Here, we report that the large-sized K+ is riveted in the prismatic Na+ sites of P2-Na0.612K0.056MnO2 to enable more thermodynamically favorable Na+ vacancies. The Mn-O bonds are reinforced to reduce phase transition during charge and discharge. 0.901 Na+ per formula are reversibly extracted and inserted, in which only the two-phase transition of P2 ↔ P’2 occurs at low voltages. It exhibits the highest specific capacity of 240.5 mAh g−1 and energy density of 654 Wh kg−1 based on the redox of Mn3+/Mn4+, and a capacity retention of 98.2% after 100 cycles. This investigation will shed lights on the tuneable chemical environments of transition-metal oxides for advanced cathode materials and promote the development of sodium-ion batteries.Layered transition-metal oxides have attracted intensive interest for cathode materials of sodium-ion batteries. However, they are hindered by the limited capacity and inferior phase transition due to the gliding of transition-metal layers upon Na 0.612K0.056MnO2 with large-sized K+ riveted in the Na-layers to enable 0.9 Na+ (de)insertion with a reversible phase transition of P2-P’2.High-capacity and structural stable cathode materials are challenges for sodium-ion batteries. Here, the authors report a layered P2-Na The battery performance and cost are mainly governed by the applied cathode materials7. A wide range of host materials, such as transition-metal oxides, phosphates, Prussian blue analogs, and so on, have been investigated for reversible (de)intercalation of Na+. However, the intrinsic large ionic radius of 102 pm and molar weight of Na+ over its Li+ counterpart impose penalties on the specific capacity or the structure stability of cathode materials, which limit the deployment of SIBs12. It is rational to design and synthesize cathode materials with more available sites for Na+ storage and robust crystal structure for Na+ (de)intercalation, ensuring their large specific capacities and long cycle life.Sodium-ion batteries (SIBs) have been considered as one promising next-generation battery system for large-scale energy storage and smart electric grid, due to the low cost and wide distribution of Na resources+ storage14. Sodium manganese oxide with a P2 phase (P2-NaxMnO2) is of significant importance for SIBs, which is consisted of two MnO6-octahedra slabs and sandwiched Na layers in each unit cell with the ABBA oxygen packing frameworks16. It usually delivers a capacity of ~160 mAh g−1, corresponding to insertion of 0.60 Na+ per formula for P2-NaxMnO218. The limited movable Na+ and phase transitions occurring upon its extraction and reinsertion are the major challenges. The Na+ extraction to x < 0.35 in P2-NaxMnO2 results in the irreversible P2 → OP4 phase transition, due to gliding of the MnO2-slabs, and consequently decaying capacities in cycles20. Fe-doping in the MnO2-slabs of NaxFe0.5Mn0.5O2 is beneficial for a high capacity of 200 mAh g−1 but with 35% capacity loss after 30 cycles, which is originated from the P2 → O2/OP4 phase transition21. The capacity can also be promoted to 185 mAh g−1 by co-doping Co and Ni for P2-Na0.7Mn0.6Ni0.3Co0.1O2 . Partial Mg substitution for Mn in P2-NaxMnO2 delays the formation of OP4 phase, leading to smooth electrochemical profiles and extended cycling performance with a capacity of 106 mAh g−1 . The substitution of Li or Al in NaxMnO2 can suppress the Jahn-Teller distortion and enhance its phase stability24. It is clear that finely tuning the chemical environments of the MnO6-octahedra with suitable dopants in the place of Mn allows for either more accessible Na+ sites for large capacities with limited cycles or minimized slab gliding for long cycles with moderate capacities. However, solutions to such paradox are challenging and highly desirable for the development of SIBs25.Layered transition-metal oxides with earth-abundant elements have been actively studied as host materials for Na+ of 138 pm is riveted in the edge-sharing sites of Na+ in P2-Na0.612K0.056MnO2, which leads to the small interslab distance of 0.56 nm in Na0.612K0.056MnO2 and results in high diffusion barrier of K+ and hence its immobility. The Na-O bond strength and Na+ vacancy formation energy are reduced to favor more extractable and insertional Na+. The reinforced MnO6-octahedra slabs are beneficial to the structure stability of P2-Na0.612K0.056MnO2. In situ X-ray diffraction (XRD) indicates that there is no inferior phase transition at the high charge voltages during cycles. It turns out to exhibit a high specific capacity of 240.5 mAh  g−1 with Mn3+/Mn4+ as the redox couple and the highest energy density of 654 Wh kg−1 with good cycle stability and rate capability.Here, we show that the large-sized Kx0.67−KxMnO2 (0 ≤ x ≤ 0.2) was synthesized via a solid-state reaction at 900 °C in air, however, the pure P2 phase was only obtained with x ≤ 0.1, as confirmed by the XRD patterns in Fig. x = 0.1. Its exact composition was determined to P2-Na0.612K0.056MnO2 by inductively coupled plasma-optical emission spectrometry (ICP-OES) in Supplementary Table 0.612K0.056MnO2-δ was determined by iodine stoichiometry titration. As listed in Supplementary Table 26. The XRD pattern and Rietveld refinement of Na0.612K0.056MnO2 and Na0.706MnO2 are displayed in Fig. P63/mmc. The refined structure of Na0.612K0.056MnO2 is depicted along the y and z axes in Fig. +, which are sharing faces (Naf) and edges (Nae) with the MnO6-octahedra in the two adjacent slabs, respectively, as depicted in Fig. + is located at the Nae site ) is enlarged to 3.58 Å in Na0.612K0.056MnO2 (3.42 Å in Na0.706MnO2) due to the larger ionic radii for K+ versus Na+. On the contrary, the dMnO2 is reduced to 2.02 Å in Na0.612K0.056MnO2 (2.28 Å in Na0.706MnO2), which is resulted from the enhanced MnO2-slabs. Thus, the distances of adjacent MnO2-slabs (d-spacing) are reduced in K-doped compounds even with larger-sized K+ in the Na layers.Na0.612K0.056MnO2 in Fig. 0.706MnO2 (1.897 and 2.901) in Supplementary Fig. 6-octahedra slabs. X-ray adsorption near-edge structure (XANES) spectra in Supplementary Fig. 0.612K0.056MnO2 and Na0.706MnO2. The fitted Mn K-edge Fourier-transformed extended X-ray adsorption fine structure (FT-EXAFS) spectrum of Na0.612K0.056MnO2 is shown in Fig. 0.706MnO2 in Supplementary Fig. 0.706MnO2) and Mn–Mn . This is in good agreement with the PDF results and the d-spacings of 0.56 and 0.57 nm between the (002) facets in Na0.612K0.056MnO2 and Na0.706MnO2, respectively, as observed by high-resolution transmission electron microscopy (HRTEM) in Supplementary Fig. Synchrotron X-ray pair distribution function (PDF) analysis was performed to study the local structure of Na0.612K0.056MnO2 in Fig. 0.706MnO2 has similar morphology and layered structure with Na0.612K0.056MnO2. Electron energy loss spectroscopy (EELS) in Fig. 0.612K0.056MnO2.Scanning transmission electron microscopy (STEM) equipped with the detectors of high-angle annular dark field (HAADF) and annular bright field (ABF) are employed to study the atomic structure of Na0.612K0.056MnO2 in Fig. 0.612K0.056MnO2. Similar trends for Na, Mn, and O are observed in the SIMS spectra of Na0.706MnO2 in Supplementary Fig. 39K nuclear magnetic resonance (NMR) spectrum of Na0.612K0.056MnO2 with K0.67MnO2 as a reference is shown in Fig. + located in the K-layers in the crystal structure of K0.67MnO2. The other two weak peaks at 70 and −70 ppm appear symmetrically and are assigned to the spinning sidebands due to the rotation of sample tube29. A negative chemical shift is observed for Na0.612K0.056MnO2, which is attributed to the smaller electronic shielding effect of Na and O in Na0.612K0.056MnO2 than K and O in K0.67MnO2, and a larger d-spacing of 0.64 nm in K0.67MnO2 than in P2-Na0.612K0.056MnO2 with the similar layered structure, as shown in Supplementary Fig. 9. These imply that K+ exists only in the Na layer, not in the MnO2-layer, and has a uniform distribution in the bulk particles of Na0.612K0.056MnO2.Dynamic secondary ion mass spectroscopy (SIMS) is performed to analyze the distribution of elements from the surface to the bulk of Na0.612K0.056MnO2 upon Na+ extraction and intercalation is monitored by in situ XRD in a home-made battery within a voltage window of 1.8–4.3 V. As shown in Fig. 0.612K0.056MnO2 is identified by the characteristic diffraction peaks of the (002), (004), (100), (102), and (104) planes in Fig. 0.612K0.056MnO2 is sustained in the first charge process without phase transition, and 0.498 Na+ per formula is extracted to give the fully charged state (I) of Na0.114K0.056MnO2, which agrees well with the ICP-OES in Supplementary Table + is inserted for Na0.65K0.056MnO2 (II). Due to the slight distortion of the MnO6-octahedra slabs, the P’2 phase starts to appear with the co-existence of P2 phase, and becomes dominant in the fully discharged state (III) of Na1.015K0.056MnO2. In the reverse charge process, the P’2 phase is gradually transformed to the P2 phase (IV), and then it is sustained at high charge voltages till 4.3 V (V). It involves the reversible two-phase transition of P2 ↔ P′2 with total insertion or extraction of 0.901 Na+ in the cycles of discharge and charge, as schematically illustrated in Fig. 0.706MnO2 involves large gliding of the MnO6-octahedra slabs for the phase transition of P2 → OP4 upon extraction of 0.455 Na+ at high charge voltages; it then converts from the OP4 phase into P2 and P′2 during discharge with total insertion or extraction of 0.671 Na+ in cycles, as shown in the schematic phase transitions of P2-Na0.706MnO2 in Fig. 0.612K0.056MnO2 and P2-Na0.706MnO2  can be extracted at the high charge voltages and re-inserted from/into P2-Na0.612K0.056MnO2 than P2-Na0.706MnO2, due to the reduced Na-O bond strength and reinforced MnO6-octahedra slabs for suppression of the inferior OP4 phase.The structure evolution of Na, (004), 0, (102),0.612K0.056MnO2 is further unveiled by X-ray photoelectron spectroscopy (XPS) and EELS. The XPS Mn 2p spectrum of Na0.612K0.056MnO2 in Fig. p3/2 and Mn 2p1/2 at 640.8/642.2 and 652.4/654.9 eV, respectively, suggesting the co-existence of Mn3+ and Mn4+ with a ratio of 2:1. Upon charging, the intensity of Mn4+ peaks increases with decrease of the intensity of Mn3+ peaks, indicating oxidation of Mn3+ to Mn4+ in Fig. 3+ is 0.08 per formula at the charge finish of 4.3 V, which is consistent with the phase transition in Fig. 4+ to Mn3+ in the following discharge process is evidenced by the gradual intensity increase of Mn3+ peaks till the disappearance of Mn4+ peaks, as shown in Fig. 30. In addition, the EELS spectra of the charged, discharged, and pristine states of Na0.612K0.056MnO2 in Fig. 3+/Mn4+ with the reversible shifts of the two peaks of Mn-L2 and Mn-L3. It is notable that there is no peak shift or broadening in both of the O pre-peak and K-edge, which suggests no valence change of O2-. Therefore, the redox of Mn3+/Mn4+ is responsible for the charge compensation in the extraction and reinsertion of Na+ from/into Na0.612K0.056MnO2 in the discharge and charge processes. Also, no oxygen evolution was detected by in situ differential electrochemical mass spectrometry (DEMS) in the charge process  from 0 to 5 ps reflect the diffusion capability of each element. It shows that Na+ is more facile to migrate in Na0.500K0.055MnO2, and K+ is immobile as riveted between layers. Its migration is energetically unfavorable as displayed in Supplementary Fig. + in Na0.500K0.055MnO2 (refs. 37). Its diffusion coefficient is accordingly five orders of magnitude lower than that of Na+ based on Arrhenius equation. This is consistent with ICP-OES and elemental mappings in Supplementary Table 0.500K0.055MnO2 in Supplementary Fig. + located in the Na layers. It can be sustained in the optimized structure of P2-Na0.220K0.055MnO2 at the fully charged state with similar structure parameters in Supplementary Fig. 0.612K0.056MnO2 for extraction and insertion of Na+. It is robust to allow for the extraction of 0.498 Na+ in the first charge without phase transition, as depicted in Fig. Density functional theory (DFT) calculations using Vienna Ab-initio Simulation Package (VASP) are performed to optimize the crystal structures of the pristine and charged cathode materials0.500K0.055MnO2 and Na0.555MnO2 have been investigated by DFT calculations. As shown in the density of states (DOS) of Na0.220K0.055MnO2 in Fig. 0.500K0.055MnO2 in Fig. 0.555MnO2. Schematic pDOS of Na0.500K0.055MnO2 and Na0.555MnO2 are shown in Fig. +, in good agreement with the analysis of FT-EXAFS in Fig. 0.500K0.055MnO2 is found to be reduced by COHP analysis in Supplementary Table xy-plane by the substitution of K. A slight fluctuation of Mn–O bond strength is found for different valence state of Mn. These lead to more thermodynamically favorable Na+ vacancies  is higher than in the Naf site (−3.500 eV) of Na0.500K0.055MnO2, which indicates the higher thermodynamic stability of the Nae site for K+. Furthermore, the reinforced Mn–O bonds play an important role to suppress the phase transition during Na+ extraction and insertion and favor the cycle stability.The electronic structures of Na0.612K0.056MnO2 is evaluated against Na metal anodes with the electrolyte of 1.0 M NaPF6 in propylene carbonate (PC) and 5 wt.% of fluoroethylene carbonate (FEC) in coin cells. Figure 0.612K0.056MnO2 and Na0.706MnO2 at 20 mA g−1. Na0.612K0.056MnO2 delivers a high initial charge capacity of 125.1 mAh g−1, corresponding to 0.498 Na+ per formula. In the following cycles, the reversible capacity of Na0.612K0.056MnO2 reaches up to 240.5 mAh g−1, which is much higher than that of Na0.706MnO2 (167.2 mAh g−1). This means that 0.901 Na+ per formula is inserted and extracted into/from Na0.612K0.056MnO2. It agrees well with the restricted phase transition in Fig. 0.612K0.056MnO2 presents higher discharge voltages than Na0.706MnO2 does, which is consistent with the cyclic voltammetry (CV) curves in Supplementary Fig. 0.612K0.056MnO2 is as high as 654 Wh kg−1, in comparison with the reported cathode materials in Fig. 0.72Li0.24Mn0.76O2, which works with both the O redox and Mn3+/Mn4+ to deliver the highest specific capacity in refs. 45.The electrochemical performance of Na0.612K0.056MnO2 is tested at 50 mA g−1, and the selected charge/discharge profiles of the 1st, 5th, 10th, and 100th cycles are presented in Supplementary Fig. −1, corresponding to capacity retention of 98.2% with only 0.018% loss per cycle in Fig. 0.706MnO2. This implies that the structure of Na0.612K0.056MnO2 is robust upon extraction and insertion of Na+ in cycles, as confirmed by in situ XRD in Fig. The cycle stability of Na+ diffusion coefficients are estimated by Galvanostatic intermittent titration technique (GITT) in Supplementary Figs. DNa+) for the single-phase regions of both NaxK0.056MnO2 and NaxMnO2 are in the range of 10−12–10−10 cm2s−1, and they exhibit decreasing trends as the sodiation proceeds near the two-phase regions in Fig. DNa+ in Na0.612K0.056MnO2 and Na0.706MnO2, as presented in the gray region in Fig. 0.612K0.056MnO2 exhibits faster Na+ transport with the content change of Na during charge and discharge, though there is K+ riveted between the layers. Enhanced rate capability is presented in Na0.612K0.056MnO2 in Supplementary Fig. −1 to 211.6 and 110.2 mAh g−1 at the currents from 50 mA g−1 to 100 and 2000 mA g−1, respectively. The demonstrated good cycle performance and rate capability of Na0.612K0.056MnO2 are attributed to the reinforced structure stability and facile Na+ transport.The Na0.612K0.056MnO2 cathode and presodiated hard carbon (HC) anode was demonstrated. The performance of HC anode was firstly investigated in half battery in Supplementary Fig. 0.612K0.056MnO2, implying reversible electrochemical reactions. The full battery delivers a specific capacity of 230.6 mAh g−1 based on the mass of active cathode material at 50 mA g−1 in Fig. −1 based on the total mass of active anode and cathode materials. It also exhibits good cycling stability and rate performance in Supplementary Fig. 0.612K0.056MnO2 will be of great potential for the high-performance SIBs.A full battery with P2-Na0.612K0.056MnO2 has been successfully obtained with the K+ riveted in the prismatic edge-sharing sites of Nae between the MnO6-octahedra slabs. The reinforced MnO6-octahedra slabs and reduced Na+ vacancy formation energy enable more active Na+ sites and favor its facile diffusion. This leads to the highest reversible capacity of 240.5 mAh g−1, corresponding to 0.901 Na+ per formula, and energy density of 654 Wh kg−1 based on the redox of Mn3+/Mn4+ among the reported transition-metal oxide cathodes for SIBs, and good capacity retention of 98.2% after 100 cycles. The reinforced interaction in the MnO6-octahedra slabs restricts the layer gliding and ensures the robust crystal structure of P2-Na0.612K0.056MnO2, in which two-phase transition of P2 ↔ P’2 occurs only at low charge and discharge voltages, favoring the structure stability and cycle performance. This strategy to finely tune the local chemistry for high energy density and long cycle life will shed lights into the design of electrode materials for SIBs and beyond.In summary, a cathode material of P2-Nax0.67-KxMnO2 was synthesized via a solid-state reaction. For the synthesis of Na0.612K0.056MnO2, NaCH3COOH , K2CO3 , and MnO2  were mixed in a planetary ball mill (QM-3SP2) at 400 rpm for 12 h in the presence of a small amount of acetone. The as-synthesized precursors were dried at 80 °C for 12 h, and then pressed into a pallet under 20 MPa. It was heated at 500 °C for 2 h and 900 °C for 10 h in air at 5 °C/min. Then, it was cooled down to 200 °C at 5 °C/min and naturally to room temperature, and stored in an argon-filled glovebox. Similar processes were applicable to the synthesis of Na0.706MnO2 except the addition of K2CO3.P2-NaN-methyl-2-pyrrolidone onto an Al foil. The loading mass of active material in each electrode pellet of 12.0 mm in diameter was ~3.0 mg cm−2. For half battery, sodium foil was used as the anode, and 1.0 M NaPF6 in PC containing 5 wt.% of FEC was employed as the electrolyte with glass fiber filter papers as the separator. The charge/discharge tests and Galvanostatic intermittent titration technique (GITT) were carried out on a Land CT2001A battery test system  in a voltage range of 1.8–4.3 V at 25 °C. GITT was measured by applying the repeated current pulses for 30 min at a current density of 10 mA g−1, followed by relaxation for 1.5 h. CV experiments were performed within the same potential range at 0.1 mV s−1. CV was carried out on a Solartron 1470E electrochemical workstation in a voltage range of 1.8–4.3 V. The full battery was assembled with P2-Na0.612K0.056MnO2 as the cathode, HC as the anode, and 1.0 M NaPF6 in PC containing 5 wt.% of FEC as the electrolyte, respectively. In order to supplement sufficient Na+ in HC before assembling full battery, the HC electrode was obtained by disassembling the half battery in the glove box after half-discharging at 50 mA g−1. The mass ratio of Na0.612K0.056MnO2/HC is around 1.2, corresponding to 4% excess of anode capacity. The charge/discharge and CV tests were carried out at 25 °C in a voltage range of 1.0–4.3 V.Coin cells (CR2032) were assembled in an argon-filled glovebox for all the electrochemical tests. The working electrode was prepared by coating the mixture of active material (80 wt.%), Super P (10 wt.%), and polytetrafluoroethylene (10 wt.%) in 46. The PDF analysis was carried out based on X-ray scattering data measured at the 11-ID-B beamline of the Advanced Photon Source at Argonne National Laboratory. High energy X-rays (λ = 0.1173 Å) were used to collect data to high values of momentum transfer (Q ~ 30 Å−1). Diffraction images were integrated within fit2D to obtain one-dimensional diffraction data47. Pair distribution functions, G(r), were extracted from the data within PDFgetX348, after correcting for background and Compton scattering. Mn K-edge X-ray absorption fine structure (XAFS) results were collected at the beamline 8C in Pohang Accelerator Laboratory (PAL). The XAFS data were collected by fluorescence mode and processed through Demeter software package49. The extended x-ray absorption fine structure (EXAFS) region was k3 weighted and Fourier-transformed in k-ranges of 3.0–10.0 Å−1. Schematic representations of the synthesized samples were obtained by using the VESTA software50. The morphology was examined by SEM . The HRTEM images, SAED patterns and EDS mappings were recorded by TEM . HAADF and ABF images were obtained on a JEM-AR200F STEM with an accelerating voltage of 200 kV with cold filed-emission gun and double hexapole Cs correctors . Dynamic secondary ion mass spectroscopy (SIMS) was carried out on a Cameca IMS-5FE7 system using Cs+ primary ions with a beam current of 3 nA. Solid-state 39K NMR spectra were obtained on a JNM-ECZ600R widebore spectrometer.The chemical composition of the electrodes was determined by an inductively coupled plasma-optical emission spectrometer . XRD was collected on Rigaku SmartLab equipped with a Cu Kα radiation source. Rietveld refinement was conducted with the Fullprof software0.500K0.055MnO2, Na0.220K0.055MnO2, and Na0.275MnO2 were simulated using Vienna ab-initio simulation package (VASP). The generalized gradient approximation (GGA) of exchange-correlation energy, Perdew-Burke-Eznerhof (PBE) functional, was used with its corresponding projector augmented-wave pseudopotentials34. A Hubbard U term of 3.9 is adopted to better describe the strong correlation of Mn 3d electrons. The slight difference between the simulated formulas and the obtained composition is due to that the number of K and Na atoms can only be integral in a unit cell, and such difference can be reasonably ignored. All optimization calculations adopted a 4 × 4 × 4 Monkhorst-Pack51 k-point grid and a cutoff energy of 400 eV, and a higher cutoff of 450 eV and a 6 × 6 × 4 k-point grid were applied for DOS and charge calculations for enhanced accuracy. First-principle molecular dynamics calculations, which adopt an NVT ensemble at 1000 K, were carried out using a single Γ point to balance the accuracy and cost. Atomic positions and cell vectors were fully optimized until all force components were less than 0.02 eV Å−1. The trajectory of MD simulation was further analyzed using Rigorous Investigation of Networks Generated using Simulations (R.I.N.G.S.) code to obtain MSD statistics. The COHP calculations were carried out using Local-Orbital Basis Suite towards Electronic-Structure Reconstruction (LOBSTER) code. The diffusion energy barriers are calculated using climbing-image nudged elastic band (CI-NEB) method with a single Γ point.The crystal structures of NaSupplementary Information


These novel histologic findings are based on the pathology of cardiac sarcoidosis. We also discussed the principal histologic differential diagnoses of cardiac sarcoidosis, such as tuberculosis myocarditis, fungal myocarditis, giant cell myocarditis, and dilated cardiomyopathy.Sarcoidosis is a rare disease of isolated or diffuse granulomatous inflammation. Although any organs can be affected by sarcoidosis, cardiac sarcoidosis is a fatal disorder, and it is crucial to accurately diagnose it to prevent sudden death due to dysrhythmia. Although endomyocardial biopsy is invasive and has limited sensitivity for identifying granulomas, it is the only modality that yields a definitive diagnosis of cardiac sarcoidosis. It is imperative to develop novel pathological approaches for the precise diagnosis of cardiac sarcoidosis. Here, we aimed to discuss commonly used diagnostic criteria for cardiac sarcoidosis and to summarize useful and novel histopathologic criteria of cardiac sarcoidosis. While classical histologic observations including noncaseating granulomas and multinucleated giant cells  are the most important findings, others such as microgranulomas, CD68 Sarcoidosis is a systemic inflammatory disease of unknown etiology and is characterized by noncaseating granulomas that may involve any organs of the body, particularly the lungs, lymph nodes, skin, liver, spleen, and eyes . CardiacCardiac sarcoidosis typically occurs in the 20s to 60s age group and is more common in women than in men ,6. The iCardiac sarcoidosis is discovered at autopsy in 15% to 90% of cases of systemic sarcoidosis, and it can affect anywhere in the heart ,7,8,9,10It is often difficult to recognize cardiac sarcoidosis via antemortem endomyocardial biopsies. It is typically diagnosed when endomyocardial biopsy demonstrates noncaseating granulomas; however, endomyocardial biopsy has low sensitivity (approximately 20% to 30%) for the diagnosis owing to the small size and high sampling error rate associated with biopsy specimens ,13. In aHere, we introduce commonly used diagnostic criteria for cardiac sarcoidosis and discuss the significance of endomyocardial biopsy. Thereafter, we summarize useful and novel histopathologic criteria of cardiac sarcoidosis. Finally, we discuss the principal differential diagnoses of cardiac sarcoidosis.The most commonly used diagnostic criteria for cardiac sarcoidosis are the 2014 Heart Rhythm Society Expert Consensus Statement on the Diagnosis and Management of Arrhythmias Associated with Cardiac Sarcoidosis see  16] and and16] aMost pathologists diagnose cardiac sarcoidosis depending on the presence of sarcoidal granulomas. However, the diagnostic sensitivity of endomyocardial biopsy is low because sarcoidal granulomas are rarely detected in many cases. Here, we emphasize the histologic findings of cardiac sarcoidosis based on the temporal evolution and significance of the morphologic alterations.+ helper T cells are predominantly recruited, expanded oligoclonally, activated, and polarized into the T helper 1 (Th1) phenotype, which results in a CD4/CD8 T-cell ratio of approximately 10:1 ..59].Giant cell myocarditis and cardiac sarcoidosis are sometimes notably difficult to distinguish because they have many similar clinical manifestations, histopathological findings, and pathoetiology. Although there is some controversy regarding whether giant cell myocarditis and cardiac sarcoidosis are different phenotypes of a single disease or two distinct entities, some histological differences have been identified . Granulo+ helper T cells are dominantly recruited in cardiac sarcoidosis as described above, whereas CD8+ cytotoxic T cells more often occur in giant cell myocarditis [Moreover, some reports indicate that immunohistochemistry may be helpful for diagnosis; although both are caused by T cell-mediated autoimmune inflammation, CD4carditis .The histologic alterations in dilated cardiomyopathy are nonspecific because they do not point to a specific etiology. Dilated cardiomyopathy occasionally exhibits similar clinical symptoms and histopathology to those of cardiac sarcoidosis. In particular, when typical granulomas are not formed in endomyocardial biopsy, it is occasionally difficult to distinguish between the interstitial phase of cardiac sarcoidosis and dilated cardiomyopathy. A fibrotic pattern is one of the histologic findings that can distinguish between both conditions. Fibrosis is pronounced in the myocardium with a mosaic pattern in cardiac sarcoidosis, whereas diffuse fibrosis in extensively injured cardiomyocytes is often seen in dilated cardiomyopathy ,36.Additionally, fatty infiltration can occur within loose connective tissue around granulomas in cardiac sarcoidosis, whereas there is no fatty infiltration within fibrous granulation tissues in dilated cardiomyopathy . It is uC. acnes components (detected by PAB antibody), confluent fibrosis, and fatty infiltration may be helpful to improve the sensitivity of endomyocardial biopsy. These novel histologic findings are based on the pathology of cardiac sarcoidosis.Sarcoidosis is a rare disease of isolated or diffuse granulomatous inflammation. Cardiac sarcoidosis is sometimes a fatal disorder, and it is important to accurately diagnose it to prevent sudden death due to fatal dysrhythmia. Although endomyocardial biopsy is invasive and has limited sensitivity for detecting granulomas, it is the only procedure that provides a definitive diagnosis of cardiac sarcoidosis. While classical histologic observations including noncaseating granulomas and multinuclear giant cells  are the most crucial, other findings such as microgranulomas, M1 macrophage accumulation, lymphangiogenesis, Cardiac sarcoidosis is an involvement of a major organ and should be thoroughly studied and treated to avoid the risk of arrhythmia and of sudden death. Therefore, pathologists should carefully observe for granulomas and other histologic findings that indicate cardiac sarcoidosis, in order to achieve a timely and accurate diagnosis of the condition.


The G peak thermal lineshift rates dω/dT showed that the H2 intercalation significantly weakened the pinning effect in epitaxial graphene. Furthermore, the G peak dω/dT value showed a perspicuous pinning effect disparity of QFSBEG samples. Additionally, the anharmonic phonon effect was investigated from the Raman lineshift of peaks. The physical mechanism responsible for dominating the G-mode temperature-dependent behavior among samples with different substrate coupling effects was elucidated. The phonon decay process of different samples was compared as the temperature increased. The evolution from in situ grown graphene to QFSBEG was determined. This study will expand the understanding of QFSBEG and pave a new way for its fabrication.Epitaxial graphene on SiC without substrate interaction is viewed as one of the most promising two-dimensional (2D) materials in the microelectronics field. In this study, quasi-free-standing bilayer epitaxial graphene (QFSBEG) on SiC was fabricated by H ComPb, etc. ,7,8,9, Hrication ,11,12.2 intercalation condition, and how the evolution proceeds from the in situ graphene and QFSBEG have not been determined. The extent of H intercalation at the interface of the SiC substrate and the graphene has a direct influence on transport properties. Therefore, it is of vital importance to conduct a study on the progress from in situ graphene to QFSBEG.The difference between the structure and properties of in situ grown monolayer graphene and QFSBEG has been researched elsewhere ,12,13,14 407 GHz . However2 intercalation condition using a Hall instrument. However, the sample can be damaged by the probe contact even under the simple four-probe Van der Pauw configuration. It is well known that Raman spectroscopy is one of the most effective and nondestructive methods for graphene characterization. In addition to the layer number and structure characteristics of graphene, the Raman spectra can also provide the anharmonic phonon effect by analyzing the temperature-dependent spectra. The anharmonic phonon effect is the collective effect of thermal expansion and electron-phonon and phonon-phonon interactions, which play an important role in electronic transport properties [For the characterization of QFSBEG, a mobility test is the most intuitive way to reveal the Hoperties ,18,19,202 intercalation conditions in the temperature range from 213 K to 663 K. The substrate effect was compared through the line shift rates of G mode. Furthermore, the anharmonic phonon process was studied by conducting theoretical model calculations.In this study, the Raman spectra were measured as a function of temperature on in situ graphene and QFSBEG obtained under different H2 etching procedure to obtain a regular and micron-scale step structure in a 5 sccm flow of H2 at a temperature of 1500–1600 °C and pressure of 800–900 mbar for 20 min. Graphene growth was conducted at 1600–1700 °C, 800–900 mbar in 20 sccm argon atmosphere for 60 min. The sample was then cooled to room temperature, corresponding to in situ graphene. For H2 intercalation, two samples were selected for H2 intercalation processing in a hydrogen flux of 50 sccm at a temperature of 1600 °C and pressure of 800–900 mbar. The operation time was selected as 30 min and 60 min to study the evolution process from in situ grown graphene to QFSBEG, corresponding to samples QFSBEG-1 and QFSBEG-2, respectively.Epitaxial graphene was grown on home-made 4 inch SiC wafers in a self-developed furnace. The (0001) face was processed by chemical mechanical polishing (CMP) and followed by cleaning with acetone , ethanol , and deionized water cleaning to remove surface contamination. The SiC substrate was subjected to a HFor the Raman measurement, HR 800 by Horiba Jobin Yvon  was used in the backscattering geometry with a 532 nm laser for excitation. For surface enhanced Raman spectra testing, the 100× objective was used with a 600 diffraction grating with the assistance of Ag particles at room temperature. For the temperature variation measurement, a Linkam  hot/cold stage was used and cooled with liquid nitrogen. The 50× long-focus objective was used with a 600 diffraction grating restricted by the distance to the sample stage. The samples were measured from 213 K to 663 K in steps of 15 K. In the case of temperature interference on SiC substrate for substrate subtraction, a SiC substrate was placed in the hot/cold stage and its spectra were synchronously collected under different temperatures. Before each measurement, the samples were stabilized for about 5 min to ensure that the temperature had reached the set value and the optical system was refocused to avoid the thermal effect. For the accuracy of data, the positions of four samples  were recorded, and after each measurement, the stage was set at the same coordinate to ensure that the spectrum was taken at the same spot.2 at different time periods were studied systematically. First, in order to get intuitive understanding of graphene after H2 intercalation, the surface enhanced Raman spectra of these three samples were measured with the assistance of Ag particles. as shown in −1 appeared in the Raman spectra of samples after H2 intercalation. It was considered that the perpendicular stretch mode of the Si-H bond was excited by the near-field plasmonic effect, thus, the H2 intercalation could be proved straightforward in comparison with the conventional method [2 was intercalated successfully or failed, but they did not provide information about the inner condition of QFSBEG.The in situ grown graphene and the samples intercalated by Hl method ,21. HoweRaman spectra of three samples were measured in the temperature range of 213K–663K, and the spectra obtained in three typical temperatures  are shown in dω/dT are summarized in dω/dT rate of −0.048 cm−1/K, which was consistent with the reported in situ grown epitaxial graphene on SiC rate of −0.043 cm−1/K [2 intercalation, the value declined dramatically and varied by the time duration of H2 intercalation. The dω/dT rates of samples QFSBEG-1 and QFSBEG-2 were −0.035 cm−1/K and −0.022 cm−1/K, respectively. It is clear that the dω/dT rate is an effective parameter for characterizing the intrinsic state of the substrate pinning effect under different H2 conditions. As the time extended, the substrate effect would be further released but beyond the reach of free-standing graphene at the rate of ~−0.010 cm−1/K [For the analysis of the G peak, the Raman spectra of graphene were subtracted by the SiC Raman spectra simultaneously taken at a specific temperature. The subtraction process was shown in 3 cm−1/K . After H0 cm−1/K .2 intercalation deepening from It is worth noting that, with the extent of Hsp2 carbon hybridization and resulted from the degenerate phonon mode 2gE at the center of Brillouin zone. The D peak is contributed by a one-phonon process assisted by the defect near the K point of the Brillouin zone, thus it is sensitive to the crystalline defects. The 2D peak is an overtone of the D peak, which originated from a second-order phonon process and is sensitive to the number of graphene layers. For a Raman spectrum taken at room temperature, the domain size or the layer of graphene could be extracted through these three peaks using an empirical equation. Temperature-dependent Raman spectra are an excellent tool to investigate anharmonic effects. According to the ab initio calculation, the Raman peak position is the real part of the phonon self-energy, and the Raman linewidth is the imaginary part of the self-energy. The Raman shift of G peak at a certain temperature can be expressed using the following formula [γ is the Gruneisen parameter of graphene for the Raman G mode; here, we take Gγ = 1.99 [For a typical Raman spectrum of graphene, it is of vital importance to study the G, D, and 2D peaks. The G peak is representative of  formula ,29.(1)ΩG = 1.99 ; βT is twas used .(3)βT=αM1 and M2 are constants representing the three-phonon and four-phonon processes to the Raman shift, respectively. The term wHere, the simple Klemens model was used to clarify the Raman shift variation caused by the pure temperature effect. In this model, the effect can be considered as three-phonon and four-phonon interactions, and can be expressed as follows:Ω0, M1, and M2 are shown in Ω0 of the QFSBEG samples exhibited a redshift, in contrast with the in situ graphene. In consideration of the phonon process, the reported studies showed that the four-phonon effect was the dominating factor over the thermal expansion and three-phonon process G mode with the temperature [1/M2 value of QFSBEG-2 decreased in contrast with that of QFSBEG-1, which indicated that the three-phonon effect decreased with the extent of H intercalation deepening.perature ,32. In cperature . Indeed,tructure . It can Usually, the phonon damping rate is proportional to the linewidth Γ(T) of the Raman peak. Hence, the phonon lifetime τ can be calculated from the peak linewidth Γ through the energy-time uncertainty relation, and the equation can be expressed as follows :(6)Γћ=1τ−1∙ps. Based on the G mode, the peak width was extracted by Lorentz fitting and the phonon lifetime was calculated by the above equation. The linewidth of the Raman G peak and phonon lifetime of these three samples are shown in the Here, τ, it was affected by two main factors, and the equation can be expressed as [For the measured phonon lifetime essed as ,36:(7)12 processing, the Si-C bonds were broken and the Si dangling bonds were partially saturated by H atoms, as shown in Based on the experimental data, the transformation evolution process from in situ graphene to QFSBEG is summarized in 2 intercalation time periods were fabricated. Temperature-dependent Raman spectra were collected in the temperature range of 213–663 K. In contrast to the in situ grown graphene that the G peak dω/dT value at changed at the rate −0.048 cm−1/K, the G peak thermal lineshift rates dω/dT declined to −0.035 cm−1/K after 30 min H2 management and −0.022cm−1/K after 60 min H2 management, respectively. The anharmonic effect analysis showed that the three-phonon process was the dominant decay pathway, and the ratio of the three-phonon process to the four-phonon process dropped from 4.74 to 3.69 as the H2 treatment time increased. At the conclusion, the evolution process of transformation was demonstrated.In conclusion, in situ grown graphene on SiC and QFSBEG under different H


Accurate and imperceptible monitoring of electrophysiological signals is of primary importance for wearable healthcare. Stiff and bulky pregelled electrodes are now commonly used in clinical diagnosis, causing severe discomfort to users for long-time using as well as artifact signals in motion. Here, we report a ~100 nm ultra-thin dry epidermal electrode that is able to conformably adhere to skin and accurately measure electrophysiological signals. It showed low sheet resistance , high transparency, and mechano-electrical stability. The enhanced optoelectronic performance was due to the synergistic effect between graphene and poly  polystyrene sulfonate (PEDOT:PSS), which induced a high degree of molecular ordering on PEDOT and charge transfer on graphene by strong π-π interaction. Together with ultra-thin nature, this dry epidermal electrode is able to accurately monitor electrophysiological signals such as facial skin and brain activity with low-motion artifact, enabling human-machine interfacing and long-time mental/physical health monitoring. Novel dry electrodes with enhanced mechano-electrical stability and conformability are attractive for long-time electrophysiological signal monitoring. Here, the authors report polymer-covered CVD-grown graphene with enhanced optoelectronic performance for biopotential monitoring. Skin-contact electrophysiology, such as electromyography (EMG), electrocardiography (ECG), electrooculogram (EOG), and electroencephalogram (EEG), relying on the recording of biopotential, are important clinical measurements to assess health status. The ability to accurately and imperceptibly monitor electrophysiological signals is the most basic requirement for wearable electronics8. Skin-contact electrodes in clinical use are mainly pregelled Ag/AgCl11. Although they provide good signal quality, the electrolytic gel may irritate the skin and will dry in a certain period, making it impossible for long-term monitoring and diagnosis. Moreover, with the gel interfacing on skin, electrode displacement with respect to skin will cause severe motion artifact13. Specifically, when detecting facial electrophysiological signals, gel electrodes are easily detached from skin, resulting in poor electrophysiological signal resolution. This is because superficial portions of the face muscles are uneven, full of wrinkles, and inevitably large-shape changed. To achieve high signal quality with low-motion artifact for long-time monitoring, dry electrodes being ultra-conformably adhesive to skin and sufficiently mechano-electrical stable are highly desirable for biopotential monitoring.Epidermal electronics refer to electronic devices that are mounted on the epidermis and able to acquire information about the body precisely or to provide therapy to the body. It can open new pathways for wearable healthcare18 and gold film17 are two frequently studied dry electrodes for electrophysiological recordings. As a representative conducting polymer, PEDOT:PSS blended with waterborne polyurethane and D-sorbitol (~20 µm) was reported to be adhesive on skin and electrically stable upon strain, being able to measure electrophysiological signals19. But the detected signal quality  is not good, because thick material is difficult to be ultra-conformal on skin and blending with polymer results in poor conductivity (545 S/cm)  was proved to be self-adhesive to the arbitrary 3D biological surface and able to monitor surface biopotentials with less motion artifact2. However, Au-film dry electrodes are expensive and visible on the skin, which may bring light-induced photoelectric artifact when applied in certain scenarios. To make dry electrodes more suitable for epidermal electronics, low-cost, optical transparency, ultra-conformability to skin and mechano-electrical stability are essential characteristics to be introduced into dry electrodes for biopotential monitoring.Conducting polymer32 (Supplementary Table 8. Poly(methyl methacrylate) (PMMA)—assisted monolayer CVD graphene electronic tattoo, with a total thickness of 463 ± 30 nm, the optical transparency of ~85%, the sheet resistance of 1994.33 ± 264 Ω/sq and stretchability of ~40%, was applied on skin to detect electrophysiological signals. To fabricate dry electrodes that are easy to handle, polymer thin film (herein PMMA) was utilized. However, the PMMA polymer carrier layer would introduce contamination, cracks, or defects into the CVD-grown graphene during transfer, leading to a significant decrease in the conductivity of final electrodes34. Therefore, to improve the quality of electrophysiological recordings for graphene-based ultra-conformal dry electrodes, a modified PEDOT:PSS layer, not only capable of handling graphene as well as enhancing its superior electronic performances, was proposed in this work.As the thinnest electrode material, graphene demonstrates superior optical transparency, high electrical conductivity, excellent mechanical property, and low electrochemical reactivityHere, we elaborated a strategy for the fabrication of transparent, highly conductive, and ultra-thin dry electrodes by using PEDOT:PSS thin film as a polymer carrier layer to handle and enhance the electrical performance of CVD-grown graphene for biopotential monitoring. To enable a thin, uniform, continuous and highly conductive PEDOT:PSS layer covered on top of graphene, surfactants and ionic compounds were added since pristine PEDOT:PSS solution did not wet graphene and exhibited limited conductivity. Optimized dry electrodes of PEDOT:PSS transferred CVD graphene film (PTG) demonstrated a total thickness of ~100 nm, a sheet resistance of ~24 Ω/sq (4142 S/cm), high optical transparency, and a sufficient mechano-electrical stability for skin electronics. The dramatic increase in electrical conductivity is not merely a consequence of a parallel connection of two conductive units, but a synergistic effect of improved molecular packing of PEDOT and charge transfer on underlying graphene. This observation has been recognized by Grazing-incidence wide-angle X-ray scattering (GIWAXS), Raman, electron spin resonance (ESR) spectroscopy, and so on. PTG exhibited outstanding capability for electrophysiological signal recording such as facial skin and brain activity in motion, enabling human-machine interfacing and long-time mental/physical health monitoring. Our synergistic strategy towards ultra-conformal, transparent, and dry conductive electrodes will pave a new way for epidermal electronics.35 method on copper foil and used conducting polymer PEDOT:PSS 38, as it is a safe additive that has been widely used in detergents for laundry  sulfonimide lithium salt as an example, BSL in short), was also introduced into PEDOT:PSS solution to further improve the electrical conductivity and mechanical stretchability of PTG as skin electrodes. As-obtained PEDOT:PSS (with SDS and BSL)/graphene/Cu foil was annealed at 120 °C for several minutes to enhance the interaction between PEDOT:PSS and graphene, and subsequently immersed in (NH4)2S2O8 solution to etch the underlying Cu foil, followed by rinsing in deionized water to remove possible residues . The resultant PTG thin film was then floated on water and scooped onto arbitrary substrates, such as a thin elastomer, glass, decal transfer tattoo paper and even skin, working as a dry electrode immediately.To achieve dry electrodes with superior electronic performance for biopotential monitoring, we synthesized graphene via CVDPSS Fig. a, b as tPSS Fig. . Adding dry Fig. . In addiDue to the amphiphilic nature of these additives, the addition of SDS and BSL to PEDOT:PSS solution improves its wettability on the hydrophobic surface and helps to form a continuous and uniform PEDOT:PSS conductive layer on graphene. To verify the improvement in electrical conductivity of PTG with the addition of SDS and BSL, we measured 4-probe sheet resistance and transmittance on a variety of PTGs as a function of SDS and BSL. The sheet resistance of PTG thin film was clearly reduced with increased SDS concentrations in the PEDOT:PSS solution Fig. . This is41. Thus, to compromise the transparency, conductivity and mechano-electrical stability of PTG, our optimal condition for following epidermal electrophysiology is bi-layer graphene transferred by PEDOT:PSS at both SDS and BSL concentrations of 1.0 wt%.Next, the mechano-electrical stability of PTG was evaluated (see “Methods”). We placed PTG onto poly (PET) and bent it at various angles. Almost no resistance change was observed when PTG was at a bending radius from 12 to 2 mm Fig. . To eval2/Si substrates are similar  in PEDOT:PSS solution, PEDOT:PSS only (~185 nm) film is ~2.3 times thickness of PTG  film  was shifted to 1599.5 cm−1 in PTG. Such blue-shifts indicated that there is strong π–π interaction between PEDOT:PSS layer and graphene, which will engender the delocalization of π electrons, thus increasing the charge carrier mobility of PTG thin film.To Figure out the synergistic effect between PEDOT:PSS and graphene, we performed a comprehensive study on the structure of PTG. Raman spectroscopy is a powerful tool to characterize carbon-based materials. For the PEDOT:PSS film only, there was a vibration peak at 1434.5 cmThe charge delocalization of π electrons was also confirmed by UV–vis-NIR spectroscopy. Compared with pristine PEDOT:PSS film, a strong absorbance occurred in PTG over a broad range, both at the ~800-nm peak and the free-carrier tail of the near-infrared region. This observation proved the charge delocalization between graphene and PEDOT:PSS, leading to the formation of polaron pairs, that is, bipolarons. The π-π interaction effect was then studied by electron spin resonance (ESR), which probed the presence of unpaired electrons Fig. . The strz and qxy directions were obtained from the 2D GIWAXS patterns , 3.46 Å (PEDOT:PSS with SDS), 3.47 Å (PEDOT:PSS with SDS and BSL) 3.40 Å (PTG with SDS) and 3.37 Å (PTG with SDS and BSL). With the addition of SDS, BSL, and mainly the underlying graphene, the π-π stacking distance of PEDOT decreases from 3.54 to 3.37 Å  and graphene on the molecular packing of PEDOT was further investigated by Grazing-incidence wide-angle X-ray scattering (GIWAXS), an advanced technique to analyze thin-film morphology. Fig. 7 Å Fig. , which r44. The more ordered PEDOT packing in PTGs can be explained by the strong π–π interaction between graphene and PEDOT:PSS. Overall, the synergistic enhancement between PEDOT:PSS and graphene will induce a higher crystallinity in PEDOT, so as to achieve a high conductivity for PTG to accurately detect electrophysiological signals.The highly ordered molecular packing in PTGs can also be reflected in the surface morphology. We used atomic force microscopy (AFM) to observe the morphological evolution of pure PEDOT:PSS and PTGs Fig. . With thh represents thickness, 45. Transferred onto a human ring finger full of wrinkles, PTG indeed exhibited intimate contact on it  pulses to drive the servo motor  recording is to measure the electrical activity of the brain, providing helpful information for the diagnosis or treatment of the brain disorders. Due to the very weak signal strength , long-time monitoring of EEG is extremely challenging, in particular for the status of exercising. Normally, for EEG monitoring, the volunteer sits in a comfortable position peacefully and experiences brain stimulus, such as doing math quiz, watching videos or listening to music. As PTG is ultra-thin and the attachment is comfortable, we, therefore, applied PTG on a volunteer for long-time monitoring of EEG without taking off the electrode pairs. During the monitoring of 12 h, the volunteer slept, exercised, and recovered to calm with two PTG electrodes mounted in the frontal region (Fp1 and Fp2) of the brain and one reference electrode behind the earely Fig. . The intely Fig. . The capIn conclusion, by leveraging two well-known skin electrodes together, we have demonstrated a transparent, highly conductive, and ultra-conformal dry electrode, which is PEDOT:PSS thin film-transferred CVD-grown graphene (PTG). Graphene transferred by PEDOT:PSS without the step of removing the carrier layer was confirmed to be electron doped with higher conductivity; on the other hand, the PEDOT backbones showed improved molecular packing due to the underlying graphene with strong π–π interaction. Thus, the synergistic enhancement effect between graphene and PEDOT:PSS resulted PTG in a very high conductivity (4142 S/cm). Together with the ultra-thin feature, PTG dry electrode exhibited excellent capability of long-time and low-motion artifact electrophysiological signal recording. For facial muscle evaluation, it showed superior advances in lower-motion artifact electrophysiology and less disturbance to laser speckle contrast imaging than commercial Ag/AgCl gel electrodes, and being able to stably and accurately control the robotic hand. For the very weak EEG signals, the user can continuously wear the PTG for 12 h monitoring regardless the extension/contraction and sweat secretion of skin. Our proposed ultra-thin, transparent, and conductive dry electrodes will pave a new way for epidermal electronics.PEDOT:PSS solution (Clevios PH 1000) was purchased from Han Feng, the sodium dodecyl sulfate (99%) was obtained from Innochem, the SEBS and BSL were obtained from Aladdin Industrial Corporation. The decal transfer paper was purchased from The Magic Touch Ltd., UK.2 and 35 sccm CH4. After 40 min growth at 1000 °C, both sides of Cu foil were covered with monolayer graphene. The SDS of 0.25 wt%–1.5 wt% was added to PEDOT:PSS solution and stirred vigorously for 2 h. The BSL was then added to PEDOT:PSS/SDS solution and stirred for another 2 h. Before spin-coating, the prepared PEDOT:PSS solution was filtered by 0.45 µm PES syringe filter. And the PEDOT: PSS solution was spin-coated on the Cu foil with graphene at 2000 rpm for 60 s, forming a PEDOT:PSS/graphene/Cu structure. The PEDOT:PSS coated graphene on Cu foil was annealed at 120 °C. Subsequently, Cu foil was etched by ammonium persulfate [(NH4)2S2O8] solution about 2 h, then the PEDOT:PSS/graphene film was washed in deionized water for 3 times, The film can be transferred to arbitrary substrates, such as silicon wafer, glass, decal transfer tattoo paper, and flexible substrates.Cu foil was put in the quartz tube under a constant pressure of 20 sccm (standard cubic centimeter per minute) HThe sheet resistances were measured using a four-probe resistance measuring meter, with a Keithley 2400 Source Measure Unit. The transparency of the thin film was tested by an ultraviolet-visible spectrometer (UV-2450). The AFM images were obtained in tapping mode . The Grazing-incidence wide-angle X-ray scattering (GIWAXS) measurement was performed on a laboratory beamline system (Xenocs Inc. Xeuss 2.0) with an X-ray wavelength of 1.54 Å and sample to detector distance of 17.1 cm. An incidence angle of 0.2° was used. Diffraction images were recorded on a Pilatus 1 M detector (Dectris Inc.) and processed using the Nika software package, in combination with WAXS Tools. The Raman spectra were obtained at 532 nm excitation (LabRAM HR Evolution). The electrochemical impedance spectroscopy (EIS) measurements were conducted using a CHI600e electrochemical workstation (CH Instruments).Two working electrodes were placed on forearm, and the reference electrode was placed on the back of wrist for sEMG. The electrodes were connected with commercial signal recording equipment (DueLite) and sEMG signals were digitally filtered in MATLAB by a six-order Butterworth filter. And the low cut-off frequency is 10 Hz while the high cut-off frequency is 450 Hz. For ECG, two working electrodes and reference electrodes were placed on arm by using the Einthoven triangle principle. During EOG measurement, the working electrodes were placing on the forehead and lower eyelids, respectively. For easy handling and stable connection to electrophysiological recording instrument, SEBS or decal transfer tattoo paper were sometimes assisted as an additional substrate, and evaporated Ag films were aided in the connection between soft-hard interfaces.The electromechanical vibrator was placed near the working electrodes (about 2 cm) to induce skin vibration, which is similar to vibration in arm movement.Two working electrodes of PTG were placed on levator labii superioris, and one reference electrode was placed behind the ear. Ag/AgCl and PEDOT:PSS electrodes were placed on the same position on the other side of face as control. Laser speckle contrast imaging was recorded in SIM BFI-WF System . The monitor detector was set 25–30 cm above the volunteer’ face, with a maximum field of view of 14.8 × 14.8 cm so that the entire facial area could be illuminated. The frequency used for the blood perfusion image was 1 s, and the size of the image was 512 × 512 pixels. To control the robotic hand, three pairs of PTG were attached on zygomaticus, risorius and corrugator to extract sEMG signals of the index, middle finger, and wrist of the robotic hand, respectively.Brain activity was evaluated by EEG signals using a commercial signal recording equipment (DueLite). The EEG signal was measured by mounting two working electrodes (PTG on tattoo paper with Ag as connections) in the frontal region of the brain and one reference electrode behind the ear. During the monitoring of 12 hours, the volunteer slept for 1 h, and exercised for 0.5 h, finally recovered to calm without taking off PTG electrodes. The EEG signal was analyzed by Matlab and Fast Fourier Transformation.Supplementary InformationPeer Review FileDescription of Additional Supplementary FilesSupplementary Movie 1Supplementary Movie 2


Biocompatible neural guidance conduits are alternatives to less abundant autologous tissue grafts for small nerve gap injuries. To address larger peripheral nerve injuries, it is necessary to design cell selective biomaterials that attract neuronal and/or glial cells to an injury site while preventing the intrusion of fibroblasts that cause inhibitory scarring. Here, we investigate a potential method for obtaining this selective cellular response by analysing the responses of rat Schwann cells and human dermal fibroblasts to isoleucine-lysine-valine-alanine-valine (IKVAV)-capped dendrimer-activated collagen films. A high quantity of nanoscale IKVAV-capped dendrimers incorporated onto pre-crosslinked collagen films promoted rat Schwann cell attachment and proliferation, and inhibited human dermal fibroblast proliferation. In addition, while pre-crosslinked dendrimer-activated films inhibited fibroblast proliferation, non-crosslinked dendrimer-activated films and films that were crosslinked after dendrimer-activation (post-crosslinked films) did not. The different cellular responses to pre-crosslinked and post-crosslinked films highlight the importance of having fully exposed, non-covalently bound biochemical motifs (pre-crosslinked films) directing certain cellular responses. These results also suggest that high concentrations of nanoscale IKVAV motifs can inhibit fibroblast attachment to biological substrates, such as collagen, which inherently attract fibroblasts. Therefore, this work points toward the potential of IKVAV-capped dendrimer-activated collagen biomaterials in limiting neuropathy caused by fibrotic scarring at peripheral nerve injury sites. Nerve autografts, which are nerve tissues transplanted from one part of the body to another in the same individual, are the gold standard for severed peripheral nerve reconstruction ,2. Howev2 and gold surfaces [Attaching small biochemical motifs to a larger scaffold is an attractive method to engage a specific subset of cells on biocompatible materials for neural applications . Small psurfaces . It is usurfaces , while ssurfaces . UnderstAn important consideration when activating collagen scaffolds with small molecules is to ensure that appended biochemical motifs are physically accessible to cells. Recent research has looked into the use of poly ε-lysine dendrimers, which are branched peptides, to activate different scaffold materials such as collagen ,17,18,19Given the potential of dendrimer-activated biomaterials for improved neural repair, the objective of this work is to integrate nanoscale IKVAV-capped poly ε-lysine dendrimers onto collagen surfaces and study the response of rat Schwann cells (RSCs) and human dermal fibroblasts (HDFs). By adding IKVAV-capped dendrimers to non-crosslinked collagen films, pre-crosslinked collagen films or by crosslinking collagen films after dendrimer addition (post-crosslinking), the response of Schwann cells and fibroblasts to different surface presentations of dendrimers can be studied. Previous studies have shown that non-crosslinked collagen films possess a lower tensile modulus compared to 1-ethyl-3-(3-dimethylaminopropyl)carbodiimide hydrochloride (EDC)/ N-hydroxy-succinimide (NHS) crosslinked collagen films, thus providing different environments for cellular attachment and proliferation . What is® CRL-2765™, Manassas, VA, USA). Unless stated otherwise, all other reagents were analytical grade as obtained from Sigma-Aldrich, Gillingham, UK.Insoluble bovine dermis type I collagen powder (product code: 01AWB003) was obtained from Devro PLC, . 16-arm IKVAV-capped poly ε-lysine dendrimers  were obtained from Tissue Click Ltd., Brighton, UK. Human dermal fibroblasts (HDFs) were sourced from Sigma Aldrich . RSC-96 rat Schwann cells (RSCs) were obtained from the ATCC cell biology collection (product code: ATCCw/v) suspension of collagen powder in 50 mM acetic acid at 4 °C for 3 days. Using a commercially available blender, the hydrated powder was blended for 2 min, rested for 2 min, and blended again for another 2 min at 22,000 rpm. Air bubbles were removed from the suspension by centrifuging at 2500 rpm for 30 s. After cooling the suspension at 4 °C overnight, films of micron-scale thickness were cast by reverse pipetting 200 µL of slurry per well into a CytoOne TC-treated 48-well plate  and drying for 48 h in a laminar flow cabinet [Collagen slurries were prepared by swelling a 0.5%  = 708.11 Da). Atto 565 nm NHS ester stock solution was prepared by dissolving 1 mg of powdered ester dye in 0.5 mL of dry dimethyl sulfoxide (DMSO). One mg of dendrimer powder was suspended in 1 mL of phosphate-buffered saline (PBS). For dye conjugation, it was assumed that a 1:1 molar ratio of dye to free primary amine would tag each terminal amine with Atto dye. To adjust to pH 9, 1M NaHCO562) immediately after incubation using a Fluostar Optima plate reader . The absorbance values of the control samples were subtracted from the absorbance values of the experimental samples to correct for background dye absorption.To remove non-conjugated dye, the agitated solutions were dialysed for 24 h in Slide-A-Lyzer dialysis cassettes . The concentrations of purified dye-conjugated dendrimers were measured colourimetrically with a Pierce Bicinchoninic Acid (BCA) Protein Assay Kit,  using a Bovine serum albumin (BSA) protein standard. Prior to measurement, the mixtures were agitated on a rocker for 30 s covered in foil and then incubated at 37 °C for 40 min (78% dye-conjugated dendrimers) or agitated on a rocker for 20–30 min at room temperature (8% dye-conjugated dendrimers). Absorbance values were read at 562 nm (ADye-conjugated IKVAV-capped dendrimers were added to non-crosslinked films (A), pre-EDC/NHS crosslinked films (B), or non-crosslinked films that were crosslinked with EDC/NHS agents after dendrimer addition (C). The three varieties of dendrimer-activated collagen films were simultaneously fabricated over three days to ensure that they were exposed to the same experimental conditions. To coat the collagen films, which were used as dried, with dye-conjugated dendrimers, 120 µL of dendrimer solution (µg of dendrimer/mL of PBS) was added to non-crosslinked  or crosslinked (B) collagen films. The films were incubated for 1.5 h on a rocker at room temperature covered in foil. After incubation, excess dendrimer solution was removed, and the films were washed with 3 × 200 µL PBS. The films were dried in a dark fume hood overnight.v/v) ethanol was added and incubated for 1.5 h. These conditions were deemed optimal by Olde Damink and co-workers [To chemically and covalently crosslink the collagen films, a 5:2:1 molar ratio of EDC/NHS/COO-(Col) in 75% (-workers  for stanv/v) in deionised water, 0.1% Tween (v/v) in PBS, saturated saline, or 0.1 M citrate buffer (pH 3). Collagen films were analysed at 0 h (pre-wash), and after 2 h, 1 day, 2 days, 3 weeks, and 7 weeks. Fluorescent images of the collagen films were taken with a Zeiss Axio Observer Z1 microscope fitted with an Axiocam 530 camera  after each time interval. Fluorescent images were taken at 10× magnification and at 565 nm and 488 nm to detect the dye-conjugated dendrimers and collagen autofluorescence, respectively. Prior to imaging, the samples were washed with 3 × 200 µL deionised water, and all images were taken of collagen films that were hydrated in 200 µL of either fresh PBS or deionised water. In between imaging time points, the samples were incubated in their respective washing solutions on a rocker at room temperature covered in foil. ImageJ software was used to quantify the dendrimers bound to the collagen films after each time interval.To probe dendrimer-collagen interactions, all samples were washed with various detergents for increasing lengths of time. The samples were washed with PBS, 75% ethanol (EtOH) (v/v) foetal bovine serum and 1% (v/v) streptomycin/penicillin antibiotics in a humidified incubator with 5% CO2 at 37 °C. RSCs were passaged 1:20 and HDFs were passaged 1:4 every 3–4 days. Cell layers at ~75% confluence were harvested for analysis by detaching the cells from T75 flasks with 0.05% (w/v) trypsin/0.02% (w/v) EDTA, centrifuging the cells at 180 g for 3 min, and resuspending the cells in serum-free DMEM.RSCs and HDFs were cultured in Dulbecco’s modified Eagle’s medium (DMEM) containing 10% (v/v) BSA solution in PBS and washed with 3 × 300 µL PBS to prevent non-specific cell binding. The cells were seeded at a density of 5 × 105 cells/mL for RSCs and 3.06 × 105 cells/mL for HDFs. These cell densities were determined using a haemocytometer. Three hundred µL of appropriately diluted cells were added to the films and tissue culture plastic (TCP) controls and incubated for 1.5 h in a humidified incubator with 5% CO2 at 37 °C. Loosely bound cells were gently removed with 3 × 300 µL PBS washes.Before seeding the cells, the dendrimer-activated collagen films were blocked for 1.5 h with 300 µL of 2% (v/v) in deionised water was added to all the samples and incubated at 4 °C overnight. According to manufacturer recommendations, 100 µL of LDH substrate  was made and was added to the samples and left to incubate until a significant colour change was apparent (20–30 min). Absorbance values were read at 490 nm (A490) using a Fluostar Optima plate reader .The number of cells bound to the collagen films were quantified colourimetrically by measuring the amount of lactate dehydrogenase (LDH) released from lysed cells into the media ,24. To lv/v) foetal bovine serum at a density of 1.46 × 104 cells/mL and 2.2 × 104 cells/mL for RSCs and HDFs, respectively. Eight hundred µL of these cell suspensions were incubated on the collagen samples for 72 h in a humidified incubator with 5% CO2 at 37 °C to allow for ~75% confluent proliferation.Dendrimer-activated collagen films were blocked with BSA solution and seeded with cells as for cell attachment analysis, except the cells were resuspended in media containing 10% (w/v) glutaraldehyde directly to the cell media ) for 30 min at room temperature. After the glutaraldehyde solution was removed, the films were washed 3× with sterile deionised water. The cells were stained with 200 µL of 5 µM DAPI  in water. The DAPI was incubated for 10–20 min at room temperature and then removed. The collagen films were washed 3× with sterile deionised water before they were imaged with a Zeiss Axio Observer Z1 microscope fitted with an Axiocam 530 camera. Fluorescent images were taken at 10× magnification and at 358 nm excitation wavelength to detect the DAPI stain. The number of cells per field of view were quantified with CellProfiler .To image the cell number on dendrimer-activated collagen films, the cells were fixed by adding 200 µL of 25% , a 2-way analysis of variance (ANOVA) was still conducted when sample sizes were roughly equivalent. Mean values were compared using 2-way ANOVA with Tukey’s post-hoc tests to determine statistical significance (p ≤ 0.05). Bonferroni comparisons were used for multiple comparisons. The values reported are means of triplicate measures ± standard deviations, unless stated otherwise. Error bars indicate standard deviations from the mean. N/S denotes p ≥ 0.05.The data were analysed for normality and equal variances using Shapiro-Wilk (or Kolmogorov-Smirnov) and Levene’s tests, respectively, with To visualise and quantify the IKAV-capped dendrimers integrated onto collagen substrates, the dendrimers were first conjugated with Atto 565 nm fluorescent probes. The efficiency of Atto dye conjugation was investigated and the dendrimer:dye ratio was optimised. The percent conjugation was defined as the ratio of the molecular weight of the free 16-arm dendrimer  ,15 to thDye-tagged IKVAV-capped dendrimers were incorporated onto collagen substrates using three different methods to yield three different types of collagen surface environments. Dye-tagged IKVAV-capped dendrimers were coated onto (1) non-crosslinked collagen films, (2) pre-crosslinked films that were initially crosslinked with EDC/NHS, and (3) non-crosslinked films that were subsequently crosslinked with EDC/NHS after the addition of dendrimers to yield “post-crosslinked” films. It was thought that crosslinked surfaces would be more hydrophobic given the ablation of exposed free amines (lysine-based) and carboxylates (glutamate and aspartate-based) upon EDC/NHS treatment ,25. Conv3+ groups) than 78% dye-tagged dendrimers, and it was hypothesised that differences between their interactions with the collagen films could be attributed to differences in steric bulk and overall charge.Dendrimers with low extent of Atto 565 nm dye-conjugation (8%) and high extent of dye-conjugation (78%) were investigated in this study. Of the dye-tagged dendrimers, 8% have a lower molecular weight and higher overall charge (fewer dye-consumed NHv/v) in deionised water, 0.1% Tween (v/v) in PBS, saturated saline, and 0.1 M citrate buffer (pH 3). These solvents were chosen to probe different types of chemical interactions that could exist between the dendrimers and the surface of the collagen films. To probe non-ionic interactions, 75% EtOH was used, and saline and PBS solutions, which are salt-based buffers, were used to probe ionic interactions. To probe ionic and non-specific, non-covalent interactions, 0.1% Tween in PBS—a salt-based buffer containing a non-ionic detergent—was used. Citrate buffer (pH 3) was used to probe interactions at an acidic pH. It was thought that by incubating the films in wash buffer for an extended period of time and taking fluorescence measurements of the films at various time points, the relative rates of dendrimer dissociation from the collagen substrates could be compared. These trends would provide insight into how bulky, charged nanomaterials, such as dendrimers, interact with different collagen surface chemistries long-term, and how this could ultimately impact cell binding and proliferation.Long-term interactions between charged, dye-tagged IKVAV-capped dendrimers and non-crosslinked, pre-crosslinked, and post-crosslinked collagen films were probed by studying how dendrimers dissociate from the films when exposed to various solvents. As mentioned in the methods above, the solvents selected were PBS, 75% ethanol (EtOH)  to dye-conjugated dendrimers (565 nm), was initially measured at each time point to ensure that collagen substrate breakdown was not a confounding factor for dendrimer dissociation . After cp ≥ 0.05) unlike the other wash buffers. For the non-crosslinked films activated with 8% dye-tagged dendrimers, none of the wash buffers significantly dissociated dendrimers from the films (p ≥ 0.05). These results suggest that the overall charge on the dendrimers affects how they interact with non-crosslinked collagen substrates over time.p ≤ 0.001), with limited dissociation for the remaining incubation time period. On the other hand, there was no significant dissociation (p ≥ 0.05) of either 8% or 78% dye-tagged IKVAV-capped dendrimers incorporated into post-crosslinked films when washed with any of the buffers. These results provide evidence that dendrimer-collagen interactions in the pre-crosslinked films are easily disturbed with a variety of aqueous solvents, and that dendrimer-collagen interactions in the post-crosslinked films are too strong to be disturbed by aqueous, ionic wash buffers.Unlike the non-crosslinked films, for which dendrimer dissociation was dependent on both the extent of dye-conjugation and also the type of solvent used to dissociate the dendrimers, dendrimer dissociations from pre-crosslinked and post-crosslinked films were unaffected by these parameters. When dendrimer-activated pre-crosslinked films were washed with any of the wash buffers, there was drastic dissociation of both 8% and 78% dye-tagged dendrimers in <2 h   via elecAfter the mode of chemical interaction between the dendrimers and the various collagen films were probed, cellular responses to the various dendrimer-activated collagen films were studied. Rat Schwann cells (RSCs) were used to model how glial cells, which nourish neurites during regrowth , respondv/v) Triton-X 100 in water. As all cells were lysed, the quantity of bound cells could be determined given a linear relationship between LDH detection and cell density [To probe cellular adhesion, a modified LDH assay, as previously reported, was employed ,24. WhilWhile fibroblasts had similar binding affinity to collagen films, irrespective of both the degree of crosslinking and extent of dendrimer activation, RSCs only had low binding affinity to non-crosslinked and crosslinked collagen films with 0.12 µg/mg or no dendrimer activation . Increasp ≥ 0.05) by dendrimer activation. This data supports that dendrimer-activated collagen surface properties affect HDF proliferation, and that a relatively large quantity of untethered dendrimers integrated into pre-crosslinked films provides the desired ablation of fibroblast proliferation.In addition to the cellular binding trends analysed after one day of incubation on the collagen substrates, cellular proliferation was probed after three-day incubation on the collagen surfaces. Cellular proliferation was calculated as a fold-change at day 3 from the initial loading density on day 0. A fold-change of 1 indicated no change in cell density. For non-crosslinked and post-crosslinked films, lower dendrimer quantities (0.12 and 0.36 µg/mg) did not reduce HDF proliferation relative to films without dendrimer activation . DespiteDAPI-stained fluorescent images of HDFs after 3 days in culture were obtained to show the cell distribution . This shRSC proliferation was dependent on the collagen substrate and peptide coating. Intermediate dendrimer coating concentrations  increased RSC attachment to both non-crosslinked and pre-crosslinked films. However, only the highest dendrimer coating concentration (0.72 µg/mg) increased RSC proliferation, relative to non-dendrimer-activated films, for all three collagen substrate varieties . This inAs for HDF proliferation, DAPI-stained fluorescent images of RSCs on the different collagen substrates with various extents of dendrimer activation were obtained . Unlike In this study, IKVAV-capped dendrimers were integrated into three substrate varieties: non-crosslinked collagen films, pre-crosslinked collagen films or initially non-crosslinked collagen films that were crosslinked after dendrimer addition (“post-crosslinked collagen films”). Prior to studying the cellular responses to these three substrate varieties, the chemical interactions between the dendrimers and the collagen films were probed.It was predicted that dendrimers with higher MW and lower charge (78% dye-tagged dendrimers), due to their hydrophobicity and steric bulk, would preferentially associate with similarly hydrophobic, crosslinked collagen substrates. On the other hand, it was anticipated that dendrimers with lower MW and higher charge (8% dye-tagged dendrimers), which are more hydrophilic, and less sterically encumbered, would favour ionic or electrostatic associations with non-crosslinked collagen. However, both 78% and 8% dye-tagged IKVAV-capped dendrimers had higher initial association to post-crosslinked films than non-crosslinked films, even though both substrates were non-crosslinked at the point of dendrimer addition . Unlike While the association studies suggested that dendrimer interactions with non-crosslinked films are dependent on local environments, the dissociation assays elucidated the dendrimer binding mechanisms for pre-crosslinked and post-crosslinked films, in addition to further supporting that dendrimers make ionic associations with non-crosslinked films.p ≤ 0.001) to a significant extent. Ethanol may wash away non-covalently associated dendrimers like 0.1% Tween solutions. In addition, while water molecules stabilise collagen triple helices via hydrogen bonded hydration shells [For dendrimer-coated non-crosslinked films, the salt containing solutions, saline, PBS and 0.1% Tween (in PBS) dissociated a large quantity of 78% dye-conjugated dendrimers . It is ln shells , larger Interestingly, citrate buffer (pH 3) did not significantly dissociate dendrimers from the non-crosslinked films. As an acidic buffer, citrate buffer was expected to increase dissociation of dendrimers from their collagen substrates as negatively charged glutamate (pKa 4.25) and aspartate (pKa 3.65) residues  become pp ≤ 0.05, not denoted in Furthermore, the dissociation assays also suggested that dendrimers are physisorbed onto pre-crosslinked films. Dendrimers, which are sterically bulky, likely entangle into pre-crosslinked substrates via weak non-covalent interactions, as these bioactive motifs quickly and completely dissociate when washed with any of the solvents. The lack of strong chemical interactions also explains why the dissociation of dendrimers from pre-crosslinked films is not dependent on the initial quantity of associated dendrimers . While 0Finally, dendrimers may be more permanently bound to post-crosslinked films than either non-crosslinked or pre-crosslinked films as there was little dissociation for all wash buffers. The detergents used in these dissociation studies targeted ionic, non-ionic, and non-covalent interactions, showing that these are not the only bond types present between the peptide and collagen after EDC/NHS crosslinking. Although covalent interactions cannot be probed directly, as detergents such as SDS that disrupt all non-covalent interactions in proteins can also breakdown collagen films, these results suggest that covalent linkages may form between the peptide and collagen post-crosslinking. Like for non-crosslinked films, these stronger chemical interactions also explain why dissociation was proportionally increased when higher dendrimer quantities were initially integrated into the collagen films. These studies show that while EDC/NHS crosslinked collagen has been shown to have superior mechanical stability relative to non-crosslinked collagen, albeit at the expense of reduced bioactivity ,22,33,34In designing biomaterials for neural applications, it is important to not only consider the response of neurite forming cells, but to also consider the responses of helper glial cells, such as Schwann cells, and inhibitory cells, like scar-forming fibroblasts. In this study, human dermal fibroblast (HDF) and rat Schwann cell (RSC) adhesion and proliferation on various IKVAV-capped dendrimer-activated collagen substrates were studied to better understand how various cell types interact with laminin-derived peptides, such as IKVAV, which have been thought to be promising for neural applications .IKVAVSADR coated onto tissue culture plastic was found to increase fibroblast adhesion at or above the mass of IKVAV used per sample in this study [For biomaterials targeted at neural repair, the response of fibroblasts is extraordinarily important given that scar-forming fibrotic tissue is one of the greatest hinderances to attaining full-functional recovery after injury . The datis study . As showin vivo [Although there was an increase in HDF adhesion to non-crosslinked films with greater dendrimer-activation, this increase was minimal relative to the already high affinity of HDFs to both non-crosslinked and crosslinked collagen films without dendrimer activation. When non-crosslinked collagen films, with or without IKVAV-capped dendrimers, are used, HDFs may bind exposed biochemical motifs on native collagen such as GFOGER, or laminin-derived peptides such as IKVAV, via β1-type cell-surface integrins . This isin vivo  and humain vivo . While tin vivo . HoweverUnlike HDF adhesion, HDF proliferation, as well as both RSC attachment and RSC proliferation (proliferation defined here as fold-change in cell number over 3 days in culture), were affected not only by the surface properties of collagen but also the presence and quantity of IKVAV motifs present. As discussed above, dendrimers can bind ionically/non-specifically, non-covalently, or covalently to non-crosslinked, pre-crosslinked and post-crosslinked films, respectively. Ionic or non-specific interactions, and non-covalent physisorption do not alter the functional groups on IKVAV motifs. However, covalent binding reduces primary amines to amides on lysine or terminal valine residues on IKVAV. Higher RSC adhesion  and prolp ≤ 0.05) and HDF proliferation on non-crosslinked films, relative to pre-crosslinked films, even though both substrate varieties had high quantities of unaltered IKVAV motifs. These lower-than-expected cellular responses may be affected by the decreased stiffness of non-crosslinked surfaces [When lower dendrimer quantities were used, there was lower than expected RSC attachment  was incorporated into pre-EDC/NHS crosslinked collagen films, rat Schwann cell attachment and proliferation were promoted while human dermal fibroblast proliferation was inhibited, as desired. These results show the potential of IKVAV-capped dendrimer-activated collagen biomaterials for neural applications since glial Schwann cells help neurite organisation and reduced fibroblast activity can limit detrimental fibrotic scarring at peripheral nerve injury sites. The different cellular responses to pre-crosslinked and post-crosslinked films also highlighted the importance of having fully exposed, non-covalently hindered biochemical motifs in attracting the desired cell type.


Polyamide 46 (PA46) is used in various automotive parts because of its excellent heat resistance and mechanical properties. This study aims to improve the frictional properties of PA46 using the lubricating ability of graphene. Nanocomposites are prepared via two mixing methods: Graphene powder is compounded directly with PA46 pellets through a twin-screw extruder, or PA46 powder is added to graphene dispersion for self-adsorption, and subsequently, it is dried and compounded with PA46 through the twin-screw extruder. Application of the nanocomposite in the friction field is evaluated via the pin-on-disk method. The coefficient of friction of the nanocomposite prepared by self-adsorption is lower than that of the nanocomposite prepared by direct compounding. The mechanical properties of the nanocomposite fabricated by self-adsorption are superior to those of other materials. This can be attributed to the uniform dispersion of graphene and the strong attractive force between the PA46 matrix and graphene. Polyamide (PA) is used in the automotive industry as a representative engineering plastic. Polyamide 46 (PA46) exhibits higher crystallinity, melting temperature, and thermal stability than the commonly used polyamide 6 (PA6) and polyamide 66 (PA66) because of its high amide content per unit length chain length and its symmetrical chain structure. Further, PA46 exhibits excellent wear properties because of its high crystallinity, which incentivizes its application in highly heat-resistant frictional parts, such as valves and gears ,2,3. RecGraphene is a two-dimensional single layer of graphite, and it is an emerging two-dimensional carbon-based nanomaterial. Graphene exhibits excellent thermal properties, electrical conductivity, mechanical strength, moisture barrier properties, and wear resistance. Consequently, graphene-polymer nanocomposites are expected to have numerous potential applications ,6,7,8,9.In this study, graphene was composited in various ways to improve the wear resistance of PA46. A chemically modified graphene (CMG) dispersion was used to increase the dispersibility of graphene, and composite particles were prepared and processed using PA46 nanoparticles and by exploiting self-adsorption behavior. For comparison, a composite was prepared by directly compounding graphene powder with PA46 pellets. The frictional properties were compared and analyzed based on the manufacturing method.® TW341 grade from DSM  was used to produce PA46.Graphene powders of BGF-H grade  and BGF-R grade  were provided by BESTGRAPHENE . For the production of functionalized graphene, 150 μm grade graphite flakes purchased from Graphene Supermarket were used. Stanyl2SO4/H3PO4 (3600 mL:400 mL) and KMNO4  at a 9:1 ratio and stirred for 12 h. The reaction mixture was cooled to room temperature (23 °C) and diluted by immersing it in ice water (4000 mL) containing 30% H2O2 (3 mL) after the reaction was completed. The non-oxidized graphite was removed by filtration, whereafter the oxidized graphite was precipitated by centrifugation at 4000 rpm for 2 h. The precipitate was washed a minimum of two times in a mixed solution of DI water (2000 mL), 30% HCl (2000 mL), and ethanol (2000 mL), and subsequently, it was freeze-dried to obtain fine non-agglomerated particles. Distilled water (2000 mL) and oxidized graphite (2 g) were mixed and sonicated for a minimum of 4 h using an ultrasonic grinder to obtain an aqueous solution of GO.Chemically modified graphene (CMG) was designed to substitute -OH, -COOH, and epoxide groups of graphene oxide with -CN, amide, and amine through functionalization to enable molecular structural interaction with PA46. An improved, recently developed, and widely used method was used to produce graphene oxide (GO) . At 50 °+) of the substituted 4,4-oxidianiline generated an electrostatic repulsive force between the graphene monolayers to yield a chemically modified graphene colloid with high dispersibility.A graphene surface-modification reaction was performed using 4,4-oxidianiline to obtain a chemically modified graphene colloid that can undergo physicochemical interactions with PA46. A 500 mL aqueous solution of 0.1% GO was prepared by sonicating for a minimum of 2 h using an ultrasonic homogenizer. The aforementioned solution was mixed with a solution of 20 g of 4,4-oxidianiline in 500 mL of N,N-dimethylformamide (DMF). This mixture was stirred at 90 °C for 40 h. Thereafter, the mixture was cooled to room temperature (23 °C), a precipitate was separated by centrifugation, washed three to five times with acetone, and filtered to obtain graphene. The as-obtained graphene was immediately immersed in ethanol, and the positive charge (-N50 = 385 μm) PA46 powder was mixed with a graphene dispersion, precipitated through chemical self-adsorption, and dried  to prepare a master batch . The MB was injected into the main feeder of the twin-screw extruder with PA46 pellets. The graphene content was adjusted to 0.05, 0.10, and 0.15 wt% by adjusting the ratio of 0.5 wt% MB and PA46 pellets (sample name S05). The mixing ratios are listed in Nanocomposites were fabricated using two methods. In the first method, PA46 pellets were injected into the main feeder of a twin-screw extruder , graphene powder was injected into the side feeder, and the two feeds were compounded at 300–330 °C. Samples with graphene powder  contents of 0.05, 0.1, and 0.2 wt%  were prepared. In the second method, finely ground , transmission electron microscopy , zeta potentials measurement , and FT-IR . A pin-on-disc tester  was used to determine the frictional properties of the material. The testing was conducted at 25 °C, a motor speed of 400 RPM, a test load of 10 N, and by using a stainless-steel ball of 6 mm diameter as the pin. Following the frictional evaluation, the frictional surface was analyzed using an optical microscope . The tensile strength, flexural strength, and flexural modulus were measured using a universal testing machine .−1 are assigned to the amide vibrations. Peaks between 3330 and 3300 cm−1 were attributed to the secondary amide group. The C-N stretching vibration gave rise to bands in the region 1200–1100 cm−1.−1) and G band (near 1580 cm−1) peaks of graphene can be identified [−1) was identified in the spectrum of CMG that underwent reduction and surface-functionalization. The D/G ratio decreased to 0.927 for CMG compared to 1.049 for GO. The smaller D/G ratio of CMG than that of GO indicates that reduction and functionalization decrease defects in the CMG structure, which increases crystallinity. This suggests the formation of a crystalline structure that possesses the functionality of graphene. Since the 2D peak near 2700 cm−1 is caused by the double resonance of a phonon with the energy of 1350 cm−1, the particular peak appears around 2700 cm−1, which is double the energy of 1350 cm−1. In the case of multilayered graphene, various scattering processes take place because of the higher number of energy bands, and this increases the number of 2D peaks and thereby leads to a broad 2D band [The structural characteristics of the as-synthesized graphene oxide and CMG were identified through their Raman spectra. In the Raman spectrum of graphene, a D peak is caused by defects within the crystal and the edge of the sample, whereas a G peak is attributed to the vibration of the hexagonal crystal structure composed of sp2 carbon atoms commonly seen in graphite-based materials. entified . The 2D  2D band . As showThe number of layers of graphene is the main parameter in basic research conducted to improve frictional properties using graphene. Interactions between graphene layers occur through van der Waals forces. The mechanical properties of graphene change significantly as the number of graphene layers increases because van der Waals forces have low bonding strength compared to the covalent bonds connecting carbon atoms in individual planes . PreviouThe effect of the friction coefficient can be examined from The friction coefficient of the S05 sample compounded through the self-adsorbing powder was smaller than that of the GH and GR conditions directly compounded with the polymer matrix, as shown in The uniform dispersion of the nanoadditives and the attractive force between graphene and the matrix are critical to the mechanical properties of nanocomposites. The frictional and mechanical properties of graphene-reinforced polyamide 46 were investigated. The uniform dispersion of the nanoadditives and attractive force between the graphene and matrix are critical to the frictional and mechanical properties of nanocomposites. The dispersibility and compatibility with the matrix were higher when the graphene was self-adsorbed onto the PA46 surface. Thereafter, it was subjected to secondary melt processing when the graphene powder was directly compounded. Thus, the friction coefficient decreased by more than 50%. The findings of this study may guide the application of PA46 composite materials in friction-driving parts to reduce the weight of automotives.


For expanding applications of spongy graphene aerogels (GAs) cost-effectively, we report a marriage of the two-step hydrothermal reduction and atmospheric drying method to fabricate a spongy CNC-graphene aerogel (CNG) with oil/water selectivity and tunable mechanical strength by a low-cost and straightforward approach. The reduced graphene oxide (rGO) with CNC by the ice-templated method can give rise to forming the hierarchical structure of hybrid GAs within the PUS network. Meanwhile, the fractured structure of PUS with a pre-compressive step arouses more versatility and durability, involving its selective and high-volume absorbability (up to 143%). The enhanced elastic modulus and more significant swelling effect than pure sponge materials give it a high potential for durable wastewater treatment. Assembling graphene to a three-dimensional 3D) porous monolith while keeping the intrinsic advantage of the building blocks is of great promise to achieve the practical application ,2. CelluD porous Additionally, 3D graphene/cellulose monoliths utilizing APD showed excellent hydrophobicity and elasticity, emerging as promising large-scale and low-cost absorbent materials in water treatment ,12,13. XHerein, porous cellulose graphene aerogel (CNG) was prepared using a two-step reduction process without lyophilization, where the cellular walls of aerogels enclosed by macroscopic-scale sponge porous scaffold led to a primarily enhanced mechanical performance. Even when the impact load with maximum (99%) strain is applied, the CNG exhibits resilience capacity without fracture and degradation. In addition, the cooperation of GAs and sponge porous cell walls endows a water droplet “sticky” hydrophobicity, which can absorb n-heptane over 1.3 times its volume and 120 times its weight and reuse manually about 1000 times. More importantly, these graphene sponges display a noticeable swelling effect, which means that adsorbed oil extended storage in more volume space of the nanocomposite. The CNG enabled a combination of visible polymer framework and microscale graphene scaffold, which can be successfully used for efficient and cost-effective water purification. These properties also demonstrate that, if graphene sheets are assembled into polymer skeletons where graphene nanosheets are kept, novel and combined properties will be observed.http://sciencek.com/, accessed on January 2018) and its length nanoscale is approximately 100 to 200 nm size range, with the length–diameter ratio being 10–20. Graphite powder (CP), L-Ascorbic acid (AR), sulfuric acid (98% ), sodium hydroxide (AR), potassium permanganate (AR), and hydrogen peroxide (AR) were supplied by Shanghai Aladdin Bio-Chem (China) and used without further purification. Graphene oxide as a raw material was synthesized according to a modified Hummer’s method [The recycled PUS samples were obtained from Shenzhen Tenghuida Electronic auxiliary material Co., Ltd. The Cellulose Nanocrystals (CNCs) were obtained from ScienceK (s method . The radCNG was characterized with a Nicolet 6700 Fourier transform infrared spectroscopy (FT-IR). Scanning electron microscopy (SEM) images were obtained by a Hitachi SU8010 at 20 kV. The X-ray diffraction (XRD) analysis is done with an X-ray source of CuKA high efficient and energy-saving wet shaping method under atmospheric pressure was demonstrated in this work. The typical fabrication procedure was depicted in Moreover, these steps towards a more facile and green route, CNC/GA as an essential building block, are assembled into the 3D architecture of PUS. When the intermediate is frozen, GO nanosheets are expelled by the grown ice crystals, also known as ice templating, and for the formation of the micro-pore network in each sponge nest. The freezing temperature and the amount of oxygen-containing groups of GO are adjusted via reduction time . These eame time . The uniThe XRD diffraction patterns of CNC, the prepared GO and the rGO/PUS and after the CNC adding (CNG) are presented in s nature . After aarticles . These cTypical architectures of CNG sponge were prepared by the hydrothermal reduction and frozen by dry ice, as shown in Comparable results were obtained when different concentrations of the dispersed GO solution were tested. Previous research has shown the mechanical properties of the pre-treated sponge relied mainly on the skeleton contact area and being verified that the fractured sponge network shows an increase of the contact area with the compressive stress in contrast to the network without pre-compression treatment . Aside fIn contrast to what is observed in the CNG specimens, the micropores structure strongly affects the stress–strain response without a plateau stress region such as deformation over the above strain region. This phenomenon can be attributed to the high modulus of whisker-like CNCs nanoparticles and changes induced by the dual network, which combined polymer networks and graphene cellular porosity. The previous research also revealed that the compressive yield stress of porous graphene materials increases linearly with an increase in porosity . In ordeOwing to the hydrophobic porous structure and ultra-lightweight, the oil adsorption capacity  of cellulose/graphene aerogels up to 198 times their weight . PristinIn particular, the volume ratio of absorbed liquid is much higher than lightweight aerogel and other graphene-based foams as shown in In summary, spongy CNC/graphene aerogels have been successfully prepared via an ambient-drying method with a pre-compaction treatment. The corresponding hybrid structure allows CNGs to have unique compressive strength to afford compression cycling numbers and maintain a linearly mechanical property. Moreover, optimizing the freezing and drying process and controlling the internal microstructure with hierarchical pores simultaneously led to a high volumic oil-absorption capacity (up to 142%). The composite developed here also demonstrated a swelling effect with easy recyclability. Significantly, this energy saving synthesis process would be a key progress for realizing large-scale commercial production, and this graphene-based hybrid material provided access to assembling a monolith by embedding porous structures of graphene aerogel inside another porosint. The networks’ interactions changed the material’s bulk properties to find broader promising applications for graphene aerogels, including pressure sensors and petroleum contamination remediation.