Ultrathin-graphite foam with high mechanical resilience and electroconductibility fabricated through morphology-controlled solid-state pyrolysis of polyaniline foam

Accepted Manuscript Ultrathin-graphite foam with high mechanical resilience and electroconductibility fabricated through morphology-controlled solid-state pyrolysis of polyaniline foam Hua Xiao, Pu Xie, Shou Ji Qiu, Min Zhi Rong, Ming Qiu Zhang PII:

S0008-6223(18)30659-6

DOI:

10.1016/j.carbon.2018.07.017

Reference:

CARBON 13298

To appear in:

Carbon

Received Date: 3 May 2018 Revised Date:

5 July 2018

Accepted Date: 8 July 2018

Please cite this article as: H. Xiao, P. Xie, S.J. Qiu, M.Z. Rong, M.Q. Zhang, Ultrathin-graphite foam with high mechanical resilience and electroconductibility fabricated through morphology-controlled solid-state pyrolysis of polyaniline foam, Carbon (2018), doi: 10.1016/j.carbon.2018.07.017. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Ultrathin-graphite foam with high mechanical resilience and electroconductibility fabricated

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through morphology-controlled solid-state pyrolysis

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of polyaniline foam

Hua Xiao‡, Pu Xie*,‡, Shou Ji Qiu, Min Zhi Rong* and Ming Qiu Zhang*

Key Laboratory for Polymeric Composite and Functional Materials of Ministry of Education,

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DSAPM Lab, School of Chemistry, Sun Yat-sen (Zhongshan) University, Guangzhou, 510275,

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P. R. China

ABSTRACT: Three-dimensional carbon-based foam (3D-CF) possesses many fascinating properties such as giant conductivity, high thermal insulation, outstanding mechanical performance, and excellent chemical and thermal stabilities. However, most 3D-CFs acquired by the previously reported fabrication methods have some limitations, such as uncontrollable morphology, plenty of inherent defects, difficult to realize self-standing after removing sacrificial template, and overlapping joint flaws among graphene layers. Although significant

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research efforts have been devoted to develop 3D foam, it is still a challenging task to obtain 3DCF with both high conductivity and outstanding mechanical performance, simultaneously. To address this issue, a novel and easy fabricating approach for 3D ultrathin-graphite carbon foam

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(3D-UGF) was explored by a morphology-retaining pyrolysis of polyaniline foam precursors. The typical 3D-UGF graphitized at 2800 °C exhibited stronger compressive strength and electrical conductivity under the compressive strain of 90%, and remained constant even after

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2000 cycles, which are comparable to the previously reported state-of-the-art 3D-CF foam with the same density. Cycling performance of the lithium-sulfur and lithium-air batteries with the

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3D-UGF as binder-free cathode is as high as 1590 mAh g–1 and 920 mAh g–1 after deep discharge/charge 100 cycles, respectively.

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1. Introduction

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Ultralight weight material

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Keywords: Ultrathin-graphite, Polyaniline foam, Flexible graphite carbon, Solid-state pyrolysis,

Three-dimensional carbon-based foams (3D-CFs), such as graphene foam [1–7], amorphous carbon foam [8–12], carbon nanotube foam [13–17], and graphite foam [18–22], possess a unique combination of ultralight weight, open pore structure, high electrical conductivity, thermal and chemical robustness, and good mechanical properties, not available in other foam materials. Therefore, these foams are currently being used in industrial application in various fields including catalysis, filtration, chemical sensors, sorbents, electronics, energy storage, and

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energy conversion. So far, 3D-CF can be fabricated via various methods including chemical vapor deposition [5, 13, 18–21], template-mediated assembly [5, 8, 23, 24], 3D printing [25, 26], and freeze-casting [5–8, 14]. However, most of these fabrication methods offer a few limitations

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owing to the complexity of manufacture process, relatively weak mechanical and electrical performance due to the lapping defects among graphene layers of the 3D-CF, and environmental pollution caused by exfoliation of graphite with strong oxidant. Thus, it still remains a challenge

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to explore a facile manufacturing method which can acquire 3D-CF with controllable

the removal of sacrificial template.

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morphology and through-hole wall thickness, and retain their self-standing properties even after

Graphene, as 2D single sheets of carbon atoms arranged in a hexagonal network, has obtained substantial attention owing to its excellent mechanical, thermal, and electrical properties. The ideal graphene must combine the following advantages: highly ordered structure,

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outstanding surface area (2630 m2 g−1), high Young’s modulus (1 TPa), high thermal conductivity (5000 W m−1 K−1), strong chemical durability, and high electron mobility (2.5 x 105

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cm2 V−1 s−1) [27, 28]. Although, this interest in graphene-inspired materials has also become evident in the foam community, synthesis of graphene-inspired foam tends to be more

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challenging because of hard removing sacrificial template, uncontrollable pore topography, and less/no availability of water-soluble precursor (for e.g., GO aqueous solution) [5–8, 29]. Therefore, in many cases, new strategies are adopted to realize these next-generation foam materials.

Conducting polymers as precursors for carbonization and graphitization have recently attracted significant attention because of their long conjugation chain length, high carbon

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content, and theoretical carbonization yield after the heat treatment [30–32]. Among these conducting polymers, polyaniline (PANI) is the most promising polymer because of its environmental stability, ease of synthesis, cost effectiveness, and controllable morphology [33,

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34]. Improving the mechanical and electrical properties of graphitic foam materials produced using conjugated polymeric precursors is of considerable importance. One of the most promising approaches for enhancing the electrical conductivity and mechanical property of 3D ultrathin-

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graphite carbon foam (3D-UGF) is the use of 3D PANI foam with ultrathin and integrity pore

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wall as presursor.

Herein, a facile and universal fabricating approach was reported for manufacturing flexible graphite carbon foam, composed of ultrathin graphite (less than 20 graphene layers) network, by the morphology-retaining pyrolysis of PANI foam precursor. The as-prepared 3D-UGF is ultralight, superelastic, and superflexible with a density of 5.2 mg cm−3. It also displays high

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elasticity under compressive stress of 0.59 MPa and compressive strain of 90%, and thus maintains its mechanical properties even after 2000 cycles. Notably, in this novel carbonization

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method, poly(methyl methacrylate) (PMMA) foam mainly acts as a microstructure-controllable sacrificial template and PANI enacts as a shape-fixable precursor to prepare graphite foam with

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various morphologies and wall thickness. Compared to classical fabrication methods, this preparation technology can help in achieving 3D-CF with less flaw of lapping among graphene layers. Therefore, the resulting 3D-UGF exhibits excellent mechanical performance and remarkable conductivity.

2. Experimental

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2.1 Materials Aniline and hydrochloric acid (HCl) were purchased from Guangzhou Chemical Reagent Technology Co., Ltd. Ammonium persulfate was obtained from Aladdin. All chemicals

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employed were of analytical grade and used as received without further purification. 2.2 Fabrication of three-dimensional carbon-based foam

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The 3D-UGF was prepared by graphitizing PANI foam. PANI foam with different wall thickness and morphology was obtained according to the previously reported method (see

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supplementary method) [27]. The as-prepared PANI foam was carbonized at 1000 °C for 5 h with a heating rate of 2.5 °C min−1. Then the foam was heated to a high temperature of 2800 °C for 5 h with a heating rate of 10 °C min−1 in a graphite furnace under protection of argon flow, followed by cooling down slowly to room temperature. The foam (3D-T-UGF) was prepared by

2.3. Characterization

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pyrolysis of the PANI foam together with PMMA template.

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The morphologies and microstructures of the 3D-CF were characterized by field-emission scanning electron microscopy (FESEM, S-4800, 10 kV) and transmission electron microscopy

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(TEM, FEI Tecnai G2 F30, 300 kV). The degree of graphitization of 3D-CF was obtained by powder X-ray diffraction (XRD, D-MAX 2200 VPC) with Cu Kα radiation (λ = 1.541 Å). Raman spectroscopy was performed using a Laser Micro-Raman Spectrometer (Renishaw inVia). The chemical states and surface compositions of the as-synthesized sample were analyzed by X-ray photoelectron spectroscopy (XPS, ESCALab250). The compression test of 3D-CF was performed using electronic universal testing machines (SANS, CMT6103). The

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electrical conductivity of 3D-CF was measured on an electrochemical workstation (CHI750E, C16076). The electrochemical performances of the no-binder cathodes (3D-GF/S) with 81.7 wt%

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active material were conducted using the 2032-type coin cell. The electrolyte was a solution of 1 M LiClO4 and 0.1 M LiNO3 in 1,3-dioxolane (DOL) and 1,2-dimethoxymethane (DME) (1:1 V/V). Besides, the 3D-GFs were uniformly adhered to stainless steel mesh to prepare the air

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electrode (cathode), a lithium plate acted as the anode. Both the electrodes were assembled into a home-made Swagelok-type hybrid Li–air battery, and 1 M KOH aqueous solution and organic

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electrolyte (1M LiClO4 in EC/DEC) were used as the electrolyte. They are separated by a water stable ceramic film of lithium superionic conductor. Galvanostatic discharge/charge measurements were conducted using a multi-channel current static instrument Land-CT2001A battery test system (Wuhan, China). All of the specific capacity values were calculated based on

3. Results and discussion

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the mass of cathode active material.

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The previously reported preparation methods usually have some limitations. Accordingly,

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the 3D-UGF was fabricated by the morphology-retaining pyrolysis of PANI foam precursors and the process is shown in Fig. 1. PMMA foam with through-holes structure and controllable morphologies was obtained by supercritical CO2 foaming (Fig. 2), followed by immersion in the

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Fig. 1. Schematic illustration of fabrication process of graphite foam.

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Fig. 2. PMMA foam template with different through-holes structure by extracting PS phase from the foamed PMMA/PS specimen under different immersion time in supercritical CO2: (a) 2 h (tubular structure), (b) 5 h (combined lamellar and tubular structure), and (c) 10 h (lamellar structure).

aniline solution. The PMMA template was pre-treated with aqueous alkali solution to carry negative charges on the surface for the uniform adsorption and dispersion of the monomer

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aniline cations on its surface. With the participation of ammonium persulfate (APS), aniline can be polymerized to form PANI layer at the skeleton surface of the PMMA foam (Fig. S1). Aniline

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monomers are tightly packed and assembled; therefore, they are polymerized in situ to obtain a non-defective polymer layered structure. The adsorption of nucleates at PMMA solid surfaces immersed in the reaction mixture led to PANI nanofilms and coatings of PMMA foam substrates. The resulting foam was then extracted in methylene dichloride to remove PMMA foam template, leading to the formation of the PANI foam (Figs. S2, S3a, S4a, and S5a), which was further carbonized at 1000 °C for 5 h. In this process, PANI was converted into a carbon layer and retained its original structure (Figs. S3b, S4b, and S5b). Then the foam was heated to a

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high temperature of 2800 °C for 5 h in a graphite furnace under protection of argon flow. Based on this, various 3D-UGFs (Figs. 3a, 3c, 3d, and 3e) with superelastic, ultralight, polyporous,

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on different-morphology PMMA templates (Fig. 2).

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high-insensitive strength structure toward stress and electroconductibility were prepared based

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Fig. 3. The morphologies and microstructures of the 3D-CF characterized by SEM and TEM: (a) 3D-UGF with lamellar and tubular cell walls (ρ = 5.2 mg cm−3, d = 5 nm) and its surface structure, and cross section of cell walls in (b), (c) 3D-UGF with lamellar cell walls, (d) 3D-UGF with tubular cell walls, (e) Morphology of 3D-T-UGF (pyrolysis of the PANI foam together with PMMA template), and (f) HRTEM images and electron diffraction pattern of 3D-UGF with lamellar and tubular cell walls.

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The SEM images reveal that the 3D-UGF exhibits three types of microstructures: lamellar and tubular cell walls (Fig. 3a), lamellar cell walls (Fig. 3c), and tubular cell walls (Fig. 3d) owing to the differences of the microstructure of the PMMA templates (Fig. 2). The foam with lamellar-tubular cell walls is constructed by connected ultrathin graphite lamellas and tubes, and cell walls are formed by graphene layers stacking tightly layer by layer (Fig. 3b). HRTEM (Fig. 3f) analysis shows that the cell wall of the 3D-UGF consists of ~20 layers of well-ordered graphene and the d-spacing of the graphene layer is 0.34 nm, indicating that the graphene layers

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are stacked closely via π–π interactions. Moreover, the electron diffraction pattern of the cell wall of foam also revealed high degree of graphitization and few defects with two types of crystal forms. Furthermore, noteworthy, the BET surface area is very small for the 3D-UGF with

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lamellar and tubular cell walls, about 99.52 m2 g−1 (Fig. 4), about one-twentieth of the theoretical surface area of graphene (2600 m2 g−1). This value corresponds to an average number of flexible

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observation of the thickness of aligned walls.

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graphite pore wall form of ∼20 layers graphene, which agrees well with the SEM and TEM

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Fig. 4. Isotherm plot and BJH pore distribution (inset) of 3D-UGF with lamellar and tubular cell walls (ρ = 5.2 mg cm−3 and d = 5 nm).

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Fig. 5. SEM images of the cross section with a thickness of (a) 20 nm and (b) 50 nm of 3D-UGF with lamellar and tubular cell walls; and (c, d) low and high magnification views of the surface structure of the 3D-T-UGF.

By controlling the wall thickness of PANI foam precursor at 20, 100, and 200 nm (Fig. S6),

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respective graphite foams were obtained with different wall thickness, for e.g., 5 nm (Fig. 3b), 20 nm (Fig. 5a), and 50 nm (Fig. 5b), which could significantly influence their mechanical properties and conductivity. The other two foams consisting of either lamellar cell walls (Fig. 3c) or tubular cell walls (Fig. 3d) possess relatively weak mechanical and electrical properties than the one with both lamellar and tubular cell walls. Moreover, the foam (3D-T-UGF) was prepared by pyrolysis of the PANI foam together with PMMA template. Detailed studies by SEM (Figs.

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3e, 5c, and 5d) reveal that the external morphological structure of tubular cell walls of 3D-TUGF displays features such as annular ridge and short lines. This is attributed to the fact that the PMMA template became soft and collapsed during the thermal decomposition process; therefore,

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the self-assembled PANI layer was forced to form wrinkles and ridges.

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Fig. 6. (a) XRD patterns of 3D-UGF and 3D-T-UGF, (b) Raman spectra of 3D-UGF and 3D-T-UGF, (c) XPS spectra of the full-scan of 3D-UGF, and (d) high resolution C 1s of 3D-UGF.

XRD pattern shows that the as-prepared 3D-UGF has a sharp (002) peak at 26.6° which is stronger than that of the 3D-T-UGF, indicating well degree of graphitization (Fig. 6a). In the Raman spectra (Fig. 6b) of the 3D-UGF, the strong G band at 1579 cm−1 is ascribed to the inplane vibration of sp2 carbon atoms. The IG/ID ratio of strong G band (1579 cm−1) and weak D

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band (1353 cm−1) is 35.36, indicating high quality of the 3D-UGF with only a few defects. Moreover, the G band at 2726 cm−1 has a wider half peak width and splitting of the D mode, showing that the foam consists of non-planar graphite [36]. In contrast, for the 3D-T-UGF,

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strong D band and the IG/ID ratio of 2.01, manifest less degree of graphitization and more defects. XPS was also used to characterize the degree of graphitization of 3D-UGF (Figs. 6c and 6d). The C and O contents of the foam are 98.12 and 1.97%, respectively. Furthermore, the C/O mole

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ratio of 54.82 demonstrates that oxygen-containing functional groups (C–OH, C–O–C, C=O, and C(O)–OH) were largely eliminated during the process of the graphitization. The C 1s peak of the

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3D-UGFs can be fitted into four components at 284.8, 285.4, 286.3, and 290.7 eV (Fig. 6d), which are ascribed to sp2 hybridized carbon atom, aromatic C–C bonds, C–O bonds, and π–π* interaction, respectively.

The mechanical resilience and electrical properties of the 3D-CF during the cyclic

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compression processes were characterized by the compression test (Figs. 7, S7, and S8). The 3DUGF with the density of 5.2 mg cm−3 shows excellent mechanical strength with a maximum

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compressive stress of 0.5 MPa at 90% compression strain (Fig. 7a) and remains unchanged even

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after 2000 cycles (Fig. 7b). Moreover, outstanding mechanical resilience of 3D-UGF emerged

Fig. 7. (a, c, e) The stress–strain curves at different compressive strains of 10–90% (a, c, e) maximum strain of 90% for 2000 cycles (b, d) and corresponding electric conductivity-time curve of 3D-UGF with lamellar and tubular cell walls (ρ = 5.2 mg cm−3, d = 5 nm) (a, b), 3D-T-UGF (c, d) and 3D-UGF with lamellar and tubular cell walls (d = 50 nm) (e), (f) typical stress–strain curves at maximum strain of 90% for 2 cycles of 3D-UGF with lamellar and tubular cell walls and a wall thickness of 50 nm.

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from compressive stress–strain curves at different compressive strains from 10 to 90%, which indicate that every next cycle always go through the maximum stress of the previous cycle revealing outstanding resilience (Fig. 7a). Furthermore, the electrical conductivity-time curves of

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3D-UGF in the top left corner of Figs. 7a and 7b also exhibit the eminent restorability which hold the same electrical conductivity of 8 S cm−2 after the compression is released during 2000 cycles. Compared to the 3D-UGF, the 3D-T-UGF (Figs. 7c and 7d) exhibits relatively weaker

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mechanical strength, electrical properties, and resilience owing to the existence of wrinkles and flaws. Wrinkles and ridges at the surface of cell walls tend to generate unrecoverable

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deformation and defects, leading to the reduction of the mechanical strength during the cyclic compression tests. More seriously, the initial electrical conductivity of 3D-T-UGF is 3 S cm−1, which decreases gradually during the compression deformation from 10 to 90%. Besides, the electrical conductivity reduces to 10−3 after 5 cycles, indicating complete breaking of the inner

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construction (Fig. 7d). Moreover, the other two foams possessing lamellar cell walls (Figs. S7a and S7b) or tubular cell walls (Figs. S7c and S7d) display poor mechanical strength and resilience, and their electrical conductivities are lower than that of the 3D-UGF consisting of

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both lamellar and tubular cell walls. Comparatively, the lamellar-cell walls foam composed of a large area of lamellar structure have better self-supporting property than the tubular-cell walls

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foam. Therefore, they can acquire a higher stress of 0.29 MPa under 90% strain, which remains constant after 2000 cycles. The foam is so weak that the contact between foam and electrode is poor at first; however, it becomes tighter after certain cyclic compression. As a result, its electrical conductivity increases a little after 2000 cycles. The tubular-cell walls foam is softer and weaker with relatively low compression strength and electrical conductivity. Furthermore, the 3D-UGF with greater wall thickness of 20 nm and 50 nm became brittle because of the

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probable increase in the defects in the cell walls and decrease in the deformability of cell walls with increasing thickness. Although they have high strength, they would be smashed into pieces during the compression tests. Fig. S8 demonstrates that the 3D-UGF with 20 nm-wall thickness

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is squashed gradually after first cycle, leading to the dramatic decrease of cyclic compressive performance and electrical conductivity after 2000 cycles. The 50 nm-wall thickness 3D-UGF yielded at the strain of 35% due to the wide wall thickness during the first cycle and crushed

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after 90% strain (Figs. 7e and 7f). All the above mentioned results confirm that the wall

thickness of 3D-UGF plays an important role in deciding mechanical resilience and electrical

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conductivity. However, the PANI foam precursor with wall thickness less than 20 nm is not complete and self-standing after removing PMMA template. Therefore, it is difficult to acquire 3D-UGF with wall thickness less than 5 nm.

To demonstrate the example application of the 3D-UGF, cycling performance of the

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lithium-sulfur and lithium-air batteries with the 3D-UGF as binder-free cathode is given in Fig. 8. In the case of lithium-sulfur battery (Fig. 8a), the residual specific capacity is as high as 1590 mAh g–1 after 100 cycles, which is much higher than the highest value under the same

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circumstances reported so far (i.e. 1000 mAh g–1) [37-42], and achieves Coulombic efficiency of

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100%. Clearly, the 3D-UGF ensures good electrical contact with sulfur, thus stabilizing the electrochemical reaction within the cathode region and providing reduced interface contact resistance. As for lithium-air battery (Fig. 8b), the cycling number can reach at least 100 at a controlled capacity of 1000 mAh g–1, characterized after deep discharge. It is almost four times the latest record [43-47]. In addition, both the discharge and charge potentials do not evidently change up to 100 cycles. More detailed analysis of the application of the 3D-UGF in electrochemistry will be discussed in another paper.

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a

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Fig. 8. Cycling performance and Coulombic efficiency of (a) lithium-sulfur and (b) lithium-air batteries with 3D ultrathin graphite framework as binder-free cathode.

4. Conclusions

In this study, the three-dimensional ultrathin-graphite carbon foam (3D-UGF) was obtained by the morphology-controlled solid-state pyrolysis technology with special-designed PANI foam as precursor, and this method is conceptually different from the existing methods. The

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fabrication method provides a versatile platform for exploring distinctive 3D macroscopic curvable graphite foam with controllable microstructure, through-hole wall thickness, stress, and electrical conductivity. Compared to the previously reported state-of-the-art 3D-CFs, including

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graphene foams, our 3D-UGF with a density of ~5 mg cm−3 (Fig. 9), lamellar and tubular cell

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walls, and wall thickness of 5 nm displays the highest mechanical strength of 0.59 MPa and electrical conductivity of 11 S cm−1 when it is compressed to the strain of 90%, as well as exhibits stronger compressive strengths and electrical conductivity under the compressive strain of 90%, and the value remains constant after 2000 cycles. Cycling performance of the lithiumsulfur and lithium-air batteries with the 3D-UGF as binder-free cathode is as high as 1590 mAh g–1 and 920 mAh g–1 after deep discharge/charge 100 cycles, respectively. In addition, the discharge and charge potentials, and Coulombic efficiency do not evidently change up to 100

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cycles. These excellent performances of 3D-UGF make them have a broad range of potential applications in sensors, sorbents, energy storage, and conversion fields. In short, the method of solid-state pyrolysis of PANI foam employed in this study integrates various advantages such as

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safe, morphology-controllable, environmentally friendly, easy to scale up, and no metal ion contaminations, which opens a new route for the preparation and application of 3D carbon-based

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superlight foam.

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Fig. 9. Comparison of density, ultimate stress, and electric conductivity of 3D-UGF with other 3D-CFs under 90% strain.

Acknowledgment The authors greatly acknowledge the financial support from the National Nature Science Foundation of China (Grant No.: 51333008), Scientific and Technological Program of

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Guangdong (Grant No.: 2017A090905001) and Young Teacher Training Program of Sun Yatsen University (Grant No.: 17lgpy86).

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Author Contributions All the authors have equally contributed in the preparation of the manuscript. All authors have given approval to the final version of the manuscript. ‡These authors contributed equally.

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Appendix A. Supplementary data

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Supplementary data related to this article can be found at https://doi.org/

References

[1] A.H. Lu, G.P. Hao, Q. Sun, Design of three-dimensional porous carbon materials: from static

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to dynamic skeletons. Ange. Chem. Int. Ed. 52(2013) 7930-7932. [2] V. Presser, M. Heon, Y. Gogotsi, Carbide-derived carbons-from porous networks to nanotubes and graphene. Adv. Funct. Mater. 21(2011) 810-833.

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[3] Q. Zhang, X. Xu, D. Lin, W. Chen, G. Xiong, Y. Yu, et al., Hyperbolically patterned 3D graphene metamaterial with negative Poisson’s ratio and uperelasticity. Adv. Mater.

AC C

28(2016) 2229-2237.

[4] C. Li, D. Jiang, H. Liang, B. Huo, C. Liu, W. Yang, et al. Superelastic and arbitrary-shaped graphene aerogels with sacrificial skeleton of melamine foam for varied applications. Adv. Funct. Mater. 28(2018) 1704674. [5] V. Chabot, D. Higgins, A. Yu, X. Xiao, Z. Chen, J. Zhang, A review of graphene and graphene oxide sponge: material synthesis and applications to energy and the environment. Energy Environ. Sci. 7(2014) 1564-1596.

18

ACCEPTED MANUSCRIPT

[6] Z. Chen, W. Ren, L. Gao, B. Liu, S. Pei, H.M. Cheng, Three-dimensional flexible and conductive interconnected graphene networks grown by chemical vapour deposition. Nat. Mater. 10(2011) 424.

aerogels. Adv. Mater. 15(2013) 2219-2223.

RI PT

[7] H. Hu, Z. Zhao, W. Wan, Y. Gogotsi, J. Qiu, Ultralight and highly compressible graphene

aerogels. Adv. Mater. 25(2013) 2554-2560.

SC

[8] H. Sun, Z. Xu, C. Gao, Multifunctional, ultra-flyweight, synergistically assembled carbon

[9] C. Chen, E.B. Kennel, A.H. Stiller, P.G. Stansberry, J.W. Zondlo, Carbon foam derived from

M AN U

various precursors. Carbon 44(2006) 1535-1543.

[10] D. Lee, J. Lee, J. Kim, J. Kim, H.B. Na, B. Kim, et al., Simple fabrication of a highly sensitive and fast glucose biosensor using enzymes immobilized in mesocellular carbon foam. Adv. Mater. 17(2005) 2828-2833.

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[11] M. Inagaki, T. Morishita, A. Kuno, T. Kito, M. Hirano, T. Suwa, et al., Carbon foams prepared from polyimide using urethane foam template. Carbon 42(2004) 497-502. [12] S. Chen, G. He, H. Hu, S. Jin, Y. Zhou, Y. He, et al., Elastic carbon foam via direct

EP

carbonization of polymer foam for flexible electrodes and organic chemical absorption. Energy Environ. Sci. 6(2013) 2435-2439.

AC C

[13] M. Mecklenburg, A. Schuchardt, Y.K. Mishra, S. Kaps, R. Adelung, A. Lotnyk, et al., Aerographite: ultra lightweight, flexible nanowall, carbon microtube material with outstanding mechanical performance. Adv. Mater. 24(2012) 3486-3490. [14] S. Nardecchia, M.C. Serrano, M.C. Gutiérrez, M.L. Ferrer, F. del-Monte, Modulating the cytocompatibility of tridimensional carbon nanotube-based scaffolds. J. Mater. Chem. B 1(2013) 3064-3072.

19

ACCEPTED MANUSCRIPT

[15] K.J. Zhang, A. Yadav, K.H. Kim, Y. Oh, M.F. Islam, C. Uher, et al., Thermal and electrical transport in ultralow density single-walled carbon nanotube networks. Adv. Mater. 25(2013) 2926-2931.

RI PT

[16] S. Nardecchia, D. Carriazo, M.L. Ferrer, M.C. Gutiérrez, F. del-Monte, Three dimensional macroporous architectures and aerogels built of carbon nanotubes and/or graphene: synthesis and applications. Chem. Soc. Rev. 42(2013) 794-830.

SC

[17] Z.Y. Wu, C. Li, H.W. Liang, J.F. Chen, S.H. Yu, Ultralight, flexible, and fire‐resistant carbon nanofiber aerogels from bacterial cellulose. Angew. Chem. Int. Ed. 125(2013) 2997-

M AN U

3001.

[18] H. Ji, L. Zhang, M.T. Pettes, H. Li, S. Chen, L. Shi, et al., Ultrathin graphite foam: a threedimensional conductive network for battery electrodes. Nano Lett. 12(2012) 2446-2451. [19] F. Arand, J. Hesser, Quantitative morphological analysis and digital modeling of

TE D

polydisperse anisotropic carbon foam. Carbon 136(2018) 11-20.

[20] I. Kholmanov, J. Kim, E. Ou, R.S. Ruoff, L. Shi, Continuous carbon nanotube–ultrathin graphite hybrid foams for increased thermal conductivity and suppressed subcooling in

EP

composite phase change materials. ACS Nano 9(2015) 11699-11707. [21] M.T. Pettes, H. Ji, R.S. Ruoff, L. Shi, Thermal transport in three-dimensional foam

AC C

architectures of few-layer graphene and ultrathin graphite. Nano Lett. 12(2012) 2959-2964. [22] S. Garlof, M. Mecklenburg, D. Smazna, Y.K. Mishra, R. Adelung, K. Schulte, et al., 3D carbon networks and their polymer composites: fabrication and electromechanical investigations of neat aerographite and aerographite-based PNCs under compressive load. Carbon 111(2017) 103-112.

20

ACCEPTED MANUSCRIPT

[23] H.P. Cong, X.C. Ren, P. Wang, S.H. Yu, Macroscopic multifunctional graphene-based hydrogels and aerogels by a metal ion induced self-assembly process. ACS Nano 6(2012) 2693-2703.

RI PT

[24] W. Lv, Y. Tao, W. Ni, Z. Zhou, F.Y. Su, X.C. Chen, et al., One-pot self-assembly of threedimensional graphene macroassemblies with porous core and layered shell. J. Mate. Chem. 21(2011) 12352-12357.

SC

[25] J. Kuang, Z. Dai, L. Liu, Z. Yang, M. Jin, Z. Zhang Synergistic effects from graphene and carbon nanotubes endow ordered hierarchical structure foams with a combination of

Nanoscale 7(2015) 9252-9260.

M AN U

compressibility, super-elasticity and stability and potential application as pressure sensors.

[26] C. Zhu, T.Y.J. Han, E.B. Duoss, A.M. Golobic, J.D. Kuntz, C.M. Spadaccini, et al., Highly compressible 3D periodic graphene aerogel microlattices. Nat. Commun. 6(2015) 6962.

TE D

[27] K.S. Novoselov, V.I. Fal, L. Colombo, P.R. Gellert, M.G. Schwab, K.A. Kim, Roadmap for Graphene. Nature 490(2012) 192-200.

[28] A. Bianco, H.M. Cheng, T. Enoki, Y. Gogotsi, R.H. Hurt, N. Koratkar, et al. All in the

EP

graphene family–a recommended nomenclature for two-dimensional carbon materials. Carbon 65(2013) 1-6.

AC C

[29] S. Chandrasekaran, P.G. Campbell, T.F. Baumann, M.A. Worsley, Carbon aerogel evolution: Allotrope, graphene-inspired, and 3D-printed aerogels. J. Mater. Res. 32(2017) 4166-4185.

[30] S. Matsushita, K. Akagi, Macroscopically aligned graphite films prepared from iodinedoped stretchable polyacetylene films using morphology-retaining carbonization. J. Am. Chem. Soc. 137(2015) 9077-9087.

21

ACCEPTED MANUSCRIPT

[31] G. Ćirić-Marjanović, I. Pašti, N. Gavrilov, A. Janošević, S. Mentus, Carbonised polyaniline and polypyrrole: towards advanced nitrogen-containing carbon materials. Chemical Papers 67(2013) 781-813.

RI PT

[32] U. Male, P. Srinivasan, Improved electrochemical performances of polyaniline by graphitized mesoporus carbon: Hybrid electrode for supercapacitor. J. Appl. Polym. Sci. 132(2015) 42540.

SC

[33] J. Stejskal, I. Sapurina, M. Trchová, Polyaniline nanostructures and the role of aniline oligomers in their formation. Prog. Polym. Sci. 35(2010) 1420-1481.

M AN U

[34] G. Ćirić-Marjanović, Recent advances in polyaniline research: Polymerization mechanisms, structural aspects, properties and applications. Synth. Met. 177(2013) 1-47. [35] P. Xie, M.Z. Rong, M.Q. Zhang, Moisture battery formed by direct contact of magnesium with foamed polyaniline. Angew. Chem. Int. Ed. 128(2016) 1837-1841.

TE D

[36] P. Tan, S. Dimovski, Y. Gogotsi, Raman scattering of non–planar graphite: arched edges, polyhedral crystals, whiskers and cones. Philos. Trans. R. Soc. London, Ser. A 362(2004) 2289-22310.

EP

[37] Gnan-kumar G., Chuang S. H., Raj-kumar, T., Manthiram A, Three-dimensional graphene−carbon nanotube−Ni hierarchical architecture as a polysulfide trap for

AC C

lithium−sulfur batteries. ACS Appl. Mater. Inter. DOI: 10.1021/acsami.8b06054. [38] Chong W. G., Huang J. Q., Xu Z. L., Qin X. Y., Wang X. Y., Kin J. K., Lithium–sulfur battery cable made from ultralight, flexible graphene/carbon nanotube/sulfur composite fibers. Adv. Function. Mater. 27(2017) 1604815.

22

ACCEPTED MANUSCRIPT

[39] Zhang K., Xie K. Y., Yuan K., Lu W., Hu S., Wei W., et al, Enabling effective polysulfide trapping and high sulfur loading via a pyrrole modified graphene foam host for advanced lithium–sulfur batteries. J. Mater. Chem. A, 5(2017) 7309-7315.

RI PT

[40] Liang Y., Zhang W. Wu D., Ni Q. Q., Zhang M. Q., Interface engineering of carbon-based nanocomposites for advanced electrochemical energy storage. Adv. Mater. Interfaces 2018, 1800430

SC

[41] Sun Z., Zhang J., Yin L., Hu G., Fang R., Cheng H. M., et al, Conductive porous vanadium

Nat. Commun. 8(2017) 14627.

M AN U

nitride/graphene composite as chemical anchor of polysulfides for lithium-sulfur batteries,

[42] Liu B., Xu W., Tao J., Yan P., Zheng J., Engelhard M. H., et al, Enhanced cyclability of lithium–oxygen batteries with electrodes protected by surface films induced via in situ electrochemical process, Adv. Energy Mater. 2018 1702340.

TE D

[43] Lacey S. D., Kirsch D. J., Li Y., Morgenstern J. T., Zarket B. C., Yao Y., et al. Extrusionbased 3D printing of hierarchically porous advanced battery electrodes. Adv. Mater. 30(2018) 1705651.

EP

[44] Lin Y., Moitoso B., Martinez-Martinez C., Walsh E. D., Lacey S. D., Kin J. W., et al, Ultrahigh-capacity lithium−oxygen batteries enabled by drypressed holey graphene air

AC C

cathodes, Nano Lett. 17(2017) 3252-3260. [45] Z. Wang, X. Shen, M. Akbari Garakani, X. Lin, Y. Wu, X. Liu, et al., Graphene aerogel/epoxy composites with exceptional anisotropic structure and properties. ACS Appl. Mater. Inter. 7(2015) 5538-5549.

23

ACCEPTED MANUSCRIPT

[46] X. Xu, H. Li, Q. Zhang, H. Hu, Z. Zhao, J. Li, et al., Self-sensing, ultralight, and conductive 3d graphene/iron oxide aerogel elastomer deformable in a magnetic field. ACS Nano 9(2015) 3969-3977.

RI PT

[47] Y. Wu, N. Yi, L. Huang, T. Zhang, S. Fang, H. Chang, et al., Three-dimensionally bonded spongy graphene material with super compressive elasticity and near-zero Poisson’s ratio. Nat. Commun. 6(2015) 6141.

SC

[48] Q. Zhang, X. Xu, H. Li, G. Xiong, H. Hu, T.S. Fisher, Mechanically robust honeycomb graphene aerogel multifunctional polymer composites. Carbon 93(2015) 659-670.

M AN U

[49] S. Barg, F.M. Perez, N. Ni, P. do Vale Pereira, R.C. Maher, E. Garcia-Tunon, et al. Mesoscale assembly of chemically modified graphene into complex cellular networks. Nat.

AC C

EP

TE D

Commun. 5(2014) 4328.

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