 m2 bilayers exhibit wide spatial distributions: the intrinsic frequencies of essentially strain-free graphene (G0, 2D0) at (1,584 cm−1, 2,675 cm−1) and the pre-tensile mode of (1,592 cm−1, 2,690 cm−1). Across the entire area, the peak position varies on the length scale of several tens of microns (Figure 2e). This distribution in G0 and

2D0 is attributed to the possibility of pre-tension in the graphene transferred onto the parylene surface due to the lattice mismatch. The Raman spectra obtained from the micro-rolls are plotted as a function of . Irrespective of curvature radii, we measure symmetric Lorentzian lines with (G, 2D) at (1,592 cm−1, 2,690 cm−1) (Figure 2d). The blue-shifts in G ~8 cm−1 and in 2D ~16 cm−1 from the original modes (G0, 2D0) stem from the distorted graphene lattice and the altered interatomic distance. Besides, there is neither splitting in the G and 2D bands into two subbands11 nor a radial breathing mode near 224 cm−1 by analogy with carbon nanotubes (CNTs),29 owing to the three orders of magnitude differences in the curvature radius. It is interesting to note that D peaks ranging at ~1,340 cm–1, which reveal the existence of defect-induced breathing modes of sp2 rings, are attenuated after self-folding, which importantly demonstrates the elimination of defects in the micro-rolls (Figure S7b). The multipoint spectra per micro-roll exhibit the spatial homogeneity for G at 1,590 cm−1 and for 2D at 2,690 cm−1 with few spatial pixel-to-pixel variations (Figure 2e). Since all characteristic peaks are largely preserved after self-folding, there is no significant exfoliation or bond breakage of the folded graphene.


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The blue-shift denoted as ΔG or Δ2D determines compressive strain leading to phonon hardening compared with flat graphene. The data sets from micro-rolls with any curvature fall on a single line with a slope of approximately 2.0 (Δ2D/ΔG) relative to (G0, 2D0). This slope agrees well with the aforementioned value for graphene stretched on flexible polymers with a uniaxial strain. In contrast, this blue-shift contrasts with the tensile state of graphene stretched on PMMA30 and polyethylene terephthalate,11 but is similar to that of graphene stretched on PDMS.31 Previous studies have shown that a high tensile strain leads to the high mobility of parylene chains and uniformly compressed graphene unit hexagonal cells thanks to the high Poisson’s ratio.31 Notably, since the self-folding experienced by graphene/parylene bilayers lacks any external force, the graphene on the parylene will experience not only longitudinal elongation but also transverse contraction with uniaxial tension. The strain is estimated using the Grüneisen parameter,  with G or 2D32: ∆𝜔

1

(2)

𝜀 = 2(1 ― 𝜈)𝜔 ∙ 𝛾

where  is the Poisson’s ratio of parylene.33 The Grüneisen parameter, which is calculated to be 2.7 using first principles,11 provides a compressive strain of 0.184% for Δ2D of 16 cm-1. This estimated strain within the graphene is almost in agreement with the predicted value in the 0.160.18% range. When both elongated and compressed distortion is considered, the strain calibration of the hexagonal shape will be altered. Here, the Raman spectroscopic shift indicates that the self-folding-driven stress and strain on the graphene lattice is significant in terms of changing the electronic structure in contact with parylene. Apart from the Raman shifts, the strain-driven structural alteration of graphene tends to tune its electrical properties. In subsequent experiments, the change in graphene conductivity before


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and after self-folding was measured at room temperature under both dry and ambient conditions. To avoid the rupture of graphene that accompanies its shape change, the stacking order of parylene and graphene was reversed (Figure 3a). We newly micro-patterned hinged bilayers that were partially pinned down, as this prevents the folded structures from free-floating and enables self-folding in the opposite direction. Hence, the hinged graphene was engineered so that it adopted various curvatures, whereas the tapered hinges remained pinned and processed so that they stayed flat for testing (Figure S8). The Au pads were connected to two-point probes to measure drain current (Id) –voltage (Vd) characteristics for flat and folded 200×800 m2 graphene/parylene bilayers. The flat graphene exhibits linear behavior at 17.9 k sheet resistance: 617.1 /□) in dry and wet conditions with water (Figure 3b). This linear plot confirms that the - adhesion between graphene and parylene is noncovalent without compromising the graphene conductivity. In contrast, the selffolding process greatly alters the electrical property of loaded graphene and it becomes nonlinear. The resistance increases from 17.9 to 38.0 k, where it is measured at 0 mV. The Id/Vd curve remains a non-linear after replacing the EDTA with water. Figure 3c compares the Id–Vd characteristics as a function of micro-roll curvature. The resistances of graphene in the flat 200×800 and 200×400 m2 bilayers were measured and normalized to 17.9 and 21.6 k, respectively. While the non-linear electrical transport is common to any folding crease dimensions, the non-linearity increases with a smaller . Note that the parylene layer above graphene functions as an insulator that eliminates interlayer tunnelling. The self-folding process never leads to the exfoliation or rupture of graphene, and the sp2 hybridization and electrical properties of graphene are retained. We attribute the altered charge carrier densities even at room temperature to the uniformly compressed graphene unit hexagonal cell.


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By applying back-gate voltages (Vg), the Id corresponding was measured for various Vd values. Both positive and negative back-gate voltages modulate the non-linear resistance of micro-rolls (Figure 3d). When Vg decreases from +40 to -5 V, Id increases piecemeal and the curve reduces the linearity. In contrast, less than -6 V of Vg induces the saturation of Id. With Vg in the -40 to +40 V range, Id in the flat graphene decreases almost linearly with increasing Vg at Vd=400 mV (Figure 3e). After self-folding, the graphene displays a transfer curve with a significant increase in the corresponding Id, which is consistent with semiconducting graphene nanoribbons with substantial band gaps.34 With higher Vg, higher defect densities in the folded graphene suppress conductance. Although Id is caused by the surrounding molecules including EDTA and Na-alginate,35,36 the modulatable property makes it possible to realize switching transistors. These results also indicate that the graphene in the micro-rolls behaves as a p-type semiconductor in both the flat and folded states. It is worth noting that non-linear electrical transport has been observed particularly in folded nanochannels37 and nano-gaps.38 Previous studies suggest that the nanometer-scale bridges, ribbons, and gaps in graphene cause the junctions to undergo electrical breakdown such as the formation of barriers and increased defect density. However, there have been few reports stating that graphene with micrometer-scale curvature exhibits a defect-induced constriction.13 Since the graphene in micro-rolls with any  value is exposed to 0.18% strain, an increase in the proportion of strained graphene would increase the formation of nanoscale gaps, which in turns generates tunnelling barriers across this gap.39 Because the generation of such tunnelling barriers in graphene remains the subjects of intense debate, further investigation is required to determine the effect of the strain and resulting electronic structure in folded polycrystalline graphene.


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Tunable self-folding process for 3D geometries Self-folded 3D shapes are tunable by optimizing such aspects of graphene elements as 2D design, thickness, and crystalline direction. First, 3D spatial geometries are rationally characterized depending on the micro-patterned 2D design (Figure 4a). For instance, flowershaped and auxetic micropatterns are transformed into caged-shaped grippers that are useful for the encapsulation and manipulation of living cells. The meshed micro-patterns are also curved and shaped into see-through and hollow gauze-like structures, which enable carbon-based architectures in association with CNTs, nano-belts,40 and fullerene41. Unfortunately, the folded films are not joined end-to-end; thus, the characteristic breathing modes are not observed with Raman spectroscopy. In addition, parallelogram micro-patterns incline the angle of self-folding orientation, consequently transforming the films into coiled hollow cylinders 114.3 m in diameter. The optimization of the interior angles of the parallelograms and parylene thickness can tune the programmable number of turns and pitches of helices, which is potentially applicable to stretchable electronic components. Hence, by incorporating various 2D micropatterns, it is possible to produce any types of complex 3D spatial geometries. Another significant characteristic that stipulates self-folded final 3D shapes is the number of transferred graphene layers. Sequentially stacked multi-layered graphene create a steeper gradient of membrane stiffness, and provides significant control over the  value of 200×400 m2 micro-rolls. When tp is 85 nm,  gradually decreases from 32.6 to 9.8 m for monolayer and four-layer graphene (Figure 4b). In spite of the increase in graphene thickness, the micro-rolls still retain sufficiently high transparency to expose inner structures. No matter how thin the parylene layer is, stacking multi-layered graphene makes the micro-rolls much tighter (Figure


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4c). The minimum value of allowable curvature radii is around 5.0 m, regardless of graphene thickness and tp. Raman spectroscopy characterizes that 2D at 2,690 cm-1 exhibit neither a rednor a blue-shift, as the graphene layers are transferred one after another (Figure 4d). In contrast,

G at 1,592 cm-1 for the monolayer graphene experiences a red-shift to 1,584 cm-1 for the increasing numbers of graphene layers with any curvature radii. The red-shifts in G suggest that an increase of graphene thicknesses causes an alleviation of strains especially in outer graphene, partly because of the randomly oriented hexagonal lattices in stacked multilayer graphene.42 This result supports the evidence for the absence of the delamination of multilayer graphene and a tight π-π stacking interaction between adjacent graphene layers. Multi-layered graphene generally contributes to not only the increases in conductivity, but also determining the amount of doping, disorder, and edges, which is applicable for integration with modulatable resistor and transistor.43 A final significant characteristic is the folding orientation. Provided that the domain size is larger than one sheet of micro-patterned film, the crystalline orientation of graphene would certainly affect the orientation of self-folding (Figure 5a). To verify this effect, a millimeter-scale single-crystal graphene layer was synthesized on a Cu substrate using an atmospheric pressure CVD (APCVD) process (Figure 5b, S9).44 The single-domain of the synthesized graphene assumes a hexagonal shape with various sizes whose edges possess zigzag-edged nanoribbons. The single-crystalline with a domain size exceeding 1 mm2 is selectively transferred to the 76 nm thick parylene. Single-crystalline graphene-laden 100×100 m2 parylene films are self-folded into tubular shapes in a similar manner to polycrystalline graphene. Whereas with the original hexagonal boundary of graphene there is a sharp distinction between folded and flat films, two adjacent hexagonally shaped-single graphene domains cause the subsequent unidirectional self-


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folding to form a tubular shape with a uniform curvature (Figure 5c, Supplementary movie SM3). The folding angles of nine samples of single-crystalline graphene-laden squares and polycrystalline graphene-laden squares as a control group are mapped in Figure 5d. Given monolayer graphene, the angles of polycrystalline graphene-laden parylene are randomly distributed with s.d. values of 59.2 degrees, partly owing to its multi-domain. In contrast, the s.d. values of single-crystalline graphene-laden parylene are less than 19.0 degrees, namely three times smaller than those of polycrystalline graphene. Raman spectroscopy characterizes the ΔG and Δ2D values of single-crystal graphene as +4 and +8 cm−1, respectively (Figure 5e). Although the blue-shift is constant with that of polycrystalline graphene, both ΔG and Δ2D are twice smaller. These results imply that the difference between the uniaxial strains of single- and poly-crystalline graphene originates from the lower elastic modulus as the domain boundary and contributes the reduction of hexagonal symmetry of graphene under the same strain. It is also noteworthy that when the sides of the domain are parallel to those of square films, they inevitably tend to be folded as two facing sides become closer. These results confirm that the folding angle is unidirectionally controllable with the crystalline direction, and is equivalently limited to the armchair-edge direction. Hence, an atomic-scale molecular configuration that includes well-controlled crystal orientations possesses the intrinsic force needed to regulate macroscopic phenomena such as the film self-folding process.

Conclusions


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We have demonstrated that monolayer graphene can induce 3D transformation of the polymeric thin film with a simple transferral method. The adhesion of graphene to parylene is mainly attributable to intermolecular force including π-π stacking, which ensures a conformal and stable contact within bilayers. The strain gradient in homogeneous films initiates self-folding into micro-rolls. Thanks to these traits, the self-folded 3D structures contain few wrinkles or deflections and are robust against highly reactive chemicals and harsh conditions. Remarkably, the 2D micropatterned design and layer thickness are decisive as regards the final 3D macroscopic structure of self-folded micro-rolls, which is theoretically rationalized with the bimorph-beam theory. In essence, a fixed amount of in-plane tensile strain in graphene functions as a driving force that induces bilayer films to bend or form 3D structures. Micro-scale curved graphene exhibits a previously unachievable non-linear electrical behavior without any fracturing or insulation. While the strain is constant in any configuration, the smaller curvature radius is attributed to the increase in resistance non-linearity. This ability to alter the electronic structure makes it possible to employ strain engineering to produce a specific electronic structure or possibly introduce a band gap. In combination with the 3D reconfigurable structure with electrical conductivity, this method could be extensively used for the multifunctional integrated circuits needed for flexible electronics, and field effect devices where the electrical conductivity of graphene is tunable with self-folded geometries. Furthermore, the ability to fold the atomically thin sheet into 3D shapes allows the realization of multi-functional micro-rolls with diverse chemical, electrical, spintronic, and optical features.

Experimental Section


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CVD growth of graphene A polycrystalline monolayer graphene CVD-grown on Cu foils (100 mm×100 mm square) was purchased from Graphene Platform Corporation (No.7, Japan). Large scale single-crystal graphene was synthesized with an atmospheric pressure CVD (APCVD) process following previously reported protocols.44 Briefly, a quartz tube furnace was filled with argon gas under atmospheric pressure. The Cu foils (HA Type, 35 m in thickness, JX Nippon Mining & Metals) were placed in the quartz tube, heated to 1075 °C, and pre-annealed under 1000 sccm Ar. Thereafter, a mixture of 1% CH4 (1 sccm) and H2 (35 sccm) was introduced to induce the graphene growth on the Cu foils. Finally, the as-grown samples were baked on a hot plate at 160°C to visualize graphene boundaries under the microscopes. Figure S9 shows the typical optical images of graphene domains on the Cu foils.

Device fabrication Figure S1 shows the process used to fabricate micro-patterned graphene/parylene bilayers. A 1 wt% solution of sodium alginate was spin-coated on SiO2 wafers 100 m thick (24×32 mm2, Matsunami) at a maximum speed of 2,000 r.p.m. for 50 s. The wafers were immersed in 100 mM calcium chloride for the gelation of the Ca-alginate. Subsequently, they were coated with parylene-C 27 – 322 nm thick using a CVD process (SCS LABCOATER PDS2010). Inside the deposition system, dichloro-di(p-xylylene) was vaporized at 150°C, then pyrolysed at 690°C to generate chloro-p-xylylene monomer, and chloro-p-xylylene was condensed onto the Ca-alginate surface to realize uniform parylene membranes. After the parylene deposition, both singlecrystalline and polycrystalline monolayer graphene were released from Cu foils by floating them


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on Cu etchant, and transferred to the parylene surface by the conventional polymethyl methacrylate (PMMA)-assisted method. The stacked graphene was covered with photolithographically micro-patterned positive photoresist (S-1813G) to protect against damage during subsequent RIE steps. The triple-layered film composed of Ca-alginate, parylene, and graphene was patterned via RIE with oxygen plasma (Mory Engineering) to create a micro-patterned film array. Finally, the arrays were released from the SiO2 wafer via immersion in aqueous EDTA.

Characterization and observation of graphene The graphene transferred on various substrates was characterized by using a confocal Raman microscope (Renishaw, Invia) at an excitation wavelength of 532 nm and with a standard grating (1800 lines/mm). For Raman characterization, the micro-rolls were suspended in water in glass bottomed dishes (D11140H, No.1S, =27mm, Matsunami). The micro-rolls were viewed through a 50× lens with a long working distance and a numerical aperture of 0.5. The exposure time and laser power was set at 10 sec and 2 mW, respectively. Spectra in the 100 to 3100 cm-1 range were measured and accumulated 5 times for each sample. The measurement accuracy was veriﬁed against on Si reference at the wavenumber of 520 nm. For electrical characterization, the graphene samples were transferred to a 120 nm SiO2/Si substrate with coated Ca-alginate sacrificial layer. A BCT-21 MDC prober station with a 4-probe head (Nagase Techno-Engineering Co., Ltd.), 100 m titanium tips arranged in a straight line 1 mm apart, combined with a digital multimeter (Keysight B1500A) were used to record Id-Vd and Id-Vg characteristics. The I-V characteristics were recorded in aqueous solution including distilled water and EDTA at room temperature. A voltage was applied to the source and backgate, and the


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corresponding drain currents were measured. The measurement accuracy was veriﬁed against a 2.5 Ω/□ ITO on glass reference. The flat or curved parylene/graphene bilayers were observed using optical microscopy and scanning electron microscopy (SEM). Optical images were acquired with inverted microscopes (ECLIPSE TE2000, Nikon; BZ-X700, KEYENCE) (Figure 1, 4a, 4b). SEM images were acquired by collecting secondary electrons on an FIB-integrated SEM (Ultra55, Auriga, Carl Zeiss) working at 5 keV (Figure 2a, S5). The micro-rolls were dehydrated by immersing them in ethanol and dried at room temperature using a freeze drier (FS-2030, EYELA). Prior to SEM imaging, the micro-rolls were gold metalized in a metal sputter coater (Ion Sputter E-1030; Hitachi, Japan).


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Figure 1. Self-foldable micro-rolls forming cylindrical shapes. (a) Schematic illustration of the micro-patterning process of sequentially laminated multi-layers. A SiO2 substrate is first coated with Ca-alginate hydrogel (a-1), and subsequently laminated with deposited parylene-C (a-2) and transferred monolayer graphene by a PMMA-assisted transferral method (a-3). After the photoresist had been micro-patterned, the multi-layered film was exposed to O2 plasma for etching and micro-patterning (a-4). (b) Phase-contrast images of an array of micro-patterned square parylene/Ca-alginate films with (left) and without (right) monolayer graphene. While the pristine parylene flattened in EDTA, the free-floating graphene-laden parylene became curved and unidirectionally folded when the structure was free-floating. (c) Time-lapsed phase-contrast


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images of sequentially self-folded micro-rolls. The removal of the sacrificial Ca-alginate layer formed creases in the films and triggered the spontaneous self-folding of the micro-rolls. (d) Schematic illustration and corresponding optical images of the variation in the final 3D structure of the micro-rolls. Scale bars are 100 m.


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Figure 2. Characterization of graphene-laden micro-rolls. (a) Low-magnification scanning electron microscopy (SEM) images of micro-rolls on a SiO2 substrate (a-1). The boxes shown by the white dotted line selected in (a-1) indicate enlarged central areas (a-2) and ends (b-3). (b) Plots of average curvature radii of self-folded 200×400 m2 films () versus various parylene thicknesses (tp). Theoretical curves derived from equation (1) are overlaid on the plots. The results are shown as the mean ± standard deviation (s.d.) values of more than fifty different micro-rolls and are less than 26.8 m. The left insets show the tp and  values of the micro-rolls. (c) Raman spectra of monolayer graphene-laden micro-rolls with different curvature radii, . Enlarged G (c-1), and 2D (c-2) peak regions for graphene. (d) Optical micrograph of flat and


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folded graphene on parylene, and overlaid Raman maps for the G-mode (magenta) and 2D-mode frequency (green) obtained in the areas specified by the white dashed boxes spanning 220×420 and 72×180 m2. The measured curvature radius of the micro-rolls was measured to be 52.4 m. Scale bars: 100 m in (a-1, d), 10 m in (a-2), and 1 m in (a-3).


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Figure 3. Electrical property of graphene-laden micro-rolls. (a) Schematic circuit diagrams and optical image of a partially foldable micropattern design, with folding portions, tapered hinges, anchor sections, and Au connection pads. A portion of the tapered hinges and anchor sections is processed so that it remains flat to allow for probe measurement, while a rectangular portion is processed so that it curves to induce strains in graphene. (b) Representative Id-Vd curves of self-folded 400×1600 m2 122 nm-thick films where  = 47.8 m in flat and folded states at room temperature. (c) Nonlinear Id-Vd curves of self-folded 200×800 m2 films with different curvature radii, . (d) Corresponding nonlinear Id-Vd curves of self-folded 200×400 m2 films where  = 60.9 m with varying Vg. (e) Id-Vg curves of self-folded 400×800 m2 films where  = 60.9 m with a constant gate voltage value (Vd).


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Figure 4. Tunability of 3D spatial geometries of micro-rolls. (a) Optical microscope image of 3D macroscopic structures of micropatterned films. For instance, flower (a-1), auxetic (a-2), mesh (a-3, a-4), and parallelogram (a-5) micro-patterns are transformed into grippers, wired cages, see-through gauze-like structures, and coiled hollow cylinders, respectively. The schematic illustrations depict the final 3D geometries depending on the 2D micro-patterns with tp=85 nm. (b) Phase-contrast images of micro-rolls with different as of graphene layers. (c) Curvature radii of the micro-rolls with different number of transferred graphene layers and with different parylene thicknesses. (d) Enlarged G (d-1) and 2D (d-2) peak regions in Raman spectra of multilayer graphene-laden micro-rolls. Scale bars represent 100 m.


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Figure 5. Crystalline-dependent self-folding orientation. (a) Schematic of the transferring and micro-patterning process for generating single-crystalline graphene-laden micro-rolls. (b) A polarizing microscopy image shows transferred mostly hexagonal graphene grains on parylene surface. A pair of adjacent grains is seen to coalesce to form larger islands. (c) Snapshots highlighting the bilayer films of single-crystalline graphene and parylene in the flat and folded states. White dotted lines indicate the grain boundaries of hexagonally-shaped graphene. (d)


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Scatter diagram and mean values plotting the folding angles of micro-rolls with single-crystalline and polycrystalline graphene in the equilibrium and stable states. (e) Enlarged G (e-1) and 2D (e2) peak regions in Raman spectra of single-crystalline graphene-laden micro-rolls. Scale bars represent 100 m.

Supporting Information. The following files of Supporting Information are available free of charge on the ACS Publications website at DOI: Supporting Information: process for fabricating self-folded graphene-laden micro-rolls; dissolution of sacrificial Ca-alginate layers; multi-angle observation of sequentially self-folded micro-rolls with inclined lenses; Control experiment using silk fibroin protein films; observation of multi-layered films by FIB-SEM; external environmental stimuli; entire and representative Raman spectra; electrical measurement of self-folded micro-rolls; Single crystalline graphene (file type, PDF) Supplementary movie SM1. Time-lapse phase-contrast observation of self-folding process (file type, mp4) Supplementary movie SM2. Multi-angle observation of self-folding graphene/parylene bilayers after the dissolution of Ca-alginate layer (file type, mp4) Supplementary movie SM3. Time-lapse phase-contrast observation of self-folding process of single-crystalline graphene-laden parylene films (file type, mp4)


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AUTHOR INFORMATION Corresponding Author * Tetsuhiko F. Teshima NTT Basic Research Laboratories, NTT Corporation. 3-1 Morinosato Wakamiya, Atsugi, Kanagawa 243-0198, Japan E-mail address: [email protected] Present Addresses † Calum S. Henderson ISIS

Neutron

and

Muon

Source,

Science

and

Technology

Facilities

Council.

Rutherford Appleton Laboratory, Didcot, OX11 0QX, United Kingdom. Author Contributions T.F.T., C.S.H. H.N., and Y.U. and designed the project and experiments. M.T., Y.O., and S.W. grew graphene samples and performed the theoretical simulation. T.F.T., C.S.H., Y.K., S.S., and T.G. fabricated the samples, and undertook the optical and electrical characterization. T.F.T., C.S.H., and Y.U. carried out data analysis. T.F.T. and C.S.H. wrote the manuscript with comments and additions from all the authors. All the authors approved the final version of the manuscript. Funding Sources This work was partially supported by Grants-in-Aid for Scientific Research (JP17H02759) from the Japan Society for the Promotion of Science (JSPS).


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Notes The authors declare that they have no competing financial interests.

ACKNOWLEDGMENT We gratefully acknowledge Prof. U. Schneider of the University of Edinburgh and N. Clément of Centre national de la recherche scientifique (CNRS) for fruitful discussions regarding graphene characterization. We are also grateful to Dr. Y. Sekine, H. Murofushi, W. Kurata, and Dr. A. Tsukada for help with the microfabrication process.

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Table of Contents Graphic

Specification: This illustration depicts the self-folding rectangular bilayer of graphene and polymeric film into cylindrical three-dimensional structures.


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Self-Folded Three-Dimensional Graphene with a ... - ACS Publications

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