This extraordinary enhancement in thermal conductivity is shown to be due to a network of continuous graphene sheets over long−length scales, resulting in low thermal contact resistance at bends/turns due to the graphene sheets being covalently bonded at such junctions. Solvent casting offers the advantage of preserving the porous structure of expanded graphite in the composite, resulting in the above highly thermally conductive interpenetrating network of graphene and polymer. Solvent casting also does not break down the expanded graphite particles due to minimal forces involved, allowing for efficient heat transfer over long−length scales, further enhancing overall composite thermal conductivity. Comparisons with a recently introduced effective medium model show a very high value of predicted particle–particle interfacial conductance, providing evidence for efficient interfacial thermal transport in expanded graphite composites. Field emission environmental scanning electron microscopy (FE−ESEM) is used to provide a detailed understanding of the interpenetrating graphene−polymer structure in the expanded graphite composite. These results open up novel avenues for achieving high thermal conductivity polymer composites.We demonstrate in this work that expanded graphite (EG) can lead to a very large enhancement in thermal conductivity of polyetherimide−graphene and epoxy−graphene nanocomposites prepared via solvent casting technique. A  One of the most promising approaches has been the percolative effect, where graphene–graphene contact is used to bypass the polymer, leading to significantly higher thermal conductivity enhancement [k values [−1−1 at just 10 wt% filler concentration. This value represents the highest k value measured in a polymer composite at this low weight fraction of filler. Presented results provide new avenues for achieving efficient thermal management in a wide range of applications.Increased thermal diss0ipation in modern electronic devices has led to a demand for thermally conductive materials with superior thermal conductivity (k) ,2. The lk) ,4. Howevm−1⋅K−1) ,6 of polm−1⋅K−1) ,11,12,13m−1⋅K−1) ,18,19,20ancement . More rek values ,25,26,27k values ,29,30, w−3) and exhibits high mechanical strength (10 MPa), thermal conductivity (25–470 W−1−1), and electrical conductivity (106–108 S−1) [Intercalation of acid molecules and oxidizing agents into graphite followed by rapid heating (900 °C) causes intercalated agents to gasify, producing enough pressure to push the adjacent graphite layers apart, resulting in expanded graphite a. Figure S⋅cm−1) . As a re S⋅cm−1) , phase−c S⋅cm−1) ,34, elec S⋅cm−1) ,36, elec S⋅cm−1) , fuel ce S⋅cm−1) ,39, batt S⋅cm−1) ,41 and s S⋅cm−1) ,43.−1−1 for EG particles. Tao et al. [−1−1 at 10 wt% EG composition. Zhao et al. [−1−1 in EG/paraffin wax composites at 25 wt% composition. Song et al. [−1−1 at 20 wt% composition in EG/MgCl2·6H2O composite. Wei et al. [−1−1 using ~25 vol% expanded graphite. Che et al. [−1−1 at ~20 wt% composition in high−density polyethylene (HDPE) composites.Significant research has been performed into the use of EG in enhancing the thermal conductivity of polymer graphene nanocomposites. Wu et al.  measuredo et al.  preparedo et al.  measuredg et al.  measuredi et al.  created e et al.  employedIn this work, we show that expanded graphite can lead to a remarkable increase in the thermal conductivity of polyetherimide composites prepared via the approach of solvent casting. Expanded graphite has a highly porous structure with interconnected graphitic walls, as seen in the SEM image of expanded graphite in the air in k~2000 W−1−1) [ink~2000 W−1−1, and outk~10 W−1−1) [The resulting interpenetrating polymer graphene structure from the solvent casting method allows for very efficient heat transfer. This can be seen by observing that graphitic nanosheets in this network are continuous for very long−length scales , as shown in m−1⋅K−1)  of graphm−1⋅K−1) .−8 K/W, which is comparable to graphene−polymer interfacial thermal resistance [k enhancement over percolation effect is that typically very high volume concentrations are required to achieve particle–particle contact for percolation. In the case of solvent−cast expanded graphite/polymer composites, however, the continuous graphitic networks are present at all volume fractions, enabling achieving high k even at low particle concentrations.A unique advantage of this continuous network can be understood by noticing that even at bends/turns, the graphitic sheets are still continuous, allowing for a very high thermal conductance  at such junctions. The above effect of low interfacial resistance at bends/turns in an expanded graphite network offers an advantage over the effect of percolation ,51,52,53The polymers chosen for this work, namely polyetherimide and epoxy, enable the advantage of superior thermal interaction with graphene ,56. BothWe further investigated the effect of sonication parameters on thermal conductivity enhancement and showed that lower sonication time and power result in larger expanded graphite particles in the composite, which allows for heat to be conducted efficiently along longer lengths scales, resulting in higher thermal conductivities. Finally, we also compare the measured thermal conductivity results with a recently introduced effective medium model that takes graphene–graphene contact into account for thermal conductivity prediction.Graphite intercalated compound (GIC) or expandable graphite (EPG), with an average particle size of ≥180 te store . Graphente store . The epote store . Commercte store  and Sigmte store . The orgThe fabrication procedure of the EG/PEID composite is illustrated in To prepare epoxy/expanded graphite composites, resin was added to 90 mL N,N−dimethylformamide (DMF) solution and stirred while heating at 150 °C to obtain a homogeneous mixture. Expanded graphite was added to this solution and stirred for one hour. The solution was then tip sonicated for 40 s followed by stirring at 150 °C until the solvent completely evaporated. After the solvent evaporated, a thick mixture of EG/epoxy was obtained. This mixture was then spread over a PTFE sheet and kept in a vacuum oven at 140 °C for 15 h to remove any residual solvent present in the mixture. On cooling the mixture to room temperature, a hardener was added to it and mixed to obtain a homogeneous viscous paste. This paste was then transferred to aluminum molds and cured at 90 °C for 20 h.k): k of EG/polymer composites was measured by the laser flash technique. A Netzsch LFA 467 Hyperflash  laser was used to measure the through−thickness thermal diffusivity of the samples. A total of 8–12 samples of 12.5 mm diameter and 0.3–0.4 mm thickness were used to measure the thermal diffusivity (α) at room temperature (23 °C). The samples were coated with graphite spray before the measurement to efficiently absorb heat from a flash lamp, and an average of 6–8 measurements was reported. This laser flash technique induces heat by a laser pulse on one surface of the sample, and the temperature rise is captured on the other surface of the sample as a function of time. α is determined by LFA using the following equation: α = (0.1388 d2)/t1/2, where α is the thermal diffusivity (mm2/s), t1/2 is the time to obtain half of the maximum temperature on the rear surface, and d denotes the sample thickness (mm). The thermal conductivity was calculated using k = αk, ρ, and Cp represent the thermal conductivity, density, and specific heat constant of the sample, respectively. In this work, the density and specific heat of the composite samples were calculated using a gas pycnometer  and differential scanning calorimetry (DSC) .Thermal Conductivity (Scanning Electron Microscopy (SEM): Morphological characterization of EG filler and EG/polymer composites was carried out by high−resolution field emission environmental scanning microscopy . This SEM was operated in secondary electron (SE) mode at an accelerating voltage of 20 kV. To prepare the samples for SEM imaging, liquid nitrogen was used to crack down the composite to image over the cross−sectional area. Montage large area mapping of EG fillers has been captured using MAPS software of FE−ESEM.Raman Spectroscopy (RS): Raman spectroscopy (RS) was performed using a DXR3 SmartRaman Spectrometer  to collect the data over the range from 3250 to 250 cm−1, laser wavelength λL = 633 nm, spectral resolution = 0.16 cm−1, and imaging resolution = 702 nm for the EG and GIC samples. An Olympus BX 41 microscope with 5× objective, 10 s exposure time for 15 accumulations, and 3 scans per sample were used to collect the spectra.X−ray Diffraction (XRD): A PANalytical Empyrean Diffractometer  produced the information regarding the crystal structure of EG compared to GIC using Bragg−Brentano focusing geometry at room temperature. A total of 3 kW Cu Kα radiation (λ = 1.5406 Å) with a scan range of 2θ = 5 to 80°and step size of 0.013°.X−ray Photoelectron Spectroscopy (XPS): X−ray Photoelectron Spectroscopy (XPS) was performed for GIC and EG samples by Thermo Scientific K−alpha XPS . A Kα gun source was used to excite the sample, and the data were collected for an acquisition time of ~70 s at 400 µm spot size. The passing energy of 200 eV was used to find the carbon (C), oxygen(O), and sulfur (S) peak in this analysis spectrum. The atomic percentage of C, O, and S were investigated using the Avantage software. To determine the atomic percentage, this software was used to perform the curve fitting in accordance with Gaussian and Lorentzian functions.−1−1 at 10 wt% filler composition. This value represents a remarkable 2770% enhancement over the k (0.23 W−1−1) of pristine PEID, providing new avenues for high thermal conductivity composites. Similarly, the measured thermal conductivity of EG/epoxy composites also shows remarkable enhancement at very low loadings of EG. At just 2 wt% EG composition, a thermal conductivity value of 1.74 W−1−1 is achieved for the epoxy composite, representing an enhancement of 1025% over pristine epoxy (0.16 W−1−1). The measured thermal conductivities of EG/PEID and EG/epoxy composites are shown in k value is achieved even at low graphene loading, as opposed to the case of percolation, where typically much higher particle concentrations are required to achieve particle–particle contact. Kargar et al. achieved around 6 W−1−1 in graphene/epoxy composites at a high graphene loading of 35 vol% through the percolation effect [−1−1 in Boron–Nitride/epoxy composites [k of 6.6 W−1−1 is achieved in this work through the use of just 10 wt% EG content. Even in percolative k enhancement, the thermal interfacial resistance at particle–particle contact can be relatively high. Konatham et al. reported a thermal boundary resistance of 5.5 × 10−8 K/W, almost as high as at the graphene−polymer contact [k value of EG/PEID composite has been achieved due to such continuous graphitic network of EG as presented in We first compare the measured values against previous results reported for the percolation effect based on establishing a graphene–graphene network. A key advantage of expanded graphite is that a large enhancement in n effect . Percolamposites  at a vol contact . Expandek achieved in this work is also due to the use of solvent casting, which offers the advantage of preserving the porous structure of expanded graphite within the composite  is higher by 572% relative to GnP/PEID composite (1.0 W−1−1). Similarly, at 2 wt% composition, k of EG/epoxy composite (1.74 W−1−1)) is higher by ~600% relative to GnP/epoxy composite (0.25 W−1−1). At the low graphene content used in this work, GnPs are well separated by polymer; high interfacial thermal resistance between GnPs and polymer then restricts the enhancement achievable through the use of graphene nanoplatelets. These results highlight the large advantage of expanded graphite over graphene nanoplatelets in enhancing composite thermal conductivity.We also compare measured k enhancement. We further investigated the effect of sonication parameters used during composite preparation on While the above images show the expanded graphite embedded in the polymer, we also show expanded graphite before it is embedded into the polymer in k values of EG/PEID composite with theoretical predictions based on a recently introduced effective medium model by Su et al. [In the next section, we compare measured u et al. . CompariThe effective medium model introduced by Su et al.  includes0 and 0k is the thermal conductivity of an interlayer surrounding graphene sheets. The role of this interlayer is to include the effect of interfacial thermal resistance at graphene−polymer and graphene–graphene contacts.In the above equations, The effective thermal conductivities In the above equations, −1−1 from our measurements (in suitable agreement with the literature). The in—plane . This interlayer thermal conductivity of 0.17 W−1−1 at graphene–graphene contact corresponds to a graphene–graphene interfacial thermal resistance of 1.17 × 10−8 m2K/W. This value is significantly smaller than the interfacial resistance predicted for regular graphene–graphene contact [−8 m2K/W), providing evidence that graphene sheets are in superior contact  at bends/turns in an expanded graphite network, compared to the contact between discrete graphene particles. The analysis points to the advantage of the continuous graphitic network achieved through the use of expanded graphite via solvent casting technique in enhancing the thermal conductivity of polymer composites. We further show the effect of lower graphene–graphene contact conductance by decreasing the graphene–graphene contact thermal conductivity in the model. It is seen that as the graphene–graphene contact thermal conductivity is decreased from 0.17 to 0.10, the effective composite thermal conductivity is decreased from 6.6 W−1−1 to 4.4 W−1−1 , indicating the importance of superior graphene–graphene contact conductance for achieving overall high composite k values. We next discuss the characterization of EG filler and EG/PEID composites.A key parameter in the above model is the interfacial resistance at graphene–graphene contact. This value was determined by fitting the measured values to the effective medium model. A suitable agreement between measured thermal conductivities and predicted values was obtained by using a graphene–graphene interlayer thermal conductivity of 0.17 W⋅K−1 see , which i contact   images enable understanding of morphological differences between graphene nanoplatelets (GnPs), Graphene Intercalation compound (GIC), and thermally expanded graphite (EG) a. ThermaX-ray diffraction (XRD) analysis was performed to determine the crystal structure and interlayer spacing of GIC and EG. −1 and ~1350 cm−1 for graphitic material [2 carbon of hexagonal ring due to in−plane tangential stretching of the carbon−carbon bonds [3 hybridized carbon [−1 and represents a second−order two−photon process [D/IG ratio is typically used to characterize the defective state of graphene. The absence of a D band in EG suggests the presence of a highly ordered defect−free graphite structure in EG. This high degree of the ordered structure of graphite, even after thermal expansion, has a strong beneficial impact on the thermal conductivity enhancement of polymer graphene composite as it preserves the intrinsic high thermal conductivity of graphene itself.Nondestructive Raman analysis was also performed to further analyze the crystal structure before and after the thermal expansion. Raman spectroscopy of material . The G bon bonds , and thed carbon . The 2D  process . Raman sXPS analysis was further performed to investigate the concentration of carbon (C), oxygen (O), and sulfur (S) elements before and after the thermal expansion of GIC, as presented in −1−1 is measured, representing an enhancement of 2770% over pristine polyetherimide. This large enhancement in thermal conductivity is found to be due to a network of continuous graphene sheets over long−length scales achieved through the solvent casting technique, which preserves the interconnected porous structure of expanded graphite within the composite. Even at bends/turns in graphene sheets in such a network, the sheets are covalently bonded, which minimizes the interfacial thermal resistance at junctions, enhancing heat transfer. Overall, the resulting structure allows highly efficient heat conduction over long−length scales along the continuous graphitic sheets. This results in the observed high thermal conductivity of the composite through the use of EG. The thermal conductivity of EG/PEID composite is also found to dramatically exceed that of graphene nanoplatelet (GnP)/PEID composites by 572% for 10 wt% filler composition. At low filler loading, GnPs are well separated by polymer and the resulting high graphene−polymer interfacial thermal resistance results in low effective GnP/PEID thermal conductivity. Presented results highlight the advantage of expanded graphite in enhancing the thermal conductivity of polymer composites and can lead to novel avenues for achieving efficient thermal management in a wide array of technologies.In summary, we demonstrate that expanded graphite (EG) can lead to a large enhancement in the thermal conductivity of EG/PEI composites prepared through solvent casting. At 10 wt% EG composition, a high thermal conductivity of 6.6 W


Polymer nanocomposites have been gaining attention in recent years. The addition of a low content of nanomaterials into the matrix improves mechanical, wear, thermal, electrical, and flame-retardant properties. The present work aimed to investigate the effect of graphene and hexagonal boron nitride nanoparticles on Kevlar and hybrid fiber-reinforced composites (FRP). Composites are fabricated with different filler concentrations of 0, 0.1, 0.3, and 0.5 wt.% by using a hand layup process. Characterizations like tensile, flexural, hardness, and impact strength were evaluated separately, heat deflection and viscosity properties of the epoxy composites. The dynamic viscosity findings indicated that a higher concentration of filler material resulted in nano-particle agglomeration. Graphene filler showed superior properties when compared to hexagonal boron nitride filler. Graphene showed optimum mechanical properties at 0.3 wt.%, whereas the hBN filler showed optimum properties at 0.5 wt.%. As compared to Kevlar composites, hybrid (carbon–Kevlar) composites significantly improved properties. As compared to graphene-filled composites, hexagonal boron nitride-filled composites increased scratch resistance. Digimat simulations were performed to validate experimental results, and it was observed that hybrid fabric composites exhibited better results when compared to Kevlar composites. The error percentage of all composites are within 10%, and it was concluded that hybrid–graphene fiber composites exhibited superior properties compared to Kevlar composites. Graphite, hexagonal boron nitride (hBN), and graphene nanomaterials are most suited for improving the mechanical and electrical properties of epoxy composites [Recent scientific advances in composites have opened the way for fiber-reinforced composites (FRP) for lightweight and imperishable structural applications like aerospace, automotive, rocket, and sports, which demand multifunctional properties. To achieve these multifunctional properties, nanoparticles are being added to polymer matrix systems. Several types of particles, such as silica, carbon nanotube (CNT), and clay have been added to FRPs. These nanoparticles enhance good mechanical properties but due to an insulating property of epoxy, thermal properties are poor in FRP composites . To addr2 on the mechanical properties of Kevlar composites and found an increase in hardness and flexural properties by 29% and 50% for 3 wt.% of composites. Srivatsava et al. [Sharma et al.  investiga et al.  investigBased on the literature review, it can be observed that several works reported on the mechanical properties of carbon fabric with MWCNT, silica, clay, graphene, etc. Very few works reported on Kevlar and hybrid fabric composites to the best of our knowledge. The novelty of the present work lies in incorporating the hybrid Kevlar–carbon fabric as a matrix into the 2D nanofillers reinforced epoxy. To the best of authors’ knowledge, such work has not been reported elsewhere in the literature. The present work aims to study the effect of different wt.% graphene  and hBN  fillers on Kevlar and hybrid (carbon as a wrap and Kevlar weft) fiber on mechanical properties, dynamic viscosity, hardness, and heat deflection response. Digimat simulations were also performed and compared with the experimental results.2/g. The functionalization of h-BN was performed as described in the literature [2/g, bulk density of 0.3 g/cc, and an average diameter of 10–15 microns with a purity greater than 98%.Bisphenol A diglycidyl ether (DGEBA), commonly named Epoxy Ly556, and Hardener Hy951 were obtained from Singhal Chemical Corporation, Delhi, India. Kevlar 200 GSM) and hybrid  were obtained from Marktech companies Pvt. Ltd., Bangalore, Karnataka, India. Hexagonal boron nitride (h-BN) was obtained from Sisco Research Laboratories Pvt. Ltd., Hyderabad,Telangana,India. Hexagonal boron nitride (h-BN) platelets had an average particle size of 70 nm and surface area of 19.4 m0 GSM andFabrics were washed with acetone to extract dust particles from the fiber’s surface and then dried for 24 h at room temperature. After drying, the fibers were subjected to moisture removal in a vacuum oven at 70 °C for 4 h. To ensure that no air bubbles were present in the epoxy, degassing was performed 60 min before the addition of the curing agent and 5 min after the addition of the curing agent. Graphene and hexagonal boron nitride particles at different wt.%  were added into the epoxy individually. Ultrasonication was performed on different filler concentrations at a temperature of 80 °C for 0.5 h. Hardener HY951 was added to the epoxy with a mixing ratio of 10:1 . Kevlar Tensile testing was conducted on a Universal Testing Machine , according to ASTM D 3039 standard, at a crosshead speed of 2 mm/min at room temperature. The test samples were cut by using a Bosch Jigsaw machine . The final dimensions of the cut tensile samples were 250 mm × 25 mm × 3 mm. For accuracy and repeatability, 5 samples were tested per each filler concentration and the average values were reported.The flexural test was performed on a Tinius Olsen H10KL, Noida, Uttar Pradesh, India, according to ASTM D7264, at a crosshead speed of 2 mm/min. For accuracy and repeatability, 5 samples were tested per each filler concentration and the average values were reported.Hardness measured by using Micro Vickers hardness tester supplied by R.S.Scientific Kolkata, West Bengal, India. ASTM E384-17 standard was used for the measure of microhardness. Hardness was measured at a constant load of 1 kg over a dwell period of 20 s. To ensure repeatability, 5 readings were taken for each configuration and the average value was reported.Viscosity measurements were carried out on the Rheometer (MCR 102) model from Anton Paar India limited, Bangalore, Karnataka, India. A frequency sweep was performed on the epoxy filled with nanomaterials (graphene and hBN).The heat deflection test is a measure of the ability of a polymer to carry a load at raised temperatures. The heat deflection apparatus was supplied by S.C.Dey & Co, Kolkata, West Bengal, India. ASTM D 648 or ISO 75 standards were used for the measuring. The impact test was performed on a Digital Izod/Charpy tester supplied by R.S.Scientific, Kolkata, West Bengal, India, according to ASTM D256, with sample dimensions of 63.5 mm × 12.7 mm × 3 mm. The prepared sample was placed 2 mm away from stoppers with a 21.7 Kg weight hammer at the center of the sample.A scratch test was performed on a Scratch Tester , according to ASTM D7027-05. Experiments were conducted under a constant load of 150 N, a ramp 30–150 N with a velocity of 1 mm/s, and scratch length of 25 mm.Digimat software was used for linear and nonlinear multiscale material modeling. Digimat software enables the prediction of the constitutive behavior of heterogeneous and/or anisotropic materials such as polymer matrix composites (PMC), rubber matrix composites (RMC), metal matrix composites (MMC), or even nanocomposites. It consists of different modules, namely, MF, FE, MX, MAP, CAE, RP, VA, HC, and AM. In the present study, MF and FE modules were used. MF indicates mean-field homogenization, and it uses Eshelby-based semi-analytical mean-field homogenization approaches and an analytical description of the material in order to compute the thermo-mechanical, thermal, or electrical properties of a composite as a function of its microstructure morphology, i.e., the inclusion of shape, orientation, volume/mass fraction, and micro  material behavior. Mori–Tanaka homogenization models were used in the present study. In this uniform model, fiber distribution assumption was taken into consideration to approach the problem within the micromechanical theory of periodically arranged heterogeneous materials. The Digimat MF module was used for the evaluation of fiber volume fraction, and Young’s modulus and Poisson’s ratio of Kevlar- and hybrid fiber-reinforced graphene filler composites were calculated.The Digimat FE module was used for the visualization of RVE geometry and the evaluation the stress–strain analysis. Fabric microstructure was generated by 3D woven interlock wrap weft phases considered for Kevlar composites and for hybrid composites with Kevlar as a wrap and hybrid fabric (carbon–Kevlar) as weft phase, with parameters shown in The generated RVE model was subjected to meshing (voxel) with a size of 100 mm, followed by uniaxial loading with 90° theta.2 compared to an all-composite variation, whereas GK0 demonstrated the least value of 390.10 N/mm2. The reason for the superior properties of GH composites is attributed to their better dispersibility and excellent load-carrying capacity [capacity . The inc2, while hBK0 demonstrated the least value of stress. The stress behavior of different composites in incremental order are as follows: hBH3 > hBH2 > hBH0 > hBK3 > hBK2 > hBK1 > hBK0. The tensile properties of different composites are shown in 2 was observed in the GH2 composite among all composite configurations.The ultimate tensile strength comparison of GK and GH composites are shown in Young’s modulus comparison of GK and GH composites are shown in RVE analysis was performed using Digmat software, and the findings were compared with the experimental data, as shown in Young’s modulus of Kevlar and hybrid composites incorporated with graphene are shown in The ultimate tensile strength of hBK composites obtained experimentally and those obtained using Digimat is shown in Young’s modulus of Kevlar and hybrid composites incorporated with graphene are shown in From the above discussion, it can be observed that 0.3 wt.% graphene exhibits superior UTS and Young’s modulus in all combinations; this may be due to the uniform dispersion of particles through the fiber surface and epoxy. Uniform dispersion ensures better interlocking between the fiber and matrix parts, which results in the improvement of the load-carrying capacity of the composite. At 0.5 wt.%, the UTS and Young’s modulus decreased because of the agglomeration caused by Van der Waals forces . An incrRVE FE analysis of hybrid composites is shown in Viscoelastic materials behave linearly up to the critical strain level, beyond which the materials behave non-linearly. Therefore, a strain sweep was initially performed to estimate the linear viscoelasticity of the epoxy-filled graphene and hBN composites. The dynamic viscosity of epoxy filled with graphene and hBN is shown in The dynamic viscosity of pure epoxy incorporated with graphene and HBN filler was observed to decrease as the angular frequency increased. The behavior of neat epoxy is similar to a Newtonian fluid at lower frequencies, but as the frequency increases, it shows a shear-thinning effect. The addition of nanofiller content linearly increased the viscosity of epoxy. At higher viscosity, it is difficult to mix the nanofillers properly into the epoxy matrix, which leads to the agglomeration of nanoparticles, leading to the poor strength of the matrix. Another problem at higher viscosity is that while carrying out the hand layup process it becomes exceedingly difficult to roll epoxy smoothly over the entire composite laminate. Up to 0.5 wt.% filler content, the viscosity of filler increases proportionately, whereas at 0.7 wt.%, the jump in the viscosity is high, which leads to a higher chance of agglomeration at 0.7 wt.%. Therefore, in the present study, we experimented with up to 0.5 wt.% of filler content.2. It should be noted that beyond GK2, the addition of filler content resulted in decreased values. All composites exhibited no delamination until peak stress but thereafter displayed a significant drop due to crack propagation inside the composites, beyond which the composite continued to bear load but never achieved peak values. The hardness values of graphene-based fiber composites are depicted in The impact test was performed according to ASTM D 256. The impact strength of various graphene-based composites was shown in The impact strength of hBN-filled composites is shown in Heat deflection curves of graphene-filled composites are shown in Heat deflection curves of hBN-filled composites are shown in Hybrid composites show a higher flexural modulus, which demonstrates their usability for compressive and bending load applications. The flexural modulus in the HBH3 composite was increased by 85%.Hardness, tensile strength, and heat deflection of graphene-based composites exhibited superior properties when compared to HBN-based composites. The optimum properties were found at 0.3% filler content.Digimat simulations were close to experimental results regarding ultimate tensile strength. Simulation results suggested that the applied load conditions are within acceptable limits.Kevlar’s impact resistance properties were greatly enhanced by the addition of fillers and thus enhancing their applicability.In the present research work, a comparative study on the effect of graphene and hBN fillers on carbon, Kevlar, and hybrid composites and their mechanical properties was evaluated. Based on the study, the following conclusions can be drawn:


Modulating near-fieldradiative heat transfer (NFRHT)with a highdynamic range is challenging in nanoscale thermal science and engineering.Modulation depths [/ × 100%] of ≈2% to ≈15.7%have been reported with matched modes, but breaking the constraintof mode matching theoretically allows for higher modulation depth.We demonstrate a modulation depth of ≈32.2% by a pair of graphene-coveredSU8 heterostructures at a gap distance of ≈80 nm. DissimilarFermi levels tuned by bias voltages enable mismatched surface plasmonpolaritons which improves the modulation. The modulation depth whenswitching from a matched mode to a mismatched mode is ≈4.4-foldcompared to that when switching between matched modes. This work showsthe importance of symmetry in polariton-mediated NFRHT and representsthe largest modulation depth to date in a two-body system with fixedgap distance and temperature. Graphene hashighly tunable surface plasmon polaritons (SPPs) related to its Fermilevel in a linear Dirac band, which makes it an ideal thermal modulatorwith external stimuli at a fixed operating temperature.23 Altering the free carrier states of a van der Waals’ heterostructurealso allows for exotic nanoscale phenomena such as tunable Mott insulator,24 nanoimaging in an infrared-waveguide,25 and efficient Fizeau drag from Dirac electrons,26 etc. Recent experimental works have demonstratedthat graphene plasmons enable giant radiative heat transfer at thenanoscale.12 Graphene withdifferent Fermi levels supports tunable SPPs and accounts for differentradiative heat flux. A high dynamic range modulation of NFRHT allowsfor a high signal-to-noise ratio (SNR) enabling potential applicationsin thermal switches and communication. Thomas et al. have shown anelectronic modulation depth of ≈2% [defined by / × 100%] with a pair ofgraphene sheets on Al2O3/SiO2 substrate.21 The Fermi levels of the two graphene sheetsare assumed to be equal since the samples are in conductive contact.Our recent work has shown that the modulation depth could reach ≈15.7%with stacked graphene layers at similar Fermi levels.6 However, the modulation effect is limited by the matchedresonance modes between the identical graphene sheets, as the grapheneon the emitter and receiver have similar Fermi levels.Dynamic modulation of NFRHTrequires changing the optical responsesof the emitter and receiver to modulate the radiative heat flux. Previousworks28 thermal communication, and is suitable for tunable plasmon- or phonon-inducedthermal or photonic regulation.Here,we present a significant modulation improvement by a pairof graphene/SU8 heterostructures at a gap distance of ≈80 nm.Back-gated tuning was employed for the graphene Fermi level control.The mismatched SPPs due to different Fermi levels of the two graphenesheets allow for a much smaller heat flux compared to the matchedSPPs case. The measured maximum modulation depth could reach ≈32.2%and is ≈4.4-fold compared to that of the matched case. Thisexperimental work represents the largest modulation depth ever reportedfor the radiative heat transfer in a two-body planar system with afixed gap distance and operating temperature. The mode-mismatch-inducedhigh-efficiency NFRHT modulation should inspire potential applicationsof thermal switches,2/Si substrates at a gap distance d is calculated by33T,ω)= ℏω/[exp(ℏω/kBT) – 1] isthe mean energy of the Planck thermal harmonic oscillators withoutzero point energy. ℏ is the reduced Planckconstant and kB is the Boltzmann constant.ξj represents the energy transmissioncoefficient between the emitter and receiver (considering both s- and p-polarization modes):kz0 is the z-component of thewave vector in vacuum (k0). rj,e and rj,r are the Fresnel reflectioncoefficients of the emitter and receiver, respectively. ξp with β > k0 represents the photon tunneling probability of the p-polarized evanescent modes.We study the NFRHT between apair of graphene-covered SU8 heterostructureson SiO32h1 is thethickness of layer one and rs(p),01 is the Fresnel reflection coefficientat the interface between layer 0 and layer 1 for s- or p-polarization modes:n = 0, 1, 2 is the numberof the layers, ε⊥n)( and ε||n)( = ε||n) and emitter (EF1) are both −0.13 eV  =  V. The SPP coupling modes are slightly deviatedfrom the ideal dispersion curves of a pair of suspended graphene sheets(dashed-dotted lines) due to the impact of the SU8 spacers and theSiO2 substrates. Typical phonon polaritons from the SiO2 substrate support the near-unity ξp aroundwavelengths of 8.56 and 19.9 μm, but have no contribution tothe heat flux modulation. The yellow-dashed lines correspond to theoccupation factor Θ– Θ in arbitraryunits. The near-unity ξp overlaps the dispersioncurves of the SPP coupling modes of two Gr/SU8 heterostructures (notshown). The coupled modes split into two branches at a lower βbut merge at a larger β, as the larger loss (at large β)in the vertical direction enables rapid attenuation of the SPP modesand prevents the interaction between the emitter and receiver. Thesematched SPP coupling modes are supported at the mid- and far-infraredregions with a larger occupation factor and are the dominant contributorto the radiative heat flux. Compared with changing the graphene Fermilevels synchronously , producingdissimilar Fermi levels is a more effective way to pursue an improvedmodulation depth. When only the graphene Fermi level of the receiverchanges to −0.22 eV  between the Gr/SU8heterostructures shows the influence of the graphene SPPs with identicalor dissimilar Fermi levels (EF2 = EF1 = −0.13 eV (case I), the spectrumattributed to the matched strong SPPs covers a broad frequency regionwith the highest value. A broader spectrum with lower spectral heatflux (compared to case I) could be observed when EF2 = EF1 = −0.22 eV(case II). However, the spectral heat flux decreases dramaticallywhen EF2 is not equal to EF1, where EF2 = −0.22and EF1 = −0.13 eV (case III).The spectrum in case III shares a similar frequency region with caseI, but it has smaller intensity due to the mismatched resonance modes.The radiative heat flux reaches 1.34 × 104, 1.19 ×104, and 0.57× 104 W/m2 forcases I, II, and III, respectively. This leads to a remarkable improvementof the NFRHT modulation depth from case I to case III, where the modulationdepth is 40.3% [calculated by (radiative heat flux of case I –radiative heat flux of case III)/(radiative heat flux of case I +radiative heat flux of case III) × 100%] and is ≈6.7-foldcompared to that from case I to case II. EF1 and EF2. Fordifferent EF1 at a fixed EF2, the radiative heat flux arrives at a maximal valuewith two similar EF. This confirms theimportance of the symmetry of the polariton-mediated NFRHT system.White-dashed line shows the cases with matched coupling SPP modes.Case I with η = 416 is quite close to the peak point at EF2 = EF1 = −0.14eV. The white-dashed-dotted line highlights the cases correspondingto the mismatched resonance modes. Finding a point with smaller heatflux based on the mismatched graphene Fermi levels allows for a highermodulation depth. In this work, the radiative heat flux from caseI to case II, and case I to case III were measured.The photon tunneling probability ξ−0.13 eV 2a, corre−0.22 eV 2b, the Si levels 2c. When 2 on Si substrates(see Supporting Information (SI) section2). Eight identical SU8 nanopillars with thickness of ≈90 nmwere fabricated on the surface of the graphene sheets for the receiver(L = 3 mm in Vg – Vn induces electron (hole) doping,41 allowing an excess-electron surface concentration of n = ηc(Vg – Vn), where Vg isthe gate-tuning voltage, Vn is the voltageat the charge neutral point, and ηc is a coefficientrelated to the back-gated structure with gate insulators composedof 90 nm thick SU8 and 300 nm thick SiO2 (see SI section 1). The graphene Fermi levels couldbe obtained by EF = sgn(n) ℏVF(π |n|)1/2/e (unit: eV), where sgn(x) is thesign of x and VF = 1× 106 m/s is the Fermi velocity of graphene.44 The emitter was pressed on the receiver in a cross shape and separatedby the SU8 nanopillars  based onthe one-dimensional linear elastic analysis6 (SI section 4). Psum detected by the heatflux sensor (HFS) is the sum heat power of Pc and Pr, where Pr is the near-field radiative heat power including thecontribution of both propagating waves and evanescent waves. The measuredradiative heat flux could be obtained by Q = (Psum – Pc)/A, where A is the active area of 3 ×3 mm2. Pc is the heat conductionfrom the eight SU8 nanopillars and is calculated based on Fourier’sLaw (SI section 5).The emitter and receiver consist of single-layergraphene-coveredSU8 heterostructures on the 300 nm-thick SiOreceiver3a. A bacopillars 3c. A totopillars 3d.6,10,1T2 = 303.15 K with temperature difference ΔT of 5 K at a gap distance of ≈81 nm. When no bias voltagesare applied to the emitter and receiver, i.e.,  =  V, EF2 and EF1 are calculatedto be −0.205 eV according to the measured radiative heat flux.According to the back-gated method, other bias voltages of ,, , and  V correspond to the graphene Fermilevels of −0.187, −0.155, −0.13, and −0.098eV respectively. The Fermi levels become closer to the Dirac pointwith higher positive voltages, indicating the hole doping of the graphenesheets. When EF2 = EF1 = −0.13 eV, the matched SPPs allow for the broadbandnear-unity ξp within the desired mid- and far-infraredregion, producing the best performance among all measured cases. Theradiative heat flux reaches 1.43 × 104 W/m2 and is ≈441-fold of the corresponding BB limit. When EF2 = EF1 = −0.205eV, the enhancement (≈ 381-fold with respect to the BB limit)is relatively robust, as the SPP modes of the emitter and receiverare still matched, despite the blue-shift to a higher frequency regionwith less occupation factor  was observed at  = V, where the graphene Fermi levels are dissimilar . When V2 increases to35 V again, the radiative heat flux can still reach a peak value of1.39 × 104 W/m2 with enhancement of ≈429-foldof the BB limit, indicating the robustness of the tuning devices. EF2, EF1) =  eV to  =  eV, but the modulationdepth is ≈26.3% when only changing EF2 to −0.205 eV. The maximum modulation depth from the matchedcase to a mismatched case reaches ≈32.2% when tuning EF2 to −0.22 eV by applying −10V bias voltage to the receiver. It could potentially be further improvedwhen other larger negative voltages are applied. Considering the breakdownof the capacitor-like samples, only −10 V was investigatedin this work. Construction of the mismatched resonance modes playsa significant role in the NFRHT modulation improvement. Switchingthe symmetry of the polariton-mediated near-field system gives promisinghigh-efficiency modulation in thermal radiation. The modulation depthof ≈32.2% is ≈4.4-fold compared to that for the matchedcases and is 16.1-fold of the previous report on graphene-based heterostructures.21 The modulation of the radiative heat flux atanother ΔT of 3 K was also investigated in V2, V1) =  V to  V repeatedly. The average modulation depthof ≈7% is similar to that with ΔT =5 K in ΔT (like that in this work) ismore likely required in potential applications such as thermal communication,since the state of the thermal equilibrium will be easier to achieveafter switching the graphene Fermi levels, hence a faster responsetime. The results show good repeatability and robustness of our devicesfor the NFRHT modulation. The sample in this work consists of a single-layergraphene-covered SU8 heterostructure, which is simpler and easierto fabricate, and allows for larger modulation depth (with mismatchedmodes) than that of the previous work6 withmultilayer structure (only matched modes are considered). The physicalmechanism in ref  =  eV, and thedashed-dotted line indicates the mismatched cases with  =  eV. The modulation depths (red-solid line) are calculatedby the radiative heat flux when switching the graphene Fermi levelsfrom  eV to eV at the corresponding d. The modulation depth canbe improved to 43.6% at d = 50 nm. The modulationdepth decreases as d increases and could still reach8.4% when d = 500 nm. In general, the modulationeffect of the NFRHT is stronger with smaller gap distances due tothe near-field effect of the graphene plasmon polaritons.The modulation depth versus variable vacuum gapdistance stigated5. The blIn conclusion, we have experimentally demonstratedan improvedradiative heat flux control with an ultrahigh modulation depth of≈32.2% with a pair of graphene-covered SU8 heterostructuresat a gap distance of ≈80 nm. Tuning the graphene Fermi levelsallows for coupling or decoupling of the SPP modes. Breaking the symmetryof the SPP resonance modes by changing the Fermi level of either theemitter or the receiver provides a much higher NFRHT modulation depthcompared to other known methods. The theoretical analysis of the gap-dependenceof the modulation depth shows that the modulation effect of the NFRHTis stronger with smaller gap distances due to the near-field effectof the graphene plasmon polaritons. This work focused on the dynamicmodulation of the NFRHT by the mode mismatching method and is importantfor the fundamental understanding of the tunable collective optoelectronicphenomena for light–matter interaction and heat transfer atthe nanoscale. The method is feasible for other van der Waals’heterostructures with resonance modes sensitive to external stimuli.The experimental results obtained in this work provide strong evidenceon mode-mismatch-induced ultrahigh dynamic NFRHT modulation with polariton-mediatedmaterials in the near-field systems, which is of fundamental interestin polaritons control and energy manipulation. This result shouldinspire potential applications such as thermal switches and thermalcommunication.


This is why the PA66/graphene film produced a more subtle secondary network with the addition of graphene, used as a charge capture site to increase its surface charge. Additionally, all the layered structures of TENGs were composed of breathable electrospun films and plain conductive cloth, with water vapor transmittance (WVT) of 9.6 Kgm−2d−1, reflecting excellent wearing comfort. The study showed that TENGs, based on all electrospinning, have great potential in the field of wearable energy supply devices.Triboelectric nanogenerators (TENGs) are considered to be the most promising energy supply equipment for wearable devices, due to their excellent portability and good mechanical properties. Nevertheless, low power generation efficiency, high fabrication difficulty, and poor wearability hinder their application in the wearable field. In this work, PA66/graphene fiber films with 0, 1 wt%, 1.5 wt%, 2 wt%, 2.5 wt% graphene and PVDF films were prepared by electrospinning. Meanwhile, TENGs were prepared with PA66/graphene fiber films, PVDF films and plain weave conductive cloth, which were used as the positive friction layer, negative friction layer and the flexible substrate, respectively. The results demonstrated that TENGs prepared by PA66/graphene fiber films with 2 wt% grapheme showed the best performance, and that the maximum open circuit voltage and short circuit current of TENGs could reach 180 V and 7.8 μA, respectively, and that the power density was 2.67 W/m With the development of Internet of Things technology (IoT), portable, flexible wearable electronic devices with various functions have been widely considered. In the IoT society, various devices offer many possibilities through data transmission and processing ,2,3,4, s2 under hand patching. Incorporating functional materials into polymer nanofibers could improve electrical output performance. Currently, most research focuses on improving output performance, such as selection of friction materials [All electrospun nanofiber films are widely used in TENGs design. The electrospun fibrous films have remarkable properties, such as inherent rough structure, large specific surface area and hierarchical porous structure ,17, whicaterials ,23,24, saterials ,26,27 anaterials ,29,30. H2. Water vapor transmission rate could be up to 9.6 Kgm−2d−1. In addition, the TENGs prepared could be placed anywhere on the body, detecting various human movements, and the mechanical energy that is wasted in life could be obtained. This energy can instead be used to power wearable electronic devices in daily life.In this work, a method for all-electrospun fiber-based TENGs, consisting of two parts, including a PA66/graphene nanofiber-film-rubbed positive base layer and PVDF-rubbed negative electrode layer, was proposed. On the one hand, a PA66/graphene nanofiber film with multi-layer micro-nano structure could not only increase the effective friction area of contact, but also graphene dispersed in the PA66 fiber could be used as a charge capture site, which could increase the surface charge on the film surface, greatly improving its triboelectric properties . In addiPolyamide66 (PA66) was bought from Aladdin Pharmaceuticals . Graphene was sold by Shenzhen Turing Artificial Intelligence Black Technology Co., Ltd. . Polyvinylidene Difluoride  was produced by Arkema . Formic Acid and Polyethylene Oxide (PEO) were bought from Macklin . N, N-Dimethylformamide (DMF) and Acetone were produced by Sinopharm . Plain weave conductive fabric was produced by Saintyear Electronic .Firstly, 0.5 wt% PEO was added to 10 mL of Formic Acid and a magnetic stirrer was used to stir for 1 h to increase the viscosity of the solution, so that it could be well electrospun. Secondly, 15 wt% PA66 particles were put into Formic Acid/PEO solution, which were fully dissolved via a magnetic stirrer for 4 h to obtain 15 wt% PA66 solutions. Finally, different mass fractions of graphene  were put into the prepared PA66 solution, respectively, and stirred to fully mix. Further, ultrasonic dispersion was applied for 6 h to make the graphene evenly dispersed in the solution.The electrospinning process was performed with a supply voltage of 20 kV, a working distance between the spray nozzle and the substrate was 15 cm and the PA66/graphene solution was injected into a 10 mL syringe with a blunt needle (0.6 mm inner diameter). The injection rate of the syringe was 0.1 mL/h. PA66/graphene nanofiber films were collected on plain weave conductive cloth (Ni-Cu) fabric electrodes, which were prepared at 25 °C and relative humidity of 50 ± 5%. The thickness of PA66/graphene fiber film was 80 ± 10 μm, measured by Talysurf Profiler .Firstly, 3 mL Acetone was added into 7 mL DMF to obtain a DMF/Acetone solution in a 7-to-3 volume ratio. Secondly, 10 wt% PVDF powder was added into the mixed solution and magnetically stirred for 4 h under heating in a 60 °C water bath. Finally, the 10 wt% clear and transparent PVDF solution was obtained.The electrospinning process was performed with a supply voltage of 12 kV, while the working distance between the spray nozzle and the substrate was 10 cm. The advance rate of the syringe was 0.4 mL/h. The PVDF nanofiber membrane was collected on the conductive fabric electrode, which was also collected under the same conditions, mentioned in Based on the electrospinning process, the composite fabric of the two friction layers of TENGs was prepared. Firstly, the double-sided conductive tape was used as the electrode lead, which was tightly pasted on the back of the conductive fabric substrate where the different films were located. Then, the required size could be trimmed, such as 4 cm × 4 cm. Finally, these breathable films could be pasted on the corresponding positions of daily clothes, collecting the mechanical energy generated by wasted human movement.SC and VOC were measured by Keithley System Electrometer .The surface morphologies of PA66 films, PA66/graphene films and PVDF films were characterized by the Field Emission Scanning Electron Microscopy . X-ray Diffraction  was used to study the elements and material structure of the electrospun fiber films. The IThe main concept of the design in this work was to use common fabric materials as the basic building unit to construct a wearable TENGs platform, so as to achieve effective harvesting of low-frequency, ubiquitous and easily wasted human biomechanical energy, such as walking and running. The all-electrospinning-based power-generating textile had multiple secondary structures and it could be better integrated with personal clothing, powering wearable electronic devices. These power-generating fabrics with excellent mechanical properties exhibited excellent breathability and cutability. It was an ideal device to make energy harvesters and sensors of a flexible wearable electronic system integrated, which guaranteed the integration of TENGs and wearable electronic devices with clothing.A detailed preparation process for the PA66/graphene nanofiber film and PVDF film are schematically illustrated in The prepared TENGs were composed of three functional components, including a PA66/graphene fiber film and PVDF film with two tribological polarities for contact electrification, and a plain weave conductive cloth for charge conduction. Based on previous studies , construTo study the elements in the films, the XRD patterns of PA66/graphene thin-film layers with different mass fractions of graphene were observed, shown in To study the effect of graphene doping on the output performance, contact-separation-mode TENGs based on PA66/graphene and PVDF films were designed. The two films were glued to two 3D-printed slabs with bosses a,b. The C(x) was the capacitance between the two electrodes at different time-varying displacements x; ε0 was the vacuum dielectric constant; S was the surface area of the friction layer; d0 was the effective thickness of the dielectric layer; σ was the triboelectric surface charge density; SCQ(x) was the amount of short-circuit charge transfer at displacement x; and f was the frequency. Therefore, according to Equations (1) and (2), we could observe that the open-circuit voltage of TENGs was independent of frequency f and increased only with larger displacement between the two contact surfaces, and the short-circuit current was proportional to the frequency f, so the measured maximum short-circuit current increased with the frequency f, shown in SCI) and open-circuit voltage (OCV) of the TENGs and an external force exerted with a frequency of 4 Hz, a higher circuit signal and output power could be obtained.According to previous studies ,34,35, wOCV and SCI were both increased by about four-times. Therefore, the TENGs made of PA66/graphene fiber film with 2 wt% graphene content were the most suitable experimental sample. Further, the piezoelectric properties of PVDF films were tested by the method reported by Yin [With the increase in graphene, the dielectric constant of the electrospun film also had an effect on the output performance. The enhanced dielectric constant of the triboelectric layer could improve the ability to transfer charge density and the surface potential also had a large impact on the performance of TENGs. A higher dielectric constant could use the electrospun PA66/graphene film as a friction layer to store more charge. However, an excessively large dielectric constant would generate a larger dielectric loss, which leads to the leakage of charge. Due to the changes in frequency and external strain during human motion, it was necessary to explore their effects on the output performance of wearable devices. As shown in d by Yin , so the P = I2R, the power of TENGs could be calculated. At 113 MΩ resistance, the maximum power density could reach 2.67 W/m2. The generated output power was sufficient to drive some low-energy electronic devices, solving the problem of sustainable power supply for wearable systems in a stable and reliable manner effectively.To study the effects of load resistance on output of the voltage and current, the output performance of nanofiber-film-based TENGs was researched, shown in m1 was the test cup mass before the test; m2 was the test cup mass after the test; S was the test cup mouth area; and the obtained WVT rate was 9.6 Kgm−2d−1, which could confirm that the all-electrospun textile film had excellent air permeability. In other words, electrospun nanofibers had excellent wearing comfort, ensuring human thermal comfort and wide personal practical applications after their integration into wearable electronic devices.Due to the striking features of the porous structure, the fabricated power generation fabric had excellent water vapor permeability, which can be visually seen in To explore the durability and stability of TENGs, the fabric current outputted after bending the two friction layers for 1000 cycles was tested and compared. It could be clearly seen that there was no significant mechanical damage to the surface, shown in To accurately evaluate the practical application of TENGs, a series of tests was carried out and the combined light chain composed of 110 LEDs in series could be lit by using the TENGs a, which To demonstrate the charging capability of TENGs, a circuit composed of a commercial capacitor and bridge rectifier was designed to test its charging efficiency. The charge in the capacitors  to 1.2V with 4 s, 36 s, 180 s, respectively, is shown in To further demonstrate the application of TENGs based on all-electrospun nanofilms, they could be easily combined with textiles to harvest the mechanical energy generated from the human body, As shown in 2, respectively. In addition, TENGs maintained excellent performance after 1000 cycles under cyclic pressure, keeping its advantages in durability and stability. TENGs with good comfort and high flexibility could obtain biomechanical energy from various human motions. There are great prospects in the system integration of wearable electronic devices.In this work, an electrospun PA66/graphene nanofiber film with a secondary nano-network structure, as a tribo-cathode layer, was fabricated, and it was assembled with the electrospun PVDF film as the tribo-anode layer to form all-electrospun film-based TENGs. The rough surface of the two friction layers not only increased the effective contact area, but also improved the breathability. Furthermore, the porous fiber network in each functional layer guaranteed the air permeability of the TENGs. Compared with the pure PA66 fiber film, the introduced graphene increased the effective contact area of the PA66 nanofiber film and graphene could act as a charge trapping site, enhancing the electrical properties of the TENGs. The fabricated TENGs exhibited high Voc, Isc and peak power density of 180 V, 7.8 µA and 2.67 W/m


Generally, the pore shape and diameter of GAs are difficult to control or the preparation process is complicated, requiring a multi-step operation. Herein, a soft-template one-step hydrothermal synthesis process was proposed to produce GAs with controllable pore sizes. Cyclohexane and n-butanol were added to a graphene oxide suspension to form a uniform aqueous dispersion under emulsification by sodium lauryl sulfate. The reduction process may have occurred around the organic droplets during the hydrothermal reaction, and a large number of organic droplets became countless physical barriers inside the hydrogel. In the later freeze-drying and high-temperature calcination procedures, the droplets evaporated to form a rich pore structure. Compared to the conventional templating method, the organic template was volatilized during the drying process such that no additional process for removing the template was required. In addition, GAs prepared by the template method possessed lower density (2.66 mg cm−3) and better compression performance and, as an adsorbent material, absorbed organic matter and petroleum from wastewater more efficiently than GAs obtained by the traditional one-step hydrothermal method; Q for n-hexane reached 116, and Q for xylene reached 147; also, the GAs prepared by the soft template method can absorb all crude oil in water samples within 30 s.Graphene aerogels (GAs) were synthesized  via a one-step hydrothermal method.Graphene aerogels (GAs) were synthesized  Therefore, three-dimensional (3D) graphene aerogels (GAs) have attracted more attention. GA has an ultrahigh specific surface area,2,12 high electrical conductivity,13 very low density,8,13,14 abundant pore structure,15–18 and certain hydrophobic and lipophilic properties,19 and it not only effectively retains the properties of graphene but is also easy to assemble into parts for use in macroscopic materials.Graphene has excellent electrical and mechanical properties and can be widely applied in supercapacitors,20 chemical crosslinking,21 and polymerization.22 Graphene oxide (GO) contains many oxygen-containing functional groups and is often used as a precursor for the construction of graphene aerogels. GO can be uniformly dispersed in a solvent, such as water, using ultrasound. In a hydrothermal process23 using a solvent or a chemical reduction with a reducing agent,24 GO loses most of its oxygen-containing functional groups and is converted to reduced GO (rGO). Because of its low hydrophilicity, rGO cannot be stably dispersed in water and the resulting layers stack spontaneously to form a stable hydrogel under the action of π–π bonding interactions. However, due to the special 2D structure of GO, with a very high aspect ratio, rGO is more inclined to stack via self-assembly rather than chaotic aggregation. During the stacking process, rGO sheets undergo a hydrophobic action that captures large amounts of solvent to form a graphene hydrogel (GAH). The density, pore size, and wall thickness of the GAH are randomly synthesized and cannot be precisely controlled and, thus, a template method has been proposed. There are two mature ideas for the synthesis of GAHs and/or GAs by the template method; one is a chemical vapor deposition (CVD) process and the other is a solution assembly method. The CVD method mainly deposits graphene directly onto commercially available foamed nickel25 or a special functional template.26 The process maximizes the excellent properties of graphene and allows for custom shapes through using a stencil, but the CVD process is relatively expensive and can only produce small quantities in the laboratory, with, in most cases, the original stencil etched away. In contrast, self-assembled templates are relatively simple to use and GA can be produced in mass and, thus, are widely used.In recent years, researchers have developed a variety of methods to assemble 2D materials into 3D forms in more efficient and applicable ways, including chemical reduction,27–29 is considered to be the easiest method for constructing 3D graphene aerogels. In the process of synthesizing a GAH, when the system is about to reach the gel point, the precursor is immediately frozen in liquid nitrogen or in a refrigerator and ice crystals formed during the freezing process are used as templates. Before the gel point, the GO has been partially deoxidized into a hydrophobic material, such that the ice crystals divide the system into a porous structure. The GAH formed by the reheating reaction has a special pore structure, which can be controlled by the freezing temperature, freezing time, and freezing direction. At the same time, however, the pore size of the GA synthesized by the ISISA method is generally nonuniform, mainly due to the difference in temperature gradients at different locations. Porous solid monoliths have also been used by many researchers as templates for GA synthesis, such as some common polymeric porous materials,30,31 silica,32 and even everyday sponges. Different from the CVD method, these monoliths can accommodate large amounts of GO suspension. And GO is then reduced to rGO using an appropriate method and the template is etched by calcination or strong acid to form a 3D porous GA. The porous monolithic template can adjust the pore size and distribution very conveniently according to needs and the prepared GA has good mechanical strength, chemical properties, and thermal stability. However, the etching process is time consuming and environmentally unfriendly. Although no high temperature treatment is needed during the etching process, further high temperature reactions are required to obtain better electrical performance compared to the CVD method. Pickering emulsions are also a popular method for preparing GA materials. A Pickering emulsion utilizes solid particles33–35 as an emulsifier and has better stability than conventional emulsions. As an amphiphilic particle, GO can form a stable oil-in-water (O/W) emulsion. The monomer is polymerized as a foreign phase to form a skeleton and then the solvent is removed to obtain a polymer/graphene monolith. However, to obtain a pure GA, it is still necessary to etch away the already-formed polymer, thus limiting its application. Shi36 proposed a better and simpler method for controlling the pore structure. By dispersing a certain amount of n-hexane into the GO suspension, GO is reduced and assembles into a 3D structure in the hydrothermal reaction. Droplets of n-hexane do not disperse GO and, being incompatible with water, form a spherical physical barrier, such that the pore size of the formed GAH is affected by the oil droplet size, which can be controlled by adjusting the amounts of n-hexane and stirring, and thus the GA pore size can be predetermined. A further advantage is that the removal of n-hexane as a template does not require an additional step, with n-hexane being removed along with water during dialysis and freeze-drying. However, n-hexane droplets are easily separated from water without the aid of an emulsifier, resulting in an uneven pore diameter in the GA upper and lower positions, thus the method needs to be further improved.An ice template process, called ice-segregation-induced self-assembly (ISISA),n-hexane systems were improved by introducing cyclohexane and n-butanol into GO suspensions and forming more stable aqueous cyclohexane–butanol/GO emulsions with the help of emulsifiers and dispersers. The prepared GO emulsions were then used to synthesize graphene hydrogels by a one-step hydrothermal method and the obtained hydrogels, named GAH-Ts, were then freeze-dried to form graphene aerogels, named GA-Ts. To obtain better graphene performance, residual oxygen atoms were removed from GA-Ts by calcination to form aerogels (FGAs) with improved graphene structures. By adjusting the ratio of aqueous to organic phases (cyclohexane and butanol), the pore structures of these hydrogels and aerogels, as well as the final FGA structure, were adjusted.In this study, unstable water/2.2.12SO4), cyclohexane, n-butanol, HCl, KMnO4, ethylenediamine , hydrogen peroxide (H2O2), phosphoric acid (H3PO4), and sodium dodecyl sulfate (SDS) were all analytical grade and purchased from Aladdin Chemistry Co. Ltd . All chemicals were used as received.Flake graphite (325 mesh) was supplied by Unigram Carbon Graphene Materials Co., Ltd, . Concentrated sulfuric acid (H2.2−1) was mixed with cyclohexane (50 ml) and n-butanol (37.5 ml) in a 150 ml beaker and then SDS (0.9 g) as emulsifier and 250 μl of EDA as reducing agent were added to the mixture. The mixture was dispersed by subsequent treatment with a high-speed homogenizer at 13 500 rpm for 5 min to obtain an off-white uniform emulsion. After the emulsion was degassed under vacuum, 10 g of the emulsion was fed into a hydrothermal kettle with a Teflon lining. Graphene hydrogels (GAH-Ts) were then obtained by keeping the hydrothermal kettle in a 150 °C muffle furnace for 6 h. GAH-T was carefully purged with 20% ethanol solution and freeze-dried for 48 h to form absolutely dry aerogels (GA-Ts). The product was further calcined at 600 °C for 5 h under N2 protection in a high temperature tube furnace such that GA-Ts were completely reduced to form the porous aerogels (FAGs).GO was prepared by a modified Hummers' process and the detailed steps can be found in the ESI.via the conventional hydrothermal method, in which only GO suspension (6 mg ml−1) and EDA (reducing agent) were used in the hydrothermal reaction without a soft template (cyclohexane and butanol). These obtained hydrogels and aerogels were termed GAH and GA, respectively.For comparison, graphene hydrogels and aerogels were also prepared 2.3α radiation (λ = 0.15418 nm) at a scan rate (2θ) of 5° min−1. Raman spectra of the samples were recorded using a laser Raman spectrometer  with an excitation wavelength of 532 nm. Surface element composition analysis was performed by using Thermo ESCALAB Model 250XI X-ray photoelectron spectroscopy , in which the binding energy was calibrated to the C-1s peak (284.8 eV) with a reference.37 Thermal gravimetric analysis (TGA) was performed using a thermal analyser  under a nitrogen flow of 20 ml min−1, at a heating rate of 10 °C min−1, and in a range from 25 to 800 °C. The surface roughness and morphology of GO were measured by using atomic force microscopy  in tapping mode and with 256 × 256-pixel resolution. Samples for AFM imaging were prepared by spin-coating the GO dispersion in water (1 × 10−5 mg ml−1) onto a freshly cleaved mica surface (200 rpm and 60 s) and allowing it to dry in air.The morphology of the graphene hydrogels (or aerogels) was characterized by using scanning electron microscopy . The chemical functional groups were characterized by using Fourier transform-infrared spectroscopy . The crystal structure was characterized by using X-ray diffraction , operating with Cu-K2.4n-hexane and xylene were selected as representatives of hydrocarbons and aromatic hydrocarbons, respectively, and a water/crude oil mixture as a model for spilled-oil wastewater.The solvent/oil adsorption capacity was measured according to the ASTM F726-99: Standard Test Method for Sorbent Performance of Adsorbents. Here, w1) and then placed in a 150 ml beaker, to which 100 ml of organic solvent (n-hexane or xylene) was added. The adsorbent was removed after 60 s of immersion, gently wiped with filter paper, and then immediately weighed (w2).Graphene aerogel samples (GA or GA-T) were weighed (Q) was calculated using the weight of the aerogels before (w1) and after adsorption (w2) asFor evaluation of adsorbent performance in spilled oil adsorption, 5 g of crude oil was mixed with 95 g of distilled water in a beaker to obtain the simulated petroleum-contaminated wastewater. The subsequent test method was the same as that for the organic solvent above. The adsorption capacity  and the density was very low (<10 mg cm−3), such that it could stand on a flower easily  and GO sheets were dispersed in water, the organic droplets formed a dispersed phase and the GO sheets adhered to the droplet surfaces (Under the effect of agitation by a high-speed disperser and with the emulsifier (SDS), the GO suspension was converted to a relatively stable oil-in-water emulsion with cyclohexane and surfaces . An optisurfaces . The grasurfaces .GA prepared without a soft template exhibited disordered porous structures . The gra3.2The GO transformation to the final aerogels was analyzed by using XRD, Raman spectroscopy, XPS, and TGA . The ID/IG values (ratio of the intensities of peak D to peak G) of GA, GA-T, and FGA were ∼1.00/1.05/1.06, respectively. The results showed that the chemical compositions of GA, GA-T, and FGA were very similar. However, the 3D structures (rGO sheet packing structures) of GA and GA-T were very different, which was attributed to the effects of the soft-template  of GO was 1.60/1. After the hydrothermal reaction, GO was converted to GA-T and C/O rose slightly to 2.91/1, indicating that GO reduction was not complete at the lower temperature . After cThe TGA curves of GA and GA-T are very different . In the 2 carbon (C–C 284.8 eV), carbon–nitrogen (C–N 285.4 eV), epoxy/hydroxyl (C–O 256.1 eV), and carbonyl  scale" fill="currentColor" stroke="none"><path d="M0 440 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z M0 280 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z"/></g></svg>O 288.2 eV)  . It was C, C–N, C–O, and CO bonds, but the contents of these various chemical bonds were quite different  and graphitic nitrogen  of GA-T and FGA for n-hexane and xylene were greatly improved compared to that of GA  and a richer pore structure. In organic adsorption experiments, the newly prepared graphene aerogels absorbed 90 times their weight of n-hexane and 139 times their weight of xylene. In the treatment of wastewater containing spilled oil, these graphene aerogels absorbed oil quickly and efficiently. These results show that graphene aerogels prepared with this improved soft-template method possess great potential for applications in wastewater treatment and organic matter separation and recycling from water.Graphene aerogels were synthesized by a simple one-step hydrothermal method, in which droplets of cyclohexane and There are no conflicts to declare.RA-010-C9RA10988A-s001


Background and Objectives: It is still unclear whether sarcoidosis is likely to be associated with tumors. In addition, the use of an immune checkpoint inhibitor has been reported to initiate the onset of sarcoidosis. We retrospectively analyzed tumor development before and after the diagnosis of sarcoidosis and examined the impact of having a history of tumors on the activity or the severity of sarcoidosis. Materials and Methods: We recruited 312 consecutive cases of sarcoidosis and analyzed the tumor development before and after the onset of sarcoidosis. Results: Among them, 25 cases were diagnosed with malignant tumor after diagnosis of sarcoidosis. In the analysis of the tumor-development group after diagnosis of sarcoidosis, both serum angiotensin I-converting enzyme and mediastinal lymph node size were significantly reduced at the time of malignant tumor diagnosis compared to at the onset of sarcoidosis, indicating that the decreasing activity of sarcoidosis may be partly associated with tumor development. Furthermore, we examined 34 cases having tumor history before the onset of sarcoidosis and analyzed the effect of tumor history on the severity of sarcoidosis. Cases with a malignant tumor in the past were older and had less complicated organs of sarcoidosis than cases without malignant tumors in the past. Oral corticosteroid therapy was administrated more frequently in cases without malignant tumors in the past, indicating that the history of a malignant tumor may influence the severity of sarcoidosis. Conclusion: These results indicate that tumor development may be partly associated with the activity or severity of sarcoidosis. Sarcoidosis is an unexplained systemic granulomatous disease. Occupational exposure to silica dust increases the risk of sarcoidosis in younger men . The risThe advent of ICI treatment is advancing research into the development of Th17 balance-mediated sarcoidosis. Furthermore, epidemiologically, research on silica exposure and the development of sarcoidosis is also progressing . Not onlInvestigations regarding sarcoidosis along with a history of malignant tumors have been limited. In this study, we retrospectively analyzed tumor development before and after the diagnosis of sarcoidosis and examined the impact of having a history of tumors on the activity or the severity of sarcoidosis.Consecutive patients diagnosed with sarcoidosis at Gunma University Hospital were recruited retrospectively. We applied the opt-out method to obtain consent for this study by declaring an information disclosure statement. The information disclosure statement was approved by the Gunma University Institutional Review Board. Patients were diagnosed with sarcoidosis by multiple respiratory specialists in this study.A total of 312 consecutive patients were enrolled in this study from those that were diagnosed with sarcoidosis at our hospital between January 2007 to December 2017. We excluded two cases in which malignant tumors and sarcoidosis were diagnosed at the same time to exclude sarcoid reactions. The diagnostic criteria of sarcoidosis were based on the statement of American Thoracic Society (ATS)/European Respiratory Society (ERS)/World Association of Sarcoidosis and Other Granulomatous Disorders (WASOG) ,29. BronStatistical information was collected from the medical record, including patient’s age, sex, height, weight at the time of diagnosis, hematological data on the course of treatment, imaging, allergy history, and preference history, which may potentially affect the course of sarcoidosis. Patient’s body mass index (BMI) was calculated using their height and weight at the time of analysis. Alcohol history defined patients’ daily consumption habits. A history of smoking was defined as an individual who currently smokes or has smoked previously more than 100 cigarettes in total . MethotrMost patients underwent contrast-free computed tomography (CT). It was difficult to measure hilar lymph nodes because the borderline between hilar lymph nodes and blood vessels was unclear. Therefore, we evaluated the most measurable lymph nodes, # 4 and 7, as lesions there can be measured by contrast-free CT. The measurements of lymph nodes were performed according to resist criteria 1.1 .p-value of <0.05 was considered statistically significant. All analyzes were done using SPSS version 24 .Cases with tumor development/history of a tumor (tumor group) and without tumor development/history of a tumor (non-tumor group) were divided into groups. For each factor of the tumor group and the non-tumor group, the number of cases and the ratio were calculated on the nominal and average scale, and the standard deviation was calculated on the order scale. Missing data were replaced by dummy variables. The Fisher’s accuracy test was used to analyze the nominal scale, while the numerical scale used the Mann–Whitney’s test. A First, we analyzed the relationship between the activity of sarcoidosis and the development of malignant tumors. Of the 312 sarcoidosis patients, 25 cases were later diagnosed with malignancies. In our study, lung cancer was the most common in seven cases, followed by malignant lymphoma in four cases . One casA statistical comparison of the background of patients did not show any significant differences between the tumor development group and the non-tumor development group .We analyzed the relationship between the activity of sarcoidosis and the development of malignant tumors. Serological data in patients with sarcoidosis having tumor development showed that serum angiotensin I-converting enzyme (ACE) was significantly reduced at the onset of malignant tumors compared to at the time of diagnosis of sarcoidosis. The major and minor diameters of the mediastinal lymph nodes were measured. The major and minor diameters of lymph node # 7 and the minor diameters of lymph node # 4 were significantly reduced at the diagnosis of malignant tumors compared to the size of initial lymphadenopathy at the onset of sarcoidosis . Lymph nNext, we analyzed the effect of tumor history on the development of sarcoidosis. Thirty-four cases had a history of having a malignant tumor in the past . One case had a history of two types of cancer in the past. Tumor types are shown in p = 0.009). Diabetes mellitus, smoking history, and alcohol intake were not significantly different between both groups.A statistical comparison of the background of patients was made between the past-tumor group and the non-past-tumor group . The agep = 0.061, 782 ± 351 vs. 1093 ± 889, p = 0.072, respectively). Although we recently reported that serum neuron-specific enolase levels could be a diagnosing and monitoring marker of sarcoidosis, there was no significant difference in serum neuron-specific enolase levels [p = 0.038) -2 receptor tended to be lower in past tumor groups  .n = 3) vs. 28.1% (n = 78), p = 0.013). Furthermore, oral corticosteroid therapy was administrated more frequently in the non-tumor group than in those cases that had a malignant tumor in the past (5.9% (n = 2) vs. 23.0% (n = 64), p = 0.024). There was no significant difference in the follow-up period and mortality.Furthermore, we analyzed the severity of sarcoidosis with or without tumor history. A comparison of the affected organs and the severity of sarcoidosis between the past tumor and the non-past-tumor group are shown in P. acnes or mycobacteria. Various symptoms are observed in systemic organs such as the eyes, lungs, heart, nerves, and skin, with oral corticosteroids being administrated as needed [Sarcoidosis is an unexplained systemic granulomatous disease. Sarcoidosis is characterized by the activation of the Th1 immune response induced by antigen exposure such as s needed ,21. In ts needed . Consists needed . AlthougIn this study, both the number of patients with systemic oral corticosteroids (OCSs) and the number of affected organs were significantly lower in the past-tumor group. The serum ACE and soluble IL-2 receptor, which reflects the activity of sarcoidosis, showed a decreasing trend in the past-tumor group. To our knowledge, this is first report describing the influence of malignant tumor history on the severity of sarcoidosis. As the number of diseased organs and frequency of OCSs treatment may reflect on the activity of sarcoidosis, immune status after tumor development is possibly associated with the severity of sarcoidosis after its onset.+ cells [+ T cells [There have been several reports on the relationship between sarcoidosis and tumors. Bonifaze et al. reported that skin tumors, blood tumors, upper gastrointestinal tumors, kidney tumors, liver tumors, and colon tumors are more frequently found with sarcoidosis . Askling+ cells ,37,38 (F T cells ,40. In tTh17-dominant CD4-positive T cells are abundant in BALF and granulomatous lesions of sarcoidosis ,42. Th17Interestingly, a transition from Th1- to Th2-dominant balance is thought to be particularly relevant in the development of advanced pulmonary sarcoidosis . The shiThere are several limitations to this study. Firstly, this was a retrospective study at a single facility, and as such, the number and the background of the cases were limited. A multicenter prospective study with a greater number of sarcoidosis cases will be required to clarify the relationship between sarcoidosis and tumor immunity. Furthermore, this was a cohort study of patients with sarcoidosis but not with tumors. Future analysis on a cohort of patients with tumors is also required to elucidate the relationship between the severity of sarcoidosis and tumor history.Serum ACE levels and mediastinal lymph node size were significantly reduced at subsequent tumor occurrence. Past-tumor history in patients with sarcoidosis was associated with a lower rate of OCS administration and fewer organ complications. Tumor development may be partly associated with the pathology or severity of sarcoidosis with regard to anti-tumor immunity.


Caenorhabditis elegans reproductive tract is orchestrated by the anchor cell (AC). This occurs in part through a cell invasion event that connects the uterine and vulval tissues. Several key transcription factors regulate AC invasion, such as EGL-43, HLH-2, and NHR-67. Specifically, these transcription factors function together to maintain the post-mitotic state of the AC, a requirement for AC invasion. Recently, a mechanistic connection has been made between loss of EGL-43 and AC cell-cycle entry. The current model states that EGL-43 represses LIN-12 (Notch) expression to prevent AC proliferation, suggesting that Notch signaling has mitogenic effects in the invasive AC. To reexamine the relationship between EGL-43 and LIN-12, we first designed and implemented a heterologous co-expression system called AIDHB that combines the auxin-inducible degron (AID) system of plants with a live cell-cycle sensor based on human DNA helicase B (DHB). After validating AIDHB using AID-tagged GFP, we sought to test it by using AID-tagged alleles of egl-43 and lin-12. Auxin-induced degradation of either EGL-43 or LIN-12 resulted in the expected AC phenotypes. Lastly, we seized the opportunity to pair AIDHB with RNAi to co-deplete LIN-12 and EGL-43, respectively, which revealed that LIN-12 is not required for AC proliferation following loss of EGL-43.Development of the  Summary: Using a dual depletion approach, i.e., combining auxin-inducible degradation with RNAi, we show that LIN-12 is not required for proliferation of the C. elegans anchor cell following EGL-43 loss. Caenorhabditis elegans larval development, a specialized uterine cell, the anchor cell (AC), breaches BM to contact the underlying vulval epithelium. This developmental event initiates the attachment of the uterus to the vulva, which later forms the reproductive tract of the animal. Several laboratories, including ours, have taken advantage of the animal's simple anatomy, transparent body, and genetic amenability to characterize molecular and cellular features of C. elegans AC invasion. Collectively, this has yielded important insights into the regulation of BM invasion in vivo (Cell invasion through basement membrane (BM) is essential for animal development, tissue inflammation, and cancer metastasis. During  in vivo .C. elegans, though its paralog, GLP-1, has a well-established role in promoting germline proliferation throughout post-embryonic development  system enables rapid degradation of proteins . It requradation B. Here, radation , which uradation .rpl-28 promoter and a self-cleaving T2A peptide that separates both transgenes (We co-expressed TIR1(F79G) with a small fragment of human DNA helicase B (DHB) fused to two copies of mKate2 (DHB::2xmKate2) A. Co-expansgenes . DHB::2xansgenes   larval stage when anchor cell (AC) invasion normally occurs  leads to the formation of two ACs and/or post-specification defects such as AC proliferation and failure to breach BM.The null phenotype of ethality  and L1 lethality . RNAi diinvasion . Specifiegl-43 null mutants. To explore this, we examined AC phenotypes using AIDHB with a new internally AID-tagged allele of egl-43 that targets the long and short isoforms of endogenous EGL-43 (lag-2 (LAG-2::P2A::H2B::mTurquoise2) (lam-2 (LAM-2::mNeonGreen) (egl-43(RNAi)   and lam-onGreen)  to labelonGreen) B-D. Of t43(RNAi) .lin-12 null mutant, both cells become ACs. To further test AIDHB, we combined it with an endogenous allele of lin-12 tagged at the C-terminus with mNeonGreen::AID  with and without 5-Ph-IAA. At the time of AC invasion, 30/30 auxin-treated animals and 26/29 control animals displayed the proliferative AC phenotype (n=196 versus 118). The higher total is expected for animals with two post-specified ACs that then entered the cell cycle and proliferated. Lastly, we confirmed the presence of LIN-12::mNeonGreen::AID in proliferating ACs of auxin controls after egl-43(RNAi)  , which legl-43 or lin-12 for degradation beginning in the L1 larval stage before AC specification. These experiments produced the expected AC phenotypes observed with either LIN-12 or EGL-43 depletion. Whereas loss of LIN-12 results in the formation of two ACs due to a defect in AC specification (egl-43(RNAi) results in ectopic LIN-12 expression in proliferating ACs, we did not observe localization in the nucleus, suggesting that ectopic LIN-12 may not be representative of active Notch signaling  can force the AC to proliferate . The NIC12 locus . The eme elegans  should fegl-43, and (iii) postulate that in the absence of EGL-43, LIN-12 expression is not necessary for AC proliferation. It is our hope that investigators will use AIDHB to interrogate the function of diverse proteins that may be required for cell-cycle-driven cellular behaviors.In summary, we (i) created a heterologous co-expression system called AIDHB, which we later paired with RNAi, (ii) generated a new AID-tagged allele of bmd284[rpl-28p::TIR1(F79G)::T2A::DHB::2xmKate2]; LG II: wy1514[egl-43::TagRFP-T::AID::egl-43]; LG III: ljf33[lin-12::mNeonGreen::AID] ::T2A::DHB::2xmKate2), pWZ192 (NotI-ccdB-SphI-DHB::2xmKate2) was double digested with NotI and SphI to excise ccdB and a PCR product representing rpl-28p::TIR1(F79G)::T2A was amplified from plasmid pCMH2123 using primers DQM1136 and DQM1137. pWZ259 was constructed by Gibson assembly (NEB) using the backbone from pWZ192 and the PCR product from pCMH2123. After sequence confirmation, pWZ259 was used as a repair template for insertion into the genome at a safe harbor site on chromosome I corresponding to the MosSCI insertion site ttTi4348 . pAP082 egl-43 locus was PCR amplified and purified using a PCR purification kit (Qiagen). 3 μl of 10 μM tracRNA (IDT) was incubated with 0.5 μl of 100 μM of a crRNA (IDT) targeting exon 6 of the egl-43 locus at 95°C for 5 min, followed by 25°C for 5 min. Following incubation, the mixture was incubated with 0.5 μl of Cas9 protein (IDT) at 37°C for 10 min. Repair template and a co-injection marker (pRF4) were added to the mixture to a final concentration of 200 ng/μl and 50 ng/μl, respectively. Young adult worms were transformed using standard microinjection techniques and progeny were genotyped for successful insertions . The egl-43(RNAi) feeding construct was published previously (Synchronized L1 larvae were plated on NGM plates containing 0.1 mM 5-Ph-IAA (MCE) and fed either OP50 or eviously , and it Images were collected using a custom-built spinning disk confocal microscope (Nobska Imaging), which was configured for automation with Metamorph software (Molecular Devices). This confocal consists of a Hamamatsu ORCA EM-CCD camera mounted on an upright Zeiss Axio Imager.A2 with a Borealis-modified Yokogawa CSU-10 spinning disk scanning unit and a Zeiss Plan-Apochromat 100x/1.4 oil DIC objective. Animals were anesthetized for imaging by picking them into a drop of M9 on a 5% agarose pad containing 7 mM sodium azide and secured with a coverslip.Acquired images were processed using ImageJ/Fiji . AID::GFP-values, and statistical tests used are specified.To determine the number of animals (N) needed per condition, a power analysis was performed . In each10.1242/biolopen.059668_sup1Supplementary informationClick here for additional data file.


The accurate modeling of frequency-dispersive materials is a challenging task, especially when a scheme with a transient nature is utilized, as it is the case of the finite-difference time-domain method. In this work, a novel implementation for the modeling of graphene-oriented dispersive materials via the piecewise linear recursive convolution scheme, is introduced, while the time-varying conductivity feature is, additionally, launched. The proposed algorithm is employed to design a reduced graphene-oxide antenna operating at  Since the discovery of isolated and stable graphene films , grapheny KMnO4) . The prey nature . In ordey nature . The proy nature , such asy nature .In particular, the research for efficient graphene-oxide-based sensors has, lately, escalated, owing to their remarkable properties ,24,25. Aetection ,28, althetection ,30 are p minutes ,32. Fina minutes ,34,35.Essentially, the aforementioned applications rely on the appropriate modeling of the involved media. In general, the characterization of graphene-based materials is achieved through their surface conductivity. Specifically, graphene itself is considered as a truly two-dimensional layer and its conductivity is evaluated via the Kubo formula . The latNevertheless, the analysis of modern devices is not feasible through analytical solutions; therefore, the surface conductivity features must be incorporated into powerful numerical algorithms, such as the finite-difference time-domain (FDTD) method . The latBased on these aspects, the main contribution of this work is the development of a novel scheme for the modeling of graphene-oriented conductive materials with frequency-dispersive response in the popular FDTD method. The transient nature of the latter is exploited to enable the simulation of planar materials with time-varying features and properties, which offers many possibilities for the design of complex devices for advanced applications. Moreover, the frequency dispersion is handled using an efficient PLRC algorithm that is appropriately adjusted to the surface conductivity function of the involved materials. Then, the developed implementation is applied to the design of an rGO antenna. An observation of the rGO transient response for sensing applications is performed, and some concepts for the enhancement of the device’s response are proposed by means of the time-varying rGO traits.The rest of the paper is organized as: Initially, the featured FDTD scheme for the modeling of the rGO is introduced in Two challenging tasks are encountered for the effective transient modeling of frequency-dependent conductive layers, such as the graphene, the GO, and the rGO. Initially, we consider that any planar material is characterized through its surface conductivity Here, we included the Dirac delta function, Equation , using tIt is important to mention that the Now, the second challenge appears, namely the robust implementation of the frequency dispersion for the conductive layer. In particular, the second-order accurate PLRC method is modified appropriately to model the surface current dispersion. First of all, the transition of Equation  into theJp(t)=σpσrGOEquation : (15)χi=It is obvious that the exponential term leads to the convenient computation of the recursion constant: Finally, the time-domain nature of the FDTD method enables the powerful feature of transient conductivity alteration, a feature that is absent from the majority of the commercial packages. Particularly, the conductivity of various planar materials can be adjusted externally, such as the electrostatic bias of graphene or the gas and humidity dependence of rGO . As a coThe previously described scheme is verified in terms of its accuracy via its thorough comparison with analytical solutions as well as with well-established a numerical software. Firstly, a finite rGO layer is considered and the transmission coefficient is evaluated at microwave frequencies. The conductivity of the rGO at this regime is almost constant; thus, the relaxation time is selected to be Additionally, the more complex setup of io Suite , where tFinally, the featured modeling methodology is compared to some well-established techniques in In this section, the newly introduced modeling of planar conductive materials is applied on the design of a patch antenna that includes an rGO patch. The transient nature of the method is exploited to perform an analysis on the effect of the planar material’s conductivity changes on the performance of the antenna. Moreover, it is shown how the conductivity variations can result on an enhancement of the antenna radiation characteristics, and, thus, can be used for sensing purposes.The proposed patch antenna has the dimensions depicted in er (PML) ,42,45.Firstly, we investigate the effect of rGO conductivity changes on the antenna response, both for constant and time-modulated values of Subsequently, taking avantage of the transient nature of the FDTD method, Equation . Both thOn the other hand, the variation of the modulation frequency leads to different antenna responses. Explicitly, the deepest resonant pick is reached for a modulation frequency of It becomes apparent that the optimal rGO antenna with the time-modulated conductivity exhibits a definite improvement of around Moreover, the total efficiency (gain) is calculated as As already demonstrated in A new implementation of the PLRC scheme for the accurate modeling of graphene-based materials with time-varying properties has been introduced in this paper. Using this scheme, the rGO has been efficiently modeled and integrated within a patch antenna operating at