Ultra-Lightweight Strain-Responsive 3D Graphene Network - The

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C: Surfaces, Interfaces, Porous Materials, and Catalysis

Ultra-Lightweight Strain-Responsive 3D Graphene Network Dong Hae Ho, Hyun Min Jun, Seon Ju Yeo, Panuk Hong, Min Jun Oh, Byung Mook Weon, Won Bo Lee, S. Joon Kwon, Pil J. Yoo, and Jeong Ho Cho J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.9b00630 • Publication Date (Web): 13 Mar 2019 Downloaded from http://pubs.acs.org on March 13, 2019

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The Journal of Physical Chemistry

Ultra-Lightweight Strain-Responsive 3D Graphene Network Dong Hae Ho,1† Hyun Min Jun,2† Seon Ju Yeo,4† Panuk Hong,1 Min Jun Oh,2 Byung Mook Weon,1,3 Won Bo Lee,5 S. Joon Kwon,4* Pil J. Yoo,1,2*, Jeong Ho Cho6* 1SKKU

Advanced Institute of Nanotechnology (SAINT), 2School of Chemical Engineering, 3School of Advanced Materials Science and Engineering, Sungkyunkwan University (SKKU), Suwon 16419, Republic of Korea. 4Nanophotonics Research Center, Korea Institute of Science and Technology (KIST), Seoul 02792, Republic of Korea. 5School of Chemical and Biological Engineering, Institute of Chemical Processes, Seoul National University, Seoul 08826, Republic of Korea. 6Department of Chemical and Biomolecular Engineering, Yonsei University, Seoul 03722, Republic of Korea. *Corresponding authors: [email protected], [email protected], and [email protected] authors contributed equally to this work.

†These

Abstract In this study, we fabricated three-dimensionally assembled architecture made of reduced graphene oxide (rGO) and utilized it as an ultra-lightweight strain gauge. Building units for the assembly were prepared over multi-scale starting from functionalized GO nanosheets at nano-scale to microfluidically processed solid-shelled bubbles at micro-scale. These GO solid bubbles were elaborately assembled into close-packed 3D structures over centimeter scale and then reduced by thermal treatment. Thermally reduced rGO assembly of which internal structure was spontaneously transformed into closed-cellular structure such as 3D rhombic dodecahedral honeycomb lattice during thermal reduction could manifest superior elasticity against the strain of 30% by virtue of the hierarchically interconnected network while securing a low density of about ~10 mg/cm3 and mechanical robustness, which were then applied as a strain gauge. The strain gauge with thermally-reduced 3D rGO structure exhibited the gauge factor of around 4 and excellent mechanical durability over 250 cycles, suggesting a new pathway for implementing ultra-lightweight strain-sensitive materials. Keywords: three-dimensional assembly, strain gauge, reduced graphene oxide, ultra-lightweight, closed cellular network

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INTRODUCTION Organization of individual building blocks of various shapes into well-controlled, three-dimensional (3D) architectures, such as photonic balls or superlattices, is of both scientific and technological importance since such an organization can impart remarkably improved properties to these materials in comparison to their bulk counterparts, because of which these materials can have wide applicability.1-8 In particular, well-defined 3D structures made of two-dimensional (2D) carbon nanomaterials such as graphene and graphene-related materials are expected to attract extensive research attention and to consequently possess unique physical properties, i.e., low weight, high surface area, enhanced electrical conductivity, and reinforced strength, owing to successful implementation of multiple interfacial networks.9-11 Moreover, recent advances in the design of 3D porous graphene-based architectures have opened up new possibilities for their use in a variety of potential applications, including energy storage, optical materials, catalysis, sorbents, and sensors.12-19 In general, 3D porous graphene materials can be synthesized either by template-free self-gelation methods such as hydrothermal processes, direct freeze-drying, and sol-gel reactions or by template-assisted chemical vapor deposition and 3D printing.16, 18, 20-27 Although these approaches have enabled the synthesis of large-scale free-standing structures of graphene, it remains challenging to fabricate regular and highly ordered cellular structures of graphene while ensuring that they offer advantages of low weight and outstanding mechanical and electrical properties. To overcome these limitations, multiphasic fluid mixtures such as emulsion droplets and bubbles have been investigated as alternative templates for producing 3D hierarchical graphene structures.28-30

Previous

efforts have been largely focused on an approach wherein emulsions of organic solvents stabilized by graphene oxide (GO) nanosheets are prepared by hand-shaking or using a homogenizer, which results in the formation of nonuniformly porous cellular structures.31-32 Although the fabrication of graphene aerogel spheres with a relatively regular size by the wet-spinning method has recently been reported,33 there are several limitations on their progressive utilization, primarily because of a low production yield that results from the multistep procedure, time-consuming post-processing such as freeze-drying, and a restriction on control over the size and uniformity of pores in the cellular network. Because of the remarkable mechanical strength, electrical conductivity, and durability of 3D porous graphene structures, they have received considerable attention, especially for application to mechanical or tactile sensors.12-13, 34-36 Previous studies focused on the fabrication of composite fibers, films, or foams by using graphene as a conducting filler for achieving high stretchability and sensitivity.37 Since most of the reported structures were made of nanomaterial composites in order to ensure their mechanical integrity and elasticity, they have been limited in achieving the property of lightweight density. Moreover, several additional requirements need to be addressed for the realization of high-performance strain sensors, such as a broad sensing range, high sensitivity, low weight, and structural durability and stability. To achieve these goals, in the present study, a hierarchically structured closed-cellular architecture comprising reduced GO (rGO) nanosheets was fabricated for application as a strain gauge. To use the rGO-based building units for construction of the 3D network, gas-in-oil-in-water (G/O/W) compound bubbles were generated ACS Paragon Plus Environment 2

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using a microfluidic system and the generated bubbles were readily stabilized by the adsorption of functionalized GO nanosheets from the oil phase to the gas–water interface. The stability of the GO bubbles was precisely controlled by adjusting the microfluidic processing conditions, such as the types of constituent fluids, GO concentration, and flow rates of the phases. The formed solid-shelled GO microbubbles were then assembled into 3D structures and subsequently transformed into a self-interconnected closed-cellular structure (CCS) through thermal reduction. As a result, enhanced electrical conductivity and elasticity could be achieved, which further aided the architecture in functioning as a strain gauge. The sensitivity and performance of the strain gauge were found to be dependent on the degree of reduction of the GO nanosheets, because this degree is strongly related to the variation in the tunneling distance between stacked rGO nanosheets during strain-induced deformation. The resulting strain gauge showed a gauge factor of around 4 and excellent mechanical durability over 250 cycles, thereby successfully demonstrating the feasibility of application of the synthesized structure as an ultralightweight strain-sensitive material. METHODS Fabrication of glass-capillary microfluidic device The glass-capillary microfluidic device used for the generation of gas-in-oil-in-water (G/O/W) compound bubbles was fabricated by combining co-flow and flowfocusing geometries, as shown in Figure 1d. Two round capillary tubes (1B100-6, World Precision Instruments, Inc.), which were tapered using a micropipette puller (P-1000, Sutter Instrument, Inc.) and a microforge (MF830, Narishige), were aligned in a square capillary tube (inner diameter 1.05 mm, Harvard borosilicate square tubing). In order to improve the wettability of the intermediate phase, a round capillary tube used for injection was treated with ocatadecyltrichlorosilane (OTS), which rendered the outside of the glass surface hydrophobic. The diameters of orifices used for injection and collection were 5 μm and 200 μm, respectively. Synthesis of alkylated graphene oxide Graphene oxide (GO) nanosheets were synthesized by a modified Hummers method using graphite powder as a precursor. Graphite oxides were exfoliated by tip sonication (ULH700S, Ulsso Hitech Co., Korea) for 10 min. Then, the GO suspension was centrifuged at 6000 rpm for 10 min to remove the unexfoliated graphite oxides, and the supernatants were redispersed in deionized (DI) water at a concentration of 4 mg/mL. In order to synthesize the alkylated GO nanosheets, 2 g of 1-ethyl-3-(3dimethylamino) propyl carbodiimide and hydrochloride (EDC) was first added to 400 mL of the GO solution (0.5 mg/mL) at room temperature. For functionalization, octadecylamine (ODA, 2 g) was slowly added to it and then reacted with vigorous stirring at 70 ℃ for 48 h. After completion of the reaction, the solution was rinsed with anhydrous ethanol (99.9%) and then dried at 70 ℃. The ODA-functionalized GO (ODA-GO) nanosheets were dissolved in toluene at a targeted concentration (25 mg/mL). Preparation of ODA-GO-shelled microbubbles Poly(vinyl alcohol) (PVA, 87%–89% hydrolyzed, average Mw = 13,000–23,000) and anhydrous toluene (99.8%) were purchased from Sigma-Aldrich. Nitrogen (N2) gas was used as the innermost phase, and the intermediate oil phase consisted of ODA-GO nanosheets suspended in ACS Paragon Plus Environment 3


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toluene (25 mg/mL). The outer phase was 5 wt% PVA aqueous solution, where PVA was added as a surfactant. To produce G/O/W compound bubbles, N2 gas was injected into the microfluidic device via a pressure regulator (New-Flow DPG3000, New-Flow Technologies, Inc.), and the intermediate and outer phases were injected into the microfluidic device via syringe pumps (KDS 100, KD Scientific). The flow rates of the three immiscible fluid phases could be precisely controlled to manipulate the size and shell thickness of the bubbles. In this study, we set the flow rates of the inner phase (G), intermediate phase (O), and outer phase (W) to 6.5 psi, 3 mL/h, and 90 mL/h, respectively, for the generation of stable and monodisperse bubbles with an average diameter of 115 μm. The generated bubbles were collected in a petri dish filled with 5 wt% PVA aqueous solution, and the organic solvent in the intermediate phase was evaporated at room temperature. Assembly of spherical microbubbles into 3D structures The ODA-GO-shelled bubbles were washed several times with a large amount of DI water in a petri dish to remove impurities and excess PVA. To assemble the lightweight solid-shelled bubbles into a centimeter-scale 3D structure, we developed a new system, i.e., container by using the designed stainless steel mesh (325 mesh). The stainless steel mesh had a pore size of 44 μm, which was smaller than the diameter of the microbubbles in order to retain the microbubbles within the designed container. The bubbles collected in the petri dish were assembled inside the designed container and filtered and washed with DI water to remove the residual PVA on the shell of the bubbles. Then, the container was inverted to rearrange the solid microbubbles by pouring the DI water. After the 3D assembled structures of spherical microbubbles were dried, the structures were thermally reduced at a targeted temperature for 1 h by H2/Ar (1:3) purging in a furnace, which resulted in the formation of 3D self-interconnected CCSs. Characterization Fourier transform infrared (FTIR) spectra (IFS-66/S, Bruker Instruments, Germany) of ODAGO and GO were measured in the absorbance mode in the wavenumber range of 600 cm-1 to 4000 cm-1. X-ray photoelectron spectroscopy (XPS) measurements were performed (ESCA 2000, VG Microtech, UK) with an Al Kα source. Raman spectra were measured using a Raman spectrometer with 532-nm-wavelength incident laser light (inVia Raman microscope, Renishaw). The generation of ODA-GO microbubbles was monitored using an inverted microscope (Nikon Eclipse Ti-U) with a high-speed camera (Phantom M110). Optical microscopy (OM) images of the ODA-GO microbubbles were acquired using a microscope (Olympus BX60) equipped with a charge-coupled device camera (Nikon DS-fi3). All optical images were analyzed using MATLAB to measure the diameter of the ODA-GO microbubbles. Scanning electron microscopy (SEM) images of the thermally reduced structures were acquired using a JSM-7401F (JEOL) scanning electron microscope. The mechanical properties of the 3D rhombic dodecahedral honeycomb (RDH) structures were measured using a universal testing machine (Instron 3343). The compression speed was 20 μm/s, and there was no compression pressure after the samples were compressed to 30% of their original thickness in every cycle. The electrical resistance was measured using source meters (2400 and 2635A, Keithley, USA). The strain-sensing performance was measured using a precision motor controller with an accuracy of 5 μm. ACS Paragon Plus Environment 4

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RESULTS AND DISCUSSION Figure 1a shows the synthesis procedure of alkyl-functionalized GO nanosheets used to fabricate the solid-shelled microbubbles via a microfluidic technique. Alkylated GO nanosheets were synthesized by means of carbodiimide crosslinking chemistry in order to form a dispersion in organic solvents as described previously.38 The carboxyl groups in the as-synthesized GO nanosheets prepared by exfoliating graphite were first activated with 1-ethyl-3-(3-dimethylamino) propyl carbodiimide and hydrochloride (EDC), which resulted in the formation of an active intermediate that was easily displaced by a primary amine group (e.g., octadecylamine (ODA)) through amide bond formation. In addition, alkyl chains were grafted onto the basal plane of the GO nanosheets by means of a nucleophilic substitution reaction between epoxy and amine groups. This yielded ODAfunctionalized GOs (ODA-GOs), which have good dispersibility in organic solvents because of their hydrophobic nature. The chemical functionalization from GO nanosheets to ODA-GO nanosheets was confirmed by both Fourier transform infrared (FTIR) spectroscopy and X-ray photoelectron spectroscopy (XPS). Figure 1b shows the FTIR spectra before and after the alkylation of the GO nanosheets. The strong band observed at around 721 cm-1 was relevant with CH2 rocking vibration of alkyl chains and the peaks at around 2849 cm-1 and 2916 cm-1 were attributed to the C-H stretching vibration of alkyl chains.39-40 In addition, characteristic peaks in the range of 1600–1700 cm-1 and at 1650 cm-1, which were assigned to the C=O stretching of the amide band and the N-H bond of the amide Ⅱ band, respectively, were clearly observed for the ODA-GO nanosheets. The N1s peaks at 400 eV in the XPS spectra were also verified as being evidence of nitrogen species resulting from amine coupling via EDC chemistry and nucleophilic substitution reactions (Figure 1c). In addition, detailed analysis of deconvoluted XPS spectra of C1s orbitals of ODA-GO shows a new C-N peak at 285.8 eV, which support the successful synthesis of alkylated GO nanosheets (Figure S1). Meanwhile, from the C/O ratios of the GO and ODA-GO nanosheets (see Table S1 in Supporting Information), the degree of alkylation of the ODA-GO nanosheets was estimated to be around 100%, implying that the basal plane of the GO nanosheets was sufficiently decorated with alkyl chains. Next, highly stable and monodisperse ODA-GO-shelled bubbles were fabricated using microfluidic G/O/W compound bubbles as templates (Figure 1d). The following three immiscible fluidic phases were used in a glass-capillary microfluidic device: (i) nitrogen (N2) as the inner gas, (ii) ODA-GO nanosheets suspended in an organic solvent as the intermediate phase, and (iii) poly(vinyl alcohol) aqueous solution as the outer phase. Any types of inert gases can be used to generate solid-shelled microbubbles using a microfluidic device. However, the stability of the bubbles depends on the types of inner gas. As an example, helium gas was used as a bubble source with the same condition and most of the microbubbles were collapsed during the solvent evaporation from the shell. This can be ascribed to the larger contact angle of the bubbles at an air-water interface due to the high buoyancy force of the helium bubbles. In addition, when the air-exposed surface area of the bubbles becomes larger, the capillary force between the ODA-GO shell and water surface increases to compensate buoyancy force, ACS Paragon Plus Environment 5


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which leads to the generation of unstable microbubbles. As has been reported previously,41 the resulting morphologies and dimensions of bubbles are influenced by several processing variables, including the geometry of the microfluidic device, properties of the constituent fluids, and flow rates of the constituent fluids. The variables selected in this study, i.e., the type of solvent, the concentration of the ODA-GO solution, and the flow rates, were systematically controlled to optimally generate monodisperse spherical bubbles for subsequent assembly of the 3D network structure. First, we investigated the stability and morphology of the ODA-GO-shelled bubbles as functions of the type of organic solvent used, such as dichloromethane (DCM), chloroform, and toluene. When DCM and chloroform were separately used as the organic solvents for the oil phase, the surface of the bubbles underwent mechanical buckling during evaporation of the solvent and eventually became concavely deformed. However, when toluene—having relatively lower polarity—was used as the organic solvent, the bubbles were able to retain their sphericity, which led to the generation of highly monodisperse and stable ODA-GO-shelled bubbles (Figure 1e). These contrasting behaviors of the bubble shape and stability are ascribed primarily to the difference in the dispersibility of ODA-functionalized GOs in different organic solvents.42-43 Although the ODA-GO nanosheets were intrinsically well-dispersed in all three organic solvents, a slight difference in the dielectric constants of the solvents would affect the stability of the dispersibility of the ODA-GO nanosheets in the solvents: the dielectric constants of toluene, chloroform, and DCM are 2.38, 4.81, and 8.93, respectively.44 Thus, the stronger affinity between the less polar toluene and the ODA-GO nanosheets made the ODA-GO-shelled bubbles more stable. Next, we characterized the effect of the concentration of the alkylated GO nanosheets dispersed in toluene on the stability and morphology of the bubbles. Although a concentration in the range of 10–40 mg/mL was found to have a negligible influence on the stability of the bubbles, concentrations in this range influenced the size and shell thickness of the bubbles under fixed conditions of microfluidic synthesis (e.g., fixed flow rates and types of fluids). However, an increase in the concentration of the alkylated GO nanosheets to above 40 mg/mL caused the solution to become highly viscous, which made it difficult to precisely control the size and shell thickness of the compound bubbles. In contrast, use of a solution with concentrations lower than 10 mg/mL resulted in the formation of spherical ultrathin-shelled bubbles. These bubbles were rather prone to mechanical collapse after drying (Figure S2). Finally, on the basis of the pre-obtained information discussed above, the flow rates of the inner gas phase (G), intermediate oil phase (O, with ODA-GO dispersion), and outer aqueous phase (W) were varied to achieve an optimum size and shell thickness of the compound bubbles, which could subsequently be used in the assembly of the 3D structure; achieving these optimum dimensions was necessary since bubble stability is strongly associated with the ratio of the shell thickness to the bubble radius (h/R).45 When the flow rate in the middle phases was increased from 1500 to 4000 μL/h under fixed flow rates of the inner gas and outer liquid at 6.5 psi and 90 mL/h, respectively, the stability of the generated bubbles gradually decreased as the flow rate of the middle phases decreased below 3000 μL/h, as shown in Figure 1f. The fraction of the stable bubbles with an average ACS Paragon Plus Environment 6

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diameter of 115 um and a coefficient of variation smaller than 1% is over 80% at an oil-phase flow rate of 3000 μL/h. A decrease in the flow rate of the intermediate oil phase led to a decrease in the shell thickness, but it had no significant effect on the bubble size. Therefore, as the oil flow rate decreased, the bubbles became unstable with smaller values of h/R, mainly because of the thinning of the shell. Meanwhile, a change in the flow rate of the outer phase under the condition of fixed flow rates of the inner and intermediate phases (6.5 psi and 3000 μL/h, respectively) significantly affected the monodispersity and size of the generated bubbles, as shown in Figure 1g and Figure S3. The size of the spherical bubbles substantially decreased from 140 μm to 70 μm with an increase in the flow rate of the outer fluid phase from 90 mL/h to 180 mL/h, and the stability also increased because the value of h/R became much larger. To sum up, spherical ODA-GO-shelled microbubbles with high stability and uniformity were obtained by elaborate modification of the microfluidic conditions, which were intended to be subsequently utilized for assembly of the 3D structure in the next step. We then assembled microscale spherical solid-shelled bubbles into 3D hierarchical structures for realization of a strain gauge. Because bubbles were continuously generated in a high yield (~105 bubbles per minute), the 3D hierarchical structures could be easily formed at the centimeter scale. Figure 2a shows the schematic procedure for the fabrication of 3D structures of spherical ODA-GO-shelled bubbles. Before the bubbles were used as building units for the 3D structures, they were completely washed with deionized (DI) water to eliminate impurities, including excess PVA and residual GO nanosheets, resulting from burst bubbles by vacuum filtration. The washed bubbles were then redispersed and collected in a petri dish filled with DI water. Then, to fabricate 3D assembled arrays of the spherical solid-shelled bubbles, a stainless steel mesh (325 mesh) was used as a guiding template and shaped into a cubic container; it should be noted that the size of the pores of the mesh used should be smaller than that of the bubbles to prevent the bubbles from passing through the pores. The designed mesh container was then filled with a suspension of solid-shelled bubbles. In order to induce the formation of structurally stacked arrays of bubbles, the buoyancy-assisted self-assembly process was exploited; when the mesh cube containing spherical bubbles was placed upside down in water, the bubbles floated to the upper part (originally the bottom side) of the container because of buoyancy, which enabled spontaneous rearrangement of the assembled arrays of spherical bubbles. The 3D assembled structures of spherical thin-shelled bubbles had a low density of ~10 mg/cm3 on account of the efficient construction of the hierarchically porous structures, which could easily be modulated further by varying the size or shell thickness of the bubbles. However, these structures were quite weak and fragile because the spherical ODA-GO-shelled bubbles were simply stack-assembled into a close-packed face-centered cubic (FCC) structure, which was loosely linked to the nearest-neighbor bubbles by only point contacts. The physical properties of assembled structures can be greatly improved by increasing the interfacial area between the basic assembling units. To achieve this structural configuration, we performed thermal reduction treatments of the assembled-bubble structures comprising alkylated GO nanosheets, by which the structural interconnection and formation of an rGO-based network could be achieved. In due course, the initially assembled FCC structure of ACS Paragon Plus Environment 7


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spherical bubbles transformed into a CCS with a rhombic dodecahedral honeycomb (RDH) lattice geometry of the assembling units, which is a 3D Voronoi tessellation of an FCC structure, as shown in Figure 2b. Notably, in the thermal reduction process of GO nanosheets, the reduction temperature is considered an important factor in determining the physical properties of rGO.46-47 To verify this opinion, four different samples were prepared by controlling the reduction temperature to 300 °C, 500 °C, 600 °C, and 700 °C. First, the structural changes of the samples reduced at different temperatures were observed. The intensity ratio (ID/IG) increased slightly from 0.95 to 1.05 after the alkylation of the GO nanosheets (Figure S4), indicating that structural and compositional changes occurred during the alkylation process.48 Meanwhile, the ID/IG values gradually decreased from 1.02 to 0.76 with an increase in the reduction temperature, which is summarized in Figure 2c and 2d. It is also well known that the sp2 fraction of a single GO layer and thus of alkylated GO nanosheets increases with an increase in the reduction temperature, mainly because of recovery of the carbon backbone with a decrease in the number of oxygen-containing moieties.49 Consequently, the variation in the structural nature of constituent materials is expected to induce changes in the resistivity of RDH structures. Indeed, the resistivity of the 3D CCS gradually decreased from 108.9 Ω∙cm to 3.3 Ω∙cm with increasing reduction temperature (Figure 2e) and the conductivity of the structure reduced at 700 °C was measured to be as high as 0.303 S/cm at a density of 12.7 mg/cm3, which is higher than that of carbon nanotube aerogels and comparable to those of other cellular foams of graphene.50-51 This unique property was attributed to the presence of a number of conduction pathways through interconnected graphene sheets in the hierarchically assembled structure. It should be noted that all the samples showed similar structural morphologies and dimensions despite being reduced at different temperatures (Figure S5), and the density of the 3D RDH structures was determined to be in the range of 12–15 mg/cm3. More interestingly, the mechanical properties and elasticity of these structures were also highly dependent on the reduction temperature, which could have been attributable to the intrinsic mechanical properties of the constituent materials. The assembled structure reduced at 700 °C possessed a very high Young’s modulus of 700 kPa with a density of 12.7 ± 0.3 mg/cm3, whereas that of the structure reduced at 300 °C was observed to be only 80.5 kPa at a density of 15.2 ± 0.7 mg/cm3 (see Figure 2f and 2g). In addition, the RDH structures thermally reduced at higher temperatures (>500 °C) showed complete structural recovery upon application of a strain of 30%. Such high elasticity could be ascribed to the hierarchically interconnected network of rGO nanosheets containing a large sp2 fraction. After repeated cycles of mechanical compression and recovery, the initially high Young’s modulus decreased because of the high energy absorption capability of the structures. It is well known that macro-sized porous structure can effectively absorb energy. Few research groups have simulated energy absorption of the porous structure and shown the possibilities of using as a shock absorber. It is likely that some rGO nanosheets that have a high modulus in the sample get irreversible damage from external stress while remaining rGO parts keeps the overall shape of the sample.52-55 Subsequently, the Young’s modulus became stable after about 10 cycles (Figure S6), thereby indicating the superelasticity and durability of these structures. As the results, Young’s modulus of the rGO structure reduced at 700 °C were decreased to 45.8 kPa ACS Paragon Plus Environment 8

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(10 cycles) and 52.2kPa (20 cycles). In contrast, RDH structures thermally reduced at a relatively lower temperature of 300 °C showed low elasticity and a poor recovery response at a compressive strain of up to 30% (Figure 2g); therefore, these structures reduced at lower temperature may not be suitable for the realization of flexible and superelastic strain sensors. Owing to the excellent compression recovery of the thermally reduced 3D RDH structures over a wide strain range, they showed promising potential for application as a strain gauge. Figure 3a shows the resistance of the RDH structures prepared at various reduction temperatures as a function of the applied compressive strain. The change in the resistance of the strain gauge was recorded by the two-terminal test method during the deformation process. Both the top electrode and the bottom electrode were made of silver paste. The resistances of all the strain gauges decreased with increasing compressive strain. The variation in the resistance of the 3D RDH structure was governed mainly by both the conduction pathways and the contact resistance among the rGO nanosheets. With the application of compressive strain, the density of the 3D structure increased and the rGO walls subsequently started to come into close contact with each other. Therefore, the formation of additional conduction pathways resulted in a decrease in the resistance. After release of the compressive strain, the resistance recovered completely to the initial level with negligible hysteresis. Figure 3b shows the normalized resistance ( R / R0 ) of the strain gauge, where

R0

is the initial resistance and R is equal to

R  R0 .

Interestingly, the

sensitivity of the 3D graphene structures improved greatly with an increase in the reduction temperature. The rGO reduced at 700 °C showed the gauge factor of ~4. In case of common steel or stainless steel like metal type commercial strain sensors, commonly have the gauge factor of 2 or 3 which are lower than our case. Furthermore, the range of the maximum strain is limited to 2% or less. In order to elucidate the mechanism of strain sensing of the 3D cellular network structures of rGO, we first investigated the electrical transport properties of an initial building block, i.e., a single GO nanosheet. As described above, the sp2 fraction and density of topological defects such as single and/or double vacancies and Stone–Wales defects increased and decreased, respectively, with an increase in the reduction temperature as revealed by Raman spectra (Figure 2c).56 Therefore, it is reasonable to suggest that an rGO flake treated by thermal reduction would recover sp2 carbons, which would be accompanied by the annihilation of topological defects. These effects of the reduction temperature are collectively reflected by the increase in the electrical conductivity of the rGO nanosheets. Under the assumption that there are no spatial correlations among the domains of sp2 carbons in a single GO nanosheet, we can describe the conductivity by using a simple percolation model. In particular, the percolation model employs an sp2 carbon domain as a conducting entity such that

   0  x  xC  ; here, 



and

0

are the overall electrical conductivity of a single GO nanosheet and the

intrinsic conductivity of the sp2 carbon domain, respectively; flake and the percolation threshold of the flake; and



x

and

xC are the fraction of sp2 carbon in the GO

is a phenomenological percolation exponent. The


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percolation threshold of an rGO nanosheet subjected to thermal reduction at 220 °C has been reported to be

xC

~ 0.6.57-59 Therefore, in the case of our GO nanosheets reduced at a temperature higher than 300 °C, the conducting is common for all the samples: sp2 carbon. This indicates that the intrinsic conductivity dominated by

entity

sp2 carbon is beyond the threshold value, and therefore, sp2 carbon does not play a critical role in determining the overall resistance of the compressed 3D structure. Instead, the overall resistance is governed by the interflake resistance and the CCS of the sample. In particular, when the nanosheets are assembled in 100-nm-thick shells, the conductivity of the nanosheets reduces drastically mainly because of the presence of vacancies among the flakes, misorientation of the carbon lattices of neighboring GO flakes, topological defects, and nonflat morphologies of a single GO flake.60 These factors work together to reduce the conductivity of the flake assembly. In other words, the effect of the flake assembly on the GO nanosheet conductivity can be described by introducing a tunneling resistance among the flakes,

Ri , as follows: 4  2m  1  8 hd  L Ri   exp  d   2  i ,   2 2  n  3 A e  A h

1/2

,

1

where n is the density of pathways for interflake conduction, which is equal to N/L (where N and L denote the number of conduction pathways and the path length, respectively); h is Planck’s constant; d is the tunneling distance; A2 is the effective cross-sectional area; e is the electron charge;

i

is the tunneling resistivity; m is the

electron mass; and  is the potential barrier.61 The tunneling resistance is sufficiently greater than the intrinsic conductivity of a single rGO nanosheet, and thus, it is reasonable to consider Ri as the intrinsic resistance of the shell materials of the 3D cellular structures. When compressive strain is applied, the shell materials undergo deformation, which results in a change in the tunneling resistance such that Ri    n0 1    exp   d 0   1, Ri     Ri     Ri 0 , Ri 0  Ri  0  ,  Ri 0 n  

 2

where we assume  to be constant. Next, we need to consider the change in the resistance caused by the deformation of the 3D CCS. By employing a simple Kirchhoff model to represent the CCS, we can describe the overall resistivity of the cellular structure,

 T , by using the shell resistivity  S T   S 1    ,

 3



where



as follows:

is the porosity of the cellular structure and  is close to 2/3 for the 3D CCS.62 Under the assumption

that the cellular materials are composed of the shell materials such that

 S  i , the overall resistance of the CCS,

RT, can be expressed using eqs.(2) and (3) as follows: RT     Ri    1      



,


 4

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and     can be expressed as 1       1  0 , 0    0  by the use of the volume conservation relation of 1  the entire 3D CCS structure. Using eq.(4), we can calculate the gauge factor as GF 

1 R    , which can be  R0

approximated as follows:

GF 

1 R     R   / R0    R0 

  0

n0 5 1   d 0     1   d 0    d 0 . n 3

 5

In eq.(5), we assumed that the pathway density n is independent of the compressive strain. From eq.(5), it is evident that GF monotonically increases with an increase in the initial tunneling distance d0. According to experimental observations, GF increases with an increase in the reduction temperature; therefore, we can conjecture that d0 increases with an increase in the reduction temperature. This tendency can be understood by considering the geometry of the rGO nanosheets; the recovery of the sp2 nature along with a reduction in topological defects with increasing reduction temperature leads to a flatter geometry of a single rGO nanosheet and a denser local assembly. Then, the interdistance among the local assembly of GO flakes increases with increasing reduction temperature, which results in an increase in d0 with an increase in the reduction temperature. Using eq.(5), we can compare d0 at various reduction temperatures. To this end, we obtain GF values by numerical differentiation of

R( ) / R0 with respect to the strain in the vicinity of the zero-strain point. As depicted in the

inset of Figure 3c, d0 for the 3D cellular network structure of rGO reduced at 700 °C is 1.7 times and 3.2 times those of the structures reduced at 600 °C and 500 °C, respectively. From the fact that rGO shows saturation behavior with an increase in the reduction temperature, we can also conjecture that

d 0 approaches a limiting

value, which, in turn, leads to a limiting value of GF. During thermal reduction, a GO flake is partially reduced to rGO in the initial stage. As the reduction temperature increases, the rGO percolation network develops and electrical interconnections spread over the entire surface, as shown in Figure 3d. Figure 4a shows photographic images of the 3D RDH structure reduced at 700 °C realized as a strain gauge, which were captured during finger loading and unloading processes. The strain gauge was compressed completely and released using a finger. The images reveal excellent elastic performance of the strain gauge, with instantaneous complete structural recovery. In order to investigate the dynamic response of the strain gauge, it was subjected to a series of random strains and the variation of resistance was measured. First, 5 different compressive strains (1.6%, 3.5%, 5.7%, 7.5%, and 9.0%) were applied, after which the strain gauge was restored to the initial state; the strain gauge was held in each state, i.e., at each strain, for around 10 s (Figure 4b). The normalized resistance of the strain gauge responded immediately to the applied external strain, and it increased negatively to -0.08, -0.16, -0.24, -0.30, and -0.35 at these strains, respectively. A stable current signal was maintained at each strain, and the strain gauge was restored completely to the initial state at all pressures after removal of the external applied strain. Second, the strain gauge performance was tested by applying compressive ACS Paragon Plus Environment 11


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strains in small increments (Figure 4c). The device exhibited a stable response not only when the strain was applied but also when it was retained in a given state. After the highest strain was applied, it was decreased to the initial level in the same increments. The resistance decreased in a step-by-step manner to the initial level without any signal loss. These results indicated both the outstanding stability and the high elasticity of the strain gauge at different strains. Third, the variation of the resistance under application of strains with different histories was also monitored, as shown in Figure 4d. For example, strain of 0 to 0.05 (i.e., 0.1, 0.2, or 0.3) was applied to the device and then released to 0. The resistance was perfectly synchronized with the applied strain without any hysteresis. The response did not depend on the strain history, indicating the highly reliable performance of our strain gauge. Furthermore, the cyclic stability of the strain gauge was verified by repeating loading–unloading cycles at a strain of 10% (Figure 4e). The strain gauge showed a durable and highly reliable response over the repeated loading– unloading cycles, which lasted for 1000 s (250cycles). Strain gauge exhibited the stable signal output over 250 cycles and the good mechanical durability until 350 cycles. After 350 cycles, it showed the slight increment of the base signal. Finally, the minimum detectable strain level and response time were measured. Figure 4f shows a discrete increment of the normalized resistance value from 0 to 0.017 with a short response time of

Ultra-Lightweight Strain-Responsive 3D Graphene Network - The

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