Compared with graphene oxide ink/gel/resin based manners, laser‐induced graphene (LIG) is unique for facile and scalable assembly of 1D and 2D structures but still faces size and shape obstacles for constructing 3D macrostructures. In this work, a brand‐new LIG based additive manufacturing (LIG‐AM) protocol is developed to form bulk 3D graphene with freeform structures without introducing extra binders, templates, and catalysts. On the basis of selective laser sintering, LIG‐AM creatively irradiates polyimide (PI) powder‐bed for triggering both particle‐sintering and graphene‐converting processes layer‐by‐layer, which is unique for assembling varied types of graphene architectures including identical‐section, variable‐section, and graphene/PI hybrid structures. In addition to exploring combined graphitizing and fusing discipline, processing efficiency and assembling resolution of LIG‐AM are also balanceable through synergistic control of lasing power and powder‐feeding thickness. By further studying various process dependent properties, a LIG‐AM enabled aircraft‐wing section model is finally printed to comprehensively demonstrate its shiftable process, hybridizable structure, and multifunctional performance including force‐sensing, anti‐icing/deicing, and microwave shielding and absorption. In this work, a laser‐induced graphene based additive manufacturing (LIG‐AM) is developed to print bulk 3D graphene with freeform shapes such as identical‐section, variable‐section, and graphene/PI hybrid structures. By investigating particle‐sintering and graphene‐converting dual‐mechanism as well as process dependent properties, an aircraft‐wing section model is finally printed to demonstrate its shiftable process, hybridizable structure, and multifunctional performance. As one of the most promising nanotechnologies, graphene has attracted tremendous attention for assembling smart devices and structures with various geometries from 1D to 3D based on its hexagonal‐latticed molecules, embeddable/connectible/modifiable 2D nanostructures, as well as extraordinary electrical, mechanical, thermal, and optical properties. For example, graphene‐enabled 1D filaments, fibers and yarns, 2D thin films, papers and fabrics, and 3D foams, hydrogels, and millispheres have been explored extensively by endowing the assembled macrostructure with free‐standing skeletons, hierarchical microporosities, ultralow density, and interconnected high‐conductive networks for enabling critical components in wearable electronics, thin‐film batteries, artificial skins/muscles, biomimetic robots, self‐sensing composites, etc.Nowadays, 3D printing has become the most powerful technology for rapid designing, prototyping, and manufacturing structures, components, and products with multidimensional geometries and random complexities in almost all industries.1, 2 and Shi et al. respectively grew graphene on macroporous nickel and seashell templates to form freestanding foam‐like structures with the size of 20 × 20 × 2 mm3 after the etching process. Without introducing issues of high temperature and rigorous atmosphere control, Xu et al. and Yu et al. respectively applied hydrothermal and directional‐freezing approaches to cast and reduce GO suspension into self‐assembled graphene hydrogels or aerogels with the cylinder‐shaped volume around 9 cm3. Although bulk formation strategy could reach centimeter‐scaled architecture, it is unable to achieve arbitrary shapes due to the lack of programmable working path. As an advanced alternative, layer‐by‐layer assembly relies on computer‐aided process to create 3D graphene monoliths or patterns through different additive manufacturing manners including inkjet printing, extrusion printing, stereolithography, kirigami patterning, etc. With the feature size of micron scales, for example, Hensleigh et al. and Jiang et al. respectively applied photocurable GO resin and GO sol‐gel ink to print self‐standing 3D graphene structures with programmable microlattices. The layer‐by‐layer based methods could certainly promote the complexity and designability of structures; nevertheless, most of them rely on GO based inks, gels and resins as raw printing materials, still facing two major drawbacks. Firstly, the GO wet chemistry requires complex processing steps from suspension preparation, layer‐by‐layer stacking, freeze‐drying, to various chemical/physical treatments that dramatically prolong the working period and limit the size scalability. Correspondingly, the quality control in each step is tricky for tackling ink stability, printing rheology, GO reduction, structural shrinkage, etc. These further intensify the molding difficulty and part‐to‐part variations. Thus, a new 3D printing strategy toward swift assembly, freeform structuring, and ease of quality control is highly anticipated for construction of graphene macrostructures.To manufacture macroscopic graphene ensembles, two processing strategies have been applied majorly including bulk formation and layer‐by‐layer assembly. The former relies on non‐printing processes to either synthesize graphene on 3D porous metal templates or hydrothermally treat graphene oxide (GO) dispersion to form bulk foam structures. For example, by using chemical vapor deposition (CVD) protocol, Cheng et al.21 and By avoiding complexities of ink and catalyst preparations in traditional liquid‐assembling and CVD methods, LIG indeed has reflected its facile, low‐cost and scalable advantages for fabricating various single‐layered structures including graphene fibers, thin films, papers, and surface patterns, which have become core elements in various novel structures and devices, e.g., functional composites, supercapacitors, heaters, environmental, and biological devices. For assembly of macroscopic 3D objects, although several initial trials have emerged, LIG has still faced clear obstacles. For example, Luong et al. developed automated laminated object manufacturing (LOM) to fabricate centimeter‐sized graphene monoliths by repeatedly converting, fusing and milling Kapton films into stacked LIG layers with different shapes. With similar strategy, our group applied hot press to one‐step molding multiple pre‐impregnated LIG/epoxy papers to form 3D‐lamiated graphene composites with the size as large as 20 × 20 × 0.2 cm3. Although macroscopic dimensions are achievable, the above works require specific binding, milling, and/or 3D‐molding processes which have restricted the bulk formation of graphene with simple shapes. By raising a binder‐free strategy, Sha et al. applied sucrose and nickel particle mixtures as carbon source and catalyst to grow graphene layer‐by‐layer on the metal template. Although it is able to construct 1 × 1 × 0.8 cm3 sized foams, the post template‐etching treatment inevitably causes considerable volume shrinkage that confines the shaping accuracy. Other strategies such as LIG processed on preprinted 3D plastics, composites, and kirigami/origami patterns are also not able to construct freeform graphene structures. Thus, a novel LIG process toward the ease of freeform manufacturing of graphene 3D macrostructures is still not appeared.As a newly‐developed technique, laser‐induced graphene (LIG) has dramatically promoted the possibility of graphene assembly by introducing multiple unique characteristics including direct laser scribing, one‐step graphitic conversion, various compatible precursors, and customizable working routes.39 By Following this line of thought, in this work we developed a brand‐new LIG based additive manufacturing (LIG‐AM) protocol to form bulk graphene structures with freeform geometric shapes. Without introducing extra binder, template and catalyst, polyimide (PI) powder was applied as the only raw materials for both particle‐sintering and graphene‐converting purposes to construct 3D solids. Based on computer‐aided laser‐irradiation scribing to selectively process each layer with individually programmed slicing pattern, LIG‐AM is able to analogize the commercial selective laser sintering (SLS) technology to form varied types of bulk architectures including identical‐section, variable‐section, and even unique graphene/PI hybrid structures with random shape complexities. In addition to investigate the combined graphitizing and fusing discipline toward macroscopic assembly, systematic work has also been done to balance the processing efficiency and assembling resolution that are controlled interactively by laser power and powder‐feeding thickness. By further studying the tunable electrical, mechanical, piezoresistive, liquid‐sensing, and joule‐heating properties as well as resin‐integration reinforced strength and stretchability, a LIG‐AM enabled aircraft‐wing section model was specifically designed to comprehensively demonstrate its shiftable printing strategy, hybridizable structure, and multifunctional performance including multidirectional force‐sensing, anti‐icing/deicing, and microwave shielding and absorption.22.1Figure2 laser system serves as the power source to selectively solidify and graphitize the very top layer of powder bed with a preprogrammed slicing pattern before feeding a new layer to finally construct a printed 3D object. To specify, a thin layer of PI powder is firstly fed onto the working plate using a spacer to control the height of powder bed. Then, the laser irradiates and scribes the powder bed to selectively sinter and convert PI powder into a continuous graphene structure with specific surface pattern. Subsequent to the first LIG layer formation, a new powder layer is fed on top and followed by another irradiation process. The feeding and irradiating steps then repeat multiple times until processing all layers to present the entire structure after removing unprocessed powders.Inspired by the commercial SLS manufacturing to print 3D metallic, ceramic, and polymeric objects through layer‐by‐layer fusion of powders, the idea of using compatible powder‐based LIG raw materials is the key for constructing graphene macrostructures with freeform geometric shapes. In this work, commercially available thermoplastic PI powders with the size of ≈10 µm was selected as the ideal raw materials by competing other LIG‐compatible powders considering irradiation energy, degree of graphitization, and shape accuracy , the macroscopic LIG structure is no longer mechanically stable with considerable voids observable in the fibrous microstructure, indicating excessive ablation effect to destroy the LIG framework. More detailed scanning electron microscopy (SEM) images are displayed in Figure Based on mechanism analysis, it is reasonable to believe that the level of irradiation energy is the key to control the dual functionality of PI for both particle sintering and graphene conversion. Figure e layers.54 By  Raman spectroscopy in Figure IG/ID) and 2D to G (I2D/IG) peaks follow similar trend that the sample processed around 1 W shows both the highest IG/ID (1.37) and I2D/IG (0.41), indicating the best graphene quality with the fewest number of layers. Additionally, X‐ray diffraction (XRD) in Figure  Similarly, X‐ray photoelectron spectroscopy (XPS) in Figure In addition to optical and SEM analyses, a series of structural characterizations were further explored to support the process dependent graphene conversion. In contrast to 0.625 and 3.75 W processed samples, transmission electron microscopy (TEM) images in Figure 41 Ram2.2FigureL1) shows the fused PI structure, indicating that the transported energy is not high enough for graphitization. Accordingly, a dense LIG layer (L2) appears right above the PI layer indicating a proper energy for graphene conversion. It is also observed that fiber‐like structures are protruded out of the LIG surface to form a rather loose layer (L3), which is explained by the quick liberation of pyrolysis gases during the photothermal reaction to distort the LIG structure. According to soft nature of L3, it will merge with the new feeding layer to receive another irradiation treatment  should be critically controlled. Intuitively, LF should be smaller than L2 to ensure all new powders to be converted to graphene; otherwise, L1 and even unfused PI would appear to destroy the structural integrity. Along with the maximum thickness, the feeding layer LF also cannot be too small for causing too much thermal impact. Figure LF from 10 to 90 µm. Clearly, only when LF lies between ≈40 and ≈70 µm, uniform multi‐layered LIG can be formed; when LF < 40 µm, LIG shows both macroscopic and microscopic defects indicating strong ablation effect; when LF > 80 µm, clear delamination has been observed indicating the loss of interfacial bonding. Thus, by applying 50 µm as the proper LF, Figure LF for every processible condition of laser. Obviously, weaker/stronger laser power is compatible with thinner/thicker LF, indicating lower/higher processing efficiency but better/worse assembling resolution. To further support this, Figure LF while the second one chose precise process by using 1 W laser and 50 µm LF. Although the second one took longer time (4 h and 240 layers vs 1 h and 60 layers), it shows a much smoother surface with better graphitized structure.Thus, to construct full graphene objects through multi‐layer processing, the thickness of every powder‐feeding layer  and compressive strength (7.9±0.5 MPa) and LIG/Ecoflex composites with improved fracture strain (230%). Similarly, LIG‐AM structures reflect an ascending trend on both of their in‐plane (22±1.3 to 46±2.1 S m−1) and through‐plane electric conductivity (2±0.1 to 3.2±0.12 S m−1) as the increase of processing power  when subjected to cycled deformations along both tensile and compression directions, exhibiting the linear strain‐induced resistance change with high reproducibility. As the increase of laser power, the increasing trend of both tensile‐deformation determined (3.1±0.1 to 6.4±0.6) and compressive‐deformation determined gauge factor (78±2.4 to 96±2.3) indicates that higher power processed LIG‐AM structures are more sensitive to external disruptions, which coincides with previous finding that high irradiation could induce large numbers of microfissures/cracks in the conducting network. In addition to deformation sensing, the porous networks also endows the graphene structures with liquid‐sensing capability. Figure R/R0 = 1.8%) to 8 µL (maximum ΔR/R0 = 30.5%). The inset plot reveals the resistance‐change cycles with processes of fast rise and gradual recovery, relating to dynamic change of liquid rinsing and evaporation. By virtue of electrothermal characteristics, Figure  Other functional properties, including hydrophobicity, oil absorption, supercapacitance, and microwave shielding and absorption, were also evaluated and discussed in Figures By understanding the printing discipline for constructing freeform structures, processing dependent properties of LIG‐AM enabled 3D graphene were further investigated. By using a high‐precision balance for density determination, r Figure , confirmr Figure . The insAttributed to freeform 3D construction, selective graphene formation, as well as tunable multifunctionality, the proposed LIG‐AM could become a rather powerful 3D printing protocol for designing, developing and manufacturing various structures, components, and devices for meeting the need of multi‐scenario applications. Figure FigureR/R0) measured by horizontal electrodes varies accordingly depending on the force direction. Specifically, forces from horizontal directions (0° and 180°) introduce minimum ΔR/R0 around 58% and forces from vertical directions (90° and 270°) introduce maximum ΔR/R0 around 105%. Other forces introduce moderate ΔR/R0 depending on the specific angle. Secondly, the smart leading edge  has reached 44.7 dB for the smart skin with 2.5 mm thickness, which exceeds the SETotal of commercial EMI shielding materials (20 dB) and other reported LIG (24.8–32 dB). Besides, the smart skin with a nickel treatment has shown much improved reflection losses (RL) with −15.3 dB absorption peak at 9.1 and 2.5 GHz bandwidth (from 8.0 to 10.5 GHz) under −10 dB, representing that 90% of the electromagnetic wave energy has been absorbed. Along with detailed information in Figure Total and RL has proved the possibility of LIG‐AM technology for EMI shielding and stealth.To demonstrate a specific example for fully exhibiting various unique advantages of LIG‐AM, an aircraft‐wing section model with integrated smart graphene structures was designed, printed and tested. As displayed in e Figure  was desie Figure  with tun3In summary, we demonstrated a brand‐new LIG‐AM protocol to form bulk 3D graphene with freeform structures without introducing extra binders, templates and catalysts. By triggering both particle‐sintering and graphene‐converting processes layer‐by‐layer, LIG‐AM creatively irradiates PI powder‐bed to assemble varied types of graphene architectures including identical‐section, variable‐section, and unique graphene/PI hybrid structures with random shape complexities. In addition, processing efficiency and assembling resolution of LIG‐AM are also balanceable through synergistic control of lasing power and powder‐feeding thickness. By virtue of outstanding processing advantages, a series of LIG‐AM enabled smart devices/structures, especially the LIG‐AM enabled aircraft‐wing section model, are successfully printed and demonstrated their multifunctional properties to meet the need of potential multiscenario applications.42 laser with wavelength of 10.6 µm and a controlled XYZ‐movement was used for irradiating polymer powders . Typical values of laser power were varied from 0.25 to 5 W with the pulses per inch (PPI) fixed at 500 and the scribing speed at 50 mm s−1. The powder spreader  with the resolution of 10 µm was employed to control each layer thickness. After layer‐by‐layer printing, diverse LIG‐AM enabled graphene architectures were performed under ambient conditions.A laser platform  equipped with a CO−1 under argon atmosphere. Differential scanning calorimetry (DSC) tests were evaluated with TA Q5000IR instrument at a ramping rate of 10 °C min between 40 and 400 °C. Specific surface area was measured with a Quantachrome 2000e BET surface analyzer.SEM was performed with JEOL JSM 7001F at 10 kV to examine the morphologies of various samples. TEM images were taken with a high‐resolution JEOL 2100F operated at 200 kV. Raman spectra were collected with a Horiba HR800 Raman microscope using a 532 nm laser with a power of 5 mW. XRD characterization was performed on Rigaku D/max 2550. XPS measurements were performed with a Thermo ESCALAB 250XI instrument equipped with the energy standard of C1s = 284.8 eV binding energy for charge correction. Thermogravimetric analysis (TGA) was carried out with a NETZSCH STA 449F3 instrument at 10 °C minMechanical properties were evaluated through a dynamic mechanical analyzer  and an E44.104 tensile machine (MTS Systems). Electrical conductivity was quantified using a Keithley 2450 Sourcemeter. Joule heating temperature was monitored by a FLIR camera with the input DC current provided by a Keithley 2260 power source. Supercapacitive properties were collected by CHI608E electrochemical workstation. Microwave shielding/absorption properties were measured by Vector Network Analyzer  at room temperature. Water contact angles were obtained by water drop angle tester (OCA15 Pro).3) were carried out by a mechanical–electrical coupling method at a rate of 2 mm min−1 and the gauge factor (GF) was calculated by the equation GF = (R − R0)/(R0 × ε), where R, R0, and ε are instantaneous resistance, initial resistance, and strain, respectively. Liquid‐sensing tests of 3D graphene sensors (10 × 10 × 2 mm3) were applied by monitoring the real‐time resistance change through dripping microvolume acetone from 0.5 to 8 µL. Electrothermal tests of 3D graphene heaters (10 × 10 × 2 mm3) were performed by loading electrical input powers ranging from 0.25 to 2 W and deicing test of smart leading edge was conducted by recording the volume change of ice cube with the input power of 1.2 W. 3D graphene based‐supercapacitors (10 × 10 × 2 mm3) were fabricated using a three‐electrode system with 1 m H2SO4 electrolyte, where the counter electrode was a Pt wire and the reference electrode was Ag/AgCl. Adsorption capacity tests of 3D graphene removers for organic solvents and oil were calculated by the ratio between the maximum adsorbed adsorbate quantity M and the original mass m using the balance analyzer. High‐temperature resistance test of smart skin was implemented by observing the structural integrity under an alcohol lamp burning. Ni decorated LIG (LIG/Ni) was obtained by pulse electrodeposition in a standard three‐electrode system controlling the deposition time for 10 min. Electromagnetic waves shielding and absorbing tests of LIG and LIG/Ni microwave absorbers were carried out by waveguide method and coaxial wire method for S‐parameters and EM parameters (complex permittivity and complex permeability), respectively.By integrating Ecoflex0030  or epoxy resin , LIG/Ecoflex and LIG/epoxy composites were prepared respectively after the curing process at 60 °C for 6 h. Mechanical‐sensing tests of smart composites (30 × 10 × 2 mmThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.Supporting Video 1Click here for additional data file.Supporting Video 2Click here for additional data file.Supporting Video 3Click here for additional data file.Supporting Video 4Click here for additional data file.Supporting Video 5Click here for additional data file.


Ab Initio Molecular Dynamics (AIMD) simulations to study graphene island alignment evolution at two temperatures. The results shed light on our experimental observations and those reported in the literature and point to the effectiveness of controlled thermal cycling in producing high quality graphene sheets on transition metal catalyst surfaces.Repeated thermal cycling by using an organic precursor is shown to be a successful technique for growing graphene on metal substrates. Having control on this process is of vital importance in producing large areas of high quality graphene with well-ordered surface characteristics, which leads us to investigate the effect of temperature on the microscopic mechanisms behind this process. Apart from being an important factor in the dissociation of the organic precursor and promoting the reactions taking place on the surface of the catalyst, temperature also plays a major role in the structure of the catalyst surface. First, we used eight thermal cycles to successfully grow graphene on the surface of Cu (111). Then, we employed  Repeated thermal cycling by using an organic precursor is shown to be a successful technique for growing graphene on metal substrates. Of these methods, mechanical exfoliation, although capable of producing high quality graphene on small scales, is the least scalable and appropriate method for electronics applications and mass production where larger areas of graphene are required, whereas the CVD method, typically catalytic, has emerged as the most powerful and dominating method in the industry.8,9Graphene, a single atomic layer of graphite and the building block of carbon nanotubes, has been proven to possess many unusual properties such as extraordinary chemical and optical characteristics, tensile strength and thermal and electrical conductivity that find applications in many areas such as ultrathin membranes, plasmonics, high-end composite materials and electronics.3 Growth mechanism on these substrates are shown to be highly dependent on temperature.10 Among these substrates, copper has become one of the most widely used catalysts for the synthesis of graphene using the CVD method. Large high-quality graphene films were grown on copper, which exhibit transferability and unique plasmonic characteristics.2,11,12 Low carbon solubility13,14 and weak Cu–C interactions15,16 along with low surface diffusion energy barrier of C atoms on Cu17 all contribute to surface-driven mechanisms present in the nucleation and growth of graphene domains.3,18–20 These factors may also explain why Cu produces predominantly single layer graphene.3,21 Low surface diffusion energy barrier contributes to highly mobilized carbon atoms that can travel freely on the Cu surface and attach to existing graphene islands. Note that surface Cu atoms have high mobility at high temperatures which can facilitate the mobility of C atoms and graphene islands at such temperatures and lead to defect healing.20,21 Therefore, surface morphology of Cu substrate at elevated growth temperatures is an important issue.Ni, Ir and Cu are three of the most common transition metals that are used as catalyst substrate for graphene growth.et al. investigated how the type of organic precursor may have effect on graphene growth and dictate its structure.22 Ethylene, acetylene and methane are organic precursors that are widely used as the source gas of which methane is known to require higher temperatures to decompose on the Cu surface. Ethylene23 and acetylene,24 on the other hand, has shown to produce continuous graphene coverage at considerably lower temperatures. In other studies, environmental conditions such as temperature, pressure and hydrogen that may have effect on graphene growth were investigated.10,25,26 The copper catalyst surface has also been subject to many studies as the surface plays a major role in the epitaxial graphene growth and achieving a perfectly clean surface prior to growth is highly desirable.27–30 Yu et al. reported their simple approach of preparing the copper surface by growing it on single-crystal sapphire. The resulting copper layer is reported to be extremely flat, chemically clean and can be easily peeled off to use for graphene growth. The graphene grown on such a copper surface exhibited very large domain sizes.31Reported research aimed at improving the quality of graphene grown over Cu surface were focused on several aspects of this process. As an example, Wassei 3,18 In the surface-driven mechanism of growth on copper, it is shown that graphene islands grow when temperature is high enough to facilitate the diffusion of islands. At such temperatures some islands may ‘stitch’ together to form larger graphene domains.32–36 Consequently, the overarching hypothesis is that defects observed in the graphene overlayer arise from the misalignment of graphene islands constituting a domain; various islands nucleate and grow with different rotational alignments across the catalytic substrate. Several studies have been aimed at investigating what dictates these island orientations. Studies in this area have concluded that domain shapes and sizes themselves are controlled by growth conditions such as temperature.37,38 Upon closer examination, two possible routes have been identified that would lead to the transition of islands to domains; either the islands migrate and coalesce into larger islands, or islands progressively grow in size and form domains through addition of C atoms.35,39,40 In the former, with the foregoing discussion, temperature would have great influence on rotation and migration of islands and domain growth, as well as on nucleation density, as evidenced earlier.38,41 While in the latter case, surface diffusion of C adatoms would be most likely to have a greater impact on growth of domains. Thus, the identification of the prevailing case will help in optimizing growth parameters. Although this may be a challenge since many factors are involved in the competition including carbon source gas decomposition rate, nucleation rate, surface diffusion and growth rate from carbon adatom addition.High resolution electron microscopy provides valuable insights into defects and domain (grain) boundaries in graphene produced from the CVD method. These domain boundaries have been subject to intense studies and may hold the key to understanding the growth mechanism. Two main mechanisms have been proposed for the growth of graphene on transition metal surfaces: segregation growth , and surface growth (Cu).ab initio methods, we investigate how temperature and time cycles  may affect the alignment between the Cu lattice and a pre-existing graphene fragment. Our focus is to study the interplay of the Cu substrate and the graphene overlayer in conjunction with temperature at early stages of growth.In this paper, our particular interest is to understand the role of nucleation and growth parameters, such as temperature, on graphene domain orientation. In this regard, the catalyst surface may have a more active role than thought in the determination of domain boundaries. In the present work, using experimental and −10 mbar) equipped with a variable temperature scanning tunneling microscope  chamber, and a separate preparation chamber, which includes an ion sputter gun for sample cleaning, a LEED (Low Energy Electron Diffraction) system, and a directional gas doser. The Cu (111) single crystal  was polished on one side and mounted on a standard tantalum sample plate from Omicron with an 8 mm diameter hole on its rear end to facilitate electron bombardment heating. Cu (111) was cleaned by 8 cycles of sputtering  and annealing . Temperature above 600 °C (873 K) was monitored by a pyrometer . Successful cleaning of the sample was confirmed by STM and LEED. Large continuous graphene islands were produced on Cu (111) surface through exposing the clean Cu (111) substrate to 1 × 10−5 mbar C2H4 at room temperature for 5 minutes, with subsequent annealing at 1000 °C (1273 K) for 5 minutes. This cycle was repeated 8 times.64 Successful graphene preparation was confirmed by scanning tunneling microscopy (STM). All STM images were collected at room temperature by using electrochemically etched W homemade tips with a sample bias of 0.7 V and a constant tunneling current of 0.4 nA. All the post annealing images were generally acquired 2 hours after the experiment to ensure that the sample had returned to room temperature. To ensure that the images shown here are consistent across the crystal face, for any given surface condition, images were collected at multiple locations  on the surface. All the STM data were processed using WSxM 5.0 software.42All the experiments were carried out in an ultrahigh vacuum system (base pressure approximately 1 × 1043–45 with Perdew–Burke–Ernzerhof (PBE) exchange-correlation functional and in the generalized gradient approximation (GGA) formulation.46 We treated the electron-ion core interactions using the projector-augmented wave (PAW) pseudopotentials45,47 available in the VASP database. We constructed a four atomic layer thick p (7 × 7) Cu slab in Materials Studio48 by cleaving along the (111) plane. A vacuum of 20 Å thickness was applied on the slab along the z direction to prevent interactions of the slab with its upper/lower periodic images. The bottom layer was kept fixed and the remaining three layers were allowed to relax. Conjugate gradient algorithm with a Gaussian smearing width of 0.05 eV was used for all relaxations. Convergence tests were carried out, and a 1 × 1 × 1 k-point mesh with a plane wave energy cutoff of 400 eV was selected for the sampling of the Brillouin zone. The graphene fragments used in this study consist of three neighboring hexagonal rings. Graphene fragments were initially taken from Materials Studio structures database and then relaxed. Relaxed structures were then used on the Cu surface. Ab initio molecular dynamics (AIMD) were performed using VASP with the foregoing computational details. AIMD simulations were carried out in the NVT ensemble with a 1 fs time step and at two temperatures of 900 K and 1200 K.We performed our density functional theory (DFT) calculations using the Vienna Ab initio Simulation Package49–54 Gao et al. showed that epitaxial growth of graphene on Cu (111) by thermal decomposition of ethylene (C2H4) at 1000 °C results in formation of single-layer graphene platelets and sheets.33 Similarly, to produce large continuous graphene islands on Cu (111) surface, we undertook a controlled thermal cycling approach using the organic precursor, (C2H4). Successful preparation of large continuous graphene film over Cu (111) surface was confirmed by STM as shown in I/dV) image) were captured after the 7th cycle of thermal annealing. As shown in these two images, graphene islands and C atoms are both observed in one single image before final controlled annealing. Note that graphene islands look smoother and darker in differential conductance images. However, as shown in I/dV) image), after final cycle of thermal dissociation of C2H4 (8th cycle), complete graphene coverage over Cu (111) was observed. The domain boundaries are visible in these STM images. Note that the entire differential conductance image has a uniform color, which indicates that either all the terraces are Cu (111) or graphene/Cu (111) and 49 As shown in One promising approach to prepare large graphene sheets with low defects on metal substrates is by thermal dissociation of an organic precursor on the metal surface.6 rings and graphene islands on different adsorption sites has been studied elsewhere using DFT.4,55 Compared to these earlier studies, the cell used in this work is larger to accommodate end effects, and a GGA functional used (PBE) which is more accurate than the local density approximation (LDA). Using DFT optimizations, we allowed a C6 ring to relax onto the Cu (111) surface. We tested three main possibilities; (a) the ring surrounds a top Cu atom and the C atoms of the ring occupy fcc and hcp hollow sites (fcc–hcp), (b) the ring surrounds an hcp site and the C atoms occupy fcc and top sites (fcc-top), and (c) the ring surrounds an fcc site and the C atoms occupy hcp and top sites (hcp-top), alternately. Results, illustrated in Fig. S1 of ESI,Next we focused on a first-principles analysis of the graphene growth mechanisms. The adsorption of C13C atoms constituting 3 rings, to relax on the Cu surface. We studied 7 different orientations of the fragment with respect to the Cu lattice. These orientations were from 0° to 60° with 10° increments. Results are illustrated in Fig. S5 and S6 of ESI.38,56,57 Here, our focus was to determine the most favorable orientations of a graphene fragment rather than a graphene sheet on the Cu surface. This represents the initial nucleation stages observed experimentally on the Cu (111) terraces as discussed in relation to In addition, we also allowed a pre-relaxed graphene fragment, composed of 33 as well as our own. Prior to running AIMD simulations, the fragment was relaxed onto the Cu surface with an initial misalignment of 24° with respect to the Cu (111) lattice as shown in In the experimental part of this work, the technique of low and high temperature cycles was applied to grow the graphene sheet. Reportedly, graphene coverage can be controlled by the number of thermal cycles. Here we emulated several aspects of this process by using AIMD simulations to understand the effect of these cycles on the stability of the graphene fragments with respect to the underlying substrate. First, we studied the motion and rotation of an isolated fragment at two temperatures of 900 K and 1200 K. The selection of temperatures was largely dictated by other experimental studiesTo understand the effect that neighboring graphene fragments may have on a single fragment, we added a second fragment to the vicinity of the isolated fragment of the previous section. We note that the orientation of the second fragment with the first fragment was not set arbitrarily. We tried four different orientations of 0°, 30°, 60° and 90° between the two fragments. The most energetically favorable was the 60° orientation. This structure resulted in 24° misalignment of both fragments with the Cu (111) lattice and we started the AIMD simulations with this initial structure. †To study the effect of proximity on additional graphene islands in the vicinity of the first one, we placed the second graphene fragment of the previous section, closer (by 0.2 nm) to the initial fragment and relaxed the structure before running AIMD simulations. Results are shown in 58–60 Fe61 and Ni61,62 surface atoms in graphene growth was clearly observed in this part of the work. Both fragments were able to bond to two surface Cu atoms situated between the two fragments and by doing so, raise those atoms from the surface  surface when only small graphene islands have been formed. Hereafter, the islands may align with respect to one another and form a uniform graphene sheet or each may grow about their own specific alignment with the Cu lattice and produce various domains. We closely examined the effect that temperature may have in choosing either of these paths. We found that higher temperatures generally lead to the alignment of an island on the Cu surface. An isolated graphene pallet at high temperatures (above 900 K) is able to correct an initial misalignment and align itself perfectly with the Cu surface. This is attributed to the higher mobility of the Cu catalyst surface layer which allows graphene to rotate freely. As growth proceeds at lower temperatures and the number of graphene fragments increases, metal-bridging of surface Cu atoms which was observed unifies graphene islands. Therefore, increased growth rate at lower temperature combined with increased mobility at higher temperatures lead to production of large high quality graphene films with the perfect alignment of graphene islands with Cu lattice.There are no conflicts to declare.RA-008-C8RA05478A-s001


The Dark Triad has been found to be associated with malevolent creativity (MC) in terms of trait level, and its underlying mechanism remains unclear. Based on the cognitive–affective processing system theory and the existing studies, the current study aimed to explore the internal mechanism between the Dark Triad and MC behavioral tendencies/performance. The results revealed that the Dark Triad is positively related to MC behavioral tendencies through trait aggression and general creativity behavioral tendencies. Regarding MC performance, the Dark Triad is positively related to the originality of malevolent ideas through MC behavioral tendencies, but this effect is only significant at low-to-medium levels of moral identity. In line with moral identity theory, a higher moral identity may prevent individuals from acting immorally due to their desire to maintain their moral image, which may further suppress malevolent idea generation. Therefore, cultivating moral identity may be an effective approach to weaken the Dark Triad–MC performance association. Creativity is the ability to produce novel and useful work . TraditiGiven that creativity is, by nature, unexpected and rule-violating, the products of MC can be unpredictable and extremely dangerous . Thus, tThe cognitive–affective processing system (CAPS) has proposed several types of mediating units  that are activated by situational stimuli and further determine behaviors . In addiAccording to the CAPS, the mediating units interact and produce different behavioral responses . IndividThe current study aimed to answer the following question: How is the Dark Triad related to MC at the trait and behavior levels? Based on the above-mentioned studies, we hypothesized that (H1) behavioral tendencies of general creativity and aggression are significant mediators between the Dark Triad and MC behavioral tendencies and that (H2) MC behavioral tendencies mediate the relationship between the Dark Triad and MC performance , while moral identity moderates the MC behavioral tendencies–performance path.A total of 217 Chinese college participants  were recruited in the study. The recruiting poster was sent to the open college WeChat group, and anyone interested in the experiment could sign up through a QR code. Three participants were removed due to incomplete data. None of participants had a history of mental or neurological illness. The prior power analysis revealed that the sufficient sample size was 166 to obtain reliable results . It was used to measure MC performance by requiring participants to solve open-ended realistic problems in creative and malevolent ways . MCT conThe MCBS was used to measure malevolent creative behavioral tendencies , MCT, RIBS, BPAQ, DD12, and MIM .Before performing the MCT, participants were explicitly informed that all situations were fictitious in this experiment, and that malevolent ideas generated in this experiment were not related to their personality or level of morality. After completing the experiment, participants were asked whether they were all right (none of them claimed to feel unwell). Besides this, participants were also informed that psychological consultation was available and free for them if they needed it.r = .30, p < .001), aggression , and Dark Triad ; MCT harmfulness was positively correlated with MC behavioral tendencies  and creativity behavioral tendencies . Moreover, the correlation coefficient between the Dark Triad and MC behavioral tendencies was significantly higher than that between the Dark Triad and creativity behavioral tendencies  and other variables see . The resZ scores. Hierarchical regression analysis using gender , age, and other scale scores as independent variables was performed on MCT originality. The results of Step 3 revealed that MC behavioral tendencies  positively predicted MCT originality , age, and other scale scores as independent variables was performed on MCT harmfulness. The results of Step 3 revealed that none of these variables significantly predicted MCT harmfulness after controlling for age and gender see . The varb = .57, t = 10.24, p < .001; b = .40, t = 7.28, p < .001; b = .10; 95% confidence interval, 95%CI = ) and creative behavioral tendencies  significantly mediated the relation between the Dark Triad and MC behavioral tendencies. The VIF ranged from 1.08 to 1.26, which meant that multicollinearity did not seriously affect the result.Parallel multiple mediation analysis using the Dark Triad as an independent variable and BPAQ and RIBS as mediators was performed on MC behavioral tendencies see . PROCESSb = .16, t = 2.30, p = .022; b = −.03, t = −.41, p = .684; b = −.09; 95%CI = ; b = −.16; 95%CI = ). Specifically, MC behavioral tendencies significantly mediated the relation between the Dark Triad and MCT originality when moral identity was one SD lower than the mean  or equal to the mean ; the mediation effect of MC behavioral tendencies was insignificant when moral identity was one SD higher than the mean . The variance inflation factors (VIFs) ranged from 1.07 to 1.58, which meant that multicollinearity did not seriously affect these results.A moderated mediation model using the Dark Triad as an independent variable, MC behavioral tendencies as a mediator, moral identity  as a moderator, and MCT originality as a dependent variable was established See . Model 1The current study investigated the internal mechanism between the Dark Triad and MC . We observed that MC behavioral tendencies were positively correlated with the Dark Triad, aggression, and general creative behavioral tendencies. Moreover, the correlation coefficient between the Dark Triad and MC behavioral tendencies was higher than that between the Dark Triad and general creative behavioral tendencies. The results further revealed that the Dark Triad was related to MC behavioral tendencies through both aggression and general creative behavioral tendencies. Regarding MC performance, only MC behavioral tendencies positively predicted MCT originality after controlling for age and gender. Mediation analyses revealed that MC behavioral tendencies significantly mediated the relationship between the Dark Triad and MCT originality; the mediation effect of MC behavioral tendencies was only significant in low-to-medium levels of moral identity.The results showed that the Dark Triad was significantly correlated with general creativity and MC behavioral tendencies, which was consistent with the previous findings . Both thParallel multiple mediation analysis found that the relationship between the Dark Triad and MC behavioral tendencies was mediated by both aggression and general creative behavioral tendencies. This result partly supports our hypothesis . Researchers have demonstrated that individuals with higher aggression tend to generate more harmful ideas in general creativity tasks . A higheAt the behavioral level, MC behavioral tendencies positively predicted MCT originality (after the effects of age and gender were controlled for) and fully mediated the relationship between the Dark Triad and MCT originality. In addition, moral identity significantly moderated the path of MC behavioral tendencies–MCT originality in the mediation model. This supports our hypothesis . The Dark Triad refers to self-centeredness, manipulation, and callousness . These ‘In summary, the findings of the current study support CAPS and moral identity theory. Individuals with higher levels of Dark Triad personality traits tend to have higher aggression and general creativity behavioral tendencies, which then further cultivate their MC behavioral tendencies. At the behavioral level, MC behavioral tendencies may be closer to the originality than the harmfulness of malevolent ideation. The Dark Triad promotes MCT originality by fostering MC behavioral tendencies, but this mediation effect is only significant with low-to-medium moral identity. Based on the above-mentioned results, cultivating moral identity may be an effective way to prevent MC performance. Beyond the laboratory, MC performance in natural settings should be further investigated in the future. However, several limitations should be mentioned: (1) almost 76% of participants were female, so future studies should include more male participants; (2) only explicit aggression was explored, but implicit aggression could also be an effective source of malevolent creation; and (3) in addition to moral identity, moral reasoning and moral emotions should be further investigated in future studies.


Graphene doped by alkali atoms (AC Studying acoustic plasmons (APs) in single-layer ,2,3,4,5,As alkali atoms are added to graphene deposited on various substrates, they intercalate between the graphene layer and the substrate and metalize, i.e., they arrange in a two-dimensional (2D) crystal lattice, which supports a partially filled parabolic band ,19,20,21l doping ,23,24,25l doping ,34,35,36l doping ,38,39,40x) and the electrostatic doping. In other words, it is substantially unstable, and the aim of this research is to explore the mechanism of the (in)stability of the AP in the ACIn this paper, we focus on the AP, firstly because of its linear dispersion, but also because of the still rather unclear mechanism of its occurrence in the ACtime ago ,42,43. Itime ago . HoweverIn order to do that, we investigate two model systems of doped graphene, KCWe study electronic excitations in the ACapproach  and a reximation ,45. The The paper is organized as follows. In In this section, we derive the spectral function of the electronic excitations in the self-standing KCsumption  that thez direction is perpendicular to the crystal plane, the graphene layer occupies the We shall briefly describe the method of calculation of the surface electronic excitation spectra at the arbitrary position crystals ,48,49,50T. The matrix elements (or the charge vertices) areWe start from the 3D Fourier transform of the noninteracting (free) electron response functionz and z and L. We now need to determine the screened response function The Fourier expansion of the free electron response function in the Since the integrations in  are perfxpansion , and a sEquation , it agaiThe solution of Equation  has the z and After solving Equation , the nonF areThe propagator of the induced dynamically screened Coulomb interaction can be calculated from the response function  as 50]50](12)xpansion  and Equaxpansion , the intThe spectral function, which defines the intensity of the energy loss by an external perturbation to the excitation of the Up to this point, we derived the expressions for the self-standing q2D systems, but including a dielectric substrate polarization is now straightforward. It is obtained simply by replacing the bare Coulomb interaction  with theThis also means that the matrix  used in In the long-wavelength or optical limit σμinThe interband optical conductivity is determined from the optical limit of the nonlocal interband conductivityThe nonlocal interband conductivity is (24)σμinWe assume that the dielectric media is vacuum ϵFor the doped graphene, we follow the method proposed by Gjerding et al. , and wriin Refs. ,55, whic in Ref. , describG layers . We use We note that the relevant parameters for the KCgraphene , and by perature  as 4πEFgThe KC package . The cor package ,60. For nctional . The grok K-mesh  and the (PZ) LDA . The grog=2.45 Å , while fz) direction, the crystal local field effects are taken into account only in the z direction and neglected in the The ACunctions  and condunctions –25) are are8 resuctivity –(25) is k-point mesh and the band summations The ground state electronic density of the bulk Alfunction  of the Aalue 3.4 . However energy) . HoweverTo demonstrate that these effects are not driven by the dielectric substrate, in z direction), where the interband screening will be modified by changing the number of valence bands participating in the interband screening.To investigate this mechanism, we take advantage of the reduced model which clearly distinguishes between the different interband contributions to the dynamical response in the two systems. As discussed in n which denotes the number of omitted bands. For example, To explore how the screening coming from the interband excitations beyond the Fermi energy influences the AP, we omit one or more unoccupied bands from the calculation, to determine exactly how each band influences the excitation spectra. As we can see from the band structures shown in n. It is important to point out that we also attempted to keep the lowest unoccupied band in the calculation while omitting any number of the other unoccupied bands, above that one, and that did not lead to the occurrence of the AP in the CsCAnother interesting point is that the difference in the intensities of the We analyzed the origin of AP instability in graphene doped by alkali metals on two prototype systems, KC


By calculation and comparison, the large absorption loss  contributes to reducing secondary radiation, which is quite beneficial for stealth uses. Thus, this work demonstrates a promising design method for the preparation of green high-performance composites for EMI shielding and stealth applications .Because of serious electromagnetic pollution caused by the widespread use of radio frequency equipment, the study of electromagnetic interference (EMI) shielding materials has been a long-standing topic. Carbon fiber and graphene composites have great potential as EMI shielding materials due to their unique microstructure and electrical conductivity. In this work, a novel kind of core-shell composite is fabricated based on the pore-rich pine needles-derived carbon fibers (coded as PNCFs) core and the graphene shell. The pore-rich PNCFs are created by KOH activation, and the integration between the pore-rich PNCFs and the graphene relies on a plasma-enhanced chemical vapor deposition (PECVD) method. The conductivity of the pore-rich PNCFs@graphene core-shell composite reaches 4.97 S cm With the rapid development of electronic information technology, radio frequency equipment has been widely used in various fields such as communications, biological medicine, and remote sensing. Radio frequency equipment brings us convenience, but it also brings serious electromagnetic pollution. Electromagnetic pollution will not only disturb the normal operation of precision instruments, but also cause harm to people’s health. Therefore, there is an urgent need to develop high-performance electromagnetic interference (EMI) shielding materials ,2,3,4.EMI shielding materials can be roughly divided into three categories: carbon materials, metals, and conductive polymers. Among them, carbon materials have suitable electrical conductivity, strong oxidation resistance, and superior thermal stability. They are quite suitable for electromagnetic shielding materials. As people’s awareness of environmental protection continues to increase, more and more attention has been paid to green chemistry and materials ,6. Graph2 V−1 S−1 [−1 and electromagnetic shielding effectiveness is 55.0 dB at 12.4–18 GHz. The porous structure will greatly increase the specific surface area of the materials, and the specific surface area of the shielding materials is also a key parameter that affects the EMI shielding performance. For instance, Chen et al. [Meanwhile, carbon nanomaterials  have attracted the interest of outstanding scholars in different fields in the past few years due to their superior various performance. For example, Zhang et al.  prepared V−1 S−1 ,21,22,23 V−1 S−1 . Silkworn et al.  preparedn et al. , we haveSEA) and reflection shielding effectiveness (SER) of the pore-rich PNCFs@graphene account for 90.8% and 9.2% of the total effectiveness, respectively, revealing an absorption-dominate EMI shielding mechanism. In addition, a possible mechanism analysis on the EMI shielding process is proposed.Herein, we use pine needles as raw materials and carbonize them into carbon fibers (coded as PNCFs) to integrate with graphene after an activation treatment. The activation treatment is conducted by a combined process of KOH immersion and calcination. The growth of graphene is carried out on the surface of pore-rich PNCFs by virtue of plasma-enhanced chemical vapor deposition (PECVD) ,28. FiguKOH (85% purity), ethanol (98% purity), distilled water, hydrochloric acid, pure paraffin, high-purity hydrogen (99.999% purity), argon (99.9% purity), and methane (99.999% purity) are provided by Hunan FaLanKeEr Chemical Co., Ltd., Changsha, China and directly used without further purification.2 atmosphere and heated to 500 °C at a heating rate of 5 °C min−1, and heated at this temperature for 1 h. Subsequently, the material was heated to 1000 °C min−1 at a heating rate of 5 °C min−1, and maintained at this temperature for 2 h. Finally, the temperature was reduced to 500 °C min−1 at a cooling rate of 5 °C min−1, and finally cooled to room temperature naturally.Green and fresh pine needle raw materials were washed with distilled water and then placed in an oven at 60 °C for 12 h. Then, the dried pine needles were placed in a tube furnace under N−1 under the protection of argon gas (flow rate 250 scc/min) and then kept at 800 °C for 5 h. After that, the sample was naturally cooled to room temperature. Then the sample was washed with 10% HCl until the pH value approached 7. Then the sample was dried at 60 °C.A combined method of KOH immersion and calcination was used to make pores for the pyrolyzed pine needles. First, 50 mL of KOH solution was used to soak the PNCFs for 24 h (the mass ratio of PNCFs to KOH is 1:4). After being dried at 105 °C to remove moisture, the PNCFs and powder-like KOH were transferred to a nickel ark for heat treatment. The temperature was first raised from room temperature to 800 °C at a heating rate of 10 °C min2CO3, K2O, etc. Subsequently, these products further react with PNCFs. These reactions not only produce abundant new micropores but also enlarge the specific surface area [The activation mechanism using KOH as the activator can be described as follow. During the heating process, the KOH crystallizes into small particles and intercalates into PNCFs. When the temperature rises to 800 °C, the activator KOH reacts with PNCFs, leading to the production of Kace area . The relThe chemically activated samples were placed in a 250 mL WC lined vial of the planetary carrier and ground using a 10 mm diameter WC ball with a ball-to-powder mass ratio set at 4:1. Milling was performed at 300 rpm for 1.5 h. Then sieve the obtained powder to remove powder agglomerates larger than 325 microns in diameter (sieve size = 0.045 mesh) in order to obtain particles of micron size. Most of the resulting pore-rich PNCFs particle sizes are distributed between 50 and 325 µm.4/Ar flow rate ratio and the total gas pressure were set as 20 sccm/80 sccm and 400 Pa, respectively. The substrate bias voltage was 30 V and the RF power was 200 W. The graphene was deposited at 800 °C, and the deposition time was 60 min.The growth of graphene by PECVD follows our previously reported method . For theThe paraffin was first melted at 80 °C for 20 min. The liquid-stated paraffin was mixed with the pore-rich PNCFs@graphene at a weight ratio of 1:1. After the homogeneous mixing under a gentle magnetic stirring at 80 °C, the paraffin/pore-rich PNCFs@graphene composites were transferred to a mold for compression molding and naturally cooled to room temperature. The thickness of the sample was controlled at 2.000 ± 0.020 mm.−1 and 5–90°, respectively. Raman analysis was performed on a Raman spectrometer  employing a helium/neon laser (633 nm) as the excitation source. The instrument for tensile tests was an electronic universal material testing machine . Transmission electron microscope (TEM) images were acquired using an FEI, Tecnai G2 F20 TEM  with a field-mission gun operating at 200 kV. Atomic force microscopy (AFM) images were performed by Bruker Dimension Icon scanning probe microscope . The samples were ground by ball mill (PML 10). The size distribution of synthesized particles (100 mg) using a laser particle sizer .Micromorphology observations were made on a scanning electron microscope . Before the observations, the samples were adhered to the sample stage and sprayed with gold for 45 s by using an Oxford Quorum SC7620 sputter coater at 10 mA (the thickness of the sputtered gold coatings is ~5 nm). SEM digitized photographs were taken at an acceleration voltage of 15 KeV with a magnification range between 75 and 10,000 g. The chemical composition of carbon nanomaterials grown by PECVD was analyzed by energy dispersive spectroscopy (EDS). The crystal structure was studied by X-ray diffraction  with Cu Kα (target) radiation (λ = 1.5418 Å). The scan rate and scan range were set as 4° minSeaton et al.  first pr2 adsorption–desorption tests using an accelerated surface area and porosimeter system  at 77.4 K. High-purity nitrogen (99.99%) was used. Sample density was controlled at 1.000 g/cm3. All samples were degassed at 300 °C for 2 h prior to measurements.The pore structure and surface area were analyzed by NV, given in the form of the following adsorption integral equation [V(p) denotes the volume adsorbed as a function of pressure, and ρ is the model adsorption isotherm given in terms of the adsorbate density calculated for model pores as a function of pore width (w) and equilibrium pressure (p). The pore size distribution, f(w), is the distribution of pore volumes as a function of pore width. The lower and upper limit of the integration range are the widths of the smallest and largest pores  [ρ comes from the molecular model for nitrogen adsorption on the basis of nonlocal or smoothed density approximation, and it is also the kernel file in the DFT software we use.The International Organization for Standardization (ISO) recommends the use of DFT to calculate the pore size distribution . The theequation :(4)V(p)=e walls) . In this2@77K on carbon materials ” was selected.In this study, the DFT pore distribution model was calculated by using the ASIQwin software released by Quantachrome Instruments , and the model (kernel) “NS11 and transmission parameters S21 were measured. From the S-parameters, we can calculate the reflection coefficient (R), transmission coefficient (T), and absorption coefficient (A). The reflection loss (SER), absorption loss (SEA), and total shielding efficiency (SET) can be calculated by Formulas (5)–(10).A PNA-X network analyzer  was used for EMI shielding tests in the frequency range of 8.2~12.4 GHz (X-band) by coaxial method. The measuring device consists of two coaxial adapters (sample holders) and a network analyzer. The signal was coupled and transmitted through the coaxial cable and the attenuator. The paraffin/pore-rich PNCFs@graphene mixture was pressed into a ring mold  and tightly fixed in the coaxial sample holder. Based on the coaxial transmission theory, the electromagnetic wave was subjected to multiple reflections and transmissions in EMI shielding samples, hence leading to the electromagnetic energy attenuation . After mThe microstructure of the pore-rich PNCFs@graphene composite and its precursors is observed by SEM. As shown in SE). It also involves the skin depth δ where electromagnetic wave energy decreases to 1/e of the incident wave, described as δ = (πfμσ)−1/2 . In X-band (8.2–12.4 GHz), the calculated δ is in the range of 45.0–55.4 nm, which is much lower than the thickness of the graphene shell. Therefore, the graphene shell can effectively shield electromagnetic waves.The graphene shell plays a key role in EMI shielding. As a result, we study its microstructure through SEM and AFM observations. θ = ~23.2° and 43.4°, which correspond to the (002) and (100) planes of carbon structure, respectively. By comparing the XRD patterns of pore-rich PNCFs and pore-rich PNCFs@graphene, both XRD patterns show a distinct amorphous carbon structure, and the relatively intact crystal structure of graphene after growing graphene by PECVD causes the carbon (002) and (100) peaks in pore-rich PNCFs@graphene to become narrower and sharper [To further study the microstructure of composite, the crystal structure of the pore-rich PNCFs and pore-rich PNCFs@graphene are monitored by XRD a. The XR sharper ,39.−1 and 1600 cm−1 are assigned to the characteristic D-band (defects and disorder) and G-band (graphitic) bands  between the D and G bands is often used as a measure of the defect density in carbon materials. The larger the ratio of ID/IG, the higher the degree of defects in the carbon materials. After growing the graphene film, ID/IG increased significantly, and the ID/IG value increased from 1.15 (pore-rich PNCFs) to 1.21 (pore-rich PNCFs@graphene), which is caused by the graphitization edge faces (dangling bonds) exposed on the surface during the PECVD process and the additional defects generated by the active plasma [ID/IG value of the pore-rich PNCFs@graphene composite reveals a large disorder/defect density. The residual defects/groups of graphene may be beneficial for the electromagnetic wave absorption of the pore-rich PNCFs@graphene core-shell composite.Raman spectra are employed to confirm the graphitized structure and the defect density of carbon materials. The Raman spectrum of composite shows the typical signal characteristics of graphene. The peaks located at around 1350 cmc) bands b. The G-symmetry . In addie plasma . The hig2 g−1 (pine needles) to 333.2 m2 g−1 (PNCFs). A combined method of KOH immersion and calcination was used to make pores for PNCFs, and the introduction of graphene by PECVD is to make the surface of PNCFs rougher, thus further increasing the specific surface area of the material. The specific surface area of the material increases from 333.2 m2 g−1 (PNCFs) to 498.1 m2 g−1 (pore-rich PNCFs) and eventually to 699.1 m2 g−1 (pore-rich PNCFs@graphene)  . Larger 2 adsorption–desorption isotherms  to 0.22 cm3 g−1 (pore-rich PNCFs) after pore formation (3 g−1 (pore-rich PNCFs) by 1.8 times to 0.40 cm3 g−1 (pore-rich PNCFs@graphene). This is related to the growth of graphene films and the formation of core-shell structures. Thus, the final pore-rich PNCFs@graphene has a suitable porous structure. The enhancement of the surface area and pore volume of pore-rich PNCFs@graphene and its close impedance matching with free space may accelerate the incident electromagnetic waves entering the composite material. The existence of pores in the pore-rich PNCFs@graphene can also provide channels for the penetration and transmission of electromagnetic waves to promote multiple scattering, thus further increasing the attenuation of electromagnetic waves in the composite [The calculated average pore size decreases from 6.55 nm  to 2.20 nm (pore-rich PNCFs) when the calculated total pore volume increases from 0.17 cmormation . The pheomposite . The porSE, which is expressed using decibels (dB) as a unit. As electromagnetic waves propagate to the surface of the material, part of the incident power (PI) will be reflected, part of the energy will be absorbed and dissipated, and the rest (PT) will pass through the shielding material [SET is equal to the sum of electromagnetic wave reflection SER, electromagnetic wave absorption SEA, and electromagnetic wave internal multiple reflection loss (SEM), expressed as:σ means the conductivity of the shielding material, μ is the relative permeability of the shielding material, f is the frequency of a certain frequency point, d represents the thickness of the shielding material, and σ represents the skin depth of the material. Reflection loss and absorption loss are the two main mechanisms of electromagnetic shielding. The internal multiple reflection loss is the electromagnetic radiation trapped between the two boundaries of the material; that is, the electromagnetic wave reflected from the second boundary returns to the first boundary, and continues to reflect from the first boundary to the second boundary, and so on. The larger surface area of the porous structure and the larger surface-interface space contribute to providing more sites for electromagnetic wave scattering and multiple reflections. To acquire more reflected waves, shielding materials must possess abundant load carriers (electrons and holes) to interact with each other.Theoretically, for plane waves propagating to the material surface, they will be reflected and absorbed. The remaining electromagnetic wave will penetrate the material and cause multiple reflections inside the material, thereby inducing further dissipation of electromagnetic wave energy. The EMI shielding property is measured by the EMI material . The SETSEA and SER are the main ways of electromagnetic shielding. Both SEA and SER are proportional to the conductivity of the material. So the electrical conductivity is tightly responsible for the EMI shielding property. As the conductivity increases, the ability of the material to lose the incident radiation increases. According to the free electric theory, the electrical conductivity σr in an alternating electric field could be evaluated by the following Equation (15):f is frequency, ε0 is the permittivity of free space (ε0 = 8.854 × 10−12 F/m), and ε″ is the imaginary permittivity [−1, and the conductivity of the pore-rich PNCFs rises to 1.57 S cm−1. The difference between them is mainly attributed to the construction of conductive networks. The conductive networks promote the movement of charge carriers and provide more pathways for the transport of electrons in the pore-rich PNCFs, thereby improving the electrical conductivity of pore-rich PNCFs. The unique graphene “shell” in the composite is very important in improving conductivity. Due to the high conductivity of the graphene shell in the pore-rich PNCFs@graphene, the pore-rich PNCFs@graphene reaches 4.97 S cm−1 at 8.2 GHz, 216.6% higher than that of the pore-rich PNCFs (1.57 S cm−1). High electrical conductivity facilitates the efficient mobilization of charge carriers in the pore-rich PNCFs@graphene, which leads to the promotion dissipation of incident electromagnetic waves on the surface of the core-shell composite and improves the EMI SE of the composite.According to electromagnetic shielding theory, ittivity ,50. The SET) is the summation of all attenuating functions. The pine needles present an extremely low SET value, indicating that pine needles almost have no EMI shielding ability. By contrast, the PNCFs deliver a much higher SET maximum value of ~36 dB. The maximum value of SET rises to ~52 dB for the pore-rich PNCFs. As the graphene is introduced, the maximum value of SET increases to ~77 dB for the pore-rich PNCFs@graphene  [3O4@reduced graphene oxide (rGO)/natural rubber (37 dB) [To investigate the EMI shielding property of the pore-rich PNCFs@graphene, the composite is tested in the frequency range of 8.2–12.4 GHz . Generally, the total EMI shielding effectiveness (graphene a. Such a (28 dB) , PEI/PA@ (28 dB) , honeyco (28 dB) , loofah  (28 dB) , Fe3O4@r (37 dB) , rGO/pol (37 dB) , multi-w (37 dB) , ALC bas (37 dB) , SC-Co-G (37 dB) , and cot (37 dB) , as summX% is the percentage of the difference between the energy in the shielding field and the energy outside the field relative to the energy in the field. For example, 99%, that is, the radiation intensity of the radiation source is shielded by 99%. EMI shielding effectiveness (dB) is converted into EMI shielding effectiveness (%) using Equation (16) as follows:EMI shielding efficiency is usually expressed as a percentage (%) to indicate the ability of a material to block waves. The percentage shielding effect After substituting the data into the calculation, the pore-rich PNCFs@graphene shows a shielding efficiency of 99.9999%. This shows that the composite can shield most electromagnetic waves.A (A = SEA/SET) of PNCFs, pore-rich PNCFs, and pore-rich PNCFs@graphene are further analyzed. A can be approximately expressed as [To estimate the mechanism of material dissipating electromagnetic waves more accurately, the EMI absorptivity essed as ,59:(17)ASEA and SER can be expressed by Equations (18) and (19):d is the thickness of the materials, μr is the magnetic permeability, ω is the angular frequency, σAC is the frequency-dependent conductivity. The SET, SER, and SEA values of the PNCFs, pore-rich PNCFs, and pore-rich PNCFs@graphene are presented in A increases from 83.6% for PNCFs to 85.8% for the pore-rich PNCFs and eventually to 90.8% for the pore-rich PNCFs@graphene .graphene d, shows e X-band ,61. Cleaε = ε′ − iε″) and permeability (μ = μ′ − iμ″). The real part of the electromagnetic parameters  is a measure of the polarization amount and means the capability for electric or magnetic energy storage, while the imaginary part  represents the energy loss ability. As seen in ε′ decreases with the increase in frequency. This behavior is raised by the lag of the induced electric field reversed external electric field at high frequency [ε′ increases obviously. The result indicates the remarkably improved electric energy storage capacity of the composite.The electromagnetic property of the pore-rich PNCFs@graphene is determined by complex permittivity  to 498.1 m2 g−1 (pore-rich PNCFs) and eventually to 699.1 m2 g−1 (pore-rich PNCFs@graphene). Larger surface area and interface space provide more sites for electromagnetic wave scattering and multiple reflections.By activation and PECVD treatment, the specific surface area of the material increases from 333.2 m3 g−1 (pore-rich PNCFs) to 0.40 cm3 g−1 (pore-rich PNCFs@graphene). This porous structure will facilitate the multiple reflections of the incident electromagnetic waves.After the growth of graphene with the help of PECVD, total pore volume increased 1.8-fold from 0.22 cm−1 (PNCFs) to 4.97 S cm−1 (pore-rich PNCFs@graphene).The generation of conductive networks promoted by the growth of graphene, the as-prepared core-shell structure has a greatly improved conductivity from 1.13 S cmBecause of the combination of high electrical conductivity and unique core-shell porous structure, the PNCFs@graphene composite delivers a high EMI shielding effectiveness of ~77 dB, much higher than its precursors , respectively.Due to the enhanced surface area and pore volume of porous PNCFs@graphene to promote the dissipation of electromagnetic waves in the composite, the absorption ratio increases from 83.6% for PNCFs to 85.8% for the pore-rich PNCFs and eventually to 90.8% for the pore-rich PNCFs@graphene. This absorption-based shielding mechanism will have great potential for developing stealth materials.From the results obtained in this investigation, the following conclusions can be drawn:


Bipolaris sorokiniana (B. sorokiniana) and graphene oxide (GO) were selected for materials. A combination of physiological, cytological, and biochemical approaches was used to explore how GO affects the growth and pathogenicity of B. sorokiniana. The mycelial growth and spore germination of B. sorokiniana were both inhibited in a dose-dependent manner by GO treatment. The addition of GO significantly alleviated the infection of pathogenic fungi in host plants. The results of scanning electron microscopy demonstrated that the inhibitory effect of GO on B. sorokiniana was primarily related to the destruction of the cell membrane. Our study confirmed the antifungal effect of graphene in vitro and in vivo, providing an experimental basis for applying graphene in disease resistance, which is of great significance for agricultural and forestry production.The antimicrobial properties of graphene in vitro have been widely reported. However, compared to research performed on graphene’s antibacterial properties, there have been relatively few studies assessing graphene’s antifungal properties. In particular, evaluating graphene’s pathogenic effects on host plants in vivo, which is critical to using graphene in disease control, has rarely been performed. In this study, the fungal pathogen of wheat, barley, and other plants,  Graphene is a typical two-dimensional carbon nanomaterial and is famous for its excellent electronic conductivity, high thermal stability, high surface area, fine water dispersibility, and outstanding mechanical properties ,2. It haStaphylococcus aureus (S. aureus)) and Gram-negative (Escherichiacoli (E. coli), Pseudomonas aeruginosa (P. aeruginosa), Ralstonia solanacearum (R. solanacearum), Pseudomonas syringae (P. syringae), Xanthomonas campestris pv. undulosa (X. campestris pv. undulosa), and Xanthomonas oryzae pv oryzae (Xoo)) bacteria [Aspergillus niger, Aspergillus oryzae, Fusarium graminearum (F. graminearum), and Fusarium oxysporum (F. oxysporum) were demonstrated to be inhibited by GO or rGO [3O4 nanocomposites on grapevine leaves can reduce downy mildew severity under field conditions [Graphene and graphene-based materials (graphite (Gt), graphite oxide (GtO), graphene oxide (GO), and reduced graphene oxide (rGO), etc.) have different antimicrobial activities, including anti-bacterial, anti-fungal, and anti-viral properties ,14. For bacteria ,21,22,23O or rGO ,24. AddiO or rGO , HepatitO or rGO , PseudorO or rGO , etc.) wnditions .E. coli among four graphene-based materials  [E. coli is particularly vulnerable to GO as it is growing, which is largely due to alterations in the cell envelope ultrastructure [Many factors affect the antimicrobial abilities of graphene. As reviewed by Fatima et al. and Radhi et al., its physicochemical traits, including size, physical structure, surface structure, and electronic structure, the combined physiological properties of microorganisms, and interaction conditions influence graphene’s antimicrobial properties ,23,28. TGO, rGO) . AdditioGO, rGO) ,22. Impotructure . In additructure ,32.Currently, several graphene antimicrobial mechanisms have been proposed, including sharp edge cutting, lipid extraction, oxidative stress, cell trapping, and parcel isolation ,28. For Bipolaris sorokiniana (B. sorokiniana) is generally considered a hemibiotrophic pathogen related to root rot, leaf blotching, and black embryos in wheat and barley, among other grass species [B. sorokiniana is widely distributed worldwide, bringing immeasurable losses to the yield and quality of many cereal crops. B. sorokiniana could infect wheat across the whole growth period [B. sorokinian can infect many parts of host plants, such as roots, stems, leaves, and grains [ species ,39. B. sh period . The exih period . B. sorod grains . In pracd grains ,44. Mored grains . In addiB. sorokiniana, a pathogen of a serious plant disease (wheat root rot), in this study. Our results demonstrate that GO could provide broad-spectrum antimicrobial activity against pathogenic fungi and pave the way for the application of GO in preventing plant diseases.To further understand the mechanism behind graphene’s inhibition of fungal pathogens and analyze the effects of GO on host infection in vitro and in vivo, we selected 2, respectively. After 150 h of reaction, the graphene oxide suspension with a mass concentration of 0.5% was obtained. The GO suspension was then characterized by Raman spectra and scanning electron microscopy [GO suspension was synthesized by an electrochemical method in our laboratory, as in previous studies . Two elecroscopy ,11.B. sorokiniana Lankao 9-3 was provided by Dr. Shengli Ding’s laboratory from Henan Agricultural University, as previously described [escribed  and was B. sorokiniana was obtained after 7 d of incubation. Photographs were taken, and the diameter was recorded. The mycelial growth inhibition rate was calculated using the following formula:The growth curve of the mycelia was tested by following previously reported standard procedures . The strMycelial growth inhibition rate (%) = (Colony diameter of test treatment − Colony diameter of CK)/(Colony diameter of CK − 1 cm) × 100The mycelial biomass of the fungi was measured following standard procedures. Three pieces of fresh mycelia were transferred to a 50 mL liquid PDA medium containing 0, 50, 100, 200, and 500 mg/L GO in conical flasks. After incubation in shaking table for 4 d at 28 °C and 200 rpm, 1 mL of the culture was removed for subsequent mycelial morphology observation. The mycelia were filtered, repeatedly washed with distilled water, and dried to a constant weight of 50 °C. The antifungal experiments were performed in a completely random design with three replicates.B. sorokiniana spores were obtained as previously described [B. sorokiniana were incubated on a PDA medium for 5 d at 28 °C under alternating light and dark conditions. Then, the spores were collected by washing the mycelia with distilled water and centrifuging at 3500 rpm for 5 min. The spore suspensions were adjusted to a concentration of 5 × 104 spores per ml.escribed . The mycTo measure the germination rate and the length of the germ tubes, 100 µL spore suspensions were mixed with GO and PDA liquid medium in tubes to produce a final concentration of 0, 50, 100, 200, and 500 mg/L GO. 50 µL mixture was taken and transferred onto grooved slides at 28 °C for 7 h incubation. The incubation time was extended to 20 h for the mycelial branching observation. Three grooved slides were prepared for one mixture in each treatment, and the average values were compared. One hundred spores were observed under a microscope to analyze the germination rate, the length of the germ tubes, and the mycelial branching of different mixtures. Micrographs were taken, and the number of the germinated spores, the length of the germ tubes, and the mycelial branching were recorded. The experiments were performed in a completely random design with three replicates. The spore germination rate was calculated using the following formula:Spore germination rate (%) = The number of germinated spores/Total number of spores × 100.The culture liquid isolated from the biomass assay experiment was diluted 10 times with pure water for light microscope observation. A drop of diluted culture was placed on a concave slide and covered by the slide, then placed directly under the light microscope.B. sorokiniana spore suspensions were mixed with GO to a final concentration of 500 mg/L and incubated for 3 h at 28 °C. The control samples were mixed with distilled water. The mixtures were centrifuged at 3500 rpm for 5 min, and the spore cells were collected and fixed with 2.5% glutaraldehyde for 2 h and washed with 0.1 M pH 7.0 phosphate buffers three times. The spore cells were then dehydrated by a series of ethanol solutions  and dried in a vacuum freeze drier. The samples were then observed under the SEM.Changes in the morphology of the spores after GO treatment were observed via scanning electron microscopy (SEM). The Cell contents such as DNA and RNA are discarded in damaged spore cells, which can be monitored by UV absorption at 260 nm. A mixture of 200 μL GO in different concentrations  with 2 mL spore was cultured in stasis for 3 h at 28 °C. The mixtures were then filtered with a 0.22 µm filter membrane, and the filtrates were tested by absorbance at 260 nm.3 was added to the spore suspensions for 2 h to measure the total electrolyte loss.An electrolyte leakage experiment was performed, as previously described . ElectroTo assess the disease severity after fungal infection in wheat leaf inoculation experiments, 10 µL spore suspension was sucked onto the 10 d old wheat leaves, which were kept humid by covering them with clingfilm for 24 h at 28 °C under dark conditions. A total of 1 mL 500 mg/L GO solution was sprayed onto the inoculated leaves and 1 mL distilled water was sprayed onto the control leaves once a day for 3 d in a 25 °C constant temperature incubator. Lesions on the inoculated leaves were recorded by a camera, and the area of the disease spot was measured with Image J.B. sorokiniana and placed in a constant temperature incubator at 25 °C for mycelial propagation. After 6–7 days of cultivation, the millet medium with mycelia was naturally dried. In parallel, the blank millet medium was also placed at the same place as a control. Sterilized soil and the millet medium were mixed at a 3:1 ratio, and the mixture was placed in clean pots for plant growth. Wheat seeds were sterilized in 0.5% sodium hypochlorite for 15 min and washed three times with sterilized water. The sterilized seeds were covered with two sheets of moist filter paper and placed in 25 °C in darkness for 24 h to accelerate germination. Germinated wheat seeds were washed with water and 500 mg/L GO solution for control and GO treatment, respectively. Then, the seed was planted into the soil and irrigated with water and 500 mg/L GO every seven days. Plants were grown in a culture room at 23 °C under a 16-h-light/8-h-dark photoperiod. Three repetitions were randomly set for each treatment. The aerial part and root of plants were observed and photographed after 14 d. The plant height was measured by a ruler. The root was washed with water to analyze root morphology with WinRHIZO software  after scanning [Pot inoculation experiments were conducted following the procedures used by Shengli Ding’s lab. Millet was boiled in water for 2 min and dried until it was semi-dry, after which it was placed in a conical flask for sterilization. The millet medium was inoculated with scanning .p value < 0.05 or <0.01.Every experiment was conducted in a completely random design with three repeats. The results are displayed as mean ± standard deviation (SD). Statistical analysis was performed by SPSS 21 software. Significance differences between groups were analyzed by one-way analysis of variance and determined by a B. sorokiniana was examined by measuring the growth curve and dry weight of the fungal mycelium exposed to different concentrations of GO . As shown in B. sorokiniana.The antifungal activity of GO to Spore germination is the first crucial stage in the return of spores to vegetative growth and is key to the spread of pathogenic fungi . As showB. sorokiniana.After spore germination and the formation of the germ tube, hyphae, and other structures, the number of mycelial branches continuously increased . The effB. sorokiniana, SEM was used to observe the morphological changes of spores after GO treatment. As shown in Previous studies have demonstrated that GO could disrupt the cell membrane of bacterial and fungal spores, altering the membrane potential and membrane-associated energy transducing system . To explB. sorokiniana to wheat, a leaf inoculation experiment and pot inoculation experiment were performed. As shown in B. sorokiniana spores caused disease varying degrees of disease spots in wheat leaves. GO application significantly reduced the area of the disease spot by 53%  demonstrated the antifungal activity of GO.Most reports have investigated the antimicrobial effects of graphene in vitro ,54,56,57viticola ,27. In cThe prevailing view for the antifungal mechanism of graphene as follows: (i) destruction of cell membranes caused by mechanical cutting, (ii) environmental isolation caused by a wrapping effect, and (iii) oxidative stress . GraphenB. sorokiniana both in vitro and in vivo. In vitro experiments, the mycelial growth and spore germination of B. sorokiniana were both interfered by GO treatment in a dose-dependent manner. In vivo experiments, GO treatment significantly reduced the pathogenicity of the pathogen in host plants. It is worth mentioning that this in vivo anti-fungal effect of GO is rarely reported in the existing literature. Finally, the results of scanning electron microscopy showed that the anti-fungal effect of GO on B. sorokiniana was mainly depended on the destruction of cell membrane. However, further study is still needed. First, it is necessary to clarify the molecular effects of GO treatment on gene expression and metabolite changes of B. sorokiniana. Secondly, research on the synergistic effects of graphene with other technologies, such as pesticides, is needed to maximize the effectiveness of its disease resistance. Finally, close attention must be paid to how graphene treatment affects plant growth and development to avoid potential ecological crises.This study provides evidence that appropriate GO concentrations can exhibit excellent antifungal properties on