FeCo-Anchored Reduced Graphene Oxide ... - ACS Publications

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FeCo-Anchored Reduced Graphene Oxide Framework-Based Soft Composites Containing Carbon Nanotubes as Highly Efficient Microwave Absorbers with Excellent Heat Dissipation Ability Injamamul Arief,† Sourav Biswas,‡ and Suryasarathi Bose*,† †

Department of Materials Engineering, Indian Institute of Science, Bangalore, India 560012 Department of Chemistry, National Institute of Technology Durgapur 713209, West Bengal, India

‡

S Supporting Information *

ABSTRACT: Conducting polymer composites containing ferromagnetic graftedgraphene derivatives are already appreciated for their lightweight, flexibility, and cost effectiveness in terms of microwave absorption. To further leverage the said properties of this wonder material, we propose a highly efficient replacement by blending conducting multiwall carbon nanotube (MWCNT) and FeCo anchored covalent cross-linked reduced graphene oxide (rGO) with poly(vinylidene fluoride) (PVDF). Interconnected conducting network of MWCNTs introduces higher electrical conductivity in the blend which is essential for microwave absorption. FeCo-anchored porous interconnected rGO framework was designed via solventmediated in situ coreduction in the presence of Fe(II) and Co(II) precursors. Resulting cross-linked-rGO/FeCo displays fascinating coexistence of ferromagnetism and conducting-dielectric behavior, while largely preserving the robust 3D porous interconnected structure. Coupled with conducting MWCNTs, diamine cross-linked rGO/FeCo in a soft polymer matrix yields remarkably high total shielding effectiveness (SET) of −41.2 dB at 12 GHz, for a meager 10 wt % filler content. In addition, the composite materials display efficient heat dissipation abilities in conjunction with the trend in their thermal conductivities. This new-age microwaveabsorbing material, powered by multifunctionality and tunable magnetodielectric properties, henceforth offers an amendable, cost-effective replacement to the existing solutions. KEYWORDS: cross-linked GO-MDA, rGO-MDA-FeCo, 3D porous framework, ferromagnetism, EMI shielding, thermal conductivity

1. INTRODUCTION With the rapid advancement in telecommunication and electronic interfaces, the consequences of hazardous electromagnetic (EM) interference can no longer be ruled out.1,2 The EM pollution has triggered a surge in researches on possible implications and remedies. This includes design and fabrication of novel EM shielding materials and composite systems. These shielding materials find extensive applications in protecting electronic devices, ranging from cell phones to radar systems.3,4 Moreover, the possible repercussions to human health also contribute to the global concerns toward developing novel screening interfaces.5 For a typical EM absorbing material, low density, high conductivity, strong broadband absorption, and excellent thermal stability are the key parameters.6 The mechanism of microwave absorption categorizes the materials into two major sections: dielectric loss materials and magnetically “lossy” materials. Carbonaceous derivatives are the most widely known dielectric materials. One-dimensional carbon nanotubes (CNTs) and nanofibers are extensively used in polymer composites and blend systems owing to their superior properties in various applications.7−11 Since the past decade, MWCNTs and graphene have been at the forefront of materials research, remarkably due to their outstanding conductivity, © 2017 American Chemical Society

thermal, electronic, and mechanical properties. The 2D planar derivatives of graphene, derived from reduction of graphene oxide (GO) either chemically or by thermal exfoliation of GO sheets through sintering, often find use in energy storage,12−14 optoelectronics,15 biomedical, and microwave absorption.16−18 Moreover, the sheets of GO contain numerous oxygencontaining hydrophilic functional groups, making them suitable for additional chemical functionalization and reaction-induced exfoliation, thereby enhancing flexibility for a myriad of applications. Despite the aforementioned advantages for both graphene and its grafted derivatives, the said properties can be greatly compromised as a result of restacking of 2D layers.19 The stacking of monolayers is attributed to the van der Waals interaction or π−π conjugation between the layers. The stacking can be reduced significantly if the 2D sheets can be restructured into a three-dimensional, covalent framework of interconnected GO, utilizing various surface functional groups of planar GO sheets.16,20,21 The high specific surface area for these porous 3D frameworks can also be potentially beneficial Received: March 22, 2017 Accepted: May 18, 2017 Published: May 18, 2017 19202

DOI: 10.1021/acsami.7b04053 ACS Appl. Mater. Interfaces 2017, 9, 19202−19214

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ACS Applied Materials & Interfaces

linked 3D GO networks; second, simultaneous coreduction of GO to rGO, thereby providing giant 3D conducting supernetwork across the composites. While nanoscale particles largely prevent agglomeration and restacking by anchoring into interstitial positions, coreduction of GO ensures conducting framework for incurring substantial dielectric loss. The incorporation of MWCNTs into the composites for forming interconnected conducting network which can act as a mobile charge carrier inside the matrix. With the particular reference to EM shielding application, these novel ferromagnetic 3D-rGO based PVDF composites demonstrated strong microwave absorption, a staggering SET value of −41.2 dB for a meager 10 wt % particle loading in the presence of 3 wt % MWCNTs. Furthermore, coupling high thermal conductivities of the carbonaceous fillers in PVDF matrix with outstanding microwave absorption properties, a new age multifunctional FeCo decorated rGO/CNT/polymer composite with high degree of heat dissipation ability can be designed.

for electromagnetic (EM) wave absorption. However, as stated earlier, dielectric materials alone have proven to be inefficient in broadband shielding. Recent studies have shown that the crosslinked GO with necessary surface functionalization is also largely inefficient in terms of total shielding efficiency.22 Therefore, magnetic inclusion is inevitable to attain substantial shielding effectiveness in GO-based composites.23 A handful of recent progresses incorporating magnetic particles grafted-3D graphene and other carbonic nanocapsules in EMI applications have been reported, most notably by Jian et al.24 and Tang et al.25 Furthermore, Wang et al. reported SET value of −25 dB for 10 wt % rGO/Co3O4-based PVDF composite for the thickness of 4 mm and Pawar et al. recently demonstrated an SET of −34 dB for rGO-Co/MWCNT (10 wt % + 3 wt %) based polycarbonate (PC)−styrene−acrylonitrile (SAN) blended composites, which are listed in Table S4. In order to reproduce significant shielding effectiveness, Singh et al. employed rGOγFe2O3 nanostructures into conducting polyaniline (PANI) matrix with high weight (75 wt %) fraction and showed an SET of −51 dB for 4 mm thickness specimen (Table S4). However, high mass fractions of fillers in these reported systems come up with additional shortfalls, i.e., reduced processability and productivity. Moreover, the reported systems largely overlooked the contribution of interconnected network in the matrix. Another aspect of microwave-absorbing materials often go neglected is the associated heat dissipation mechanism. The exposure to EM radiation produces significant heating effect in shielding devices.26 Therefore, new-age multifunctional EM absorbing materials should be able to address the associated disadvantages it faces in conjunction with the practical applications in microelectronics.26 With the miniaturization of various electronic interfaces, an efficient thermal control mechanism has to be integrated with the devices.27 This not only offers multifunctionality to the EM shielding materials but ensures cost-effective replacements in terms of existing solutions. Herein we have taken into account the combination of magnetic and conducting components of the networks without compromising the cross-linked 3D structure of the amino-GO precursor. The intercalated 3D network of GO was generated in the presence of an industrially familiar epoxy-curing diamine monomer 4,4′-diaminodiphenylmethane (MDA). After isolation, the MDA-cross-linked GO was employed as a precursor for FeCo nanoparticle decoration by means of a facile, scalable hydrazine reduction. The intercalation of primary amineterminated organic moieties into the layered sheets of GO toward the formation of nanoporous graphene oxide frameworks (GOF) has been reported previously for a different application.28 Although the 3D GO-based networks have often been employed in numerous other applications, most prominently in energy storage and catalysis, the effects of nanoporous, cross-linked GO superstructures onto the EM screening capabilities were largely overlooked. Moreover, magnetic incorporation onto the amino cross-linked GObased modular networks without affecting the overall network is a prime challenge in terms of fabrication, structure−property relationship, and EM absorption. Careful control of reaction conditions is imperative to achieve ferromagnetic nanoparticleanchored cross-linked GO without affecting the overall network. This approach elucidates a few problems; first, reproducing ferromagnetic character and subsequently incurring huge magnetic loss without significant restacking of cross-

2. EXPERIMENTAL SECTION Fabrication of MDA-Cross-Linked GO. GO was initially synthesized by the modified Hummer’s method, as reported elsewhere.19 For the fabrication of covalently cross-linked GO-MDA network, 350 mg of GO and diamine monomer MDA (150 mg, Sigma-Aldrich) were dispersed and dissolved in DMF (50 mL) and sonicated for 2 h using probe and bath sonicator. The dispersion was then transferred into a round-bottom flask and refluxed for 24 h in an oil bath at 80 °C. The weight ratio were chosen such that −NH2 groups of MDA would be anchored preferentially to GO epoxy followed by carbonyl groups whether other functional groups would largely be intact. The product was isolated and dried under vacuum for 2 days at room temperature. Syntheses of FeCo-Anchored rGO-MDAs. The decoration of FeCo nanoparticles onto the diamine-GO network was performed in a mixture of GO-MDA dispersion in 1,2-propanediol (20 mL, SigmaAldrich) and Fe(II) acetate heptahydrate (99.9%, TCI Chemicals, Japan) and Co(II) acetate tetrahydrate (98%, Sigma-Aldrich) solution in the same. Two different batches of rGO-MDA-FeCo samples were prepared: rGO-MDA-FeCo1 with lower FeCo weight fraction whereas rGO-MDA-FeCo2 were with higher FeCo particle decoration by using concentrated precursors. Briefly for rGO-MDA-FeCo1 synthesis, 100 mg of the GO-MDA was taken in 20 mL of 1,2-propanediol and dispersed under bath sonication. Separately 50 mg each of Co(II) and Fe(II) salts were dissolved in 10 mL of 1,2-propanediol. The two solutions were then mixed together and further subjected to ultrasonication. The resulting stock was then treated with NaOH to maintain pH of 10. Afterward, 5 mL of 85% hydrazine hydrate (Loba Chemie) solution was added and refluxed for 2 h at 85 °C. The mixture was then transferred to a Teflon-lined autoclave for 3 h. Finally, the black residue of rGO-MDA-FeCo1 was isolated and vacuum-dried at room temperature. For rGO-MDA-FeCo2, 150 mg of total Fe(II) and Co(II) salts (equal mass) was dissolved in 20 mL of solvent and was treated as metal salt precursor solution in an otherwise similar reaction conditions. Fabrication of PVDF−MWCNT Composites. The polymer nanocomposites containing GO-MDA and its FeCo-anchored analogues are fabricated by a typical solvent-mediated mixing-cumcasting approach. The appropriate amounts of GO-MDA and its FeCo derivatives were dispersed in DMF (40 mL) by means of a probe sonicator (Heilscher UP 400S) for 5 min at 50% amplitude. Following the dispersion, it was further subjected to bath sonication for 20 min. For all the nanocomposite fabrication, GO-MDA and its derivative concentrations were fixed at 10 wt %. In a parallel approach, 3 wt % of MWCNTs (Nanocyl SA, Belgium; mean diameter and length are 9.5 nm and 1.5 μm, respectively) was dispersed in DMF (40 mL) following ultrasonication for 20 min. Afterward, two dispersions were mixed together and further sonicated for 1 h. This followed by mixing 19203


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Figure 1. Schematic representation of the 3D rGO-MDA-FeCo network formation from 2D GO sheets. with the commercial PVDF (KYNAR-761, Mw ∼ 440 000 g mol−1, 1 g) solution in DMF (20 mL) was then casted upon a Teflon flatbed tray and slowly heated for removal of DMF solvent. Finally the composite film was vacuum-dried at 80 °C for 24 h to remove the traces of solvents. Characterization. A PerkinElmer GX FTIR instrument was utilized to obtain FTIR spectra. Particle morphology, size, and shape were investigated by a field emission scanning electron microscope (ESEM Quanta, FEI and Carl Zeiss Ultra 55) with an energy dispersive X-ray (EDAX) attachment. High resolution transmission electron microscopy (HRTEM) was recorded using transmission electron microscopy (T20 G2, FEI). The crystal structures and phases of the samples were identified by powder X-ray diffraction using a PANalytical X’Pert PRO diffractometer using monochromatic Cu Kα radiation (λ = 0.1541 nm). For X-ray photoelectron spectroscopy (XPS), we used a Kratos Axis Ultra spectrophotometer. Spectra were taken using Al Kα radiation (1486.6 eV). Atomic force microscope images were acquired by a Park Systems NX10 AFM instrument. A Quantum Design MPMSXL-5 superconducting quantum interference device (SQUID) magnetometer was utilized for magnetization studies. Room temperature electrical conductivity of the composites was performed using an Alpha-N analyzer, Novocontrol (Germany). Thermal conductivity of composites was performed in an LFA 447 Nanoflash (ASTM E1461, NETZSCH, Germany) instrument. The disk-shaped samples were fabricated by compression molding at 220 °C and subsequently polished for conductivity studies. For electromagnetic (EM) shielding interference measurements, 5 mm toroidalshaped samples were prepared by hot compression-molding at 220 °C. The EM shielding measurements were performed by an Anritsu MS4642A vector network analyzer (VNA). Damaskos MT-07 was used as coaxial sample holder and was connected with VNA for measurements. Before any measurement, the setup was calibrated by SOLT (short-open-load-through). The S parameters (S11, S12, S22, and S21) were recorded in X- and Ku-band frequency for all the measurements. The total shielding efficiency and minimum reflection loss data were taken in dB units.

diffusion so as to contain all the GOs during the cross-linking reaction.14 The acidity and oxidation potential of GO are also expected to accelerate the cross-linking reaction.14,29 In order to get the covalently cross-linked GO networks, the excess of MDA was washed with copious amounts of DMF and water repeatedly until the supernatant became colorless. The FTIR spectra of unmodified GO, GO-MDA, GO-MDAFeCo1, and GO-MDA-FeCo2 are shown in Figure 2. The

Figure 2. FTIR spectra display possible cross-linking interaction between epoxy and diamine MDA, whereas in FeCo-anchored systems, the absence of most functional groups indicates a smooth spectra.

absorption peaks of GO correspond to stretching bands of hydroxyl (OH), epoxy (C−O), carboxyl (CO), carboxy (C− O), alkoxy (⟩C−O−), and aromatic (CC) at wavenumbers of 3396, 1080, 1729, 1392, 1224, and 1622 cm −1 , respectively.16,27 In comparison, the hydroxyl content in the GO-MDA reduced drastically. Additionally, the intensity of the carboxyl and epoxy peaks noticeably decreased, and the −COOH peak shifted to lower wavenumber of 1718 cm−1. Furthermore, a new peak at 1550 cm−1 was observed, implying the presence of amine −NH in MDA-cross-linked GO network. The disappearance of −OH peak and existence of −NH confirmed the possible noncovalent interaction of the −NH2 group of MDA monomer with that of −OH and an epoxy opening nucleophilic addition reaction between −NH2 and epoxy groups on the basal plane.30 The disappearance of the major peaks correspond O-containing functional groups in FeCo-decorated rGO-MDA samples confirmed the reduction of GO, while the presence (albeit with reduced intensity) N−H

3. RESULTS AND DISCUSSION The diamine monomer is crucial in forming the covalently cross-linked GO supernetwork.27 The reaction is schematically illustrated in Figure 1. The suspension of GO in water shows evidently anisotropic nature with very high aspect ratio. The GO sheets offer 2D arrays of O-based functional groups (i.e., carboxylic acid group at edges, epoxy and hydroxyl groups at the basal plane) for possible docking of foreign functional groups and elements.33 For the cross-linked GO networks, diamines are preferentially attached to epoxy groups. The MDA possesses two aromatic primary amines with high resonance stability, and therefore it can either undergo epoxide ringopening nucleophilic addition reaction or can attack −COOH groups on the basal plane (Figure S1). Despite the fact that two aromatic groups can produce steric hindrance while crosslinking two GO layers, we used excess of MDA for a faster 19204


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ACS Applied Materials & Interfaces bending peak at 1536 and 1530 cm−1 were accounted for largely intact chemically cross-linked GO-MDA network.30 The X-ray diffraction patterns of the unmodified GO, crosslinked GO-MDA, and FeCo-decorated rGO-MDAs are shown in Figure 3. The effect and extent of cross-linking in the

explained in terms of various O-containing labile functionalities present across the plane and the edges.32 However, increased interlayer spacing of 1.04 nm in GO-MDA corresponding to 2θ = 8.7° indicates a possible cross-linking phenomenon.33 Depending on the initial mass ratio of GO:MDA, cross-linking can follow a parallel or perpendicular interlayer stacking, in addition to loop linking.16,33 Loop cross-linking occurs when cross-linker concentration is sufficiently low. The end-to-end bridging between two GO sheets is also highly unlikely as MDA was not utilized in sufficiently high excess. The probability of bridged GO frameworks is higher when the GO:MDA ratio is kept at 1:2. Furthermore, since the spacing between respective GO layers increased by 0.2 nm as compared to unmodified GO, we can assume that cross-linking takes place by a perpendicular fashion; i.e., MDA units covalently “stitch” the two GO units. The anchoring of FeCo nanoparticles onto the GO-MDA and subsequent reduction of GO into rGO were further extracted from XRD spectra of GO-MDA-FeCo1 and GO-MDA-FeCo2 ferromagnetic cross-linked GO-network. The broad (002) peaks of reduced GO signify the removal of functional groups from GO-MDA, and the appearance of α-FeCo peaks (ICDD reference: 00-049-1568) of (110) and (200) further illustrates the formation of the decorated surface of the GO-MDA crosslinked structure.23 Although the GO-MDA undergoes chemical coreduction of Fe(II) and Co(II) salts and simultaneous reduction of GO, the network structures are largely unaffected throughout the process. This is shown in scanning and transmission electron micrographs of the pristine GO and its cross-linked analogues. The morphological insights into the GO cross-linked network and FeCo-anchored rGO-MDAs are derived from scanning electron microscopy images as shown in Figure 4. It is observed that GO-MDA shows 3D-cross-linked network with

Figure 3. X-ray diffraction patterns of GO and its cross-linked derivatives. The prominent peaks were indexed with standard references. The existence of body-centered cubic (bcc) phases in rGO-MDA-FeCo systems indicates formation of metal alloy.

presence of MDA can be evaluated in terms of d-spacing between the GO planes. The characteristic (001) diffraction peak of GO is spotted at 2θ = 10.8°, corresponding to an interlayer distance of 0.81 nm.31 Contrary to graphene (d002 ∼ 0.35 nm), the increased interplanar spacing in GO can be

Figure 4. Scanning electron micrographs of (A) 2D GO sheet, (B) GO-MDA, (C) rGO-MDA-FeCo1, and (D) rGO-MDA-FeCo2. Corresponding elemental (Fe and Co) mapping of the marked regions for rGO-MDA-FeCo1 (C1, C2) and rGO-MDA-FeCo2 (D1, D2). The color mapping illustrates equimolar distribution of Fe and Co across the interfaces. 19205


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Figure 5. Transmission electron micrograph (TEM) images of (A) GO 2D sheet, (B) GO-MDA, (C) rGO-MDA-FeCo1, and (D) rGO-MDAFeCo2. HRTEM images of the FeCo nanocrystals embedded on the rGO matrix in rGO-MDA-FeCo1 (E) and rGO-MDA-FeCo2 (F), showing epitaxial growth directions.

Figure 6. (A) Survey XPS scan of all the four samples. (B) N 1s of pure GO shows little/no presence of N. (C) N 1s spectra of GO-MDA can be deconvoluted into three major peaks while the peak correspond to −NH accounts for 80.2% of the total population. (D) N 1s of rGO-MDA-FeCo1 shows nearly similar distribution of −NH and imine peaks while free amine peaks are vanished.

random porous distribution. On the contrary, pure GO sheets show highly transparent 2D sheet-like morphology when

attached to the surface of Si substrate, without any sign of stacking. However, the FeCo-anchored surfaces of rGO-MDA19206


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spectra of the samples (Figure 6A). The survey scans of the samples hint at possible cross-linking phenomena, owing to the appearance of the N peak in GO-MDA and FeCo-decorated GO-MDAs.14,30,33 Table S2 shows the elemental compositions of GO and its cross-linked composites. Figure 6B evidences that pristine GO shows almost no or little nitrogen content, accounting for a small peak in XPS scan. However, intensity of nitrogen signal in GO-MDA, rGO-MDA-FeCo1, and rGOMDA-FeCo2 indicates an increase in nitrogen content. The C/ O ratio for the materials follows the increasing order upon cross-linking and coreduction and is best understood by considering reactivity of the O-containing functional groups on GO sheets. The deconvoluted N 1s peak of cross-linked GOMDA shows the existence of −NH (399.6 eV), imine (−N= , 401.5 eV), and primary amine (−NH2, 398 eV) groups with relative abundance of 80.2%, 10.1%, and 9.7%, respectively.35,36 This supports the fact that cross-linking predominantly took place by bridging mechanism, resulting in higher relative abundance of −NH linkages in the system. In the rGO-MDAFeCo1 system, the N 1s spectrum is exclusively populated by −NH (82.7%) and −N= groups (17.3%). Free amines in the cross-linked layers are nonexistent in the rGO-MDA-FeCo1 system. To proceed further with the explanation whether epoxy and/or carbonyl groups on GO surface host the incoming cross-linker and metal atoms, we analyzed the C 1s peaks postdeconvolution (Figure S7). The GO-MDA exhibits five characteristic peaks following Gaussian fitting of the C 1s; namely, C−C, CN, epoxy, CO, and O−CO correspond to binding energies of 284.8, 285.7, 286.6, 288.3, and 289.7 eV, respectively.37 The epoxy peaks in FeCo-anchored systems are nonexistent owing to the reduction of metal ions and GO layers. Furthermore, the presence of various O-functional groups with reduced intensities in GO-MDA indicates the fact that incoming −NH2 groups can proceed via both nucleophilic addition (epoxy ring-opening) and reaction with carbonyl groups at the edge. The aromatic amino groups in MDA are resonance-stabilized; therefore, reduced reactivity coupled with steric interference tends to favor reactivity simultaneously at carbonyl and epoxy sites. For rGO-MDA-FeCo1 and rGOMDA-FeCo2, the relative intensities of C−N/C−O peaks are higher than that in GO-MDA. The merging of oxygenated and deoxygenated peaks in C 1s is more pronounced in rGO-MDAFeCo2, implying the higher extent of reduction in the latter.38 Also, covalent C atoms have no affinity toward incoming groups and metal ions; hence, the C−C/CC/C−H peaks in all the samples remain largely unchanged in terms of relative intensity. The peaks at 778.5 and 793.2 eV (Figure S8) in rGOMDA-FeCo1 are assigned to be of Co 2p3/2 and Co 2p1/2, respectively. Similarly, peaks positioned at 793.1 and 779 eV in rGO-MDA-FeCo2 are associated with the spin−orbit splitting of Co 2p1/2 and Co 2p3/2, respectively.38 Despite the presence of satellite peaks, spin-energy separation between the peaks indicates the zerovalent metallic states of Co.39 Similarly for Fe 2p, the peaks positioned at 724.6 and 710.3 eV correspond to the binding energy of Fe 2p3/2 and Fe 2p1/2, respectively, indicating the presence of both metallic and oxide states in rGO-MDA-FeCo1. The appearance of shoulder peaks at the low binding energy sides of Fe 2p matches with the reference metallic Fe states in both the rGO-based FeCo specimens. The existence of metallic Fe and Co peaks indicates the formation of FeCo alloy.40 The compositions of FeCo in both systems were measured to be Fe45Co55 and Fe54Co46, respectively, for rGOMDA-FeCo1 and rGO-MDA-FeCo2. On the basis of XPS,

FeCo1 and rGO-MDA-FeCo2 showed increased roughness with intercalated network formation. The intercalated network formation followed by diamine cross-linked layers can be clearly observed from SEM images. The FeCo nanoparticles were grafted on the surfaces and gallery gaps of the interconnected network. This results in the reduction of porosity. For rGOMDA-FeCo2, due to the presence of higher metal precursor concentration, the surface appears to be neatly covered with nanoparticles of FeCo. The corresponding elemental mapping analysis of rGO-MDA-FeCo systems (Figure 4) suggests that the elemental distributions of both Fe and Co atoms across the framework are uniform, indicating bimetallic FeCo nanoparticles spreading smoothly across the GO-MDA matrix, with nearly equimolar contribution. The EDAX spectra of the rGO/ FeCo structures further bolstered this claim (Figure S2). The multilayer nature of the interconnected GO-MDA network can further be visualized in high-magnification SEM images (Figure S3). The morphology of interconnected 3D network structures is further investigated by TEM and HRTEM images as shown in Figure 5. Upon decoration of FeCo nanoparticles onto the accessible functional groups, the interactive network does not collapse, rather forming a robust covalent structure with substantial magnetization for magnetodielectric integrated applications. The intercalation of magnetic nanoparticles across the surface and gaps accounts for rigid interconnected framework with adjustable permeability. The HRTEM images (Figure 5E,F) of nanocrystals onto the FeCo-anchored derivatives are shown with measured periodic lattice fringe spacing of 0.20 nm, which corresponds to interplanar spacing between (110) planes of FeCo alloys. In order to study the height profiles and 3D-topographical features of cross-linked GO and rGOs, we performed atomic force microscopy (AFM) studies and are shown in Supporting Information (Figures S4 and S5). The topographical tapping mode AFM image of a 2D GO sheet (Figure S3A) shows a typical thickness around 2 nm, indicating a double layer GO sheet. Conversely, the thickness profile of rGO-MDA-FeCo1 (Figure S4B) indicates an increase in magnitude of thickness of rGO-MDA-FeCo1 network owing to cross-linking and subsequent reduction. The 3D topographical features of GOMDA and rGO-FeCo structures are also investigated using AFM and shown in Figure S4. The typical roughness with highly porous wrinkled morphology is associated with the 3D cross-linked features of the N-conjugated GO frameworks, thus echoing SEM and TEM findings. The porous properties of the cross-linked GO-MDA and FeCo-rGO derivatives were further analyzed by nitrogen adsorption−desorption isotherms measured at 77 K and shown in Figure S6. The isotherms indicate type II adsorption branch with prominent H3 hysteresis loops in the P/P0 range from 0.4 to 1. The BET surface area and total pore volume for GO-MDA were calculated to be 267.45 m2 g−1 and 0.85 cm3 g−1, respectively (Table S1). The effective bridging between GO sheets by cross-linker prevents stacking, accounting for higher surface area.14 On the contrary, both rGO-MDA-FeCo1 and rGO-MDA-FeCo2 have lower BET surface area and pore volume, most likely due to some degrees of restacking following reducing treatment, further corroborating SEM/TEM and AFM observations.34 In order to ascertain the mechanism of reaction between GO and diamine cross-linker MDA followed by ferromagnetic FeCo decoration, we performed XPS (Figure 6). The peaks of C 1s, O 1s, N 1s, Co 2p, and Fe 2p are observed in the XPS survey 19207


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ACS Applied Materials & Interfaces XRD, and FTIR spectra, we can thus propose a mechanism of formation. A schematic representative of the reaction pathway is shown in Figure 1. For the amino cross-linked GO-MDA, it is essentially a combination nucleophilic ring-opening reaction with epoxy and with the carbonyl group either simultaneously or in a particular order. It has been reported previously that aromatic amine cross-linkers are less prone to initiate attack to epoxy groups due to the steric hindrance. For the less electronrich primary aromatic amines, nucleophilic substitution is a viable choice. However, for GO-MDA, we can assume that following the substitution reaction, the amines are more prone to undergo ring-opening addition reaction at epoxy sites. Prior to hydrazine reduction in the presence of metal salts, GO-MDA in 1,2-propanediol offers negatively charged oxygenated surface as a substantial number of O-containing functional groups are left on the GO surface.39,41 The Coulombic attraction between the metal ions (Fe2+, Co2+) and negatively charged GO-MDA in the interstitial gallery gaps and edge positions prompts the metals to anchor readily. Since epoxy O with two sp3 C is more polar than that of hydroxyl O which is connected to a sp2 C, metal ions show greater affinity toward epoxy groups which are sitting at the interstitial positions. The alkali-mediated hydrazine reduction takes place exclusively at these “nucleating” sites, leading to a bimetallic solid solution of disordered FeCo nanoparticles.42 The strong magnetostatic coupling between the nanocrystallites favors the growth and agglomeration of FeCo onto the already reduced GO-MDA. The positively charged metal ions usually neutralize the accumulated negative charges on GO, resulting in formation of heavily stacked mass.42 The advantage of utilizing GO-MDA cross-linked network as a precursor for bimetallic decoration is that they survive the reducing atmosphere while hosting metals at various interstitial positions. Both the FeCo-grafted rGO-MDA samples show strongly ferromagnetic behavior at room temperature with magnetic saturation (Ms) values of 23.0 and 35.5 emu/g for rGO-MDAFeCo1 and rGO-MDA-FeCo2, respectively. The coercivity values are found to be 143.2 and 100 Oe for rGO-MDA-FeCo2 and rGO-MDA-FeCo1, respectively. Magnetic hysteresis curves are shown in Figure 7A. Long-range ferromagnetic ordering is reported for both the hydrazine-reduced FeCo-anchored samples with no typical bifurcations in zero field cooled−field cooled (ZFC-FC) curves below 300 K (Figure 7B).43,44 The bifurcations below RT in ZFC-FC curves are generally associated with superparamagnetic nanomagnets.45 For dc magnetization measurement by superconducting quantum interference device (SQUID) magnetometer at a nominal field (usually lower than 8 kA/m), the blocking region of magnetic nanoparticles is defined as the interval of temperatures where the system starts responding to the magnetic field during temperature ramp up from a lower temperature (as low as 5 K).46 The measurement of blocking temperature (TB) is an integral part of magnetic characterization of nanoparticles systems and is expressed in terms of effective anisotropy density (Keff) and nanoparticle volume (V) as TB(V ) ∼ Keff (V )/kB

Figure 7. (A) Magnetic hysteresis loops observed for rGO-MDAFeCo1 and rGO-MDA-FeCo2 at room temperature. (B) Temperature dependence of field cooled (FC, closed symbol) and zero field cooled (ZFC) curves at a constant field of 50 Oe.

curves below room temperature indicates that both the materials are in a magnetically blocked state, owing to strong interparticle magnetic interaction.44,47 The disordered body-centered cubic (bcc) phase of bimetallic FeCo anchored onto the modular 3D network of GO-MDA could possibly give rise to higher magnetocrystalline anisotropy, accounting for room temperature ferromagnetic (RTFM) behavior.46 The strong dipolar interaction between the nanoparticles could possibly suppress the thermal fluctuations of the spins, thereby leading to higher TB (i.e., above room temperature) and broadening of ZFC curves in both the mentioned ferromagnetic systems. Furthermore, for strong magnetically coupled system a hysteretic behavior is also evident as shown in FeCo-anchored rGO-MDAs at room temperature.46 It is imperative to note that the primary mechanism of the screening is mostly associated with the reflection of incident EM wave from the surface of the shield. Mobile charge carriers or colossal conducting supernetworks play an important role for such characteristics. Later studies showed that it is the interaction of incident EM radiation with the conducting pathway of the shield that plays a pivotal role in EM wave attenuation, not merely the electrical conductivity itself.23 The strategic dispersion and diffusion of various conducting nanomaterials into a polymer matrix generate a highly conducting interconnected architecture. Figure S9 illustrates the trend in specific electrical conductivities of the cross-linked GO and rGO/FeCo network structures. Furthermore, dis-

(1)

where kB is the Boltzmann constant. Because of size effects, Keff is different from bulk materials and depends on the volume and shape of nanoparticles. Most of the nanoparticles display room temperature superparamagnetism below a critical size as thermal fluctuation energy appears higher in magnitude than anisotropy energy. However, no superposition of ZFC and FC 19208


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radiation by dissipating the energies into heat.48−50 Furthermore, higher initial permeability of the shield stipulates stronger EM attenuation. The initial permeability is directly proportional to the square of saturation magnetization

persion of MWCNTs into the PVDF matrix enhances the charge transport property by forming giant interwoven conducting framework inside the insulating polymer matrix. However, diffusion of different cross-linked GO-based nanoarchitectures into the matrix further augments the conducting network formed by MWCNTs, resulting in an inflating order of charge transportation properties for rGO-FeCo/MWCNT grafted systems (Figure 8). The formation of giant conducting

μi = Ms 2 /(aK + bλξ)

(2)

where constants a and b are determined by the material composition, λ is magnetostrictive constant, elastic strain parameter of the crystal is ξ, and K is the anisotropy constant. As discussed earlier, complex permittivity and permeability parameters convey the ability of materials to attenuate EM radiation by absorption. Storage ability is the main source of polarization, and dissipated energy is associated with the loss. Therefore, shielding by absorption is increased with the increasing values of μr and εr where μr = μ′ − jμ″ and εr = ε′ − jε″. As dissipated energy is associated with the loss parameters, the consolidated loss parameters can be evaluated in terms of dielectric tangent loss (tan δε = ε″/ε′) and magnetic tangent loss (tan δμ = μ″/μ′). The energy loss in a material illuminated by electromagnetic waves comes through damping forces acting on polarized atoms and molecules and through the finite conductivity of a material.50 According to the Poynting theorem, the total power of a harmonic electromagnetic field with sufficient angular frequency while entering a volume through the surface encourages partially into increasing the field energy stored inside, and the other part is lost into heat.51 It is also observed that the conductive and dielectric losses are indistinguishable with respect to the heat generated. Applying external electromagnetic field to a composite material implies that the electromagnetic waves come across a variety of microscopic boundary conditions due to the heterostructure inclusions. The resulting local field variations can have a very strong effect on energy absorption at such boundaries because absorption depends quadratically on the electric field intensity. As shown in Figure 9B, the consolidated loss parameter is slightly improved by addition of GO into the composite because trivial change in permittivity is of principal concern. But decoration of FeCo nanoparticles onto GO-MDA network results in the enhancement of both permittivity and permeability parameters. The dielectric permittivity mainly originates due to generation of space charge polarization by different conductive materials, and it is enhanced by crosslinked formation of GO sheets by partial or full reduction of GO. However, magnetic permeability appears following ferromagnetic decoration onto the cross-linked GO network by in situ coreduction process. An enhanced magnetic permeability value appears to be associated with high magnetic hysteresis losses due to natural resonance and eddy current in the gigahertz frequency. It can also be noted that incorporation of dielectric and magnetic loss into the system has generated a massive electric dipole and magnetic interaction at the interface of the materials, further enhancing the virtual charge accumulation. The attenuation constant (α) which is quantitatively associated with the ability of EM absorption can be evaluated by the corresponding permittivity and permeability parameters (Figure S12). Higher magnetic and dielectric losses are indispensable to improved EM attenuation. The formation of cross-linked structure of GO-MDA enhances the permittivity value of the composite structure due to its higher dielectric loss parameters as well as saturation magnetization and magnetic inclusion by FeCo decoration directly enhances magnetic permeability, which has massive

Figure 8. An ac electrical conductivity plot with respect to frequency of various nanocomposites.

supernetwork between rGO-FeCo 3D structures and MWCNTs largely accounts for reinforced composite conductivity. As the percolation threshold of PVDF lies in between 1 and 2 wt % of MWCNTs concentration (Figure S10) so, here to reproduce sufficient electrical conductivity for charge transportation, we have blended a nominal 3 wt % MWCNTs along with the 10 wt % different cross-linked GO structures. The conducting network formation incorporating both rGOarchitecture and MWCNTs is further evident from SEM images (Figure S11). The amount of incident EM attenuation is analyzed through a total shielding effectiveness (SET) in X (8−12 GHz) and Ku band (12−18 GHz) frequency. We observe a frequencydependent shielding efficiency when the blend possesses sufficient electrical conductivity. The incorporation of MWCNTs into the blend greatly enhances the ac electrical conductivity, and the outcome corroborates with that of EM shielding efficiency (SE), while the neat blends are transparent to EM radiation. A significant improvement in SET is observed (Table S3) with increasing content of MWCNTs which ensures the formation of finer mesh structure for facilitating the charge transport property. Additionally, mesh-like structure possesses large specific surface area and therefore favors multiple scattering inside the network which may contribute to higher SE. As expected, a substantial enhancement in magnitude of SET is observed following the addition of rGO/FeCo-based nanostructures into the matrix along with the MWCNTs, echoing the trend in electrical conductivity of composite blends. A remarkably high SET of −41.2 dB is reported for rGO-MDA-FeCo2-based composite at 18 GHz (Figure 9A). The shielding mechanism is illustrated schematically in Figure 9E. The change in local field due to heterogeneous dielectric constant materials influences the absorption of the incident EM 19209


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Figure 9. (A) SET values of the composites as a function of frequency, (B) consolidated tangent loss of various nanocomposites in respect of frequency, (C) total shielding effectiveness with respect to shield thickness, (D) absorption and reflection component of total shielding at 18 GHz, and (E) schematic representation of shielding mechanism for rGO-FeCo/MWCNT composites.

where t is the thickness of the samples in millimeters. It is observed that skin depth is readily decreased from 5.6 mm (when only MWCNTs are present) to 1.2 mm after incorporating cross-linked rGO/FeCo nanostructures along with MWCNTs, reiterating improved EM attenuation property (Figure 9C). The conventional EM screening devices are subjected to overheating owing to prolonged exposure to microwave radiation.40 With the miniaturization of various electronic interfaces, an efficient thermal control mechanism has to be integrated with the devices. Combining excellent electrical and thermal conductivity of the cross-linked rGO/FeCo-based composites containing a small weight fraction of MWCNTs (3%), the designed materials offer new improved heat dissipation ability, largely owing to its high carbonic contents. It has long been known that both graphene and CNTs display excellent thermal conductivity (κ); κ of single layer graphene can be as high as 4000−7000 W/(m K), whereas for CNTs it lies in the range of 2000−6000 W/(m K).54,55 The high thermal conductivity in these nanocarbon derivatives is associated with low thermal conduction by phonons.56 For the polymer-based composites, however, the lower polymer− filler contact area often results in reduced thermal con-

influence in enhancing the attenuation constant value. Therefore, it can be concluded that synergetic effect of both magnetic and dielectric parameters of cross-linked rGO/FeCo based conducting PVDF composites played a pivotal role in exhibiting very high EM attenuation property through absorption (Figure 9D). Reflection loss (RL) of these nanocomposites is also depicted in Figure S13, suggesting a large enhancement after incorporation of cross-linked rGO/FeCo along with MWCNTs. Thickness of the shield material is an important tool in terms of application.52,53 Generally EM attenuation is scaled up with the thickness of the shielding material following a dependence on skin depth value of individual material. The skin depth (δ) depends on the material’s property as described by the expression50 δ=

1 πfμσ

(3)

where μ is the permittivity, σ is the conductivity of the blend, and f is the frequency in hertz. The shielding effectiveness due to absorption (SEA) is also related to the skin depth as

SEA = −8.68(t /δ)

(4) 19210


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Figure 10. (A) Infrared images of the composites (5 mm thickness) at different time interval during cooling following microwave irradiation for 3 s (top row: A−D correspond to GO-MDA; middle row, i.e. E−H correspond to rGO-MDA-FeCo1; the bottom row is for rGO-MDA-FeCo2 based composites. (B) Thermal conductivities of various PVDF composites. (C) Temperature−time profiles of the composites during cooling.

308% jumps in κ values for rGO-MDA-FeCo1 and rGO-MDAFeCo2 based CNT/PVDF composites are interesting for number of reasons. The higher thermal conductivity in reduced GO and graphene networks is attributed to defect-free morphology.63 The higher wt % of rGO may form a percolation network in polymer matrix. A strong synergetic effect between CNTs and reduced GO networks is responsible for significant reduction of phonon scattering at the interfaces.58 With the decoration of FeCo particles onto GO, the defect level in the rGO network decreases, owing to stronger heat dissipation. Furthermore, effective lateral size and thickness of rGO-MDAFeCo systems increase with higher magnetic load. The particle size distribution histograms from DLS measurements for GOMDA and rGO/FeCo suspensions (Figure S14) illustrate the trend of increasing lateral size post reduction and FeCo decoration. This enhances the thermal conductivity of rGOMDA-FeCo2 despite the fact that increased surface roughness (proportional to the magnetic loading in rGO) could encourage phonon transport.58 Therefore, it can be ascertained that the cross-linked ferromagnetic rGO-based composite systems in combination to small MWCNT loading provide new-age EM shielding materials with enhanced shielding effectiveness and excellent thermal properties for multifaceted applications.

ductivity.57 Conversely, the larger surface areas of cross-linked rGOs and MWCNTs offer improved interfaces within the polymer matrix, resulting in minimum thermal resistance in the composites.58 The transient temperature responses of the crosslinked GO-FeCo composites in CNT-PVDF matrix are illustrated in Figure 10A. The thermal images of various composite specimens showing heat dissipation behavior were captured prior to microwave irradiation at a constant frequency of 2.45 GHz for 3 s. The time−temperature response (Figure 10C) of the fabricated composites during cooling was evaluated directly from the infrared (IR) images (taken using an infrared camera FLIR T650SC, Sweden). It is observed that the cooling curves follow exponential decay with time for the composite systems. However, the cooling is remarkably faster for the FeCo-anchored rGO-MDA systems as compared to GO-MDAbased composite. The higher rates of heat transfer for the former are in conjunction with the trend in thermal conductivity values; i.e., the higher the thermal conductivity, the higher is the rate of heat dissipation.58,59 The room temperature thermal conductivity for the specimens is shown in Figure 10B. The thermal conductivity of neat PVDF is about 0.215 W/(m K), which is very close to that reported in the literature.60 For reduced GO, κ lies in the range of 0.14−2.87 W/(m K).61 With the addition of 10 wt % partially reduced GO-MDA, a nominal 13% increase in κ is observed. For PVDFCNT composites, the thermal conductivity is quite higher. This is in good agreement with the recently reported trends.62 However, for 3 wt % CNT-grafted rGO-PVDF systems, the conductivity values increase manifolds. Compared to neat PVDF, the increase in κ for GO-MDA/CNT is nearly 80%. The improved synergistic effect between CNTs and partially reduced GO-MDA may possibly reduce the phonon scattering at the matrix−filler interfaces. However, staggering 220% and

4. CONCLUSION In summary, we demonstrated that highly efficient, 3D rGO/ FeCo framework-based soft conducting composites show multifunctional behavior in terms of remarkably high EM shielding efficiency (−41.2 dB at 18 GHz) and improved thermal conductivity, in addition to excellent processability and cost effectiveness. In order to procreate multimodality, the 2D GO sheets were transformed into robust 3D network by means of diamine (MDA)-mediated covalent cross-linking reaction, 19211


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magnetic Interference Shielding. Adv. Mater. 2013, 25 (9), 1296− 1300. (2) Li, N.; Huang, Y.; Du, F.; He, X.; Lin, X.; Gao, H.; Ma, Y.; Li, F.; Chen, Y.; Eklund, P. C. Electromagnetic Interference (EMI) Shielding of Single-Walled Carbon Nanotube Epoxy Composites. Nano Lett. 2006, 6 (6), 1141−1145. (3) Yousefi, N.; Sun, X.; Lin, X.; Shen, X.; Jia, J.; Zhang, B.; Tang, B.; Chan, M.; Kim, J. K. Highly Aligned Graphene/Polymer Nanocomposites With Excellent Dielectric Properties for High-Performance Electromagnetic Interference Shielding. Adv. Mater. 2014, 26 (31), 5480−5487. (4) Zhang, Y.; Huang, Y.; Zhang, T.; Chang, H.; Xiao, P.; Chen, H.; Huang, Z.; Chen, Y. Broadband and Tunable High-Performance Microwave Absorption of an Ultralight and Highly Compressible Graphene Foam. Adv. Mater. 2015, 27 (12), 2049−2053. (5) Frey, A. H. Headaches from Cellular Telephones: are They Real and What are The Implications? Environ. Health Perspect. 1998, 106 (3), 101. (6) Chung, D. Electromagnetic Interference Shielding Effectiveness of Carbon Materials. Carbon 2001, 39 (2), 279−285. (7) Lu, J. P. Novel Magnetic Properties of Carbon Nanotubes. Phys. Rev. Lett. 1995, 74 (7), 1123. (8) Liu, X.; Yin, X.; Kong, L.; Li, Q.; Liu, Y.; Duan, W.; Zhang, L.; Cheng, L. Fabrication and Electromagnetic Interference Shielding Effectiveness of Carbon Nanotube Reinforced Carbon Fiber/Pyrolytic Carbon Composites. Carbon 2014, 68, 501−510. (9) Kar, G. P.; Biswas, S.; Rohini, R.; Bose, S. Tailoring the Dispersion of Multiwall Carbon Canotubes in Co-Continuous PVDF/ ABS Blends to Design Materials with Enhanced Electromagnetic Interference Shielding. J. Mater. Chem. A 2015, 3 (15), 7974−7985. (10) Yang, Y.; Gupta, M. C.; Dudley, K. L.; Lawrence, R. W. Novel Carbon Nanotube-Polystyrene Foam Composites for Electromagnetic Interference Shielding. Nano Lett. 2005, 5 (11), 2131−2134. (11) Biswas, S.; Kar, G. P.; Bose, S. Microwave Absorbers Designed from PVDF/SAN blends Containing Multiwall Carbon Nanotubes Anchored Cobalt Ferrite via a Pyrene Derivative. J. Mater. Chem. A 2015, 3 (23), 12413−12426. (12) Stoller, M. D.; Park, S.; Zhu, Y.; An, J.; Ruoff, R. S. GrapheneBased Ultracapacitors. Nano Lett. 2008, 8 (10), 3498−3502. (13) Fan, Z.; Yan, J.; Wei, T.; Zhi, L.; Ning, G.; Li, T.; Wei, F. Asymmetric Supercapacitors Based on Graphene/MnO2 and Activated Carbon Nanofiber Electrodes with High Power and Energy Density. Adv. Funct. Mater. 2011, 21 (12), 2366−2375. (14) Zhang, X.; Ciesielski, A.; Richard, F.; Chen, P.; Prasetyanto, E. A.; De Cola, L.; Samorì, P. Modular Graphene-Based 3D Covalent Networks: Functional Architectures for Energy Applications. Small 2016, 12 (8), 1044−1052. (15) Huang, Y.; Dong, X.; Shi, Y.; Li, C. M.; Li, L.-J.; Chen, P. Nanoelectronic Biosensors Based on CVD Grown Graphene. Nanoscale 2010, 2 (8), 1485−1488. (16) Wan, W.; Li, L.; Zhao, Z.; Hu, H.; Hao, X.; Winkler, D. A.; Xi, L.; Hughes, T. C.; Qiu, J. Ultrafast Fabrication of Covalently CrossLinked Multifunctional Graphene Oxide Monoliths. Adv. Funct. Mater. 2014, 24 (31), 4915−4921. (17) Yan, D. X.; Pang, H.; Li, B.; Vajtai, R.; Xu, L.; Ren, P. G.; Wang, J. H.; Li, Z. M. Structured Reduced Graphene Oxide/Polymer Composites for Ultra-Efficient Electromagnetic Interference Shielding. Adv. Funct. Mater. 2015, 25 (4), 559−566. (18) Biswas, S.; Arief, I.; Panja, S. S.; Bose, S. Absorption-Dominated Electromagnetic Wave Suppressor Derived from Ferrite-Doped CrossLinked Graphene Framework and Conducting Carbon. ACS Appl. Mater. Interfaces 2017, 9 (3), 3030−3039. (19) Yang, X.; Zhu, J.; Qiu, L.; Li, D. Bioinspired Effective Prevention of Restacking in Multilayered Graphene Films: Towards The Next Generation of High-Performance Supercapacitors. Adv. Mater. 2011, 23 (25), 2833−2838. (20) Cao, X.; Shi, Y.; Shi, W.; Lu, G.; Huang, X.; Yan, Q.; Zhang, Q.; Zhang, H. Preparation of Novel 3D Graphene Networks for Supercapacitor Applications. Small 2011, 7 (22), 3163−3168.

followed by in situ coreduction in the presence of Fe(II) and Co(II). This scalable, two-step approach ensures intact crosslinked nature of GO-MDA while conferring high permeability owing to decoration of FeCo nanoparticles onto GO edges and gallery gaps. Coupled with 3 wt % CNTs in PVDF matrix, the rGO-MDA/FeCo framework displayed a predominantly absorption-driven EM shielding mechanism. Moreover, improved synergy between the carbon nanofillers and cross-linked network of rGO was attributed to synergetic relationship between highly permeable FeCo and conducting rGO matrix. A strongly correlated rGO/CNT network resulted in high thermal conductivity and improved heat dissipation. Hence, we anticipate that by introducing ferromagnetic rGO-MDA materials in soft conducting matrix, this study will not only bolster the research in multifaceted applications toward costeffective solution to modern electronic interfaces but also promote a quantitative understanding of the structure− property correlation toward the nanofabrication of graphenebased multifunctional materials.

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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/acsami.7b04053. Possible reaction mechanism and pathway, EDAX spectra of the rGO/FeCo networks, high magnification SEM micrographs of various composites, topographical tapping mode AFM images (2D), 3D topological images of cross-linked rGO/FeCo structures, nitrogen absorption−desorption isotherms, C 1s XPS spectra of various cross-linked strictures, Fe 2P and Co 2p spectra from XPS for rGO/FeCo, electrical conductivities of GO and its cross-linked analogues, SEM images of various PVDFbased composites containing GO/rGO, attenuation constants of composites, reflection loss (RL) of various nanocomposites, average particle size distribution from DLS measurements, comparison of total shielding effectiveness of rGO composites, specific surface area and pore volume, XPS elemental analysis of various 3D GO structures, total shielding effectiveness of various composites and comparison of total shielding effectiveness of rGO composites (PDF)

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AUTHOR INFORMATION

Corresponding Author

*Tel +91-80-2293 3407; e-mail [email protected] (S.B.). ORCID

Suryasarathi Bose: 0000-0001-8043-9192 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge the financial support from DST (India). I.A. acknowledges DST SERB-National PostDoctoral Fellowship (N-PDF) program (Grant PDF/2016/ 000048) for financial assistance.

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