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Three-Dimensional Printing Hollow Polymer Template-Mediated Graphene Lattices with Tailorable Architectures and Multifunctional Properties Qiangqiang Zhang,*,†,‡,∇ Feng Zhang,§,∇ Xiang Xu,∥,⊥ Chi Zhou,*,§ and Dong Lin*,# ACS Nano 2018.12:1096-1106. Downloaded from pubs.acs.org by IOWA STATE UNIV on 01/09/19. For personal use only.

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College of Civil Engineering and Mechanics, Lanzhou University, Lanzhou 730000, P. R. China Key Laboratory of Mechanics on Disaster and Environment in Western China and the Ministry of Education of China, Lanzhou University, Lanzhou 730000, P. R. China § Department of Industrial and Systems Engineering, University at Buffalo, the State University of New York, Buffalo, New York 14260, United States ∥ Key Lab of Structures Dynamic Behavior and Control of the Ministry of Education, Harbin Institute of Technology, Harbin 150090, P. R. China ⊥ Center of Structural Health Monitoring and Control, School of Civil Engineering, Harbin Institute of Technology, Harbin 150090, P. R. China # Department of Industrial and Manufacturing Systems Engineering, Kansas State University, Manhattan, Kansas 66506, United States ‡

S Supporting Information *

ABSTRACT: It is a significant challenge to concurrently achieve scalable fabrication of graphene aerogels with three-dimensional (3D) tailorable architectures (e.g., lattice structure) and controllable manipulation of microstructures on the multiscale. Herein, we highlight 3D graphene lattices (GLs) with complex engineering architectures that were delicately designed and manufactured via 3D stereolithography printed hollow polymer template-mediated hydrothermal process coupled with freeze-drying strategies. The resulting GLs with overhang beams and columns show a 3D geometric configuration with hollowcarved features at the macroscale, while the construction elements of graphene cellular on the microscale exhibit a wellordered and honeycomb-like microstructure with high porosity. These GLs demonstrate multifunctional properties with robust structure, high electrical conductivity, low thermal conductivity, and superior absorption capacitance of organic solvents. Moreover, the GLs were utilized as a subtle sensor for the fast detection of chemical agents. Aforementioned superior properties of GLs confirm that the combination of 3D tailorable manipulation and self-organization design of structures on the multiscale is an effective strategy for the scalable fabrication of advanced multifunctional graphene monoliths, suggesting their promising applications as chemical detection sensors, environmental remediation absorbers, conductive electrodes, and engineering metamaterials. KEYWORDS: 3D graphene lattices, stereolithography, 3D printing hollow polymer, tailorable manipulation, self-organization design

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approaches to achieve tailorable manipulation and scalable fabrication of GM structures on the multiscale.12−16 Thereinto, 3D graphene aerogels (GAs) in sponge, foam, array, and lattice formats have recently become one of the most attractive GMs because of their low density (99%),7,17 high electrical conductivity (∼10 S cm−1),18 large Joule heating rate (∼20,000 °C

he three-dimensional (3D) porous graphene monoliths (GMs) as emerging carbon nanostructures have converted the intrinsic outstanding properties of pristine graphene into macroscopic applications, including supercapacitors,1 catalysts,2,3 stretchable electrodes,4 actuators,5 damping materials,6,7 thermal insulation and fire retardancy,8,9 and environmental remediation.10,11 However, the attainment of both well-ordered micromorphologies and specific engineering architectures for the multiscale fabrication of GMs is still under significant challenge. To meet further multifunctional demands, there is an intriguing interest to develop controllable © 2018 American Chemical Society

Received: August 27, 2017 Accepted: January 12, 2018 Published: January 12, 2018 1096

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Figure 1. (a) The schematic of 3D GL fabrication within HPA template. (b) Optical image of individual GO sheet with a diameter up to ∼40 μm. (c) AFM micrograph of monolayer GO sheets with a thickness profile in the range of 0.8−1.2 nm. (d) TEM image of isolated graphene microflake within GLs showing typically wrinkled morphology. The inset presents clear hexagonal SAED patterns. (e, f) SEM images of highly porous microstructures within GLs at different magnifications. (g) SEM micrograph of multilayered graphene cellular wall consisting of GL porous scaffolds with a typical wrinkled morphology.

s−1),19 and superinsulated conductivity (∼0.0126 W m−1 K−1).9 Scale-up fabrications of micro/nanosized graphene sheets into macroscopic GAs via large-scale assembly have been realized by in situ self-assembling,20 chemical cross-linking,21 and freezecasting strategies.7,22 On the microscale, most GAs have stochastic porous network or poor manipulation of the microstructures even though the orientation of graphene sheets can be locally mediated by freeze-casting strategies under the squeezing of ice crystals induced π−π stacking.18,22,23 In addition, simple cubic or cylindrical reactors are normally used for graphene production on macroscopic dimensions, which limits bulk shaped GAs to be tailored into precise geometric configurations or complex topological morphologies.24,25 This inhibits the assembly of graphene sheets into the well-ordered formats on the multiscale, hindering the applications of GAs in broad fields, such as high-performance energy storage devices, functionalized bio/chemical sensors, and stretchable electrodes. Recently, 3D printing techniques have been utilized to construct graphene oxide (GO) monoliths, which would be further reduced to 3D conductive GM architectures after thermal decomposition or chemical etching of additives.14,18 Worsley et al. used a silica additive to adjust the viscosity of the GO processor and formulated a non-Newtonian paste-like ink with shear-thinning behavior. A 3D printing technique was used to build GA/silica microlattices with an engineered architecture and elastic characteristic.16 Shah et al. reported a 3D printed biocompatible graphene composite from a printable graphenepolymer gel consisting of graphene and polylactide-co-glycolide. The fabrication process was implemented at ambient conditions via direct writing with structural features as small as 100 μm.24

Zhang et al. developed a freeze-casting-based multinozzle drop-on-demand inkjet printing approach through depositing GO ink onto a cold substrate. The 3D printing GAs with overhanging architectures demonstrated high electrical conductivity and stimuli-responsive performance to mechanical deformation.25 Hu et al. demonstrated 3D printing of all-solidbatteries by direct writing of printable GO-based electrode composite inks and solid-state electrolyte inks with shearthinning behaviors as typical non-Newtonian fluids. The stateof-the-art 3D printing techniques enable the manipulation and construction of graphene sheets on the multiscale level into specifically designed patterns with optimized pathways for high efficient transport of charge carriers (e.g., electrons) or transmission of loading force. These merits facilitate multifunctional productions of graphene derivations for diverse applications via multidimensional complex structures,14,16,18,24,25 compared with those of cubic or cylindrical shaped bulk GAs synthesized by conventional approaches with relative low controllability on graphene orientations.7,20−22 In addition, the extraordinary controllability and tailorability on the geometry and topology of the GAs allow for the construction of intricate internal structures as well as complex external architectures which can meet application-specific needs, including point-of-use filter in environmental recommendation and conformal organ repair in biomedical and health applications. Although these strategies achieve tailoring graphene-based materials via direct writing or freeze-casting, rheological studies indicate that the printability of these graphene-based inks depends on the additive components, viscosity, shear stress, and ratio of the storage modulus (G′) to the loss modulus (G″).16,24−26 Therefore, these direct printing strategies still 1097

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Figure 2. Characterization of structure and chemical composition of GLs. (a, b) XRD and Raman spectra at four stochastically selected positions. (c) TGA. (d) XPS for all elements at four different spots. (e) The deconvoluted spectrum of C1s. (f) Nitrogen adsorption− desorption isotherms. The inset contains the pore size distributions.

have some inevitable disadvantages, including intercalation of additive fillers (e.g., silica powder, polymer matrix) or weak π−π stacking among the graphene sheets. These shortcomings lead to poor mechanical robustness and low electrical conductivity because of the numerous insufficient microstructural joints formed via weak physical assembly. This is in contrast to sufficient π−π interactions that result in stronger cross-linking interfaces among graphene sheets via chemical-assisted hydrothermal strategies.4,17,20 Therefore, the scalable fabrication of GAs with truly 3D tailored macrostructures and more controllable microstructure remains a significant challenge. In this study, we highlight 3D graphene lattices (GLs) with controlled complex structures by hollow polymer architectures (HPAs)-mediated fabrication of GO/ethylenediamine (EDA) ink via hydrothermal assembly, followed by freeze-drying strategies. The HPAs were designed and printed by a maskimage-projection-based stereolithography (MIPSL) technique with optimized wall thickness. After thermally etching the polymer template, the as-obtained GLs present a truly 3D geometric configuration consisting of overhang beam and column elements with hollow-carved topological morphologies. The porous structure exhibits a well-ordered and honeycomblike microstructure with the multilayered graphene cellular walls assembled via strong π−π interaction. These GLs have a wealth of multifunctional properties such as structural robust-

ness, elastic deformation, high electrical conductivity, satisfactory Joule heating effect, and huge capacitance for organic solvent absorption. In addition, the GLs were designed as a sensor enabling fast detection of chemical agents (e.g., methane, ethylene glycol, methanol, dimethylbenzene, n-hexane, ethanol, and acetone), indicating the capability to distinguish different chemicals.

RESULTS AND DISCUSSION Microstructure and Chemical Composition Characterization. Figure 1a illustrates the four steps of 3D GL fabrication. First, we injected the GO/EDA mixture into hollow chamber of HPAs (wall thickness ∼0.2 mm) to facilitate the template-mediated tailoring of 3D GL architectures. The HPAs scaffolds were designed by Creo-Parametric and printed using a MIPSL 3D printer (Perfactory μicro) with clear photocurable resin (FLGPCL02 from MakerJuice Laboratories). The more detailed fabrication procedures, parameter setup, and design principles of HPAs are discussed in the Supporting Information (see Figures S1−S6). Next, the GLs were fabricated within HPAs after modified hydrothermal assembly (120 °C, 6 h, 3 vol. ‰ EDA) by directly placing the GO/EDA mixture filled template into an autoclave reactor. After that, the HPA template containing hydrogel-like GLs was freeze-dried under vacuum condition 1098

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ACS Nano (−60 °C, 24 h). The GLs synthesized by the appropriate hydrothermal reaction fill up all hollow space within HPAs. In contrast, excessively high or low temperatures fail to form the intact GL architectures because of severe shrinkage induced fracture or collapse of structural elements (Figures S7 and S8). Moreover, the optimized formation parameters of GLs that maintain original HPA template dimensions with a negligible shrinkage during hydrothermal and followed freeze-drying processes have been verified through previous reports (Figure S8a).4,7,17 This has little to no direct association with the HPAs template’s geometric configuration, but relies more on the reaction temperature, EDA content, GO concentration, and freezing temperature. Therefore, we can also create more complex HPAs-mediated GM morphologies, such as handshaped graphene bioscaffolds within bionic polymer templates (Figure S9). Ultimately, a pure cubic GL sample was obtained with a mechanically robust and precisely tailored engineering architecture after thermal decomposition of HPA template (Figure S8a). The GL was further reduced to recover high electrical conductivity by thermal annealing at 1000 °C. Moreover, this HPAs template-mediated design and scalable fabrication of GLs also validate the potential to functionalize scaffolds for biological tissues or bionic organs with defined and complex shapes such as 3D graphene liver bioscaffolds and mechanical gear parts (Figures S5 and S9). Specifically, the quality of the GO precursor used to construct GLs was characterized by atomic force microscopy (AFM), as shown in Figure 1b. An individual monolayer GO sheet has a thickness profile of 0.8−1.2 nm and a large surface area of ∼1200 μm2 with an average diameter up to ∼40 μm (Figure 1c), indicating an ultralarge aspect ratio (∼5 × 104). This facilitates the self-assembly of GO sheets into multilayered cellular walls via strong π−π interaction-induced stacking to form 3D frameworks.4 Transmission electron microscope (TEM) image (Figure 1d) shows the typical wrinkled morphology of the isolated graphene microflakes within the GLs. The inset of Figure 1d shows clear hexagonal patterns in the selected-area electron diffraction (SAED), indicating the highly graphitic nature of the thermally reduced GO (rGO) derivations. The scanning electron microscopy (SEM) image in Figure 1e shows highly porous microstructures within beam or columnlike elements of the 3D GLs, which consist of hierarchical pores sized from hundreds of nanometers to tens of micrometers. As shown in Figure 1f, GO sheets were first assembled into faciallink “Y-shaped” nodes with strong π−π interactions during hydrothermal process. These nodes subsequently formed a hierarchical porous network under oriented squeezing of the ice crystals during freeze-drying. This multiscale tailoring manipulation on the micro- and macro-topological structures enables multifunctionalization and controllable modification of 3D GMs. In addition, the SEM micrograph in Figure 1g demonstrates typical wrinkled morphologies of the multilayered graphene sheets that serve as the basic cellular wall to assemble the porous microscaffolds within the GLs. The structure and chemical composition of the GLs were characterized by X-ray diffraction (XRD), Raman spectroscopy, X-ray photoelectron spectroscopy (XPS), and Brunauer− Emmett−Teller (BET) analysis, while the thermogravimetric analysis (TGA) was used to evaluate its thermal stability. The XRD results (Figure 2a) peaked at 2θ = 26.32° imply that the interlayer spacing among stacked rGO sheets within GLs shifts

from 7.30 Å of original GO ink to 3.36 Å as HPAs were thermally decomposed (see Figure S12a). The Raman spectra in Figure 2b illustrate dominating D- and G-bands peaking at 1343 and 1586 cm−1, respectively, and the related intensity ratio (ID/IG) enlarges from 0.87 to 0.92 corresponding to GO processor and thermally reduced GLs (see Figure S12b). In addition, the 2D- (2713 cm−1) and D + D′- (2940 cm−1) bands corroborate the sufficient recovery of the nature graphitic structure and the significant decrease of defect densities on rGO sheets.9 Figure 2c elucidates that the weight drop of GLs at an initial moment with temperature lower than 100 °C is up to 6.9%, which is possibly because of the absorbed water vapor from ambient. The final effective weight loss within GLs is about 15.8% at 1000 °C due to the decomposition of residual polymer. Besides the highly graphitizing performance of final GLs, the evolution of PHA compositions over the thermal annealing processes was further characterized by Fourier transform infrared (FT-IR) spectrum. As shown in Figure S12c, with the increases of thermal annealing temperature and processing time, the number of absorption peaks dramatically decrease, and the related intensities became weak. Compared with primary HPAs and thermal annealing samples at 320 and 400 °C, the FT-IR curves for 500 and 1000 °C annealing processes show negligible absorption peak at 2914 cm−1 for the stretching vibration of C−H due to the bond fracture during thermal decomposition. The same phenomena also occurred at other absorption peaks ranging from 700 to 1800 cm−1, such as stretching vibration of CO at 1733 cm−1 and bending vibration of C−H at 750 cm−1. In particular, the peaks at 1624 and 1114 cm−1 characterizing the CC and C−C become weak with the proceeding of thermal decomposition, indicating the high carbonization of residual polymer as a form of amorphous carbon. XPS analysis in Figure 2d quantitatively demonstrates the chemical composition of GLs with typical peaks of the C1s, N1s, and O1s elements located at 285.3, 401.5, and 532.5 eV, respectively (Figures S10 and S11). As shown in Figure 2e, the deconvoluted peaks of the C1s spectrum at 284.6, 285.7, and 286.1 eV correspond to the C−C/CC, C−OH, and C−NH, among which the N originates from the cross-linking reaction between amides on EDA and oxygenic functional groups on GO sheets during the hydrothermal process.4 During the thermal annealing process, the atomic ratios of C1s/O1s increase from 2.5 for GO processor to 36.3 for final GLs (Tables S1 and S2; Figure S12d), indicating that a few carbon atoms are grafted with the residual oxygen-containing functional groups. Moreover, we took the XRD, Raman, and XPS measurements on four stochastically selected spots to characterize the distribution of composition and further evaluate discreteness of these spot-type measurements. Figure 2 and Table S1 show the similar signal patterns and key parameters at different locations within a same GL sample. All the characterized results jointly reveal that the GO precursors are uniformly constructed within the GLs on the multiscale as isotropic features. In summary, the sharp XRD peak, high intensity ratio of ID/IG, and low oxygenic group content jointly demonstrate that oxygen-containing functional groups are efficiently removed during thermal annealing process, leading to high graphitization and welloriented stacking among those graphene sheets within GLs. Nitrogen adsorption−desorption isotherms are shown in Figure 2f, while their corresponding Barret−Joyner−Halenda (BJH) pore size distribution curves are illustrated in the inset. 1099

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Figure 3. Absorption capacitances of GLs. (a) The absorption ratios to different chemical solvents. (b) Optical images of absorption to water, low-viscous organic solvent, and semisolid asphalt, respectively. (c) The cyclic stability of absorption capacitance for alcohol with natural absorption and fire discharge processes. (d) The water absorbing content under humid conditions at 96% relative humidity as an evidence of moisture-proof properties.

The GLs are categorized as type Π according to the international union of pure and applied chemistry classification. The BET surface area of the GLs is 15.8 ± 0.10 m2 g−1, revealing the large specific surface area of these hierarchical porous structures. The BJH pore size distribution curve elucidates that the GL samples mostly contain pores ranging from several to hundreds of nanometers. Absorption Capacitance for Chemical Solvents. The tailored geometric architecture and controllable manipulation of hierarchical microstructures offers 3D printed GLs with high porosity and specifically designed pore orientation. These GLs can be employed for the selective absorption of inorganic or organic solvents. As shown in Figure 3a, the GLs exhibit large absorption capacitance for organic solvents with different viscosities, including water, alcohol, acetone, n-hexane, nbutanol, phorbol-12-myristate-13-acetate, TritonX-100, Tween 80, and asphalt. The corresponding absorption ratios are up to 1.8 ± 0.1, 1824.4 ± 34.5, 1926.5 ± 38.3, 1476.4 ± 29.6, 1949.5 ± 39.0, 1946.1 ± 38.7, 2385.6 ± 47.7, 2470.3 ± 49.4, and 2067.4 ± 41.3%, respectively. Because of the intrinsic poor hydrophilicity of rGO sheet within GLs after thermal annealing at 1000 °C, the GLs demonstrate poor capability of water absorption. However, the organic solvents with low viscosity can be easily drawn into pore structure within the GLs by capillary tension force. In contrast, the high adhesive force of these asphalt-like semisolid solvents impedes them to be absorbed naturally except for the thinning state after heat processing. In particular, this suggests promising usages of GLs in environmental remediation for cleanup and recycling of high-viscous leakages utilizing Jouleheating function compared to those of polymer or chemical gellike sponges (see Figure 3b).27−29 Naturally absorbed and fire-discharged processes of organic solvents were conducted to evaluate the recycling stability of GLs to various adsorbates. For example, the wetted GL samples

were lighted in air after fully absorbing alcohol, as shown in Figure 3c. The fire went out after 10 min with all adsorbates discharged through in situ volatilization and combustion. The GLs then recovered their original dry and free-standing configurations without any structural collapses or fractures. During 10 charge−discharge cycles, the GLs retain a combined performance of robust structure, large absorption capacity, reusability, and high stability with negligible fluctuations in absorption. Figure 4d further assesses the moisture-proof property of GLs at 96% relative humidity. A little amount of absorbed water (109 Ω·m; see Figure S13). Therefore, it can be concluded that the wall

thickness of the HPAs does not affect the electrical conductivity of the final GLs. In respect to the functionalizing applications, decreasing electrical resistance has been verified as one of the most critical routes to maximize voltage performance for Joule heating.19 As a high conductive GO derivation featured with various complex architectures, GLs can serve as smart and Joule heating nanomaterials meeting the engineering heating demands, such as a high-rate heater15 or rapid cleanup material for crude-oil.19 Figure 4b exhibits a highly efficient Joule heating effect of the GLs (12 × 12 × 10 mm3) with temperature elevating over 90% within 100 s. The heating rates are as high as 0.38, 0.53, 0.72, 0.99, 1.35, 1.77, and 1.99 °C s−1, which correspond to heating voltages of 0.5. 0.8, 1.2, 1.5, 2.0, 2.5, and 2.8 V, respectively 1101

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Figure 5. Mechanical properties of GLs. (a) Snapshots of global structural evolution during compression processes. (b) The strain versus stress curve of compression test with maximum strain up to 65%. (c) Truss-structured geometric model for finite element simulation. (d, e) The comparative deformation image and node stress distribution under 4% compression strain, respectively.

Figure 4f shows that before applying heating voltage, the asphalt materials on the GLs sample retained their original semisolid state without drawing into the porous networks. After that, a 1.5 V voltage was applied between the two electrodes of the GLs to heat the asphalt, which was thin via Joule heating. This leads to elevation of the GLs temperature followed by a rapid absorption and cleanup of the fluidic asphalt into the porous scaffolds. Next, the higher voltage can be applied to continuously increase the temperature of the GLs. The absorbed asphalt within the porous frameworks becomes more fluidic with less adhesive force. After the gravity overcomes the constraint of the capillary forces, the absorbed asphalt flows outward from GLs to realize the on-site recycling. Some asphalt, however, remains strongly restrained by the larger capillary forces within those smaller pores of GLs. The combustion-derived deep cleaning strategy was then conducted to facilitate both recovery and reuse of the GLs for high-viscous adsorbates in this remediating application. Mechanical Property Investigations. HPA templatemediated fabrication of GLs via strong π−π interactions derived assembling of GO sheets resolves issues such as the stochastic element orientation, the possible delamination and weak structural interfaces during convectional hydrothermal processes, or direct writing based 3D printing.7,16,20−22,24−26 The as-fabricated GLs therefore demonstrate expected mechanical properties such as high strength and large Young’s modulus. As the snapshots shown in Figure 5a, the truss architecture of GL (ρ = 8.33 mg cm−3) was compressed into a “collapse” status when the compressive strain was increased up to 65%. The corresponding beam and column components in GLs present elastic bending and compression deformation with macroscopic elastic strain of 4%, and the following deformation triggers node torsion and induces sliding displacement. The GL structures then show local yielding and fracture propagation until the entire GL collapses. In contrast to pure GAs fabricated via freeze-drying or supercritical-drying strategies, the compression performance of

(Table 1). Simultaneously, the temperatures increase up to 41.58, 58.94, 78.55, 115.26, 150.52, 189.86, and 218.15 °C within 400 s. Both equilibrium temperatures and heating rates demonstrate approximately linear dependent relationships on heating voltages (Figure S14). In addition, we further validated the stable performance of the Joule heating effect through cyclic electric heating and natural cooling procedures with nearly overlapping data during repeated curves, as shown in Figure 4c. As demonstrated in Figure 4d, we fired the GLs in air to further investigate their fireproof performance. The unchanged morphologies and robust structure during 30 min of combustion indicate that GLs have a good fire-retardant capability, enabling them to serve as fireproofing materials. In addition, the porous microstructure and lightweight density of GLs provide a limited effective area or pathways for thermal conduction, and some semi- or fully closed cellular structures demonstrate the ability to constrain the thermal convection.9 Therefore, as shown in Figure 4e, the temperature-sensitive asphalt pie over the upper surface of the GLs retains its original semisolid status with high viscosity and cannot be absorbed into porous networks even though GLs are heated at the bottom surface by an alcohol lamp with the maximum temperature up to ∼500 °C. This indicates the expected high thermal insulation of GLs that can function for the engineering thermal insulator. Furthermore, as a common environmental pollution, crude oil leakages are conventionally remediated via absorption of polymer-based blankets. These traditional processes are inappropriate for solid or semisolid heavy oil or other organic materials with high viscosity, because the capillary tension forces in porous GLs fail to overcome the gravity and adhesive forces of adsorbates unless they have adequate fluidic features. Joule heating of the graphene derivations can result in a phase change of asphalt-like materials from highly viscous semisolid state to a good fluidic condition, which facilitates in situ absorption and discharging process of adsorbates in a more controllable way. 1102

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Figure 6. Detection properties of GLs used as a chemiresistor. (a) The detection circuit with LED light presenting sensing performance. (b) Initial state without chemical agent added with LED shown bright light. (c) The dim light of LED indicating the change of electrical resistance to chemical agent (e.g., acetone). (d) The schematic illustration of sensing detection of acetone molecular interacted with rGO within GLs. (e) The resistance changes for different chemical agents. (f) The resistance change of GLs under different contents of acetone by volume friction of 20, 50, and 100 vol %, respectively. (g) The cyclic stability of sensing detection to acetone over 10 times.

GLs in Figure 5a exhibits a degenerated elastic characteristic but high-strength features. The corresponding Young’s modulus (∼1.05 MPa) is much higher than that of GAs (∼28 kPa), enabling GLs more appropriate as a bulk structural material with a high fracture resistance (Figure 5b).4,30 This trade-off performance between strength and ductility on the GLs is mainly achieved by the optimized manipulation of structural construction on the multiscale. More specifically, when the compression strain increases from 4% to 50%, the hardened skeleton of the carbonized polymer results in brittle fracture responding to a fluctuating ε−σ curve. In contrast, the GLs framework remains elastic with related graphene sheets in the microscale synchronously generating a serial of waved “wrinkle” morphologies on their surfaces.4,7,31,32 When the strain increases by more than 50%, however, the residual polymer collapses and propagates cracks to GLs under shear forces and transversal expansion deformations (Figure 5a). Although these fabricated truss architectures of GLs have the simple geometric features, the optimized frameworks offer appropriate paths for load transmission. The vertically oriented columns bear the majority of compression loads, while the horizontally arranged beams serve as tensile rods to constrain the transversal expansion (Figure 5c−e). Both of them collaboratively contribute to the high strength, toughness, and robustness of GL structures, enabling their long-term stable services for recycling and multifunctionalizing applications such as absorbers for organic solvents in environmental remediation. Comparatively, several different architectures with varying geometric forms are selected to simulate the effects on mechanical properties utilizing the finite element method, as summarized in Table S3. Under the same normalized volume

density, the structures with inhomogeneous distributed components or complex load transmission routes exhibit the significant changes of compressive strength. This reveals that the different structural formats of GLs are associated with various mechanical properties, which can be specifically tuned with a different 3D printing HPA template-mediated fabrication strategy. Sensing Detection for Organic Solvents. In general, the strong oxidizing reaction of a modified Hummer’s method grafts a large amount of oxygenic groups onto GO sheets, and the thermal reduction at 1000 °C leads to loss of carbon atoms with some vacancies introduced on rGO sheets. Even though most of oxygenic groups are simultaneously eliminated after thermal annealing, the rGO sheets still retain some oxygenic functional groups from the basal plane to the edge of the sheets. This explains why the electrical conductivity of rGO is lower than pristine graphene.4 However, these defective features within rGO offer them specific improved chemical activities when compared with pristine graphene lattices, which enables the GLs to serve as chemiresistors to detect different chemical agents based on stronger electrostatic interactions between target molecules and active defect sites on rGO. The research societies have demonstrated intriguing interest in rGO or other carbon nanomaterial derivations (e.g., carbon nanotubes) as gas sensors.33,34 Here, the differences in binding energies and charge transfer for adsorbates at defect sites (e.g., carboxylic acid, vacancies, hydroxyl, amidogen) are confirmed to dominate these detecting effects.35 For example, Perkins et al. reported that carbon nanotubes offer significant sensing performance to 1103

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ACS Nano Table 2. Detection Sensitivity of GLs to Different Chemical Agents agents

n-hexane

dimethylbenzene

ethanol

acetone

methanol

ethylene glycol

DI water

polarity detection sensitivity (%)

0.06 5 ± 0.4

2.5 8 ± 0.6

4.3 13.6 ± 0.7

5.4 17.5 ± 0.85

6.6 20 ± 1.1

6.9 22 ± 1.2

10.2 0.4 ± 0. 15

sensor for quick identification of methane, ethylene glycol, and ethanol because of interactions between molecular and defective sites on rGO sheets. Combining 3D printing manipulation of architectures with controllable modification of microstructure patterns is an effective way for scalable fabrication and multifunctionalization of advanced graphene monoliths on the multiscale. They have promising applications such as sensors, environmental remediation tools, Joule heating devices, thermal insulation, fireproof materials, moisture prevention tools, and biomedical scaffolds.

a variety of chemical vapors and gases (e.g., acetone, NH3, NO2).36 The 3D printing technology offers fabrication strategies for tailorable and scalable manufacturing functional graphene architectures. Figure 6a,d shows that the truss-structured GLs have a sensitive detecting response to acetone. Before acetone solvent was attached onto GLs, a LED lamp first shines a bright light under a constant voltage (Figure 6b). Then, the light intensity of LED becomes dim as acetone is attracted into the porous GLs and interacts with rGO (Figure 6c). This is because the binding of acetone molecular to rGO sheet affects the charge transfer, which can be easily detected by the significant increase in resistance.35 The mechanism of this chemical stimuli-responsive performance is illustrated in Figure 6d. As shown in Figure 6e, when the amount of absorbed acetone is about 50 vol %, the GLs resistance increases to 17.5 ± 0.85%. In contrast, the detection sensitivities for other organic solvents under the same absorbing contents (50 vol %) are 5 ± 0.4, 8 ± 0.6, 13.6 ± 0.7, 17.5 ± 0.85, 20 ± 1.1, and 22 ± 1.2%, corresponding to n-hexane, dimethylbenzene, ethanol, methanol, and ethylene glycol, respectively. The different stimuli-responsive performance of GLs indicates a monotonously increasing dependence of detection sensitivity on solvent polarity except for water (Table 2), even though it has a large polarity of 10.2. This is because the hydrophobic features of rGO within GLs limit interactions between water molecular and defective sites. As shown in Figure 6f, the sensitivity of GLs chemiresistor is also dominated by the concentration (loading contents) of organic solvents. We selected three different concentrations of acetone (20, 50, and 100 vol %), and their maximum resistance changes are 4.7, 18.2, and 68.3%, correspondingly. This reveals that the signal intensity of the GL response for detecting the same target solvent is dependent on the content of absorption. The original resistance gradually recovered as acetone evaporated as a result of little interaction with rGO within the GLs. In addition, there are overlapping curves for more than 10 cycles, which demonstrates the stable selectivity of GLs to detect various organic solvents as a chemiresistor (Figure 6g). The performance implies that rGO can create GL chemiresistors with tunable sensitivity and identification of chemical agents through 3D printing and chemical modification-based controllable designs on the multiscale.

EXPERIMENTAL SECTION Materials. Natural graphite flakes (average diameter: 50 mesh) were purchased from Nanjing Xianfeng Nanomaterials Tech. Co., LTD (China) to prepare large-area GO precursor using modified Hummer’s approach.4,7 The photocurable polymer was purchased from MakerJuice Laboratories. High-purity silver paint (5 × 103 S cm−1) was purchased from Beijing Emerging Tech. Co., LTD (China). EDA, deionized water, ethanol, acetone, n-hexane, n-butanol, phorbol12-myristate-13-acetate, TritonX-100, Tween 80, ethylene glycol, methanol, dimethylbenzene, and asphalt were all obtained from local suppliers (Lanzhou, China) and used as received. GLs Fabrication. The detailed fabrication of GLs was conducted as the following steps: (1) The HPAs were fabricated from a mask-imageprojection-based stereolithography (MIPSL)-based 3D printing technique (curing light wavelength: 405 nm) with 3D hollow features according to a dedicated computer-aided design (CAD) model (Figures S1−S5).37 (2) GO aqueous suspension (concentration ∼5 mg mL−1) was mixed with EDA agent by volume ratio of 1:3 vol.‰ and ultrasonically dispersed for 30 min. Then, the mixture of GO/ EDA ink was injected into 3D printed HPAs until the chamber of the mold was fully filled under vacuum conditions. The polymer templatemediated GO was assembled into 3D hydrogel structures after hydrothermal reaction under 120 °C for 6 h, followed by freeze-drying process to obtain 3D GO aerogel/polymer composite lattice (see Figures S6 and S7). (3) Finally, the as-fabricated samples were thermally annealed under stepwise heating procedures and argon protection for 24 h (320 °C for 6 h, 400 °C for 6 h, 500 °C for 10 h, and 1000 °C for 2 h) with polymer templates decomposed into large quantities of gases and a small amount of carbonized residuals. The obtained GLs with tailored complex 3D geometric configurations consist of overhang beams and columns elements. Moreover, the individual elements (e.g., beams and columns) were assembled by strong π−π interaction among GO sheets with honeycomb-like porous microstructures (Figure S8). Characterization. Microstructure of GLs was characterized by a SEM (HELIOS NanoLab 600i, FEI, US) and a TEM (Tecnai G2 F30, FEI, US). The thickness of isolated GO sheet was measured by AFM (Dimension Icon, Bruker Corporation, Germany) in contact mode. The structure and chemical compositions of GLs were characterized by XRD, XPS, Raman spectra, and FT-IR spectrum. XRD analysis was carried out by an X-ray diffractometer (X’PERT PRO MPD, PANalytical, Netherlands) using Cu−Kα radiation (1.540598 Å) with a 2θ range of 10−40°. Raman spectra was recorded by a Raman spectrometer (inVia, Renishaw, UK) with Raman shifts in the range of 600−3000 cm−1. XPS investigation was conducted by an X-ray photoelectron spectroscopy (XPS K-AlHPA 1063, UK Thermo Fisher). FT-IR (Nicolet iS10, Thermo Fisher Scientific, US) data were recorded over the range of 500−4000 cm−1. Thermogravimetric analysis (TGA) was carried out with a simultaneous TGA/DSC thermal analyzer (TGA/SDTA851e, Mettler-Toledo, Switzerland)

CONCLUSIONS In summary, we fabricated graphene lattices (GLs) with 3D engineering architectures (consisting of overhang units) via implementing hydrothermal assembly of GO within stereolithography-based hollow polymer templates. The obtained GLs present a 3D geometric configuration with truly 3D tailoring structure consisting of overhang elements at the macroscale, while the microstructure within the construction elements exhibits honeycomb-like hierarchical patterns. The 3D GLs possess a sequence of multifunctional performances including robust structure, large elastic deformation capability, high electrical conductivity, and a superior absorption capacitance. In addition, the GLs can serve as a sensitive 1104

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ACS Nano with a heating rate of 10 °C/min from room temperature (RT) to 1000 °C. Measurements. Electrical conductivities of GLs were measured by a two-probe configuration using a digital multimeter (Victor 8245, Victor, China). The electrodes were painted with highly conductive silver paste to reduce contact resistance. A material test machine (MTS-810, MTS, US) was employed to evaluate mechanical compressibility with loading rate of 1 mm min−1. The temperature distribution was monitored by digital thermometers with K-type thermocouples (Center 309, Taiwan, China). N2 adsorption− desorption measurements were performed by using a Quantachrome instrument (Quabrasorb SI-3MP, US) at 77 K to analyze the specific surface area and pore size distribution.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.7b06095. The detailed design and fabrication processes of HPA template by a mask-image-projection-based stereolithography. The parameter setup of GLs via hydrothermal and followed freeze-drying strategies. The XRD, Raman, and FT-IR results of GO, GLs, and HPAs before and during thermal annealing processes. The XPS spectrum of N1s and O1s on GLs. The peaks and atomic concentrations of XPS results at different testing positions for GLs. Electrical properties of residual carbonized polymer on HPAs. The critical parameters of GL Joule heating performance. The finite element simulation comparison of GL mechanical properties for different structure formats (PDF)

AUTHOR INFORMATION Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. *E-mail: chizhou@buffalo.edu. ORCID

Qiangqiang Zhang: 0000-0002-6082-6782 Author Contributions ∇

These authors contributed equally to this work. C.Z. and F.Z. designed and fabricated the PHA samples. Q.Z. fabricated the GL samples, conducted the characterization of XRD, Raman, XPS, FT-IR, and BET, and implemented the finite element simulation of mechanical properties. D.L. performed measured mechanical, electrical, thermal, and Joule heating properties measurements. X.X. provided characterization of SEM, AFM, and TEM. Q.Z., D.L., and C.Z. collectively wrote the paper. All authors commented on the final manuscript. Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS The authors gratefully appreciate financial support from the Fundamental Research Funds for the Central Universities (PI: Q.Z., Grant: no. lzujbky-2017-k17), the Program of Introducing Talents of Discipline to Universities (111 Project, Grant: no. B14044) and the National Natural Science Foundation of China (PI: Q.Z., Grant: no. 51702142). D.L. and C.Z. thank Kansas State University and University at Buffalo, respectively, for the startup support. We would like to thank LetPub (www. letpub.com) for providing linguistic assistance during the preparation of this manuscript. 1105

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