Flexible wearable materials have frequently been used in drug delivery, healthcare monitoring, and wearable sensors for decades. As a novel type of artificially designed functional material, photonic crystals (PCs) are sensitive to the changes in the external environment and stimuli signals. However, the rigidity of the PCs limits their application in the field of biometric and optical sensors. This study selects silk fibroin (SF) and poly-N-isopropylacrylamide (PNIPAM) as principal components to prepare the hydrogel with the physical crosslinking agent lithium silicate (LMSH) and is then integrated with PCs to obtain the SF-PNIPAM dual-crosslinked nanocomposite for temperature and strain sensing. The structural colors of the PCs change from blue to orange-red by the variation in temperature or strain. The visual temperature-sensing and adhesion properties enable the SF-PNIPAM dual-crosslinked nanocomposite to be directly attached to the skin in order to monitor the real-time dynamic of human temperature. Based on its excellent optical properties and biocompatibility, the SF-PNIPAM dual-crosslinked nanocomposite can be applied to the field of visual biosensing, wearable display devices, and wound dressing materials. As the new generation of artificially designed optical material, photonic crystals (PCs) have unique periodic structures and optical characteristics—photonic bandgap (PBG)—which can alter the wavelength of reflected light when the stimuli such as pH, temperature, and pressure change the structural color of the PCs ,9,10,11.Silk fibroin (SF) is a material with good biocompatibility, breathability, and moisture permeability and has been widely used in the field of biomedicine for slow-release drugs, artificial organs, and immobilized enzyme materials ,26,27,28Here, we prepared a hydrogel interpenetrating polymer network by combining SF with PNIPAM via the physical crosslinking agent lithium silicate (LMSH), which can improve the tensile strength of SF and explore its application. The SF-PNIPAM dual-crosslinked nanocomposite with vivid structural colors was obtained by combining PCs with SF-PNIPAM hydrogel precursor, and the shape of the nanocomposite could be designed according to different molds. Changes in external stimuli such as temperature or strain led to variation in the lattice spacing of PCs resulting in shrinkage or swelling of the SF-PNIPAM dual-crosslinked nanocomposite, along with the change of the structural color. The visual temperature sensing and adhesive properties enabled the SF-PNIPAM dual-crosslinked nanocomposite to be directly attached to the surface of human skin, which also holds great promise in the field of visual wearable temperature sensing; Because of the biocompatibility and water retention, the SF-PNIPAM dual-crosslinked nanocomposite carries the potential to be used as wound dressings materials.As a hydrogel monomer, NIPAM can shrink with water loss when external temperature changes and will have cliff-like shrinkage near its phase transition temperature (33 °C). However, the poor mechanical properties of materials consisting of pure PNIPAM limit its application in wearable sensors.2O, then added LMSH , APS (initiator), and TEMED (accelerate the reaction) to prepare the SF-PNIPAM dual-crosslinked nanocomposite, which has best mechanical properties and excellent stability  on the surface of its charged nanoparticles, initiates free radical polymerization in situ to form PNIPAM macromolecular chain, meanwhile acts as a physical crosslinking point to adsorb a large number of SF macromolecular chain, finally obtaining dual-crosslinked nanocomposite based on the physical crosslinking  31,32,3,332,33.The preparation of SF-PNIPAM dual-crosslinked nanocomposite with NIPAM could be influenced by crossing temperature ,35,36. TTo observe the microstructure of the SF-PNIPAM dual-crosslinked nanocomposite film, we performed SEM characterization after drying it at room temperature (25 °C), as shown in To observe the real structure of the film in water, we freeze-dried the SF-PNIPAM dual-crosslinked nanocomposite film at −54 °C for 24 h and further characterized by SEM . The freWith the temperature-sensitive monomer—NIPAM—the SF-PNIPAM-dual-crosslinked nanocomposite can respond to temperature, which will shrink/relax with ambient temperature changes. A large extent of sudden contraction/relaxation occurs at the critical phase transition temperature (LCST), which, in turn, leads to a difference in the lattice spacing of the PCs, causing the changes in its PBG and structural color that can be visible to the naked eye. The LCST of pure NIPAM is 33 °C, and the addition of SF increases the LCST of the SF-PNIPAM dual-crosslinked nanocomposite by 2 °C, which is 35 °C. The structural color of the SF-PNIPAM dual-crosslinked nanocomposite will red-shift slightly with temperature decreasing below 35 °C, which indicates less shrinkage. Above 35 °C, cliff-type dehydration shrinkage occurs, and the structural color changes from original blue to red, which is different from normal hydrogel photonic crystal. The thermos-sensitive mechanism of the SF-PNIPAM dual-crosslinked nanocomposite is the interaction between the hydrophobic and hydrophilic groups in the system and water molecules. When the ambient temperature is lower than LCST, the hydrogen bonding between the SF-PNIPAM dual-crosslinked nanocomposite and the water molecules forms a highly ordered in the solvated shell. With the rising temperature, the hydrophobic interaction between molecules is enhanced, and the hydrogen bonds are destroyed. At a critical temperature (LCST), the SF-PNIPAM dual-crosslinked nanocomposite undergoes a phase transition and loses water; its structure changes from a loose coil structure to a compact colloidal structure. The illustration in The SF-PNIPAM dual-crosslinked nanocomposite exhibits excellent mechanical properties, which can produce stable deformation, accompanied by the change of photonic crystal structural color, so this material can be used to monitor joint movement and recovery. The structural color changes of the SF-PNIPAM dual-crosslinked nanocomposite are the result of deformation, and external stimuli such as pressure will deform it in the direction perpendicular to the force, which could increase the lattice constant of the PCs, resulting in the red shift of the reflection peak and structural color. The uneven change of structural color is caused by uneven pressure. The shape and structural color of the SF-PNIPAM dual-crosslinked nanocomposite will recover if external stimuli are removed, which indicates that our material has excellent mechanical properties. Due to the existence of a large number of hydrophilic groups, the SF-PNIPAM dual-crosslinked nanocomposite can fit well with the skin without falling off, without employing any adhesive. In order to explore the acid-base stability of the SF-PNIPAM dual-crosslinked nanocomposite, we prepared a phosphate buffer with pH = 5, 6, 7, 8, 9, and placed the material in the above buffer solution for 10 min. The structural color and reflection peak positions of the SF-PNIPAM dual-crosslinked nanocomposite with different pH environments were essentially unchanged . The ref2O and measured its swelling of each dimension in 2 h. 0, W0, and H0 were the original size of the SF-PNIPAM dual-crosslinked nanocomposite film, and L, W, and H were the size of the SF-PNIPAM dual-crosslinked nanocomposite film soaked in water for different periods. The result shows that our SF-PNIPAM dual-crosslinked nanocomposite film swelled in all three directions, and the swelling ratio in the height direction was slightly larger than that in the other two directions, which was 1.17. The SF-PNIPAM dual-crosslinked nanocomposite gradually reached swelling equilibrium after 10 h, and its volume remained constant. The addition of SF increased its hydrophilicity, and the film can absorb water and has a strong water-retention performance, which can shrink with water loss and swell again when immersed in water. The SF-PNIPAM dual-crosslinked nanocomposite was placed in the air when its volume shrinks correspondingly with the loss of water, and it can swell again when immersed in a solution, which has the inherent properties of hydrophilic hydrogels. Based on the above properties, the SF-PNIPAM dual-crosslinked nanocomposite film can be used in water-absorbing and drug loading and has the potential application for visual monitoring of healthcare.We characterized the swelling of the SF-PNIPAM dual-crosslinked nanocomposite in water to illustrate the excellent water-absorbing of the nanocomposite. The SF-PNIPAM dual-crosslinked nanocomposite film with the volume of 1 cm × 1 cm × 2 mm was placed in ultrapure HIn this paper, we used SF macromolecular from cocoons and PNIPAM macromolecular chains formed by free-radical polymerization of NIPAM monomers, lithium magnesium silicate as a crosslinking agent to prepare a purely physically crosslinked nanocomposite with good biocompatibility, flexibility, and stability. The reaction conditions were mild; the SF-PNIPAM dual-crosslinked nanocomposite obtained that is physically crosslinked has high strength and toughness. Through the optimization of the formula, the SF-PNIPAM dual-crosslinked nanocomposite had the best strength and strain sensing performance when the dosage ratio of NIPAM and SF was 5:1. The SF-PNIPAM dual-crosslinked nanocomposite had temperature sensing performance due to the use of NIPAM, and its phase transition temperature was near 35 °C. The structural color and reflectance spectrum of the PCs were changed with varying temperatures. When subjected to pressure stimulation, the structural color of the SF-PNIPAM dual-crosslinked nanocomposite gradually changes from blue to orange-red, which can realize naked-eye sensing of external strain. At the same time, the excellent acid-based stability of the SF-PNIPAM dual-crosslinked nanocomposite made it suitable in acidic and alkaline environments. Based on optical performance, biocompatibility, adhesion, and acid stability, the SF-PNIPAM dual-crosslinked nanocomposite can be used for body temperature measurement, real-time monitoring of human movement, and other areas.2O3, A.R.); cocoons, sodium carbonate , lithium bromide , polyethylene glycol , N-isopropylacrylamide , n-hexane (A.R.), lithium magnesium silicate 3Si4O10(OH)2·4H2O], LMSH, A.R.), ammonium persulfate , tetramethylethylenediamine , concentrated sulfuric acid , hydrogen peroxide .Methyl methacrylate , potassium persulfate , Alumina (Al2O were added to four-necked flask (250 mL), nitrogen was bubbled through the liquid surface for 30 min to remove O2, then heated to 80 °C and stirred at 300 rpm. The KPS solution was added to the above mixture and continued to react for 45 min, obtaining PMMA colloidal suspension. The mixture was centrifuged thrice to remove unreacted monomer and initiator. The vertical self-assembly method was used to prepare three-dimensional (3D) photonic crystals; the PMMA colloidal microspheres were self-assembled on a glass slide and obtained 3D photonic crystal array with face-centered cubic (FCC) stacking [The polymethylmethacrylate (PMMA) colloidal microspheres were synthesized by soap-free emulsion polymerization . MMA monstacking ,39.2O3 (20 g) was immediately added to the above solution, continued stirring for 5 min to remove the polymerization inhibitor in NIPAM, and filtered. n-hexane was added to the solution and placed in the dark at 4 °C for 24 h, and the NIPAM crystals were obtained by suction filtration, dried under vacuum at 40 °C for 24 h, and stored in a sealed and dark place at 4 °C.NIPAM (30 g) monomer and toluene (75 mL) were added to a single-necked flask, stirred magnetically, and heated to 45 °C for 30 min to obtain NIPAM solution; then, Al2CO3 solution (5 g L−1) for 30 min. The process was repeated thrice to remove sericin from the silk fibers and dried at room temperature. LiBr solution (9.5 M) was prepared and kept at 60 °C; then, added dried silk fibers with magnetic stirring and continued to react for 4 h. After the silk fibers completely dissolved, dialysis was performed with a dialysis bag  at 4 °C for three days to remove the LiBr and small protein and obtained SF solution. Then the dialyzed SF solution was concentrated with 10% polyethylene glycol . It was then reverse dialyzed at 4 °C for 12 h to obtain the SF solution with a concentration of about 5 wt%.We prepared a silk protein solution following the protocol . Cocoons2O and stirred magnetically until fully dissolved, placed in an ice-water mixing bath (<5 °C), and blew N2 for 10 min to avoid the presence of oxygen affecting polymerization effect of the prepolymer solution, APS (0.01 g) was dissolved in H2O (2 mL) and added to the above solution, TEMED  was added as a catalyst, continued to stir for 10 min and obtained prepolymer solution. The prepolymer solution was injected into the PCs to replace air and reacted at room temperature for 48 h after filling. Then the mold was taken out, and the nanocomposite was placed in H2O for 48 h to remove unreacted monomers. The SF-PNIPAM dual-crosslinked nanocomposite film was finally obtained.LMSH (mass ratio 20% of NIPAM), SF (0.2 g), and NIPAM (1 g) were mixed in 9 mL H


Clostridioides difficile infection (CDI) research is limited by a lack of standardized definitions for clinical response and disease outcomes, which impacts clinical drug development and results comparison between studies. We aimed to assess outcome definitions in CDI therapeutic trials to propose new versions that are clinically relevant, discrete and objective.A multidisciplinary group of CDI experts met monthly to review response endpoints from published clinical trials of antibiotic therapy for CDI. Previously published phase III or IV trials were assessed for outcome definitions. Discussions were held to reach a consensus on new clinical trial endpoints for adults and children to improve the accuracy and clinical relevance of measures of treatment success.Significant heterogeneity was noted amongst the primary endpoints in phase III and IV CDI antibiotic treatment trials. Initial clinical cure (ICC), strictly defined as < 3 unformed bowel movements/24 hour, and sustained clinical cure (SCC) were primary outcome measures for recent clinical trials. The strict ICC definition incompletely measures treatment success as assessed in clinical practice and, since ICC is necessary to achieve SCC, may lead to type II error for SCC. A set of proposed alternative outcome definitions was developed using the terms initial response (IR) and sustained response (SR) . IR allows for investigator assessment of overall improvement in CDI response more analogous to clinical practice and will lead to more patients eligible to meet SR. Achievement of SR requires both IR and no need for retreatment of CDI by day 30 after antibiotic completion and is the more relevant endpoint for CDI therapeutic development. The use of a less restrictive IR definition will more accurately capture early responses to treatment and importantly increase the validity of SR. The shortening of follow-up period by 30 days is also anticipated to reduce costs and efforts associated with conducting trials.Timeline of CDI outcome assessments for clinical trialsThe set of definitions proposed here will more accurately capture clinical success and standardize the approach to outcome assessment in trials of CDI therapeutics.Carolyn D. Alonso, MD, Cidara Therapeutics: Advisor/Consultant|Merck: Advisor/Consultant Oliver A. Cornely, Prof. Dr., Abbott: Honoraria|Abbvie: Advisor/Consultant|Actelion: Board Member|Al-Jazeera Pharmaceuticals: Honoraria|Allecra Therapeutics: Board Member|Amplyx: Advisor/Consultant|Amplyx: Grant/Research Support|Astellas: Honoraria|Basilea: Advisor/Consultant|Basilea: Grant/Research Support|Biocon: Advisor/Consultant|Biosys: Advisor/Consultant|BMBF: Grant/Research Support|Cidara: Advisor/Consultant|Cidara: Board Member|Cidara: Expert Testimony|Cidara: Grant/Research Support|CoRe Consulting: Stocks/Bonds|Da Volterra: Advisor/Consultant|DLR: Grant/Research Support|DZIF: Grant/Research Support|Entasis: Board Member|EU Directorate-General for Resarch and Innovation: Grant/Research Support|F2G: Grant/Research Support|German Patent and Trade Mark Office: German patent (DE 10 2021 113 007.7)|Gilead: Advisor/Consultant|Gilead: Grant/Research Support|Grupo Biotoscana/United Medical/Knight: Honoraria|Hikma: Honoraria|IQVIA: Board Member|Janssen: Board Member|Matinas: Advisor/Consultant|Matinas: Grant/Research Support|MedPace: Advisor/Consultant|MedPace: Grant/Research Support|MedScape: Honoraria|MedUpdate: Honoraria|Menarini: Advisor/Consultant|Merck/MSD: Grant/Research Support|Merck/MSD: Honoraria|Molecular Partners: Advisor/Consultant|MSG-ERC: Advisor/Consultant|Mundipharma: Grant/Research Support|Mylan: Honoraria|Noxxon: Advisor/Consultant|Octapharma: Advisor/Consultant|Octapharma: Grant/Research Support|Paratek: Board Member|Pardes: Advisor/Consultant|Pfizer: Grant/Research Support|Pfizer: Honoraria|Projektträger Jülich: Grant/Research Support|PSI: Advisor/Consultant|PSI: Board Member|Pulmocide: Board Member|Scynexis: Advisor/Consultant|Scynexis: Grant/Research Support|Seres: Advisor/Consultant|Shionogi: Board Member|Wiley (Blackwell): Editor-in-Chief, Mycoses Kevin W. Garey, PharmD, MS, Acurx Pharmaceuticals: Grant/Research Support|Paratek Pharmaceuticals: Grant/Research Support|Seres Therapeutics: Grant/Research Support|Summit Pharmaceuticals: Grant/Research Support Dale N. Gerding, MD, Destiny Pharma plc.: Advisor/Consultant Stuart Johnson, M.D., Ferring Pharmaceuticals: Membership on Ferring Publication Steering Committee|Ferring Pharmaceuticals: Employee|Summit Plc: Advisor/Consultant Stacy A. Kahn, MD, Lilly: Stocks/Bonds Ciaran P. Kelly, n/a, Artugen: Advisor/Consultant|Facile Therapeutics: Advisor/Consultant|Ferring Pharma: Advisor/Consultant|Finch: Advisor/Consultant|Finch: Advisor/Consultant|First Light Biosciences: Advisor/Consultant|First Light Biosciences: Ownership Interest|Milky Way Biosciences: Advisor/Consultant|Milky Way Biosciences: Grant/Research Support|Pfizer: Advisor/Consultant|Seres Therapeutics: Advisor/Consultant|Summit Therapeutics: Advisor/Consultant Larry K. Kociolek, MD, MSCI, Merck: Grant/Research Support Thomas J. Louie, MD, Artugen: Advisor/Consultant|Artugen: Grant/Research Support|Crestone: Advisor/Consultant|Crestone: Grant/Research Support|Finch Therapeutics: Advisor/Consultant|Finch Therapeutics: Grant/Research Support|Rebiotix: Advisor/Consultant|Rebiotix: Grant/Research Support|Seres Therapeutics: Advisor/Consultant|Seres Therapeutics: Grant/Research Support|summit plc: Advisor/Consultant|summit plc: Grant/Research Support|Vedanta Biosciences: Advisor/Consultant|Vedanta Biosciences: Grant/Research Support Maria Vehreschild, Prof. Dr., 3M: speaker fee|Astellas: Advisor/Consultant|Astellas: speaker fee|biologische heilmittel heel gmbh: Grant/Research Support|BioNtech: Grant/Research Support|EUMEDICA: Advisor/Consultant|Farmak International Holding: Advisor/Consultant|Ferring: Advisor/Consultant|Ferring: Speaker fee|Gilead Sciences: Advisor/Consultant|Immunic AG: Advisor/Consultant|MaaT: Advisor/Consultant|Merck: Advisor/Consultant|Merck: speaker fee|MSD: Advisor/Consultant|MSD: Grant/Research Support|MSD: speaker fees|Pfizer: speaker fee|Roche Molecular Systems: Grant/Research Support|Roche Molecular Systems: speaker fees|SocraRTec R&D GmbH: Advisor/Consultant|Takeda California: Grant/Research Support Professor Mark H. Wilcox, MD, FRCPath, GSK: Advisor/Consultant|GSK: Board Member|GSK: Grant/Research Support|Pfizer: Advisor/Consultant|Phico Therapeutics: Board Member|Seres: Advisor/Consultant|Seres: Board Member|Seres: Grant/Research Support|Summit: Advisor/Consultant|Summit: Grant/Research Support Erik R. Dubberke, MD, MSPH, Abbott: Advisor/Consultant|Ferring: Advisor/Consultant|Ferring: Grant/Research Support|Merck: Advisor/Consultant|Pfizer: Advisor/Consultant|Pfizer: Grant/Research Support|Seres: Advisor/Consultant|Summit: Advisor/Consultant|Synthetic Biologics: Grant/Research Support.


Hybrid structures comprised of graphene domains embedded in larger hexagonal boron nitride (h-BN) nanosheets were first synthesized in 2013. However, the existing theoretical investigations on them have only considered relaxed structures. In this work, we use Density Functional Theory (DFT) and Molecular Dynamics (MD) simulations to investigate the mechanical and electronic properties of this type of nanosheet under strain. Our results reveal that the Young's modulus of the hybrid sheets depends only on the relative concentration of graphene and h-BN in the structure, showing little dependence on the shape of the domain or the size of the structure for a given concentration. Regarding the tensile strength, we obtained higher values using triangular graphene domains. We find that the studied systems can withstand large strain values (between 15% and 22%) before fracture, which always begins at the weaker C–B bonds located at the interface between the two materials. Concerning the electronic properties, we find that by combining composition and strain, we can produce hybrid sheets with band gaps spanning an extensive range of values (between 1.0 eV and 3.5 eV). Our results also show that the band gap depends more on the composition than on the external strain, particularly for structures with low carbon concentration. The combination of atomic-scale thickness, high ultimate strain, and adjustable band gap suggests applications of h-BN nanosheets with graphene domains in wearable electronics. We investigate the mechanical and electronic properties of hBN nanosheets with graphene domains under strain. We find that the structures withstand large strain values and present highly adjustable band gaps, ranging from 1.0 to 3.5 eV. Various applications have been proposed for graphene in modern technologies3,4 due to its extraordinary properties, which include: high electron mobility at room temperature,5 high thermal conductivity,6 and broad light absorption spectrum.7 Moreover, graphene has motivated the search for other 2D materials arranged in a honeycomb-like hexagonal lattice. Well-known examples include hexagonal boron nitride (h-BN),8,9 silicene,10 and many transition metal dichalcogenides.11 In particular, h-BN, also known as white graphene, is an insulator material (energy gap >4 eV).8,9,12 This 2D material also presents remarkable properties, such as high thermal stability,13 high resistance to oxidation,14 and strong optical absorption in the UV region.15,16 Several synthesis methods have been used to achieve large area h-BN monolayers. For instance, h-BN has been grown through ultra-high vacuum chemical vapor deposition (CVD).17,18 We can list applications for h-BN as atomically flat substrates,19 as 2D dielectric materials,21 and in lubrication.20Produced from graphite by using micromechanical cleavage technology,22–25 It is also able to withstand tensile strains as large as 25%.26 h-BN also features high values of Young's modulus (≈0.865 TPa) and tensile strength (≈70.5 ± 5.5 GPa).27–29 Extensive literature exists describing the mechanical properties of graphene and h-BN and suggesting applications. For example, these structures have been used as reinforcing materials in nanocomposites30,31 and in highly flexible touch screens.32We can certainly also include the mechanical properties of graphene and h-BN in their list of remarkable properties. Graphene is the strongest material ever synthesized, with high values of Young's modulus (≈1 TPa) and ultimate tensile strength (≈130 ± 10 GPa).vice versa (graphene/h-BN sheets).33–39 These structures were first synthetized by Ci et al. using a thermal catalytic CVD method.40 Since then, other synthesis methods were proposed as well.41–43 Regarding the mechanical, electronic, and optical properties of these hybrid structures, first-principles calculations provided values intermediate between those found in graphene and h-BN. Interestingly, these values could also be controlled by adjusting the size and geometry of the nanodomain.33–38 Calculations also showed that the properties of the hybrid structures depend on the graphene/h-BN interface, where the less stable C–B and C–N bonds are found.33,37,38 Overall, the current results indicate that graphene/h-BN sheets may be useful for applications in optoelectronic devices, as they exhibit variable band gaps (<2 eV) and an optical absorption spectrum in the visible region.The contrasting properties of graphene and h-BN motivated investigations on hybrid 2D materials with intermediate properties. Arrangements of C, B, and N atoms were used to propose graphene-like hexagonal sheets with h-BN nanodomains and 22–25,27–29,44–50 One way to achieve this is by applying a tensile strain, with its magnitude and direction determining whether the material undergoes bending, wrinkling, stretching, or breaking.22,51–54 From an experimental point of view, strain can be applied to 2D structures through various methods.55–58 For example, the material may be deposited in an elastic substrate, which is then elongated.55 As a result of the applied strain, the atomic arrangement of the system is deformed, and many properties of the material are affected by this modification. For instance, changes in thermal conductivity,59 electronic and optical properties,22–25,27–29,44–48 and chemical activity60 have already been reported. In turn, these changes could improve or impair the performance of a 2D material for specific applications. These developments have given rise to a new field of research called “straintronics”, which aims to control electrical properties in 2D materials through mechanical deformation.61It is also possible to control the physicochemical properties of 2D materials (including graphene/h-BN hybrid sheets) through mechanical deformation.et al. found that a tensile strain reduces the band gap of h-BN, whereas a small compressive strain increases it.62 Other studies found similar results.22,27 In fact, each 2D material has a particular response to strain. For instance, in silicene a semimetal–metal transition was observed for biaxial strain values greater than 7.5%.63 For transition metal dichalcogenides, calculations and experiments investigated the effects of strain on the electronic properties, revealing a semiconductor–metal transition for large enough tensile strain values.64,65 And for graphene deposited on suitable substrates,24 strain opens the energy gap.Additionally, strain exerts a significant effect on the interactions between atomic orbitals in 2D materials, and hence it greatly affects their energy band gap. For example, Fujimoto et al.44 used molecular dynamics (MD) simulations to show that the fraction of h-BN contained in the structure determines its mechanical properties. In turn, this fraction depends on the geometry and/or the number of nanodomains in the sheet. They also found Young's modulus values intermediate between those of graphene and h-BN, with lower values for higher h-BN concentrations. More recently, Oliveira et al.48 combined DFT calculations and MD simulations to investigate whether the mechanical properties of hybrid sheets depend on their size. The authors found that the mechanical properties are independent of scale, so long as the graphene/h-BN concentration remained constant as the sheet size increased or decreased. Additionally, BxCyNz hexagonal sheets with atomic arrangements that included many B–C and N–C bonds were studied with MD49 and DFT50 simulations. The reported values of stiffness and tensile strength were lower than those reported for structures composed of h-BN nanodomains embedded in graphene sheets. Regarding the effect of strain on the electronic properties of graphene/h-BN sheets, Azevedo and Kaschny determined that the band gap of this hybrid material increased with strain, from 0.08 eV to 0.5 eV.45 However, note that the investigated sheets were composed mainly of carbon atoms. Overall, most studies in the literature concentrate on hybrid structures composed mostly of carbon atoms.44–50Strain is also often applied to materials in order to study their mechanical properties. For the hybrid graphene/h-BN sheets, Zhao 41–43 So far, theoretical investigations on BCN sheets with high h-BN concentration have focused on their electronic and magnetic properties.38,39 In the present contribution, we combine DFT calculations and MD simulations to investigate the mechanical properties as well as the effect of composition and strain on the electronic properties of h-BN sheets with graphene domains. We find that the band gaps of the investigated systems depend strongly on the graphene/h-BN concentration and moderately or weakly on the external strain (depending on the composition of the hybrid sheet). Furthermore, they can withstand strain values above 10% before permanent deformation. The combination of elasticity with a controllable band gap suggests possible applications of the investigated nanosheets as semiconductor elements on wearable electronics.On the other hand, hybrid structures with high h-BN concentration have received comparatively little attention, even though reports detailing their synthesis exist.II.Lx = Ly = L), with side lengths (L) ranging from 2 to 50 nm , which is given bynC follows the previous definition and nT is the total number of atoms. The values of γ for the investigated structures vary from 0.04 to 0.34, as shown in In the present work, first principles calculations and MD simulations were employed to investigate the mechanical and electronic properties of h-BN nanosheets containing graphene nanodomains (h-BN/graphene sheets). Some of the studied structures are shown in 0 nm see  and 2. Iains see , whereasains see . We deno66 as implemented in the SIESTA code.67,68 Norm-conserving Troullier–Martins pseudopotentials69 were used, in the Kleinman–Bylander factorized form.70 A double-ζ polarized basis set (DZP) was used, composed of numerical atomic orbitals of finite range. The exchange–correlation energy is expressed within the generalized gradient approximation (GGA),71 in the form of the Perdew–Burke–Ernzerhof (PBE) functional. All geometries were fully relaxed, with residual forces smaller than 0.1 eV Å−1. We used a grid cutoff of 150 Ry for the grid integration, with the Brillouin zone sampled using a 10 × 10 × 1 k-point mesh within the Monkhorst and Pack scheme. We adopted a convergence criterion where self-consistency is achieved when the maximum difference between the output and the input of each element of the density matrix is less than 10−4 eV, in a self-consistent field cycle. We adopted a rectangular unit cell and imposed periodic boundary conditions. A vacuum region of 100 Å was added along the z direction to avoid artificial interactions between a monolayer and its periodic images.All first-principles calculations are based on density functional theory (DFT),72 The interactions between B, C, and N atoms were described with the Tersoff potential,73 using the parameters adjusted by Kinaci et al.74 Our MD simulations were performed using a reasonably small timestep of 0.1 fs, and proceeded in three steps:The MD simulations were carried out using the Large-scale Atomic/Molecular Massively Parallel Simulator (LAMMPS) code.5 steps in the NPT ensemble, using Nose–Hoover thermostats and barostats75-77 to set temperature and pressure values to 10 K and 0 Pa.(1) We first evolved the system for 2 × 105 steps in the NPH ensemble, using the same barostat described above to set the pressure to 0 Pa. Note that the barostat was only applied to the planar direction that is not under strain from this step onward.(2) We turned the thermostat off and then evolved the system for 2 × 106 steps using a strain rate of 10−6 fs−1 .(3) In the final step we maintained the barostat configuration described above, and then elongated the system for 3 × 1026 In contrast, for h-BN, we find our results predict a lower Young's modulus and higher tensile strength when compared to the experiments.23 However, we note that our results are consistent with other theoretical investigations30We used the above method to obtain stress–strain curves for graphene and h-BN. The MD results were then compared against experimental results, as detailed in the ESI.L = 2 nm) that are shown in L ≤ 50 nm), which are shown in L is constant and the graphene/BN concentration is variable, while in model-II the side length L is variable and the graphene/BN concentration is constant. In both models, we considered different geometries . In order to determine the mechanical properties of these structures, we applied a tensile strain along one direction and then calculated the resulting tensile stress to obtain a stress–strain curve. Next, we calculated the slope of the linear region to determine the Young's modulus (Y). The tensile strength (σ) and the ultimate strain (ε) were taken at the point where the stress reaches its maximum value. The calculated values of Y, σ, and ε are presented in Our calculations were performed in two stages. In the first stage, we performed DFT calculations and MD simulations to determine the mechanical and electronic properties of the small hybrid h-BN/graphene sheets L =  nm that x) and zigzag (y) directions, as illustrated in One final remark is that the methodology used to obtain the stress is different for DFT and MD simulations. For the DFT approach, we first increased one side of the simulation box by 1% and then relaxed both the other side and the atomic positions until the calculation converged. After that, we increased the simulation box by 1% again and repeated the process. For the MD simulations, we increased one direction of the simulation box continuously at a fixed strain rate, and used a barostat to keep the other direction relaxed. Meanwhile, we allowed all atoms to freely evolve. Strain was applied along the armchair . For each structure, we obtained stress–strain curves for both the armchair and zigzag directions which are shown in 74 As a consequence, the difference between Young's modulus values obtained through DFT and MD is only ∼10%  and tensile strength (σ) values obtained for L2nm–C6, L2nm–C24, and L2nm–C54 are summarized in σ. Note that the same behavior was observed for h-BN, graphene, and other h-BN/graphene sheets, since this weak anisotropy is related to the hexagonal structure of the unit cell.22–25,27–29,44–50 Our results also show that the Young's modulus values increase with increasing concentrations of graphene, regardless of the direction of applied strain. In contrast, we find no conclusive relation between the tensile strength and ultimate strain with the number of C atoms. These results are in agreement with others previously reported for graphene/BN sheets, which found higher Young's modulus for higher concentrations of graphene.44,45,48 After the stress reaches its maximum value, for strain values of 21% for the armchair and 15% for the zigzag direction, the h-BN/graphene sheets fracture , which are shown in L = 40 nm) while we vary the domain size  while we vary the sheet size (10 ≤ L ≤ 50 nm). The goal of model-II is to investigate possible size effects . We also consider whether the shape of the domain affects the mechanical properties for both model-I and model-II. The mechanical properties of the larger sheets are summarized in Let us now discuss the mechanical properties of the larger hybrid sheets . Finally, note that, unlike the Young's modulus, the tensile strength has a weak dependence on the domain shape, being overall higher for triangular shapes when comparing structures with similar γ values. Furthermore, while we did not study the effect of the shape on the electronic properties, previous investigations on this topic found that shape affects these properties.33Moving on to the tensile strength results for model-I, they are summarized in γ, notice that the Young's modulus remains constant as the size of the nanosheet increases, for all pulling directions and domain shapes. Inspection of Y is higher for circular and lower for triangular domains. However, it is important to remark that although γ is constant for structures with the same domain shape, it does vary for structures with different domain shapes. For model-II, we used γ values of 0.19, 0.08, and 0.11 for circular, triangular, and star-shaped domains, respectively. Hence, model-II structures with circular domains are stiffer simply because they contain more graphene. With regard to the tensile strength, we find no clear relationship between this quantity and the nanosheet size in the results presented in Regarding model-II, we find that the mechanical properties of the BN nanosheets with graphene domains do not depend on the sheet size, as long as the graphene concentration remains constant as the sheet size varies. Young's modulus results for model-II are presented in We have also considered systems with lower symmetry than those discussed so far. We used MD simulations to study h-BN monolayers with (i) asymmetrical graphene domains and (ii) rotated triangular graphene domains. The results obtained for structures with lower symmetry are presented and discussed in the ESI.x-direction, we find that each thin segment is under the same strain (as the entire sheet is under the same strain). Hence, as the regions above and below the graphene domain are composed entirely of the less stiff h-BN, lower stress values produce the same strain. Consequently, fracture starts in the middle region, where stress is higher on average (t = 0 ps). However, note that strain is not particularly high at the graphene–BN interface, where the fracture always begins. This region is where we find the weaker hybrid bonds, and the C–B bonds were always the first to break. To understand this result, we note a previous DFT study that calculated first-neighbor bond energies for C–C, B–N, C–N, and C–B bonds, for hybrid BxCyNz structures.78 The authors found that the C–B bond had the lowest bond energy. Fracture then propagates perpendicular to the applied strain (t = 0.2 ps). This result is in agreement with the previously discussed DFT results. As the process continues, stress decreases (t = 0.4 ps).Eg against the strain for the armchair and zigzag directions. These values were calculated from the electronic band structures with 100 k-points along the Γ–X–Z direction. For zero strain, each structure has the same value of Eg in both directions, as expected. The gap values for L2nm–C6, L2nm–C24, and L2nm–C54 are 3.52 eV, 2.27 eV, and 1.56 eV, respectively. Comparing the smaller graphene domain with the others, we find a decline in the band gap of 35.5% for L2nm–C24 and 55.7% for L2nm–C54. This energy gap reduction for increased concentrations of graphene is in agreement with the results found by Manna et al.33 and Azevedo et al.37Ef are in general associated with C atoms. On the other hand, B and N atoms contribute to the electronic states in the conduction and valence bands, respectively. In agreement with the PDOS, the localized density of states (LDOS) results presented in z orbitals from C atoms. These results are common features of hybrid nanostructures with stoichiometry BxCyNz.12,36,37 Finally, note that we did not calculate the band gap of the larger structures presented in Let us now discuss the DFT predictions for the electronic properties of the structures shown in L2nm–C6 for strain values >16%, where we observed a sharp gap decline. The following discussion on the band gap variation with strain is concentrated on results for values below 15%, since fracture becomes an issue for some structures for larger strain values. In general, the energy gap decreases slightly more when the strain is applied to the armchair direction. Furthermore, we observe that the band gap of the considered structures varies weakly or moderately with deformation. These results can be seen more easily in 79 In contrast, the band gap of other 2D materials can vary by more than 50% for uniaxial strain values around 10%. For example, this is the case for many transition metal dichalcogenides80 . Finally, when comparing the effects of composition and strain, we find that the former has a more substantial influence on the band gap . We can also combine both effects to decrease the band gap from 3.5 eV to values below 1.0 eV.In general, we observe that the energy gap decreases gradually as the applied strain increases, except in the case of 81 and consequently, the numbers presented here are underestimated. However, we remark that other articles calculated band gaps of strained structures using different methods and found that, while the band gap values depend on the method employed, the general trends do not.79,82 The size of the structures considered here in the DFT calculations (160 atoms) hampered the use of more accurate methods.An important note is that the DFT method underestimates the band gap of semiconductors,83 Regarding requirement (i), our calculations show that the investigated monolayers can be extended by at least 10% before undergoing permanent deformation. Concerning condition (ii), although the hybrid sheets are very stiff, they are also very thin, requiring small forces to deform . Finally, regarding requirement (iii), our results show that the considered structures remain semiconductors over a wide range of strain values. They also reveal that it is possible to control the band gap of the considered monolayers by changing their composition. Semiconductors have been used in wearable electronics as photodetectors,83 as transistors,85 and as components in sensors.86 Therefore, h-BN nanosheets with graphene domains may play an important role in wearable electronics.Altogether, the calculated properties of the h-BN nanosheets with graphene domains indicate possible uses in wearable electronics. For such applications, a material needs to (i) be stretchable, (ii) be flexible, (iii) operate under large deformation.IV.In summary, we combined DFT and MD simulations to investigate the mechanical and electronic properties of h-BN nanosheets with graphene domains under strain. Regarding the mechanical properties, we found good agreement between the DFT calculations and MD simulations – the former was used to study smaller systems and the latter to study both smaller and larger structures. We found that Young's modulus values depend only on the relative concentration of graphene and h-BN. It does not depend on either the shape of the domain or the size of the structure (provided the h-BN and graphene domain are increased/decreased proportionately). Our results also revealed that the tensile strength is slightly higher for triangular graphene domains, particularly when applying strain to the armchair direction. In the case of DFT calculations, the system fractures for strain values between 15% and 22% at the junction between h-BN and graphene. In the case of MD calculations, the ultimate strain was slightly lower, but fracture also starts in the same region, even though stress is not higher at the interface. Fracture begins at the junction between h-BN and graphene because it is the location of the weaker C–B bonds.Regarding the electronic properties, we found that the band gap depends more on the composition of the hybrid sheets than on the external strain. Our results also show that the band gap of structures with more carbon atoms is more sensitive to the applied strain. However, the variation of the band gap with the external strain is still relatively small compared to other 2D materials. Furthermore, when composition and strain are combined to modify the band gap, we can obtain a broad range of values (from 1.0 eV to 3.5 eV). The proposed hybrid monolayers combine elasticity (the nanosheets can withstand considerable stress before plastic deformation), atomic-scale thickness , and adjustable band gap . Those properties indicate that the h-BN nanosheets with graphene domains could be promising materials for applications in wearable electronics.There are no conflicts to declare.RA-011-D1RA05831B-s001


Developing cheap and earth-abundant electrocatalysts with high activity and stability for oxygen reduction reactions (ORRs) is highly desired for the commercial implementation of fuel cells and metal-air batteries. Tremendous efforts have been made on doped-graphene catalysts. However, the progress of phosphorus-doped graphene (P-graphene) for ORRs has rarely been summarized until now. This review focuses on the recent development of P-graphene-based materials, including the various synthesis methods, ORR performance, and ORR mechanism. The applications of single phosphorus atom-doped graphene, phosphorus, nitrogen-codoped graphene , as well as phosphorus, multi-atoms codoped graphene  as catalysts, supporting materials, and coating materials for ORR are discussed thoroughly. Additionally, the current issues and perspectives for the development of P-graphene materials are proposed. The energy conversion technique is one of the most important parts that significantly influence the development of human society. The conversion efficiency of converting chemical energy into electrical energy is many folds higher than that of conversion of chemical energy into thermal  or kinetic energies . Accordingly, great attention has been given to the study of fuel cells and metal-air batteries as the next-generation energy conversion techniques because these devices can directly convert chemical energy into electric energy with nearly zero carbon footprint ,6,7,8,9.Graphene-based materials have attracted extensive interest in the field of electrocatalysis due to their fascinating mechanical, electronic, and thermal properties ,21,22,23Among these elements, nitrogen is by far the most studied dopant because of its similar atomic size to carbon and slightly higher electronegativity than carbon. The ORR performance can be remarkably enhanced in nitrogen-doped graphene (N-graphene), which has been demonstrated and reviewed in previous works ,28,29. B2 hybridized carbon to sp3 state. These lattice distortion-induced defects can serve as the active site for ORR. Moreover, the phosphorus in P-graphene could act as a donor and show an n-type nature because of its weak attraction with the valence electrons in the M-shell. To further confirm the electronic structure difference between P-graphene and pristine graphene, the density functional theory (DFT) simulations are conducted. As in the electrostatic potential map in Phosphorus is located on the periodic table in Period 3 and Group 5. Like its congener, nitrogen, phosphorous has five electrons in the valence shell. The geometric and electronic effects of phosphorus doping differ from nitrogen doping because of its large atom size and low electronegativity compared to both nitrogen and carbon. As seen from the STEM images in Accordingly, in this review, we discuss the recent studies of phosphorus-doped graphene, including phosphorus mono-doped (P-) and phosphorus with other elements codoped  graphene. Herein, the works reported on phosphorus-doped graphene served as metal-free catalysts, supporting materials, and coating materials are reviewed. The synthesis methods, catalytic mechanism, and ORR performance are systematically summarized. The synergistic effect of inducing defects and electron distribution on the ORR activity and durability are also discussed. Finally, the current challenges and future perspectives in the field of P-graphene ORR catalysts are proposed.It is well known that there are a large number of synthesis strategies to obtain graphene. According to the formation mechanism and the source of the carbon , the synthesis strategies can be divided into two types: top-down and bottom-up. The intrinsic properties and electronic structure of as-prepared doped graphene largely depend on the pristine graphene. Therefore, we divide these doping methods into two categories, as shown in The in situ doping is mainly based on the bottom-up pathway to produce graphene from small carbon molecules. Theoretically, the structure of the doped graphene can be fabricated on an atomic level. The doped graphene with high purity and fine structure can be obtained. However, this kind of method suffers from low yield, high cost, and difficulties in scaling up. Moreover, phosphorus is not easy to be introduced into the carbon skeleton because of its larger atom size than that of carbon. Engineering-desired, doped structures remain a challenge and detailed theoretical and experimental investigations are warranted.4/PH3/H2, and the phosphorus content is about 0.1%. In the prepared P-graphene, the phosphorus atoms are located at the edge of graphene as phosphorus oxides, such as C–P=O or P–Ox. Triphenylphosphine (TPP) [Chemical vapor deposition (CVD) is one of the most used technologies to produce low-dimensional materials, such as carbon nanotubes and graphene. It is based on the chemical reaction of the precursors in the vapor phase . As showne (TPP) ,40 can aAfter the graphene is grown on the catalyst surface, additional energy from external sources, including heating , ions imAlthough high-quality doped graphene films can be obtained by CVD, complicated procedures, such as substrate pretreatment, or protection layer coating are needed. How to transfer the doped graphene from the catalyst surface onto arbitrary substrates without damaging the graphene is still a million-dollar question. Furthermore, the precise controlling of the P–C bond configuration is desired for further study.2HPO4 was used as a phosphorus source [3/C2PO2 and C3PO structures. Similar methods were used and the tetrakis(hydroxymethyl)phosphonium chloride [2HPO4 [The P-doping can also be achieved simultaneously during the formation of graphene by stepwise solution chemistry of organic precursors. As shown in s source . After hchloride , Na2HPO4e [2HPO4 , and sode [2HPO4  were useTop-down approaches mean the doped graphene comes from the graphite and graphite derivatives. To extract the graphene from graphite and graphite derivatives, different methods, such as sonication and hydrothermal, were used. The precursor mainly determines the structure and intrinsic properties of the as-prepared doped graphene. Graphite oxide (GO) is the most used precursor for doped graphene. Thus, we devote a separate chapter to this method. These methods are always cost-effective, easy to scale up, and have high yields. Yet, it is difficult to precisely control the configuration of the doped graphene.Graphite oxide (GO), which comes from the oxidation of layered graphite mainly by a modified hummers method, is a promising precursor and versatile building block for chemically converted graphene in large volumes owing to its tailorable surface chemistry . After oP-graphene can be obtained by simply mixing the phosphorus source with GO in solution . The XPSMoreover, different combinations of phosphorus sources and heteroatom sources are used to prepare phosphorus and heteroatoms codoped graphene, such as P, F-graphene, P, N-graphene, and P, S-graphene ,66,67,683–P=O, C–P–O, C–O–P. The presence of these phosphorus configurations plays a key role in electrochemical performance. We will discuss these impacts later in this article.Because of the abundant oxygen-containing functional groups on GO, there is a considerable amount of oxygen content in the final P-graphene. Thus, there should be different P–O bonds and P–C bonds, such as C3PO4) was used as a phosphorus source and electrolyte. Ball milling of graphene stack and red phosphorus were also used to produce P-graphene [Graphite is thousands of graphene layers stacked together. The graphite was also used as the starting material to produce P-graphene by an electrochemical exfoliated method ,70. Becagraphene . Generalgraphene , up to 0graphene ).To enhance the phosphorus-doping content, fluorography was used . As showMetal-free catalyst based on carbon nanomaterials has been studied intensively as an emerging class of ORR catalysts since 2009 . Recentl3PO, C2PO2, and CPO3. Thus, the polarized phosphorus serves as a bridge between the oxygen atoms and the carbon atoms. The carbon atom that bonded with a phosphorus atom and oxygen atoms is positively charged. Therefore, the active sites are these positively charged carbon atoms. P-graphene shows a one-step 4 e− pathway and comparable ORR catalytic performance than that of the Pt/C catalyst. As shown in − and 4 e− pathways  . The phopathways c. Moreovpathways d. With tFurther theoretical and experimental works were conducted  annealed under 1000 °C is close to that of the commercial Pt/C (−0.02 V vs./AgCl) and shows better electrochemical stability . This is mainly related to the different charged sites of the graphene due to the different heteroatoms doping.Furthermore, the third element can also be introduced into graphene lattice to get tridoped graphene. As shown in The nanosized elemental metal, metal oxide, metal nitrides, and metal phosphides have potentially served as ORR catalysts. However, these nanoparticles are prone to agglomerate and aggregate during the synthesis and operation process because of the high surface energy. Besides, these nanosized catalysts always suffer from limited active surface areas and poor charge conductivity. Therefore, the substrate with a high surface area is generally used to disperse the nanoparticles. Doped graphene is a good candidate to accelerate the homogeneous dispersion. In particular, phosphorus atoms with multielectron orbital in P-graphene can lead to enhanced interactions between the catalyst and supports. The surface chemical state of P-graphene can not only affect the morphology but also modulate the electronic structure of the catalysts. A few research works have been devoted to using the P-graphene as catalyst support for ORR ,105,106.x@MnPy nanoparticles on P, N-graphene were obtained by annealing the mixture of GO and the desired chemicals. The phosphorus is (NH4)2HPO4 and Na2HPO2·H2O. The MoPx@MnPy homogeneously disperse on the P, N-graphene due to the negatively charged state of nitrogen and phosphorus codoping. There are P–C, Mn–P, Mo–P, P–O, P–O–H, and PO4 in the hybrid structure, indicating the enhanced interaction between MoPx@MnPy and P, N-graphene. Hence, the composite shows good electrocatalytic activity and stability. Besides, different CoMn2O4, cobalt, Co2P, and cobalt phosphate nanoparticles are also dispersed on the P, N-graphene [As shown in graphene ,105,106.Encapsulating the active nanosized catalyst into carbon materials is another way to resist aggregation and agglomeration by dual chemical and physical protection. Furthermore, the structural and electronic properties of the active catalyst can be modulated with the contact of the doped carbon materials. The carbon materials outside the catalyst can be called graphene or carbon shells in different literature. When the thickness of this carbon layer is very thin, it is called graphene in this review.3P, which were obtained by self-assembly, and 1-hydroxyethylidene-1,1-diphosphonic acid (HEDP) were annealed under N2 atmosphere to obtain Cu3P@P, N-graphene. The TEM image shows that the Cu3P NPs are covered with a thin carbon shell. The N and P atoms in the precursor play a vital role in incorporating Cu3P NPs into the carbon network and in situ doping nitrogen and phosphorus. A strong interaction can be found between the Cu3P and P, N-graphene because of the chemical bonding of copper with phosphorus. The optimized Cu3P@P, N-graphene shows comparable catalytic activity to that of the Pt/C catalyst because of the synergistic effect between the Cu3P and P, N-graphene shell. Later, similar strategies were used by different groups to prepare Co@N, P, S -graphene, FeCo@P, N-graphene/N-CNTs, and FeP@P-graphene [As shown in graphene ,109,110.Regardless of abundant works done on the P-graphene, continuous efforts are needed to rationally design and fabricate the highly efficient ORR catalysts. Some current challenges and future perspectives are listed below.First, the precisely controlled synthesis for heteroatom-doped graphene remains a key challenge and needs to be explored further. In present catalysts, there are many types of dopants and various defects. The electrochemical reaction is delicate, so it is difficult to identify the real factors that are responsible for enhancing the ORR activity. The synergistic mechanism is one of the most used explanations for the improvement of ORR. However, it is essential to clarify the effect of a specific chemical bond. Doping desired heteroatom(s) with a dominated specific type of chemical bond in a particular position is the first step to studying the doping effect. Currently, GO is being used as the carbon support in most of the research related to electrocatalysis. The conventional pyrolysis of GO cannot meet the demands. The changing structure and surface chemistry of GO during thermal treatment should be tuned. In addition, the controllable synthetic routes should be further developed. For example, magnetron sputtering is one of the promising physical vapor deposition techniques for the preparation of metal, alloy, and compound thin films. Under a high vacuum environment with a magnetically confined plasma, the atoms from the target can be sputtered onto the substrates. P, S, F, and N are commonly used as dopants by this technique ,116,117.Secondly, more efforts should be made on the physical property evaluations of P-graphene, such as thermal stability, wettability, and thermal conductivity, which also play important roles in ORR performance. It is also important for the synthesis routes. For example, the defects can be found in sing layer-graphene prepared by CVD at ~500 °C, while defects appear in bilayer graphene fabricated at ~600 °C . More stThirdly, identifying the types and quantifying the content of dopants accurately is still a big challenge. XPS is a frequently used technology to study the information of chemical bonding. However, even with a similar catalyst, the analysis of XPS data varies between different research groups. Other advanced characterization techniques, such as X-ray absorption spectroscopy, should be involved.Fourthly, the underlying catalytic mechanism, as well as the real active sites for ORR of doped graphene is not yet identified clearly. Although some theoretical simulations have been done, there is still a big gap between the theoretical model and the actual working state of the catalysts. Some well-designed experiments and in situ characterizations should be done to unveil the catalytic mechanism.Last but not least, the building of the multiple doped structures or the composites will play a more important role. For example, we have recently synthesized P, N-graphene dots/N-3D-graphene . We founIn summary, doping graphene with phosphorus and various heteroatoms has been proven as a promising way to enhance the ORR performance. More importantly, the experimental and theoretical studies of P-graphene are essential to study the catalytic mechanisms towards ORR, thus providing a general strategy to obtain the advanced non-noble metal catalysts for the application of energy conversion devices, such as fuel cells and metal-air batteries. In this review, the recent advances in P-doped graphene with an overview of the various synthesis methods, ORR performance, and ORR mechanism are discussed. Furthermore, we believe that these achievements can be used in other energy conversion and storage fields. Further efforts on the study of P-graphene are warranted.


Three cases are considered, each of which addresses a specific configuration of a rippled graphene sheet located on a flat substrate. The transitional case assumes that both the graphene sheet length and substrate length are constrained. The substrate constrained case assumes only the substrate has a constrained length. Finally, the graphene constrained case assumes only the length of the graphene sheet is constrained. Numerical results are presented for each case, and the interpretation of these results demonstrates a continuous relationship between the total energy per unit length and the substrate length, that incorporates all three configurations. The present model is in excellent agreement with earlier results of molecular dynamics (MD) simulations in predicting the profiles of graphene ripples.The calculus of variations is utilised to study the behaviour of a rippled graphene sheet supported on a metal substrate. We propose a model that is underpinned by two key parameters, the bending rigidity of graphene  A variational model is proposed to describe rippled graphene on a substrate, and obtain the relationship between total energy and substrate length. This two-dimensional structure endows graphene with many useful electronic and mechanical properties.16 as well as graphene on a substrate.17 Experimental studies find that the electronic properties of graphene can be affected by the range and height of these ripples.18,19 Gui et al.20 use first-principles calculations to predict the electronic properties of a rippled graphene sheet where a band gap opening is observed in the rippled graphene. They evaluate a direct band gap at 0.93 eV which indicates rippled graphene may be a highly tunable semiconductor. They also report that the band gap increases proportionally with ripple amplitude, which is the maximum distance between the graphene sheet and the substrate.However, graphene does not always remain planar. For example, ripples are observed in suspended graphene sheets,21 When a graphene sheet is placed on a substrate the amplitude of the ripples depends on the substrate properties such as roughness and interfacial van der Waals interactions.17 Moreover, the overall size of the graphene sheet has a significant impact on the amplitude of ripples, with the ripple amplitude increasing proportionally with the size of the sheet.22The pattern of ripples in graphene may be affected by a variety of factors, including temperature, substrate material, and the size of the graphene sheet. For example, suspended graphene remains flat with no ripples when the temperature is close to absolute zero, but ripples appear as the temperature increases.et al.10 evaluate the bending rigidity of a single-layer graphene at 1.44 eV. They compare this result to earlier obtained values for the bending rigidity of graphene which range from 0.80 to 1.60 eV. To cover this range of values, we will consider various bending rigidities for the rippled graphene in the present paper.The bending rigidity of graphene plays an important role in the conformation of ripples. Using density functional theory calculations, Wei et al.23 develop a model which describes a graphene sheet folded onto itself where two energies are taken into account. The first is the elastic energy that arises when bending a graphene sheet which is represented mathematically by integrating the square of the line curvature over the length of the total curve and scaled by the bending rigidity of graphene. The second energy is the van der Waals (vdW) interaction between the flat section of the graphene layers, prescribed as the integral of a Heaviside unit step function multiplied by the vdW energy per unit area. An extension of the fold model was further developed by Dyer et al.24 where either a carbon nanotube of circular or elliptical cross section is encapsulated in the graphene fold. A comparison between MD simulations and the mathematical model shows excellent agreement between the two approaches in determining the fold shape.The energies we consider when modelling ripples in graphene are the same as those considered when modelling a fold of graphene. Cox In the following section, we formulate a general mathematical model of a single graphene ripple on a substrate. In Section 3 we non-dimensionalise this model and employ the calculus of variations assuming a fixed length (isoperimetric) constraint of both the rippled graphene sheet and the substrate. In Section 4 we relax the isoperimetric constraint on the sheet length, but maintain a fixed substrate length. The substrate length is varied in Section 5 and an isoperimetric constraint is applied to only the sheet length. The solutions and associated numerical details are provided with each case. The main results and the relationships between the three cases considered here are described in Section 6. Discussion and some concluding remarks are made in Section 7.2z-direction which reduces the problem to two dimensions. Also, the reflective symmetry about the y-axis allows us to consider only solutions in the right half plane. We further divide the solution curve into three sections. The first section C1 is the curve from the point  to the point  where the line curvature is strictly negative. The second section C2 is from  to  where the line curvature is strictly positive. The third section C3 is the horizontal line from  to  where the line curvature is zero. The concatenation of these sections is denoted by C = C1 + C2 + C3. In these position vectors we use h to denote the ripple height, and δ to denote the separation distance between the substrate and the flat section of the graphene sheet.We propose a general formulation to determine the conformation of a rippled sheet of graphene on a substrate for each of the three cases mentioned above. We consider two different metal substrates: Cu(111) and Ni(111). The geometry of the rippled graphene sheet is shown in et al.,23 the dominant energies for our model are the elastic bending energy and the vdW interaction energy. The elastic energy Ee is modelled withwhere γ is the bending rigidity of graphene, s is the arc length, and κ is the line curvature of y = y(x) which is given byWith reference to the model developed by Cox γ, namely, 0.8, 1.0, 1.2, 1.4, and 1.6 eV. We model the vdW interaction energy between the graphene and the substrate aswhere ξ denotes the vdW energy of graphene–substrate interactions per unit area, and u(x − x1) is a Heaviside unit step function. Our model assumes that the elastic energy dominates in sections C1 and C2, and the vdW interaction dominates in C3. We comment that the curvature and therefore the elastic energy vanishes in C3, but discarding the vdW interaction in sections C1 and C2 is an approximation we make in the modelling. Therefore, the total energy per unit length may be expressed byx = x1 the value of x1 is not prescribed and we have a natural boundary condition on x. Here dots denote differentiation with respect to x.As mentioned in Section 1, the bending rigidity of graphene reported by various authors ranges from 0.8 to 1.6 eV as presented in h for each case of a rippled graphene sheet supported on a different substrate, and the reason for adopting these particular values will be explained in Section 4. The prescribed ripple height, h, and empirical values for the separation distance between graphene and the substrate, δ, and the vdW energy of graphene–substrate interactions per unit area, ξ, corresponding to each substrate are shown in For the purpose of comparison we prescribe a particular value of 3C, given byWe now apply the calculus of variations to determine the conformation of a rippled graphene sheet and derive a solution for which the functional x2 is fixed, but x1 varies . Therefore the new functional we wish to minimise is of the formλ is a Lagrange multiplier corresponding to the isoperimetric constraint. We nondimensionalise X = x/α and Y = y/α where α is a scaling factor, and therefore ÿ = Y′′/α and ẏ = Y′, where primes denote derivatives with respect to the nondimensional X coordinate. The substitution of these new variables into Moreover we assume that X2, and divide both sides of αξ. Thus we derive the nondimensionalised energy functional asFurthermore we subtract the fixed energy at  and μ = λ/ξ so thatY(0) = h/α and a natural boundary condition applies at X = X1. To simplify the following calculations, we use f to denote the integrand part of We let Y explicitly, then on integrating X, we obtain the first integralβ is a constant. Further, the integrand has no explicit dependence on X, this provides the additional first integralH is a constant. On considering the corresponding second-order variational problem, we may derive the standard equation for the first variation of Êtot aswhere P = fY′ − d(fY′′)/dX, Q = fY′′, and H = (Y′P + Y′′Q − f). Here the natural boundary condition, which applies when the X-coordinate at the end point X = X1 is not prescribed, requires H = (1 − μ). Therefore by combining Extremals of or equivalentlywhich is an equation relating the nondimensional curvature  PBM data was replaced with SVG by xgml2pxml:<glyph-data id="z.kgcir.i" format="PBM" resolution="300" x-size="10" y-size="11" xml:space="preserve">00000110000000100100000000000000000000000111000110001100100000100100000011100000011011000001000101000100011000</glyph-data><svg xmlns="http://www.w3.org/2000/svg" version="1.0" width="14.727273pt" height="16.000000pt" viewBox="0 0 14.727273 16.000000" preserveAspectRatio="xMidYMid meet"><metadata>Created by potrace 1.16, written by Peter Selinger 2001-2019</metadata><g transform="translate scale" fill="currentColor" stroke="none"><path d="M400 840 l0 -40 -40 0 -40 0 0 -40 0 -40 40 0 40 0 0 40 0 40 80 0 80 0 0 -40 0 -40 40 0 40 0 0 40 0 40 -40 0 -40 0 0 40 0 40 -80 0 -80 0 0 -40z M80 520 l0 -40 40 0 40 0 0 -120 0 -120 -40 0 -40 0 0 -120 0 -120 40 0 40 0 0 80 0 80 40 0 40 0 0 40 0 40 40 0 40 0 0 -40 0 -40 40 0 40 0 0 -80 0 -80 80 0 80 0 0 40 0 40 40 0 40 0 0 40 0 40 -40 0 -40 0 0 -40 0 -40 -40 0 -40 0 0 80 0 80 -40 0 -40 0 0 40 0 40 40 0 40 0 0 40 0 40 40 0 40 0 0 40 0 40 80 0 80 0 0 40 0 40 -80 0 -80 0 0 -40 0 -40 -40 0 -40 0 0 -40 0 -40 -40 0 -40 0 0 -40 0 -40 -80 0 -80 0 0 40 0 40 40 0 40 0 0 80 0 80 -120 0 -120 0 0 -40z"/></g></svg> and the first derivative of Y. In order to obtain the parametric solutions, we make the substitution of Y′ = tan θ which yieldsWhich after the substitution of whereupon we define a new parameter ψ, such thatThe above expression may be written as followsThis leads to the succinct formula = cos(θ)dθ/dX = sin(θ)dθ/dY, we derive two first order differential equations, namelyNow, using the fact that θ to ϕ such that θ = 2ϕ − ψ, thus where F and E represent the incomplete elliptic integrals of the first and second kind, respectively, with elliptic modulus k, and c1, c2 are arbitrary constants of integration.To facilitate further integration we change the parametric variable of our equations from We now define two functions which play a significant role in determining our parametric solutions, namelyϕ0 = sin−1(1/k) corresponds to  which is the boundary point between the two curves C1 and C2, which has coordinatesThe line curvature changes sign at X0, Y0) where ϕ varies over the range . Thus our solutions may be written in terms of X0 and Y0 asMoreover, the solutions are shown to have a rotational symmetry about this critical point (YC2(ψ/2) = δ/α, we may obtainNow, using the boundary condition which introduces the Y component of the critical point as the midpoint of the ripple amplitude and the flat section of the graphene sheet. We also comment that x = αX and y = αY, and so multiplying these solutions by the scaling factor  recovers the re-dimensionalised solutions.Hence we may rewrite ψ and μ. We may find their values by solving the following system of equations numericallywhere δ takes a specific empirically determined value for each substrate as given in L is the assumed fixed arc length of C. After finding the unknowns ψ and μ, the solution is fully determined and representative plots are shown in Lrip by integrating ds over the curve C1 + C2,We are left with two parameters to determine, namely θ = 2ϕ − ψ we deduce thatwhere . Hence, the total half arc length of the ripple is given byAgain using the substitution of C1 + C2, as followsand therefore, the elastic energy is given byWe may also calculate the elastic energy by integrating the square of the line curvature over the total length of Hence, the total energy per unit length for this model is given byx-component of the critical point x0 and the total half arc length of the ripple Lrip are presented in γ. The vdW interaction ξ takes a specific empirically derived value for each substrate as given in Numerical values for the 4C to allow graphene to overhang the substrate. Therefore we discard the variation of x2 and substitute λ = 0, and  into We now consider the special case, shown in μ = 0, and so by the same reasoning as in Section 3 we obtainWe note that this expression is identical to A = 2(sec ψ)−1/2, B = 0, and . Again the line curvature  changes sign at ϕ0 = sin−1(1/k) = π/4 which corresponds to the point  and has similar coordinates to that in Consequently, we redefine some of our parameters so that now X0, Y0), and the variation of ϕ over , the curves C1 and C2 still can be written asSimilarly, with the rotational symmetry of the solutions about the point  = δ, the value of ψ may be determined numerically. After determining ψ, the solution is fully determined and representative plots are shown in γ = 0.8 in the neighbourhood of the critical point becomes infinite when the ripple amplitude h ≈ 8.0 δ for the Cu(111) substrate, and h ≈ 6.7 δ for the Ni(111) substrate where δ takes a specific empirical value for each substrate as presented in h for each corresponding substrate in all plots to enable comparison between various values of the bending rigidity γ.In this case we have only one parameter to determine, namely Lrip and the total energy Etot for this case can be calculated from μ = 0 and ϕ0 = π/4. Thus, the total half arc length of the ripple is given byand the total energy per unit length is given byThe total half arc length of the ripple x-component of the critical point x0 and the total half arc length of the ripple are shown in γ. Again the vdW interaction strength ξ takes a specific empirically derived value for each substrate as given in Numerical values for the 5C, we take into account the variations of x2, and assume it varies proportionally with x1, that isIn this case, as shown in s = α(1 + Y′2)1/2dX. We now divide both sides by ξα and set . Further, the aim is to determine the shape of Y = Y(X) which minimises Êtot, and here λ has no impact on the solution, so we choose λ = ξ to simplify the calculation. The final form of the functional to be minimised isμ = 1. In a similar way as in the previous two cases, the use of calculus of variations then provides thatFollowing the nondimensionalisation described in Section 3, the total energy per unit length for this case isψ = π/2 and β acts in place of the term  as μ approaches 1. Therefore, we redefine the parameters A, B, and k in Similarly, by comparing (22)X0, Y0) represents the point of the rotational symmetry which again corresponds to ϕ0 = sin−1(1/k), and the coordinates of this point are given byThe critical point  = δ/α to rewrite the Y component of the critical point asNow, we use the boundary condition C1 and C2 may be obtained by scaling these solution by .The dimensional solutions for the curves β, and using the boundary condition YC2(π/4) = δ, the value of β may be determined numerically. After finding β, the solution is fully determined and representative plots are shown in Lrip for this case can be evaluated from ψ = π/2 and  to giveHere again we are left with only one parameter to determine, namely Etot for this case is obtained from Similarly, the total energy per unit length x-component of the critical point x0 and the total half arc length of the ripple Lrip are presented in γ. As before, the vdW interaction ξ takes a specific empirically derived value for each substrate as given in Numerical values for the 6L = 700 Å for the rippled graphene sheet on a flat substrate. The goal is to obtain a relationship between the variation of the substrate length and the total energy per unit length for the rippled graphene. The substrate half length Lsub is divided into three regimes that correspond to the point x2, where the graphene edge ends, which can be calculated by x2 = x1 + L − Lrip. Here we comment that x2s < x2g where x2s and x2g correspond to the substrate constrained case and the graphene constrained case, respectively.For the purpose of comparison, we assume a fixed half length Lsub < x2s, provided that as Lsub gets closer to x2s the total energy per unit length Etot decreases due to the vdW interactions that occur when Lsub increases. The graphene constrained case can be adapted to address the regime when Lsub > x2g, and here the total energy per unit length Etot remains constant since there are no more interactions that are involved for this section. Finally, the transitional case is used to evaluate the total energy per unit length Etot when Lsub ≈ x2. Multiple values for x2t were chosen from the interval  to fill the gap between the substrate constrained case and the graphene constrained case where x2t denotes the location of the graphene edge in the transitional case. The substrate constrained case was formulated to describe the regime where L = 700 Å will remain flat when it is placed on a flat substrate of the same half length. The substrate constrained case can be adapted to investigate the impact of reducing the substrate length on the graphene sheet. A graphene sheet of half length −1 when it is placed on copper, and at −63.5 eV Å−1 in the case of a nickel substrate. It increases when the substrate length decreases, as expected due to reduced vdW interactions between the graphene sheet and the substrate, as well as elastic energy stored in the ripple. The points marked with circles represent ) where L0sub denotes the substrate length when the gradient of the total energy per unit length curve approaches −ξ.The impact of reducing the substrate length on the total energy per unit length is presented in Starting from a flat graphene sheet laying on a shrinking substrate, the system might form a ripple or remain flat and beginning to overhang the substrate. The ripple formation pathway would follow the transitional case and the flat pathway would follow the substrate constrained case. In an ideal setting, ripples would not form because the energy require to follow the first pathway exceeds the energy lost in following the second. However, we comment that many experimental situations include confounding effects such as impurities and defects in the graphene structure which are not accounted for in this work.et al.28 They observed unexpected profiles of wrinkles with widths in the range of tens of nanometres, and heights in the range of a few nanometres. Their theoretical methods and MD simulations agree that the maximum width of wrinkles is less than 2.7 nm which is largely smaller than those experimentally observed. As they reported, the presence of interlayer molecules between the graphene and the copper substrate is the main mechanism that enlarges both width and height of graphene wrinkle. Since we don't take into account interlayer molecules, our model is comparable with their MD simulations without interlayer molecules. Their MD simulation model, as reprinted in γ = 1.20 eV and different heights are adopted to compare with the results of MD simulations of graphene wrinkles obtained by Wang et al.,28 as shown in −2 for the vdW energy of graphene–Cu(111) interactions per unit area, but their model used the value of 4.494 eV nm−2, both results are in excellent agreement in predicting the profiles of graphene ripples (or wrinkles).Recently, morphologies of graphene wrinkles on copper substrate were experimentally investigated by Wang 7y-axis allows us to consider only solutions in the upper right quadrant. We employ variational calculus to minimise the elastic energy arising from the curvature squared and maximising the vdW interaction energy between the flat section of the graphene and the metal substrate.In this paper, we develop a model for a rippled graphene sheet that is located on a flat metal substrate. The assumed translational symmetry along the ripple reduces our problem to two dimensions, and the reflective symmetry about the We account for the length of the substrate by considering three different cases. The first addresses the case when the edges of the rippled graphene sheet and the substrate coincide. The second case is when the edge of the graphene sheet overhangs the substrate edge, and the last case applies to a rippled graphene sheet for which the edge of the sheet does not extend to the substrate edge. We assume a fixed half length of the graphene sheet, and consider all three cases to obtain a continuous relationship between the total energy per unit length of the graphene and the length of the substrate as shown in ξ as a function of x and lead to an explicit dependence on x in the integrand part of Future research may address the phenomenon of ripple formation when the substrate shrinks by taking into account additional physical effects. For instance the model maybe extended to account for the change in the vdW interaction strength as the substrate length and density change. This will introduce There are no conflicts to declare.


Electrode structure is an essential factor affecting the performance of flexible humidity sensors. In this study, Cu and Cu2 + 1O electrodes were printed by the one-step method using laser direct writing technology to reduce the nano–CuO ink on flexible substrate PI and to be used for a humidity sensor. The resistance of the humidity sensors with nine various electrode structures was measured under the relative humidity (RH) of 16–78%. It was observed that all sensors showed good humidity sensing characteristics, and the sensitivity of the copper-based humidity sensor was not affected by the electrode structure under low humidity conditions but was significant under high humidity conditions. The sensor with the length of 1960 μm and the width of 120 μm shows the lowest sensitivity of 180.2 KΩ/%RH under 35% RH, and the sensor with the length of 2430 μm and the width of 180 μm shows the highest sensitivity of 1744 kΩ/%RH under 65% RH. It is expected that the results can provide an assessment of the performance improvement of the flexible humidity sensor and a reference for the research and development of intelligent wearable devices. Humidity sensors are widely used in smart wearables, environmental monitoring, healthcare, and food. With the rapid development of the smart wearable industry, the development of flexible humidity sensors with smaller feature sizes, lower costs, and better performance on flexible substrates has attracted much attention . The ele2O [The choice of a flexible substrate and conductive ink is critical in the manufacturing of flexible sensor electrodes by LDW technology. According to the characteristics of the ink absorption spectrum, a laser can be selectively absorbed or transmitted by ink, so the appropriate laser wavelength selection is necessary, which can effectively prevent the damage to the flexible substrate caused by the direct absorption laser. Polyethylene terephthalate (PET) , polyeth2O , and so 2O , is a fo2O , which pThe electrode structure is also an essential factor affecting the sensor’s performance. The interdigital electrode is a widely used electrode structure and includes three primary parameters: interdigital width, length, and thickness. The optimization electrode structure with excellent performance, such as precision and sensitivity, can be found by simulation software to simulate and analyze the electric field , capacitNano CuO ink was obtained by mixing 60 wt% CuO , 13 wt% polyethylene pyrrolidone , and 27 wt% ethylene glycol . To obtain a suitable concentration of ink for spin coating, the ink was diluted with DI water, mixed with magnetic stirring for 70 min, and treated by ultrasonic for 5 h.Commercial polyimide (PI) tape  with a thickness of 80 μm was pasted on a cover glass with a thickness of 0.15 mm. Then, the nano CuO ink was dropped to the PI film to rotate at a speed of 1000 rpm/min for 30 s and dry for 1 h in air. The LDW device, which was equipped with a continuous-wave laser diode with a wavelength of 650 nm  and the direct current regulated power supply, was used to irradiate the ink. Electrodes printed by LDW were connected to copper tape as electrical test leads, and the sensor was encapsulated with Teflon tape. The surface characteristics and element distribution of the sensor were examined by a scanning electron microscope (SEM)  with an energy dispersive spectroscopy (EDS). The specimens were coated with gold before observation, and the operating voltage was 7.0 kV. The crystal structure of the ink material was investigated by X-ray diffraction (XRD)  equipped with Cu Kα radiation (λ = 0.1540 nm), operating voltage 40 kV, and current 30 mA.To investigate the influence of a humidity sensor structure on its performance, the nine humidity sensors with various interdigital electrodes, whose interdigital widths respectively were 120 μm, 180 μm, and 240 μm and interdigital lengths were 1960 μm, 2210 μm, and 2430 μm, were designed based on 20 μm single lines. Four saturated salt solutions having specific relative humidity (RH), which were, respectively, sodium bromide (35% RH), sulfuric acid (55% RH), sodium chloride (65% RH), and potassium chloride (78% RH), were prepared as the performance test environment of the humidity sensor in the ambience . S is the sensitivity, R0 and R are the sensor’s resistance in air and humidity environments, respectively, and ΔRH is the variation in RH. Response time refers to the necessary time in which the sensor’s resistance is changed from the air to 90% RH, and recovery time refers to the essential time for which the sensor’s resistance is changed from 90% RH to the air.The sensor’s sensitivity S is defined as follows :(1)S=, (−111), (111), (−202), (202), (−113), (−311), and (220) crystal planes of CuO monoclinic crystals, respectively (PDF#80-0076), and in the laser irradiation area, the diffraction peaks at 2θ = 43.4°, 50.5°, and 74.2° belong to the (111), (200), and (220) crystal planes of Cu (PDF-#04-0836), which indicates that CuO is reduced to Cu. The diffraction peaks at 2θ = 29.6°, 36.5°, 42.4°, 61.5°, 73.6°, and 77.5° are attributed to the (110), (111), (200), (220), (311), and (222) crystal planes of Cu2 + 1O, which is the excessive metal defect of Cu2O [11), and 0 crystal of Cu2O . Cu nano of Cu2O .R0 − R)/R0), sensitivity, and response/recovery time of the sensors when relative humidity changes. The sensors show an excellent humidity response characteristic, and the normalized response value of the sensor increases with the increase in RH, as shown in The sensing characteristics of the prepared humidity sensor are evaluated and compared with commercial humidity sensors, including response recovery characteristics in the range of 35–78% RH. The humidity sensors with nine different electrode structures were numbered 1–9, and the electrodes were No. 1, 2, and 3 in the first row from left to right, respectively, and were No. 1, 4, and 7 in the first column from top to bottom in Reusability is also an important index to evaluate the performance of humidity sensors. As an example, under 35%, 55%, 65%, and 78% RH, the resistance response of the No. 3 humidity sensor, because the response of all sensors is similar, and the testing data of the commercial humidity sensor are shown in Water molecules are adsorbed on the surface of the electrode, forming a thin coating of water when the sensor works. The performance of the humidity sensor can be affected by the adsorbance of water molecules, and the adsorption efficiency of water molecules largely depends on the structure of the interdigital electrode. Under low RH, the sensitivity of all the sensors is less affected by the electrode structure, which indicates that the adsorption efficiency of water molecules is similar to the copper-based humidity sensors. However, the adsorption efficiency of the sensor with a longer interdigital length is more robust under high RH.0 is preferable to incorporate into Cu2O lattices as a dopant in an acidic medium to form n-type Cu2O [xO by LDW on a PC flexible substrate-covered Cu precursor film using the continuous wave of 808 nm. An integrated humidity sensor was prepared by combining a CuXO-rich porous sensing structure with a Cu-rich conductive structure which showed high sensitivity to human respiration. In this study, nine humidity sensors with various electrode structures were designed to analyze the influence of electrode structures on humidity sensing. As a rapid machining process, LDW has been used to manufacture conducting circuits [Copper, which has a low price compared with other precious metals, such as gold and silver, and has good electrical conductivity, is considered as an excellent material for preparing a conductive structure. Cupric oxide and cuprous oxide, as two oxides of copper, are widely used in sensing systems for their excellent photoelectric and gas responsiveness . Layeredype Cu2O . When thype Cu2O  printed circuits  and senscircuits  in the pcircuits , the redcircuits .The Cu and Cu2 + 1O electrodes were formed by one-step reduction in the nano CuO ink on a flexible PI substrate using LDW and were utilized as a humidity sensor. The results show that the narrowest interdigital width of the sensor is 120 μm, the most limited single line width is 20 μm, and the printing accuracy is ±1 μm. Nine humidity sensors with different electrode structures were prepared and the variation in their resistance was analyzed to be 16–78% RH. The sensitivity of the copper-based humidity sensor is hardly affected by the electrode structure under low RH, but is significantly affected under high RH. All sensors show excellent humidity sensing characteristics. The sensitivity of the sensor with a length of 1960 μm and a width of 120 μm is 180.2 KΩ/%RH and is the lowest under 35% RH, and the sensor with a length of 2430 μm and a width of 180 μm is 1744 kΩ/%RH and is the highest under 65% RH.


Our analysis also clarifies the tradeoff associated with the loss penalty and nonlinear benefit afforded by graphene patches integrated onto passive photonic circuits, thereby providing some guidelines for the design of hybrid integrated nonlinear devices, coated with graphene, or, more generally, any other 2D material.Recently, the nonlinear optical response of graphene has been widely investigated, as has the integration of this 2D material onto dielectric waveguides so as to enhance the various nonlinear phenomena that underpin all-optical signal processing applications at telecom wavelengths. However, a great disparity continues to exist from these experimental reports, depending on the used conditions or the hybrid devices under test. Most importantly, hybrid graphene-based waveguides were tested under relatively low powers, and/or combined with waveguide materials that already exhibited a nonnegligible nonlinear contribution, thereby limiting the practical use of graphene for nonlinear applications. Here, we experimentally investigate the nonlinear response of Si Graphene is the first and most mature of bidimensional materials that has been isolated ,2, and ion (SPM) ,13 or foon (SPM) ,14,15,16nstance, ).However, a great disparity exists in the literature regarding the nonlinear response of graphene , and theγ∼1 m−1W−1).Although Si is a mature photonic platform, which has been used with graphene in initial demonstrations ,12,13, i−1W−1 to −10 m−1W−1. Our simulations allow us to take into account the global response of the waveguide as well as to differentiate between the linear loss penalty and the nonlinear benefit induced by graphene on the whole waveguide structure. Surprisingly, the outstanding nonlinear response (−13 m2W−1 and up to −12 m2W−1) of graphene that was reported in a few papers [In this paper, we investigate the use of short graphene/Sit powers ,16. Our w papers  does notFor this study, we use Siapproach  for strasorption . A silichttps://www.graphenea.com/ onto the chip containing the waveguides with the patterned upper cladding. Although graphene covers the whole chip, this patterning allows us to restrict the interaction between graphene and the guided mode of the waveguides solely along the etched windows. This method ensures a relatively good control of both the position and length of the graphene area interacting with the waveguides, without additional post-processing steps after the graphene transfer, which might otherwise affect the graphene optical properties.Following this fabrication process, the top cladding of the waveguides was thus removed selectively see a,b in or−1 . We estimate a Fermi level of approximately 12 cm−1, expected for CVD graphene and this transfer process [Moreover, in order to ensure the quality of the graphene transferred onto the chip, Raman measurements were conducted after the transfer, and confirmed a high-quality monolayer graphene , with th process .The pump-probe experiments are performed by using a pulsed laser at telecom wavelengths that delivers 2 The interaction of the pump , after propagation along a waveguide of length L, is usually defined as [The idler signal is clearly detected at approximately 1541 fined as (1)CE′=PCE because the probe power at the output of the waveguides is decreased with respect to that at the entrance of the waveguides by the propagation loss. Equation  to numerically solve this system of equations. It allows us to change the parameters along the propagation direction according to the specific response of the local waveguide structure, while feeding the output signal of one given section as the input for the simulation of the subsequent one. CE as a function of the graphene patch length, and for different coupled pump powers.The loss at the etched interface corresponds to the estimated loss when the guided mode crosses the boundary between the clad section of the SiCE decrease primarily reflects the missing nonlinear contribution from the 1.7 From the simulated curves, the case CE measurements are relatively well reproduced considering −10 m−1W−1 < −1W−1 i.e., an absolute value almost one order of magnitude larger than that of the clad SiCE-enhanced reduction can be understood by the opposite contributions to the idler signal generation arising, respectively, from the hybrid graphene/SiAs observed on  reports ,23, and A comparison of our results with the literature on graphene-covered nonlinear waveguides is shown in −1W−1, i.e., several orders of magnitude less than the other values indicated in the Table, where up to a few thousand were reported. Considering the 2D nature of graphene covering the waveguide and assuming that the nonlinear response of graphene is a few orders of magnitude greater than the one of the underlying waveguide  of length ions see , we onlyOur analysis see  shows thMore quantitatively, we show that the 2D material patch increases the net idler generation of the overall waveguide as long as one of the following conditions are met, depending on the sign of In both cases, there is an optimum 2D material length that optimizes the nonlinear idler generation, which is found to beEventually, if the 2D material nonlinear response is strong enough to satisfy Equation , then thWe note that the constraint set by Equation , on γhybaveguide b, with raveguide a. Consisaveguide  also ten−1W−1 (if −1W−1 (if −1W−1), indicate that the benefit of graphene, in our chip, is not only overcompensated by its loss penalty, but is also not enough to compensate for the nonlinear effect induced by the surrounding sections of SiIn our experiments, One simple solution to experimentally observe a benefit of graphene on photonic circuits would be to shorten the surrounding bare waveguides or, more practically, to decrease its nonlinearity. For instance, other waveguide platforms might be used, such as SiOne oxide ,33, migh−1W−1) locally increased by an order of magnitude with respect to that of the underlying waveguide, and (2) a strong linear loss of graphene reducing the contribution of the following SiBy transferring graphene onto locally unclad Simaterial . From ou