Flexible layered reduced graphene oxide/aramid nanofiber composite membrane with high thermal conductivity and mechanical properties

Aramid nanofiber films have shown great potential for high-performance batteries and flexible electronics due to their lightweight and excellent formability. However, aramid nanofibers’ inherent low thermal conductivity severely limits their further applications in high-power electronic devices. Therefore, we attempted to prepare ANF-based composites with excellent thermal conductivity by using reduced graphene oxide (rGO) as a thermally conductive filler for ANF. In this work, we prepared rGO/ANF composite films with horizontal laminar structures by chemical reduction, vacuum-assisted filtration, and hot-press drying. Compared with the pure ANF film (λ, 0.52 W mK−1), the thermal conductivity of the rGO/ANF composite film was greatly improved, and the thermal conductivity increased with increasing rGO content. 50 wt% rGO/ANF composite film achieved an in-plane thermal conductivity of 7.45 W mK−1, 1182.7% higher than the ANF film, and the tensile strength reached 81.7 MPa. Overall, our prepared rGO/ANF composite film exhibits excellent thermal conductivity and good mechanical strength, which helps to explore the potential and development of aramid nanofibres applications in electronic devices.


Introduction
With the rapid development of miniaturization, high density, and high performance of electronic and power devices, there is an urgent need to develop thermal management films with high thermal efficiency and mechanical strength [1][2][3][4][5][6]. Based on the advantages of aramid nanofibres (ANF) in terms of flexibility, electrical insulation, mechanical strength, and high-temperature resistance, its films are widely used in high-performance battery separators [7,8], composite materials [9], flexible electronics [10], and other fields. However, the inherent low thermal efficiency of ANF limits further applications [11,12]. Recently, the synthesis of ANF composite films with high thermal conductivity and mechanical strength by filling the ANF matrix with highly thermally conductive fillers has been widely investigated [13]. Currently, the commonly used thermally conductive fillers include metals (such as Al [14], Cu [14], Zn [15], etc), inorganic nonmetallic fillers (such as graphene [16,17], carbon nanotubes [18,19], boron nitride (BN) [20], silicon carbide (SiC) [21], etc). Among them, the two-dimensional nanomaterials graphene and boron nitride are more widely used in thermal conductive materials.
The internal arrangement structure and interfacial bond strength of composites are reported to be the two most important factors affecting their thermal conductivity [22]. Firstly, an effective alignment structure facilitates the construction of thermally conductive pathways. Secondly, the stronger the interfacial bond, the less phonon scattering at the interface, and the denser the composite. Two-dimensional nanomaterials have good thermal conductivity due to their anisotropic crystal structure. The functionalized 2D nanomaterials can be bonded more tightly to the polymer-based interface, significantly improving heat transfer efficiency. For example, Ma et al [23] prepared a BNNS@PDA/ANF thermally conductive composite film using ANF and Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. polydopamine functionalized boron nitride nanosheets (BNNS@PDA). When the loading of BNNS@PDA was only 50 wt%, the composite film's in-plane and out-plane thermal conductivity reached 0.62 W mK −1 and 3.54 W mK −1 , respectively, respectively, an improvement of 181.8% and 196.2% over the pure ANF film. The in-plane and out-of-plane thermal conductivities of the composite films reached 0.62 W mK −1 and 3. 54 W mK −1 , respectively, at a loading of only 50 wt% BNNS, an improvement of 181.8% and 196.2% over pure ANF films. Lin et al [24] obtained the composite films' thermal conductivity by combining solution mixing, and An ANF/boron nitride composite film was obtained by blending solution mixing and vacuum filtration. The composite film prepared by adding 50 wt% BNNS had good out-of-plane thermal conductivity (0.6156 W mK −1 ) and mechanical strength (62 MPa). For a given BN loading, BN after surface functionalization or hydroxylation can further improve the thermal conductivity of the composite film [25]. However, the chemically inert surface of BN makes it challenging to form a thermally conductive pathway by uniform dispersion in the ANF matrix. It requires complex modification and processing before application. In addition, the oxygen-containing groups in BN increase the possibility of phonon scattering, and the thermal conductivity of ANF-based composites is weaker along the horizontal plane [26]. Therefore, BN has limitations when used to improve the thermal conductivity of ANF matrices.
Graphene is an ideal polymer-based thermally conductive filler due to its high thermal conductivity of 2000-5300 W mK −1 [27]. There are many reports about graphene improving the thermal conductivity of composites. However, poor dispersion and the low cost of pristine graphene have limited its application. Typically, surface modification methods are employed to improve graphene dispersion, but the improvement is not effective [28,29]. In contrast, graphene oxide (GO) has large oxygen-containing functional groups, high specific surface area, and good dispersibility, which can be uniformly dispersed in a polymer matrix and form a solid interfacial bond with the matrix. In addition, the rGO obtained by reducing GO as a graphene precursor has excellent thermal conductivity [30]. For example, Yang et al [31] prepared a graphene film with excellent thermal conductivity (1102.62 W mK −1 ) by wet spinning and reduction with hydriodic acid. Jin et al [22] reported the preparation of ANF/rGO thermally conductive composite films by vacuum-assisted filtration of ANF. They went by hot pressing and reduction with hydriodic acid, and when the rGO content was 40 wt%, the in-plane thermal conductivity of the composite film reached 0.4172 W mK −1 , an increase of more than 1250% over the pure ANF film. Inspired by the research on chemically reduced graphene oxide thermally conductive materials and ANF composites, the main objective of this work was to obtain thermally conductive composite films by constructing thermally conductive networks in the ANF film structure.
In this work, ANF was prepared by chemical dissociation of macroscopic aramid fibers (Kevlar 29 ® ) by deprotonation, following the work of Lv [32]. rGO was obtained by heating and reducing GO with the green reducing agent L-ascorbic acid. The suspension of ANF and rGO was then filtered by high-efficiency vacuumassisted filtration and hot-press dried to produce the rGO/ANF composite membrane. The obtained ANF was characterized by scanning electron microscopy (SEM), transmission electron microscopy (TEM), Fourier transform infrared spectroscopy (FTIR), and x-ray diffraction (XRD). rGO and rGO were compared and analyzed by x-ray diffraction (XRD), Fourier transform infrared spectroscopy (FTIR), and x-ray photoelectron spectroscopy (XPS) tests. The rGO/ANF composite films were characterized and performance tested using Raman spectroscopy, x-ray diffraction (XRD), scanning electron microscopy (SEM), thermal conductivity meter, universal testing machine, and infrared thermal imaging camera. Finally, the mechanical properties and thermal conductivity of the rGO/ANF composite films are discussed and analyzed in detail. Monolayer graphene oxide powder (99% purity, flake diameter 0.2-10 μm, thickness approximately 1 nm) was obtained from Source Leaf Biotechnology Co. Acetone, potassium hydroxide (KOH), dimethylsulfoxide (DMSO), and L-ascorbic acid (L-AA) were purchased from McLean Biochemical Technology (Shanghai, China). All of the above reagents were analytically pure and used without further purification.

Preparation of ANF dispersion
Cut Kevlar fibers were ultrasonically washed with acetone for 60 min, rinsed with deionized water, and vacuum dried for reserve use. Added 2 g Kevlar and 2 g KOH to 500 ml DMSO, sealed the resulting mixture in a roundbottomed stoppered flask, and stirred it in a water bath at 50 ℃ for one week to obtain a stable dark reddish brown ANF/DMSO solution. 3500 ml DMSO was added and diluted the ANF/DMSO solution into 0.5 mg ml −1 . Then 60 ml of 0.5 mg/ml ANF solution was measured, and 150 ml of deionized water was added to promote the protonation of the amide group of ANF. After ultrasonic treatment for 3 h, it is poured into the vacuum suction funnel, and then deionized water is added to wash it until DMSO is passed through to obtain the dispersion of 30 mg ANF in water.

Preparation of GO dispersion and L-AA aqueous solution
0.5 g GO was added into 500 ml of deionized water, homogenize in an ice bath for 30 min, and obtain a stable GO dispersion with a 1 mg ml −1 concentration. Add 1 g of L-AA to 100 ml of deionized water to give 1 wt% L-AA aqueous solution and seal for storage.

Preparation of rGO/ANF composite films
Firstly, quantitative GO dispersion (GO content is 25%, 50%, 75%, and 100% of the mass of ANF) and 1 wt% L-AA solution (the mass ratio of GO and L-AA is 1:10) were added to 30 mg of ANF dispersion and treated with a homogenizer for 15 min Secondly, the hybrid system was reduced in an oil bath at 95 ℃ for 2 h. The hybrid system was then poured into a vacuum funnel for filtration and washed several times with deionized water. Finally, the obtained wet film is sandwiched between two pieces of glass, dried at 50 ℃ for 48 h, and then prepared the rGO/ANF composite film. A series of rGO-x/ANF composite films were designed by controlling the rGO content (25, 50, 75, and 100 wt%). The fabrication diagram of the rGO/ANF composite film is shown in figure 1.

Characterizations
The morphology and microstructure of the samples were analyzed by scanning electron microscopy (SEM, Hitachi Regulus 8100, Hitachi, Japan). The films were soaked in liquid nitrogen for several hours before brittle fracture. Before the test, fix the breakable section on the conductive adhesive and conduct gold-plating treatment. The transmission electron microscope (TEM) patterns were collected on Tecnai G2 F20/TEM microscope (FEI Company of the United States). The sample needed to be dried on the copper mesh for testing. The samples' Fourier transforms infrared (FTIR) spectra were obtained using Bruker Vertex 70 (Thermo Flight Company, USA) instrument. The x-ray diffraction (XRD) pattern of the sample was tested by Ultima IV (Rigaku Company, Japan) device. Before the test, the film was cut into a square with a side length of 2 cm. The sample's x-ray photoelectron spectroscopy (XPS) analysis was conducted on an Axis Ultra DLD Kratos AXIS SUPRA instrument (Shimadzu, Japan), and the excitation source was Al Kα radiation. The thermal conductivity value of the sample was tested by Swedish Hot Disk TPS2500S equipment (Swedish AB company), and each sample was tested three times. According to the ISO1184-1983 standard, using the (CMT6104, China) universal testing machine to conduct tensile tests on rectangular samples (3 cm long and 1 cm in width) at room temperature. The testing machine had a 1000 N load cell with a 1 mm min −1 speed. The infrared thermal image of the sample was recorded with an infrared thermal imager (FLIR T540, Philips, USA).   2(d)). The diameters of the individual fibers were counted to be in the range of 5 to 30 nm, with the sizes mainly concentrated in the field of 10 to 20 nm (figure 2(e)), which is consistent with previous reports [33]. Under alkaline conditions, the intermolecular hydrogen bonds of the Kevlar fibers were dissociated, but most of the fibrous structure was retained to form aramid nanofibres [34]. respectively. Due to many hydrogen bonds and regularly arranged molecular chains, Kevlar fibers have a high degree of orientation [23]. In contrast, ANF has only a broad and weak diffraction peak at 20.2°due to a decrease in crystallinity due to the disruption of the original crystal structure after hydrogen bonding dissociation in Kevlar [35,36].
As shown in figure 2(g), the peaks at 3313 cm −1 and 1641 cm −1 are attributed to the stretching vibrations of N-H and C=O, respectively. The coupling of the N-H bond's deformation vibration and the C-N bond's stretching vibration may form peaks at 1538 cm −1 . The peak near 1305 cm −1 is a characteristic absorption peak of Ph-N, consistent with previous reports on ANF [37]. The combined analysis of the TEM, XRD, and FTIR test results demonstrated the successful preparation of ANF.

Characterization of GO and rGO
The mechanism of reduction of GO by L-AA is not well understood. According to the literature [38,39], the mechanism mainly involves two different reactions, including the decline of the epoxide group and the reduction of the collar hydroxyl group. L-AA and GO are reduced to rGO and dehydroascorbic acid by heating at 95°C, and the reaction process is shown in figure 3. The hydroxyl group on the five-membered ring of L-AA is acidic, which can dissociate and provide two protons, and the dissociated structure is a nucleophilic reagent. When L-AA reacts with GO, the proton reacts with the hydroxyl and epoxide groups in an acid-base reaction, followed by the SN2 reaction. The specific process of SN2 is that the nucleophilic reagent attacks the Sp3 carbon attached to the hydroxyl group or the Sp2 carbon of the epoxide group, forming an intermediate and producing the by-product water or hydroxyl group. Finally, the mediators are heated to undergo a redox reaction to create the end products rGO and dehydroascorbic acid.
The states of the GO dispersion and the rGO solution after 24 h are shown in figure 4(a), where the GO dispersion is hydrophilic, stable, and homogeneous. On the contrary, the rGO is superhydrophobic and settles to the bottom. GO is hydrophilic because it is rich in carboxyl and hydroxyl groups. In contrast, after reduction with L-AA, GO lost many hydrophilic groups and sank hydrophobically. Figure 4(b) shows the XRD spectra of GO and rGO samples. A typical sharp diffraction peak appears at 10.0°for GO with a layer spacing of 0.884 nm, indicating the presence of oxygen-containing functional groups and defects in GO [40]. After chemical reduction, the characteristic diffraction peak of GO disappears, and a new broad diffraction peak of rGO appears at 24.6°, corresponding to a typical (002) graphene crystal plane with a layer spacing of 0.362 nm. The results showed that L-AA successfully reduced some oxygen-containing groups in GO. In addition, the rGO formed by the reduction of GO has an increased sp 2 carbon content, reduced interlayer spacing, and an increased degree of crystal structure order.
As shown in the FT-IR spectrum ( figure 4(c)), GO has a broad band between 3700 cm −1 and 2500 cm −1 . 1732 cm −1 shows the C=O stretching vibration peak. 3405 cm −1 and 1409 cm −1 show the O-H stretching and bending peaks, respectively. 1223 cm −1 and 1053 cm −1 show the C-O stretching vibration peak. Peaks at 1223 cm −1 and 1053 cm −1 are characteristic peaks of the oxygen-containing functional groups of GO, thus demonstrating the hydrophilicity of GO ( figure 4(a)). In addition, the intensity of the characteristic peak of the oxygen-containing functional group in rGO is significantly reduced, indicating that some of the oxygencontaining groups have been successfully reduced. Therefore, the number of oxygen-containing groups in rGO is low, demonstrating the hydrophobicity of rGO in figure 4(a).
To further demonstrate the successful reduction of GO by L-AA and the formation of rGO, XPS characterization was performed (figures 4(d)-(f)). The intensity of the O 1 s peak of rGO was significantly weaker than that of GO, indicating that the C/O of rGO was substantially higher than that of GO, indicating that L-AA successfully removed some of the oxygen-containing groups from GO. In addition, (figure 4(e)) the C1s peaks of GO were fitted at 284.8, 286.8, 288.0, and 288.7 eV, corresponding to C=C/C-C, C-O, C=O, and O-C=O bonds, respectively [41,42]. By comparing the C-O peaks in figure 4(f), the peak height in rGO is significantly reduced, indicating that the number of oxygen-containing groups in the formed rGO is diminished, and the content of sp2 C is increased. In summary, L-AA effectively reduced GO and formed rGO. Figure 5(a) shows the XRD spectra of GO, rGO, pure ANF, and composite films. rGO, rGO, and pure ANF films have characteristic peaks at 10.0°, 24.6°, and 20.9°corresponding to crystal planes (001), (002), and (110), respectively. In addition, the characteristic peaks on the (002) crystal face become more pronounced with increasing rGO content ( figure 5(c)), indicating that the components of GO and rGO can be maintained after the film is prepared with ANF.

Characterization of GO/ANF and rGO/ANF
The Raman spectra of GO, rGO, GO/ANF composite, and rGO/ANF composite are shown in figure 5(b). The two characteristic peaks of GO and rGO are located at 1352 cm −1 (D peak) and 1589 cm −1 (G peak), respectively. The intensity ratios of D and G peaks (I D /I G ) can be used to characterize the defects or crystalline structure of grapheme [43]. The D peak of GO is significantly higher than the G peak with I D /I G = 1.79, indicating the presence of defects and a high sp 3 carbon content in GO. After L-AA reduction, the I D /I G of rGO was increased to 1.96, meaning that most of the oxygen-containing groups in GO underwent reduction reactions and the sp 2 structural domain was expanded [44], which is consistent with the XRD, FTIR, and XPS test results in figure 3. In addition, the 50 wt% GO/ANF and 50 wt% rGO/ANF composite films showed similar Raman signals with I D /I G of 1.73 and 2.13 for both, respectively. The I D /I G of the rGO/ANF composite film became significantly larger after reduction, indicating that the GO in the composite film is effectively reduced to rGO. Figures 5(d)-(f) show the tensile stress-strain curves, tensile strength, and elongation at break for rGO/ANF composite films with rGO content of 25 wt%, 50 wt%, 75 wt%, and 100 wt%, respectively. The tensile strength and elongation at the break of ANF films were 114.9 MPa and 22.3%, respectively. The tensile properties of ANF are superior due to the combined effect of hydrogen bonding, covalent bonding, and van der Waals forces in the internal structure [45]. As the rGO content increases, the tensile strength of the rGO/ANF composite film decreases, and the brittleness gradually increases. Compared to ANF, rGO is brittle, and rGO induces slippage in ANF, so the tensile properties of the rGO/ANF composite film decrease [46]. At 25 wt% rGO, the tensile strength and elongation at break of the rGO/ANF composite film were 95.2 MPa and 17.4%, respectively, with tensile strength retention of 82.8%. When the rGO content was increased to 50 wt%, the tensile strength of the rGO/ANF composite film was 81.7 MPa and elongation at break was 14.7% while still maintaining a certain level of mechanical strength and flexibility. Due to the proximity of the ANF and rGO fillers, the stress is effectively transferred [47,48]. Figures 6(a)-(c) show the optical photographs of ANF film, GO/ANF composite film, and rGO/ANF composite film before drying. The surfaces of the three samples are relatively complete, with no apparent defects. The diameter of the sample is about 5 cm, and the manufacturing process is simple. After vacuum drying, the surfaces of the films are smooth and free of cracks, and the compactness is also improved. The ANF film is similar to the yellow of the macrofibre, and the two composite films are opaque black. In addition, figure 6(d) shows the ship-shaped origami of the rGO/ANF composite film, indicating its high flexibility. As shown in figure 6(f), the cutting of rGO/ANF composite films into Chinese window patterns shows that they do not need molds to reduce the production cost of complex composite films. These results indicate that the rGO/ANF composite film has good mechanical strength and high flexibility (figures 6(d)-(f)).
In this work, commercially available GO was used as a precursor, and rGO, formed by the reduction of GO with L-AA, was used as a thermally conductive filler to prepare rGO/ANF composite films by vacuum-assisted filtration. To uniformly disperse the rGO filler in the ANF matrix, the GO and ANF dispersions were also subjected to sonication before chemical reduction, which was essential to achieve the desired layered structure. To analyze the relationship between microstructure and thermal conductivity of the composite films, the pure ANF films and four rGO/ANF composite films with different rGO contents were tested by SEM. As shown in figure 7, all membranes had a laminar structure in cross-section due to the self-assembly of the ANF and rGO/ ANF suspensions during the vacuum-assisted filtration process [49]. Pure ANF has a fibrous, filamentary structure in cross-section (figures 7(a)-a').
In contrast, the rGO/ANF composite membrane fibers exhibited a directionally stratified morphology (figures 7(b)-(e), (b)'-(e)'), indicating that the rGO was uniformly dispersed in the ANF matrix. In addition, the tightness of the rGO/ANF composite film increased with increasing rGO content, and the results indicated that ANF and rGO were effectively combined under vacuum-assisted filtration and hot pressing. In particular, the oriented lamellar structure of ANF and rGO contributes to creating an effective in-plane thermal conductivity path. Figure 8 shows the thermal conductivity of the ANF film and the rGO/ANF composite film with different rGO contents. The composite film exhibits superior thermal conductivity compared to the pure ANF film (λ = 0.52 W mK −1 ). The thermal conductivity of the rGO/ANF composite film increases sharply with increasing rGO content. The thermal conductivity of the rGO/ANF composite film reaches a maximum value of 12.31 W mK −1 when the rGO content is 100 wt%. Because the reduction of defects in the rGO formed by the removal of GO by L-AA effectively reduces phonons' scattering capacity. In particular, the thermal conductivity of the composite film increases to 7.45 W mK −1 at a mass fraction of 50 wt% rGO, an increase of 1182.7% over the original ANF film, when the rate of growth in thermal conductivity is most pronounced. This is because the directional layer structure formed in the rGO/ANF composite film provides an efficient conduction path for phonons. Therefore, there is an inextricable link between the oriented layer structure of the rGO/ANF composite film and the thermal conductivity.

rGO/ANF thermally conductive composite films
Thermal tests were performed on the films to verify further the thermal conductivity of the rGO/ANF composite film. The ANF film, 50 wt% GO/ANF, and 50 wt% rGO/ANF composite films were placed on a heating device. The variation of the sample surface temperature with time was recorded using an infrared thermal imaging camera. The results are shown in figures 9(a)-(b). The surface temperatures of the GO/ANF and rGO/ANF composite films were significantly lower than those of the ANF film over the same period. The   surface temperatures of the GO/ANF and rGO/ANF composite films were considerably lower than those of the ANF film over the same period, which means that the composite film has a better thermal diffusion performance. The 50 wt% rGO/ANF composite film has an average thermal diffusion rate of 8.3°C/s from 0 to 6 s, while the ANF film and 50 wt% GO/ANF are only 4.6°C s −1 and 5.6°C s −1 respectively. The rGO/ANF composite film has a higher thermal transfer efficiency than the pure ANF and GO/ANF composite films and can quickly transfer heat from the heat source to the external environment. In addition, the 50 wt% rGO/ANF composite film has the lowest surface temperature during heat dissipation, down to 48.5°C, which is 16.3°C lower than the lowest temperature of the pure ANF film. This is attributed to the directional layered structure formed by the rGO and ANF matrix, which creates an effective in-plane heat transfer path and allows rapid inplane heat dissipation, consistent with the data in figure 8.

Conclusions
In summary, ANF was successfully prepared by deprotonation and rGO by chemical reduction of GO using L-AA. rGO/ANF composite membranes with high thermal conductivity and excellent mechanical properties were successfully prepared by combining these two methods to prepare ANF and rGO suspensions, followed by a vacuum-assisted filtration-induced self-assembly process. The oriented lamellar structure formed by rGO and ANF facilitates the fabrication of thermally conductive pathways, enabling the rGO/ANF composite film to exhibit excellent thermal conductivity. At 50 wt% rGO, the in-plane thermal conductivity of the rGO/ANF composite film reaches 7.45 W mK −1 , 1182.7% higher than that of the ANF film. The thermal conductivity of the rGO/ANF composite film was verified by infrared thermography to be better than that of the GO/ANF composite film and the pure ANF film. The tensile properties of the 50 wt% rGO/ANF composite film were shown to be better, with a tensile strength of 81.7 MPa. Besides, the high thermal conductivity and good mechanical strength of the rGO/ ANF composite film were demonstrated. The excellent performance of the 50 wt% rGO/ANF composite film makes it a potential composite material for the thermal management of electrical and electronic devices.

Data availability statement
All data that support the findings of this study are included within the article (and any supplementary files).