High and low temperature direct CVD growth of graphene on these substrates including growth mechanism and morphology is discussed. Detailed discussions are also presented for Si and Ge substrates, which are necessary for next generation graphene/Si/Ge based hybrid electronic devices. Finally, the technology development of the metal‐catalyst free direct CVD growth of graphene on these substrates is concluded, with future outlooks.To fabricate graphene based electronic and optoelectronic devices, it is highly desirable to develop a variety of metal‐catalyst free chemical vapor deposition (CVD) techniques for direct synthesis of graphene on dielectric and semiconducting substrates. This will help to avoid metallic impurities, high costs, time consuming processes, and defect‐inducing graphene transfer processes. Direct CVD growth of graphene on dielectric substrates is usually difficult to accomplish due to their low surface energy. However, a low‐temperature plasma enhanced CVD technique could help to solve this problem. Here, the recent progress of metal‐catalyst free direct CVD growth of graphene on technologically important dielectric (SiO Owing to its excellent characteristics, 2D graphene sheet is regarded as a next‐generation transparent conductive electrode for applications in various electronic devices.FigureEf(Gr)) and lightly doped Si (Ef(Si)) at thermal equilibrium (dark condition) is shown in Figure Vfbias, which brings the Fermi level downward with respect to its “unbiased” condition. In this way, the Fermi level comes into close proximity to the quasi Fermi level for holes in Si, and thereby the number of accessible states for photoexcitation is significantly reduced. When reverse bias is applied, Ef(Gr) shifts to higher value and generates a large number of available energy states for holes to inject .2, ZrO2, HfO2, h‐BN, Al2O3, Si3N4, quartz, MgO, SrTiO3, TiO2, AlN, glass, and mica, and semiconducting substrates such as Si, Ge, GaN, and SiC. Merits and demerits of using high and low temperature CVD processes including growth mechanism and morphology of the graphene on these substrates have been discussed. Detailed discussions are presented for Si and Ge substrates, as they are important semiconductors, and suitable for next generation graphene/(Si/Ge) based hybrid electronic devices. Important results have been summarized in tables, and finally conclusions and outlook have been presented.Through the above discussion, it is evident that these inventions pave a way toward the production of graphene, in‐depth knowledge of the correlation among products, microscopic processes, and experimental conditions. This motivates the research for the metal‐catalyst free direct CVD growth of graphene on various dielectric and semiconducting substrates. Herein, we present a comprehensive review focused on the recent progress made toward the metal‐catalyst free direct CVD growth of graphene on technologically important dielectric substrates such as SiO22.14) on bare SiO2/Si substrates via oxygen aided atmospheric pressure chemical vapor deposition (APCVD) process.2 substrates could be achieved by using PECVD.Figure2/Si substrate was cleaned and then heated in H2  atmosphere at 1000 °C for 15 min. Three kinds of seeds were prepared onto the substrates, which include 1) mechanical exfoliation using scotch tape peel‐off graphene; 2) graphitic clusters nucleated by C2H4 (99.9%) + H2 plasma CVD  or CH4 (99.9%) + H2 plasma CVD ; and 3) patterning of nanoislands graphene by oxygen plasma etching and electron beam lithography. The position of the substrate with the seeds was at the center of the furnace. H2 plasma (H2: 250 mTorr) was generated upstream to activate the edge of the seeds at 500 °C. C2H4+H2 plasma CVD  or CH4+H2 plasma CVD  was then used for graphene growth. The PECVD growth mechanism of graphene on SiO2/Si substrate is shown in Figure 2 content, the pressure, and the growth temperature decided the growth at the edge   was found 1 nm min−1 and increased about 4.5 nm min−1 at 48 mTorr.A low temperature growth leads to compatibility and minimizes energy consumption along with cost effectiveness for bulk production in the industry. Therefore, in order to deposit a clean high‐quality graphene directly on dielectric substrates, a controllable, low‐cost, and reliable mode is important at low temperature. The catalyst‐free direct growth of polycrystalline graphene was achieved by the pyrolysis of methane  direct growth of micrometer‐scale graphene crystals on SiO2 substrates could be achieved by using PECVD technique. These graphene crystals can be directly embedded to fabricate electronic devices, thereby eliminating the conventional postdeposition transfer process. Lack of transfer process, good control of the method, excellent quality of grown graphene, and the compatibility of this process with the current microelectronics technology make it a facile approach for future use in graphene electronics.In addition to this, the metal‐catalyst free direct CVD growth of graphene on SiO2.22 substrates can be achieved at a temperature lower than 480 °C as reported by Scott et al.2 substrates using acetylene (C2H2) as a precursor. As soon as the substrate was placed, the reaction chamber was evacuated followed by the flow of Ar. The substrate was heated to a temperature of 325–650 °C. The flow of Ar was maintained for 10 min and then C2H2 was introduced into the chamber. X‐ray photoemission spectroscopy (XPS) confirmed the formation of zirconium oxide being present in the form of monoclinic baddeleyite, and being slightly oxygen deficient. The low temperature enables catalytically active oxides to form sp2 carbon. Wang et al.2 substrate using either CH4 or ethanol (C2H5OH) as a precursor via thermal CVD. First, the substrate was mounted to the central region of the quartz tube, and heated from room temperature to 1130 °C in 50 min with 50 sccm H2 and 50 sccm Ar. Second, the substrate was annealed for 20 min at 1130 °C, and then a certain flow rate of CH4 or C2H5OH vapor was introduced into the chamber to initiate graphene growth. 2D or 3D growth of graphene could be controlled by altering feedstock concentration and reaction time  plane  = 86 cm−1), D , and 2D  bands  analysis which in turn was conducted through Moiré pattern region of 0.55 nm  image confirmed the presence of a graphene/h‐BN film having a thickness of 2 nm Figure f. The im4 as a precursor via metal catalyst free CVD technique. Stacked‐layers on h‐BN substrate using hexane (C6H14) (vapor) as a precursor via APCVD technique was demonstrated by Liu et al.Similarly, Ding et al.2.54 as a precursor via near equilibrium CVD technique. Quartz having smooth face toward the downward side was placed in a high‐temperature horizontal silica tube furnace. The furnace was heated at 1180 °C and stabilized for about 30 min in 250 sccm H2 and 300 sccm Ar. During entire growth, a gas mixture (CH4:H2 = 1.9–2.3:50) was taken as the carbon source. AFM phase image of graphene grain on quartz substrate is shown in Figure4 as a precursor via remote (r)‐PECVD technique. Nanographene films having good uniformity were grown on 4 inch wafer. By controlling the growth duration, the value of transmittance on quartz was found to be greater than 92%, whereas it exhibited a low resistance of 40 kΩ sq−1 at 550 nm. The direct growth of VGs on quartz substrate can also be achieved using either CH4 or ethanol (C2H5OH) as a precursor via thermal CVD, as demonstrated by Wang et al.Owing to the high melting point along with structural stability, quartz has a flat surface having roughness about 0.414 nm. Chen et al.2 (20 sccm) and then CH4 (10 sccm) was introduced for 15, 25, 35, and 45 min, respectively. Figure Recently (2018), Chen et al. demonstrated direct growth of graphene on vertically placed quartz substrates via APCVD.2.63N4) substrates can be achieved by a two‐stage CVD process.3N4 surface, and at a later stage merged together to form a polycrystalline graphene film. These graphene films can be used to fabricate field‐effect transistors (FETs). A clean Si3N4/SiO2/Si substrate was kept into a quartz furnace at high‐temperature   . The graphene film was observed to be clearly separated from Si3N4 substrate, and formation of wrinkles took place along the boundary of the sheets  was found to be three times higher than those grown on SiO2/Si substrates as well as better than graphene grown using metal catalyst.3N4 substrates via two‐stage metal catalyst‐free CVD process where the detrimental effects of the substrate were minimized. In addition, the graphene sheets were successively grown along rough Si3N4 surface followed by merging together to form a high quality polycrystalline graphene film. The difficulties that arise during postgrowth transfer method can be eleminated by adopting two‐stage metal‐catalyst‐free‐growth technique, which finds its compatibility with current Si processing technology.The film was not distinguished properly from the underlying Sis Figure h. It wass Figure i. The grs Figure l–n. Figun Figure o illustrn Figure q confirm2.72/H2 mixture was used as the carrier gas. A rotational disorder along with wrinkles was formed onto the graphene films grown on AlN/Si. Here, temperature played an important role to improve the structural quality of the film. Again, high temperature growth might have influenced rough surface, but there was no impact on the structural quality of the graphene film as determined by Raman analysis. Moreover, the temperature for optimum growth might be higher (1350 °C), which suggests the growth of high‐quality graphene on bulk AlN substrates. Furthermore, during processing, AlN etching was minimized by the growth of graphene compared to the annealing , which enabled the growth of AlN at 1250 °C for 6 min without etching effects. It may be mentioned here that thermal treatment of nitrides films finds an ease with this method to enhance the quality of crystal, and favors the activation of doping at the time of ion implantation.Aluminum nitride (AlN) is a promising material due to its versatile applications such as microelectronic and optoelectronic devices, short wavelength emitters, electronic packaging, and acoustic wave resonators.2.8FigureMgO is considered to be advantageous for the growth of nanographene and few‐layer nanographene (nFLG) directly using CVD technique. However, calibration of the reaction time or temperature needs to be undertaken to grow nFLG and nanographene. The growth can be done at temperatures of 325 °C using acetylene as a precursor.2H4 as a precursor via free radical assisted CVD. Zhao et al.Ferrari2.92) leads to better gate modulation, improved gate capacitance, and reduced gate leakage.2 are carrier scattering due to charge fluctuations and surface roughness. These results are obstruction for further integration of FETs based on SiO2‐gated graphene.3 (STO), a transparent, high‐k perovskite dielectric may have great thermal stabilities and potentials.2/CH4:100/50/2.5 sccm).Figure−1, respectively. However, the 2D peak of graphene grown on the STO substrate did not appear due to the substrate screening effect . XPS analysis  sites on the surface  values at infinity due to the insufficient graphene growth  was demonstrated by many researchers. Chen et al.4 as a precursor via near equilibrium CVD technique, whereas Song et al.4 as a precursor via metal catalyst‐free APCVD. Zhang et al.4 as a precursor by PECVD technique. A study pertaining to graphene crystal onto sapphire substrate was accomplished using C2H4 as a precursor using PECVD technique.The graphene is considered to play a pivotal role in optoelectronic devices as a transparent conducting film (TCF). However, the manufacturing process available in the literature utilizes graphene as a TCF which follows transfer procedures. These steps are time‐consuming and sometimes chemical contaminations are encountered during processing which may impose detrimental effects to optoelectronic devices. Therefore, the metal‐catalyst free direct CVD growth of graphene on crystalline Al2.114 as a precursor.4 into different species such as CHx, C2Hy, C3Hz, and atomic hydrogen. These radicals could play a pivotal role during nucleation followed by the growth of nanographene. Apart from the above discussions in Sections The metal‐catalyst free direct growth of graphene on glass and mica substrates can be achieved via r‐PECVD system at a substrate temperature of ≈550 °C using pure CH2.122) has gained tremendous research interest due to its green impact on environment.2 surface results in the contamination free graphene‐TiO2 interface. Liu et al.2 surfaces such as (001) and (100) faces. Both substrates were cleaned using wet chemical etching followed by flattening into atomic smoothness.Figure2/Si substrate, the sharp G‐bands and D‐bands of the graphene became prominent  plane for the same growth parameter. Similarly, Bansal et al.2 substrate using CH4 as a precursor via APCVD/LPCVD technique. Moreover, graphitic nanostructures were grown on titania nanowire aerogel at 750 °C via CVD using ethylene. Three layer thick graphitic nanostructures gave a clear indication that titania nanowire surface enabled the graphitization to few‐layer graphene. The growth mechanism of few‐layer graphene on nanostructured metal oxides paved a way toward facile and controllable processing of metal oxide‐nanocarbon fiber–shell composites.TableRecently, photocatalysis effect based on titania (TiO001) and 0 faces. 001) and 0 faces. 001) and 0 faces. 001) and 0 faces. 3Apart from Ge, there are only few reports available on the metal‐catalyst free direct CVD growth of graphene on important semiconducting substrates such as Si, GaN, and SiC. To the best of our knowledge, there are no reports avalaible on the metal‐catalyst free direct CVD growth of graphene on other important arsenide and phosphide based semiconductors. Therefore, in the successive sections, direct CVD growth mechanism and morphology of graphene on Si, Ge, GaN, and SiC substrates will be discussed. Detailed discussion is being presented for Si and Ge, as they are important semiconductors for the next generation graphene/Si and graphene/Ge based hybrid electronic devices.3.13.1.12 substrates by using photoemission‐assisted plasma enhanced CVD system.4. Particles were closely connected to each other, and shared some graphene sheets between them. The advantage of using this system is that the DC discharge plasma assisted by photoelectrons emitted from the substrate under ultraviolet (UV) light irradiation could be generated close to the substrate with a controllable volume. Similarly, Zhang et al. reported the catalyst‐free direct graphene growth on various substrates including Si and SiC.r‐PECVD system at a relatively low temperature of ≈550 °C by using pure CH4. This graphene growth process was quite unconventional as compared to high temperature (≈900–1000 °C) CVD graphene growth.4), for 3 h under the plasma power of 100 W. Figure4 can dissociate itself into various reactive radicals, such as CHx, C2Hy, C3Hz, and atomic H2. These species play a key role during the nanographene nucleation and growth stages. The carbon‐containing radicals continuously get adsorbed onto the substrate surface, and bonded with each other by diffusion and collision to form graphene nanoclusters with H‐terminations. Nucleation on the as‐grown graphene was much easier as compared to clean substrates. Moreover, due to atomic H2, strong etching effect was there which helped in the suppression of formation of amorphous carbon. Similarly, some other low temperature (500 and 780 °C) direct CVD growth of multilayer graphene films and graphene nanowalls on Si substrates with the help of microwave surface wave plasma CVD and PECVD were reported by Adhikari et al. and Zhou et al., respectively.Takami et al. reported the catalyst‐free direct growth of networked graphite on Si and SiO3.1.22H2 and Ar with flow rates of 25 and 50 sccm, respectively, were introduced. Graphene growth at different temperatures was carried out by maintaining the pressure in the range of 2–100 Torr for 1 h. Figure −1. This was higher than the typical large scale graphene or graphite (1580 cm−1), which confirmed the growth of NG.Kim et al.2 atmosphere with a flow rate of 6 sccm at 300 mTorr and then kept for 20 min to activate the substrate surface, followed by 35 sccm of CH4 for 30 min at desired growth temperatures (870–970 °C). After growth at 920 °C, two kinds of surface features were observed: 1) the rectangular black dots highlighted by the solid red box, and 2) the large white area highlighted by the dashed blue box , but they were in the form of small dots and patches instead of continuous uniform graphene films.ox Figure12a. Fi2 was not necessary to grow graphene on SOI and just replaced 6 sccm of H2 with 50 sccm of Ar at 920 °C. Small black dots were observed and the SOI surface seemed to be less reactive. When the growth was carried out at 970 °C, the size of the black dots increased and stronger graphene bands were observed. Based on the above results, they proposed a growth mechanism for graphene on nonmetal or semiconducting surfaces as surface reaction, adsorption, decomposition, and accumulation  domains, concave BLG domains, and bulging FLG domains on the upside‐down placed single crystalline Si substrates at 900–930 °C for 1 h using the composition of CH4 or ethanol as the carbon feedstock in a conventional thermal APCVD system. It was also established that the concentration of active carbon species in CVD system exerts significant impact on the growth mode of graphene. The carbon precursor flow rate and reaction time during the CVD process controlled the growth dynamics, which ultimately led to a well‐controlled morphology of the obtained carbon material, such as 2D graphene film or VAGNs. The growth morphology was found to be independent of substrate as well as carbon source precursor. This growth strategy suggested that VAGNs can be fabricated on variety of substrates using different carbon precursors in a thermal APCVD system without using plasma, and offers a new insight for intrinsic growth mechanism. System was heated to 1130 °C in 50 min with 50 sccm H2 and 50 sccm Ar. The substrates were annealed for 20 min at 1130 °C, and then CH4 or ethanol (vapors) was introduced into the chamber to initiate graphene growth. The graphene growth morphology could be controlled by varying the precursor concentration and/or reaction time. The VAGNs growth process is schematically depicted in FigureRecently, Wang et al. demonstrated the direct synthesis of uniform VAGNs on Si substrates via catalyst‐free thermal CVD for the first time.4 flow rate between 7 and 14 sccm, whereas planar growth was observed for less than 7 sccm CH4 flow rate. The carbon precursor decomposition under low CH4 flow rate synchronized with the adsorption of carbon atoms on active sites and the growth of graphene. Therefore, the diffusion of carbon species to edge of graphene proceeded under a thermodynamic equilibrium condition, and had enough time to reach the desired positions with minimum energy to form stable crystalline phases.4 flow rates, the effective diffusion of carbon species was limited on the substrate or graphene surface. Therefore, a multiple graphene nucleation and simultaneous enlargement took place, which was induced by the supersaturation of active carbon species concentration, and finally led to the shrinking of available surface on the substrate into narrow channels. In this way, surface diffusion of carbon species reduced drastically and the direct deposition of carbon species started at the graphene edges. Hence, the edge reaction occurred very fast, and led to the vertical growth instead of boundary coalescence when two graphene domains approached each other.Initially, graphene flakes were randomly nucleated on the substrate surface and enlarged to form continuous film, as time progressed. Once a buffer layer formed, initial planar graphene growth eventually altered to upward growth. Consequently, the carbon atoms from precursor thermal cracking were continuously incorporated into the open edges to make the unique vertical growth proceed. VAGNs were obtained with the CH3.1.3Figure2 and SF6 in an ICP reactive ion etching system. Second, the VG was directly grown on the SNCs using CH4, H2, and Ar in the ratio of 1:5:45 at 1000 °C for 3 min. The chamber pressure was maintained at 2.5 kPa, and a bias voltage was applied between the filament and the substrate to facilitate the growth of graphene. Finally, SNC‐ PG electrode was fabricated by patterning the SNC‐G electrode with the help of UVL and RIE techniques. The tilted angle SEM images of the SNC electrode and SNC‐G electrode are shown in Figure Instead of using flat Si surfaces for the metal‐catalyst free direct CVD growth of graphene, researchers have used textured Si substrates to grow graphene.−1 belongs to Si, whereas the characteristic peaks of few‐layered graphene marked as D, G, and 2D were observed for the SNC‐G electrode.2 (2.04%) was due to the surface oxidation. The SNC‐G electrode contained carbon along with Si, and a very small amount of O2 (0.79%). Similarly, Son et al.4, CO2, and H2. CO2 was used as a mild oxidant which helped in achieving robust and uniform growth of MLG around each Si nanoparticle by generating catalytic sites.It was observed that the graphene nanosheets (GNs) were vertically grown on the SNCs and resembled to petaloid clusters. The average height and half‐width of the vertically aligned SNCs were found to be 2 µm and 100 nm, respectively, and the average half width of the SNCs with VG coating was 200 nm. The GNs grew vertically after covering the SNCs surface and formed 3D petaloid structures due to the internal stress and the applied bias voltage.3.22 determines the number of grown graphene layers, as H2 balances the production of reactive hydrocarbon radicals and etching of graphitic carbon during a CVD process. Therefore, the optimized graphene growth was carried out with H2:CH4 = 50:0.1 sccm at 910 °C for 100 min to obtain single‐layered graphene. Figure−1  at 900–930 °C for 5–120 min. Initially, the graphene islands were uniaxially aligned along theCatalytic growth of a single‐crystalline graphene on a solid substrate surface can be achieved by growing a single grain to a size as large as possible from a single nucleation site. Recently, a centimeter‐sized single‐crystalline graphene was obtained from a single nucleus.ID/IG ratio of the single‐crystal graphene (<0.03) was much smaller than that of the polycrystalline graphene (≈0.4), which indicated that the extended grain boundary defects in the single‐crystal graphene on Ge (110) were absent. Hence, it can be inferred that the hydrogen‐terminated Ge (110) surfaces are an ideal substrate for the catalytic growth of single‐crystalline monolayer graphene. On the other hand, Dai et al. also demonstrated the metal‐catalyst free direct CVD growth of graphene on Ge (110) wafers instead of Ge (110) buffer layers on Si (110).However, 2O (18 MΩ cm) at 90 °C for 15 min. The cleaned substrates were loaded into a horizontal tube and the system was evacuated to ≈10−6 Torr. It was filled to atmospheric pressure with a mixture of Ar (99.999%) and H2 (99.999%) at a constant total flow rate of 300 sccm. The Ge (001) samples were annealed for 30 min at 910 °C, and then CH4 (99.99%) was introduced to initiate the growth for 1–18.25 h. The GNRs were found to be self‐aligned 3° from the Ge 〈110〉 directions, and self‐defined with smooth armchair edges. They were having tunable width to <10 nm and aspect ratio to >70. It was also observed that, to obtain highly anisotropic GNRs, the growth rate in the width direction should be very slow (<5 nm h−1).After the successful direct CVD growth of single‐crystalline and polycrystalline graphene monolayer, the direct CVD growth of graphene nanoribbons (GNRs) on Ge was reported by Jacobberger et al. for the first time.It was inferred that the energy barrier associated with the graphene nucleus rotation significantly increased, as it became larger, which resulted in fixing the orientation of the NR lattice during the subsequent growth.4, H2, and Ar at 3.6–4.6, 100, and 200 sccm, respectively. After ultra high vacuum (UHV) annealing to 700 °C, both the Ge (110) and the Ge (111) surfaces restructured into domains, and exhibited in‐plane ordering underneath. The Ge(111)/graphene interface was strongly affected by UHV annealing. STS analyses revealed significant differences in electronic interactions between graphene and Ge (110)/Ge (111). Raman spectra indicated that the graphene was considerably strained after the growth with more point‐to‐point variation especially on Ge (111). Finally, the extreme strained case was observed for graphene/Ge (001), which resulted in the reorganization of the Ge surface intoKiraly et al. studied the metal‐catalyst free direct CVD growth of graphene on differently oriented Ge substrates in order to understand the nature of the graphene–Ge interfaces.4 gas at 900–930 °C2/CH4 flow ratio, and the growth time in order to understand the growth mechanism and morphology.2 and Ar, respectively, whereas the CH4 flow rate varied between 1 and 10 sccm. Once the growth temperature reached to 930 °C, CH4 was introduced into the chamber, and the total pressure was set to 100 mbar. It was possible to tune the growth in order to obtain different graphene structures, such as GNRs, monolayered and multilayered graphenes by simply varying the CH4 flow and growth time. At a flow rate of CH4 (F) = 1 sccm, GNRs oriented along the 〈110〉 directions on the surface  concentration led to strong supersaturation. The formation of carbon growth species was much faster as compared to their total consumption, which allowed the further nucleation of graphene domains.4, poorer quality graphene films with absence of the Ge nanofaceting were obtained. The time dependent growth at F = 2 sccm and H2/CH4 = 100 for 60 min was carried out several times. The growth proceeded in a layer‐by‐layer regime, as the carbon adatom species (CHx) concentration before nucleation was just above the critical supersaturation level. The adsorbed carbon species depleted as the nucleation and growth of graphene grains started. Eventually, their concentration was reduced to a level where the nucleation rate could be negligible and only monolayered domain enlargement could take place. Moreover, the initial supersaturation condition was restored and the second graphene layer formation started to take place. These results can be attributed to the similarities between the C–Cu and C–Ge alloy systems.Despite the fact that the direct CVD growth of graphene on Ge (100) results in highly strained graphene overlayers, Pasternak et al. demonstrated the direct LPCVD growth of large area high‐quality graphene films on Ge (100) overlayers supported on Si (100) substrates by using CHce Figure17a sim4 and Ar/H2 mixtures, while the total system pressure was maintained at 700 mbar.In continuation to the metal‐catalyst free direct CVD growth of graphene on Ge substrates, Lukosius et al. and Lupina et al. reported a complementary metal oxide semiconductor (CMOS) technology compatible good quality graphene on 200 mm Ge(001)/Si(001) wafers, respectively.ID/IG ratio (≈0.1), which was the indication of a low concentration of defects. Thus, a good quality, uniform, and large area graphene was obatined. Graphene‐induced Ge faceting was also observed in this case  surfaces without grain boundary defects.2/Si or HfO2/Si wafers to fabricate FETs.4 and rest of the process was same as described in ref. 2 (15 nm)/Si or HfO2 (15 nm)/Si substrates to fabricate FETs.Dabrowski et al.FigureLch) of 25–120 nm were defined via e‐beam lithography. Cr/Pd/Au (0.7/10/8.3 nm) contacts were thermally evaporated and patterned at random locations across the substrate so that the ribbons could be perpendicular to the source and drain electrodes. A schematic diagram of the FET architecture is shown in Figure Lch of ≈25 nm. The direct CVD growth of aligned, narrow, and semiconducting GNRs on Ge (001) could overcome many challenges, like control over the polydispersity in ribbon width, length, and location. Thus, it may enable significant advances in state‐of‐the‐art semiconductor electronics.The GNRs were peeled off with the help of a sacrificial multilayer stack of Au/poly(methyl methacrylate) (PMMA)/thermal release tape, and stamped onto the desired substrate Figure18a–e. 3.32H2 and 1000 sccm NH3 for 5 min, while the total pressure was maintained at 750 mbar. As GaN can dissociate at high temperatures, an overpressure of NH3 was maintained to protect the GaN surface. NH3 plays a dual role, as it compensates the loss of nitrogen from GaN during growth as well as releases H2, which is helpful in graphene‐CVD process. The thickness of the grown carbon thin films was in the range of 2–4 nm and they were continuous, uniform, and scalable.Silicon carbide (SiC) and gallium nitride (GaN) are two important wide bandgap semiconductors which can be integrated with graphene for electronic and optoelectronic device applications. There are only few reports available on the metal‐catalyst free direct CVD growth of graphene on SiC and GaN substrates. Unfortunately, no reports are available on direct CVD growth of graphene on other wide bandgap semiconductors. Direct CVD grown graphene‐based coatings as transparent electrodes can be utilized in optoelectronics, especially for fabricating GaN‐based light emitting diodes. So far, this has been achieved by transferring CVD grown graphene from metal substrates to GaN surface.4 (2 sccm) and H2 (20 sccm) with the total pressure maintained at 10 mTorr. The power for discharging the gas mixture was 50 W, and CH4 was effectively dissociated into various species, such as CHx, C2Hy, H, and H2. The dissociation rate was about 34%, which was much higher as compared to ≈0.0002% in thermal CVD,Ea ≈ 2.0–2.6 eV) in thermal CVD.2‐hybridized carbon atoms in as‐grown graphene , G (≈1580 cm−1) and 2D (≈2680 cm−1) confirmed the formation of graphene film with structural disorder. Figure Since PECVD is a low temperature process, therefore direct growth of uniform graphene films on GaN via PECVD seems to be an ideal option to fabricate optoelectronic devices in order to prevent thermal degradation of GaN and other active layers. Kim et al. demonstrated the direct integration of polycrystalline graphene into GaN‐based light emitting diodes (LEDs) via PECVD.Hence, GaN‐based blue LEDs were fabricated with directly grown graphene through a batch process for centimeter‐scale LED substrates, and their schematic diagram and a scanning TEM image with a graphene electrode are shown in Figure 4 were demonstrated by Zhang et al. and Ghosh et al., respectively.2 (4–32 sccm), and CH4 (0.5–8 sccm) for 1–6 min. Before each sample growth, substrates were heated to 1600 °C for 3 min under a 100 mbar pressure to sublimate the Si from top of the SiC surface, which exposed the inner carbon and provided the spontaneous nucleation site for external carbon species. All the important results on low and high temperature metal‐catalyst free direct CVD growth of graphene on Si, Ge, GaN, and SiC are summarized in TableSimilarly, low‐temperature direct PECVD growth of island type nanographene and vertical graphene nanosheets on SiC substrates using CH42, ZrO2, HfO2, h‐BN, Al2O3, Si3N4, quartz, MgO, SrTiO3, TiO2, etc., and semiconducting substrates, such as Si, Ge, GaN, and SiC. It was observed that direct CVD growth of graphene on dielectric substrates is difficult to achieve because of their low surface energy. Also, it is difficult to obtain good quality uniform graphene films on dielectric substrates via high‐temperature CVD methods. However, a low‐temperature PECVD technique could solve this problem. Apart from Ge, there are limited reports available on the metal‐catalyst free direct CVD growth of graphene on other important semiconducting substrates including wide bandgap semiconductors, such as Si, GaN, and SiC. Extremely low carbon diffusivity on Si surface and relatively high carbon solubility at high temperatures could hamper direct CVD growth of high‐quality monolayer graphene on Si substrates. Less low‐temperature direct CVD growth of graphene films on Si using different PECVD systems was attempted. However, the quality of grown graphene was not good, and large area uniform films could not be obtained. These PECVD methods mostly yielded island type growth and vertical graphene nanosheets (VGNs). The temperature of Si substrate should be lower than 1000 °C for a high‐temperature direct CVD growth of graphene. The high temperature (>800 °C) direct growth on Si substrates was attempted by using thermal APCVD, LPCVD, and HFCVD, respectively. These methods also yielded different morphologies of graphene, such as triangular nanographene, SLG, and FLG domains as well as VGNs. VGNs were grown on flat and textured Si substrates, respectively, but high quality and large‐area uniform graphene films could not be obtained. On the other hand, Ge substrate is far better than Si substrate due to its high catalytic activity and surface diffusivity, and a very low carbon solubility at its melting point (<108 atoms per cm3). Hence, direct CVD growth of graphene on large Ge wafers was carried out by using thermal CVD. Large area, high quality, and uniform graphene films were successfully produced. Single‐crystalline Ge substrates are also available for direct CVD growth of single‐crystalline monolayer graphene to fabricate graphene/semiconductor heterostructures based electronic devices. Furthermore, epitaxially grown large‐area single‐crystalline Ge layers on Si wafers are also available for direct CVD growth of graphene. Moreover, direct growth of uniform graphene films on GaN via PECVD seems to be an ideal option to fabricate optoelectronic devices in order to prevent thermal degradation of GaN and other active layers. As GaN can dissociate at high temperatures, therefore an overpressure of NH3 must be maintained to protect the GaN surface. NH3 plays a dual role and compensates the loss of nitrogen from GaN during growth as well as releases H2, which could be helpful in high‐temperature graphene‐CVD process. Direct CVD growth of graphene on Si and wide bandgap semiconducting substrates further needs to be explored to directly integrate graphene on these substrates for graphene/semiconductor based hybrid electronic and optoelectronic device applications. Attempts could be made toward growing high quality and large area graphene films on these substrates  by exploring new carbon precursors, and by designing novel CVD units. These growth methods exhibit great potentials pertaining to technology and scientific aspects, as some of the facile growth techniques have already been adopted. A detailed understanding of the metal‐catalyst free direct CVD growth mechanism of graphene on dielectric and semiconducting substrates is pivotal to achieve a controlled graphene growth on them, which is crucial for graphene based electronic and optoelectronic device fabrication.Recent progress made toward the metal‐catalyst free direct CVD growth of graphene on various dielectric and semiconducting substrates is summarized in this review. The focus was on the metal‐catalyst free direct CVD growth of graphene on technologically important dielectric substrates, such as SiOThe authors declare no conflict of interest.


The composition and structure were studied using a range of material characterization techniques. No covalent bonds were observed between molecules, and the self-assembled driving force was the only intermolecular interaction: Van der Waals’ force in the intra-layer and π–π interactions between layers. The arranged-rearranged structures provided a range of lithium ion shuttle channels, including the space between layers and diffusing through the nanosheets, which decrease the diffusion distance of lithium ions remarkably and reduce the irreversible capacity of the battery.This study examined the characteristics of small molecular structure nano-graphene in a dynamic hierarchical self-assembly and found that graphene is rearranged under its own pressure during dynamic aggregation and water ripples are formed by the The online version of this article (10.1186/s11671-019-2903-4) contains supplementary material, which is available to authorized users. The development of green alternative energy sources has received considerable interest. Recently, nano-graphene and graphene composites attracted interest for used as lithium ion anodes –3. In ad2 carbon atoms. Each layer has a hexagonal honeycomb structure called a nano-graphene sheet  is a representative example of nano-graphene that has been well studied –30. The eet Fig.  31]. Wh. Wh2 card-spacing generated via the dynamic self-assembly at the molecular level and the relationship between the clusters of nano-graphene, an in-depth analysis of the formation factors inside of nano-graphene was analyzed further.This paper introduces the dynamic hierarchical self-assembly structure–function relationship of hexabenzocoronene. By observing the n-butylammonium perchlorate (TBAP). Deionized water is used for all experiments.Hexabenzocoronene was synthesized according to a previously reported procedure –35. All The morphology and lattice fringe were observed using a scanning electron microscope , transmission electron microscope , and high-resolution transmission electron microscope .n-butylammonium perchlorate and 0.1 mM ferrocene electrolyte solution were scanned at a scan rate of 0.1 mV s−1.Electrochemical measurements were performed on the Shanghai Chenhua CHI660e system. A three-electrode system is used, a platinum wire for the counter electrode, a platinum plate with a fixed of electrode, and a saturated calomel electrode for the reference electrode. The concentration of the supporting electrolyte TBAP was 0.1 mol/L, and the analytical pure solvent was acetonitrile (ACN). Firstly, polish the platinum carbon compound electrode vertically on the circular gauze on the glass brick ; secondly, rinse off the white aluminum with distiller water and then use ultrasonic for 1 min by acetone; and finally, use ear ball washed and blown dry. Then, the suspension of the hexabenzocoronene sample was dropped on the surface of the glassy carbon compound electrode, and the solvent was naturally evaporated to dryness. Then 0.1 M tetra-Hexabenzocoronene is carbon–carbon material combined with significant π–π conjugate chemical bonding. A procedure for the preparation for hexabenzocoronene consisted of a series of reactions, such as Sonogashira, Diels-Alder reaction, Lewis catalyst-based cycle reaction, and deprotonation under basic conditions to give intermediates in unsatisfactory yields –38. The HBC has been widely used, but in the study of self-assembled system, it needs to be further understood. Although studies of the same or similar anode materials have been mentioned in the reported literature, the HBC study is still insufficient. Therefore, the focus of the work is on the detailed research of the self-assembly system, and put it on one by one to understand the internal dynamic distribution of aggregation and the induction and to improve the supplement of the lack of content anode materials.The small-molecule nano-graphene self-assembled dynamically to form regular thin sheets, which were sequentially and systematically stacked to form intermittent sheet nano-graphene fragments that were held tightly to each other . On the To explain the abovementioned rearrangement/change caused by its own weight and whether it will affect the material properties, scanning electron microscopy (SEM) was performed to determine if the particle size had changed. As shown in Fig. d-spacing between the layers highlight the performance of LIB anode materials.Transmission electron microscopy (TEM) showed that the hexabenzocoronene molecule exhibits structural features with a coherent layer spacing and a molecular layer spacing from 0.34 nm. High-resolution TEM (HRTEM) indicated that nanoparticles bind to each other  was further undertaken to understand the lithium storage behavior. The CV curves hexabenzocoronene were measured at the same scan rates (0.1 mV s−1) and display redox peaks with slight shifts with increasing scan rates, thereby showing a rectangular shape with increasing scan rates, as shown in Fig. V) can be separated into oxidation increases (V1), standard oxidation effects (V2), and standard reduction effects (V3) (Eq. (Cycle voltage (CV) was performed at the high potential of the lithium-ion batteries to determine the long-term stability and potential energy Fig. a. Accord rates 0. mV s−1 a+ incorporated into the hexabenzocoronene. The charging process (Li+ transfer) in hexabenzocoronene anodes requires stabilization. The calculated stabilization HOMO energy of the hexabenzocoronene radical anode ranges from 5.592 V, as shown in Fig. The anion/radical anion with an electron donating functional group leads to a homogeneous/uniform electron distribution throughout the flake, which is beneficial for maximizing the number of Lid-spacing between the layers for hexabenzocoronene was examined. This paper revealed a multi-diffusing process of lithium ions as the dynamic structure providing dynamic diffusion paths. TEM showed that lithium diffuses between the layers and has the ability to pass through the sheets, which greatly increases the lithium ion (yellow spot) diffusion efficiency; the Additional file Va/cm3 (STP) g−1 value is 110.47 and 96.62. According to adsorption-desorption isotherm, there is no hysteresis loop in the isotherms of HBC. Moreover, Additional file Vm is 18.647 cm3 (STP) g−1, and as,BET is 81.16 m2 g−1. The TEM image revealed self-assembled structures that were disorganized in the center of the fingerprint, and then they were arranged more regularly into a fingerprint-like structure. In the process of the self-assembly of graphene sheets, graphene sheets arrange in a stacked manner and self-assemble into a layered two-dimensional structure in a head-to-head manner. Moreover, the bonding force between molecules is weak with no strong chemical bonds. The self-assembled structure is a dynamic process involving the angular-rearrangement of self-assembled layers of graphene nanosheets under the action of energy. Moreover, TEM image showed that lithium ions have different diffusion modes between the graphene sheets, which can diffuse between the layers and pass through the layers, from the inner layer to the outer layer diffusion. Therefore, nano-graphene exhibits strong lithium ion diffusion properties and surprising lithium ion storage capacity.The inset in Fig. d-spacing in the self-assemblies led to a significantly enhanced LIB anode charge capacity and cycling stability. These results revealed a structure–property correlation between the nature of functional groups and Li storage capacity. Nevertheless, identifying the mechanism for how nano-graphene hierarchically assembles and dominates the overall battery performance will be an important research topic. Through these studies, a more rational and effective application of nano-graphene will be realized. Observing the characteristics of the internal architecture from a microscopic perspective and analyzing the dynamic hierarchical self-assembly properties of a nano-graphene sheet one by one will be the subjects of a future study.HBC shows good structure durability and stability. The electron density with the optimal Additional file 1:Supporting information (DOCX 436 kb)


N = 1,098 participants  without a diagnosed autism spectrum disorder, most of whom were university students. The IRI scale “Perspective Taking” and the EQ were negatively related to the AQ in male and female participants, while the IRI scale “Empathic Concern” was negatively related to the AQ in females only. Moreover, the AQ was positively related to the SQ-R in females only. Lastly, the SQ-R and a number of the empathy scales were significantly associated: For example and surprisingly, the EQ correlated weakly and positively with the SQ-R in both male and female participants. The results from this study illustrate how standard self-report measures of empathy, autistic tendencies, and systemizing tendencies are associated with each other in a large sample not diagnosed with an autism spectrum disorder. Additionally, some potential gender-specific effects are revealed.The aim of the present study was to investigate associations between four highly used self-report measures assessing empathy , autistic tendencies, and systemizing tendencies. Participants in this study completed the following self-report measures: The Interpersonal Reactivity Index (IRI) and the Empathy Quotient (EQ) to measure empathy, and the Autism Spectrum Quotient (AQ) and the Systemizing Quotient–Revised (SQ-R) to assess autistic and systemizing tendencies, respectively. The final sample consisted of  Empathy can be understood as an important concept contributing greatly toward successful human social interaction –3. HowevIn line with the various definitions of empathy, a range of self-report measures assessing individual differences in empathy exist . ThesInternational Classification of Diseases 11th Revision (ICD-11) and the Diagnostic and Statistical Manual of Mental Disorders, Fifth Edition (DSM-5). It is—among other things—characterized by deficits in social functioning and communication . ASD is/will be classified (including diagnostic criteria) and divided into various subtypes in the nication , 19. Accnication –24. Howenication , 25, 26.nication , 24, 27.nication .Aside from reporting lower levels of empathy, samples of participants diagnosed with an ASD also show higher systemizing tendencies compared to control samples , 21, 24.Although higher autistic tendencies are robustly associated with both lower empathy and higher systemizing tendencies, empathy and systemizing tendencies are not robustly correlated with each other. Correlations between scores in different EQ and SQ(-R) versions have been reported to lie between −0.06 and −0.21 in general population samples, and at around −0.29 in a sample of people diagnosed with an ASD , 30, 31.Taken together, the aim of the present study was to investigate the relationships between empathy, autistic tendencies, and systemizing tendencies. Previous studies investigating these relationships typically used versions of the EQ, AQ, and SQ(-R). The findings previously shown with these scales should be replicated. In addition to this, the present study also sought to examine these relationships using the IRI, which, contrary to the EQ , assesses empathy as a multidimensional construct. In line with results from previous studies, we expected the following associations:Positive correlations between the EQ and the IRI scales PeT and EmCNegative correlations between the AQ and both the EQ and the IRI scales PeT and EmCPositive correlations between the AQ and the SQ-RPositive correlations between the AQ and the IRI PeD scaleNo significant correlations between both the EQ and the IRI scales with the SQ-Rhttps://osf.io/q2arp/. Data will be made available upon reasonable request.This study is registered at https://www.ckannen.com/). The data collection took place at Ulm University, Ulm, Germany. Therefore, most participants tended to be younger and were students. More specifically, 1,249 participants took part online in the present study, which is part of the Ulm Gene Brain Behavior Project (UGBBP). However, 11 participants were excluded due to missing data. Hence, 1,238 participants  were retained for subsequent analyses. One participant reported suffering from Asperger syndrome. This participant was excluded from further analyses, as he/she was identified as an outlier in terms of the AQ score.Participants took part in an online study including various self-report questionnaire measures presented using the SurveyCoder-Tool (th-Quantil − (1.5x(75th-Quantil − 25th-Quantil))] or higher than [75th-Quantil + (1.5x(75th-Quantil − 25th-Quantil))] on at least one of the scales under investigation, or for age, were treated as outliers and excluded from further analyses , 140 participants were excluded due to their categorization as an outlier on at least one of the scales under investigation, or because of their reported age. All participants who scored lower than Consistent with most of the literature, the EQ correlated most strongly with the IRI scales PeT and EmC for both males and females , 16, 17.The AQ was significantly and negatively related to the EQ and the IRI scale PeT for both males and females with medium effect sizes . Additior = 0.18) could only be considered small to medium . Some of the IRI scales were also significantly related to the SQ-R depending on the sample . This is in contrast to previous results and assumptions indicating that empathy and systemizing tendencies are (weakly) negatively related and therefore potentially distinct and independent phenomena  and the Affective Neuroscience Personality Scales (ANPS), including its FEAR scale (assessing facets of anxiety using a trait approach) –45. The Next, the sole use of self-report questionnaires as measures for the constructs could be criticized. Clearly, this methodology has several disadvantages . The speIn conclusion, the results from this study provide insights into the relationship between empathy, autistic tendencies, and systemizing tendencies in a large, mostly student, sample not diagnosed with an ASD. The results generally support the notion that autistic tendencies are negatively related to empathy, but positively related to systemizing tendencies, although in males the association between autistic and systemizing tendencies was weak and nonsignificant. Therefore, the results suggest that future studies should report findings with these measures separately for males and females, and also support the notion that future studies should focus on facets of autistic tendencies and empathy to examine these relations in more fine-grained detail.All subjects gave electronic informed consent in accordance with the Declaration of Helsinki. The protocol was approved by the local ethics committee of Ulm University, Ulm, Germany.CS and CM planned and implemented the study and collected the data. CS implemented the statistical analyses and wrote the manuscript. CM worked over the manuscript to improve it. AC gave helpful advice and also worked over the manuscript to improve it.The position of CM is funded by a Heisenberg grant awarded to him by the German Research Foundation .The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.


The signaling of brain-derived neurotrophic factor (BDNF) has been suggested to be impaired in Alzheimer’s disease (AD), which may compromise the function of BDNF upon neuronal activity and survival. Accordingly, decreased levels of BDNF and its tropomyosin-receptor kinase B-full-length (TrkB-FL) have been detected in human brain samples of AD patients. We have previously found that neuronal exposure to amyloid-β (Aβ) peptide, a hallmark of AD, leads to calpain overactivation and subsequent TrkB-FL cleavage leading to decreased levels of TrkB-FL and the generation of two new fragments: a membrane-bound truncated receptor (TrkB-T′) and an intracellular fragment (TrkB-ICD). Importantly, we identified this TrkB-FL cleavage and TrkB-ICD presence in human brain samples, which indicates that this molecular mechanism contributes to the loss of BDNF signaling in humans. The exact role of this TrkB-ICD fragment is, however, unknown. Here, we used a human neuroglioma cell line and rat cortical primary neuronal cultures to track TrkB-ICD intracellularly. Our data show that TrkB-ICD is a relatively stable fragment that accumulates in the nucleus over time, through a phosphorylation-dependent process. We also found that TrkB-ICD has tyrosine kinase activity, inducing the phosphorylation of nuclear and axonal proteins. These findings suggest that TrkB-ICD may lead to a dysregulation of the activity of several proteins, including proteins in the nucleus, to where TrkB-ICD migrates. Since TrkB-ICD is formed by Aβ peptide-induced cleavage of TrkB-FL, the present data highlights a new mechanism that may have a role in AD pathophysiology. Alzheimer’s disease (AD) is a slow progressing neurodegenerative disease, leading to atrophy and neuronal loss of particular brain regions, in particular the hippocampus, therefore leading to cognitive impairments . After 16 h of transfection with TrkB-ICD vector, H4 cells were treated with cycloheximide , an inhibitor of protein biosynthesis, for 8 h and 24 h. TrkB-ICD levels were quantified at 0 h, 8 h and 24 h after CHX treatment; a time-dependent gradual decrease of TrkB-ICD expression levels was detected (p < 0.0001) towards near 50% of the value at time 0 , which codifies the nuclear import of proteins  revealed the presence of two bipartite NLS. In this algorithm, scores of 9/10 codify proteins exclusively present in nucleus, whereas scores of 1/2 indicate proteins only present in cytoplasm. In our case, both predicted sequences are characteristic of proteins that could be localized either in nucleus or in cytoplasm  that specifically detects phosphotyrosine-containing proteins and evaluated, through western-blotting and immunofluorescence assays, phosphotyrosine immunoreactivity of 24 h-transfected primary neurons. As shown in Considering that TrkB-ICD contains a TrkB-FL tyrosine kinase domain and also that phosphorylation is a central step for many biologic processes, we evaluated whether this fragment could We then evaluated if TrkB-ICD nuclear translocation was dependent on its kinase activity. To do so, immunofluorescence assays were performed using 24 h-transfected primary neurons incubated with K252a (200 nM), which inhibits Trk-FL kinase activity . These data indicate that PKA and PKC belong to the signaling cascade operated by TrkB-ICD that lead to protein phosphorylation  for 2 h, 8 h and 24 h, the overall levels of protein phosphorylation markedly decreased, being already nearly absent after 2 h of incubation with the tyrosine kinase inhibitor  is a relatively stable protein; (II) phosphorylates several proteins; and (III) accumulates in the nucleus .Previously, we described that TrkB-FL could be cleaved by Aβ peptide accumulation , having therefore a protective role. Although we do not exclude that hypothesis, it is unlikely that TrkB-ICD could have this putative protective role, since TrkB-ICD only has the anchorage site for PLCγ and we do not detect (results not shown) any activation of the three pathways. Furthermore, these mediators of TrkB-FL signaling are in close vicinity to the cytoplasmic membrane and our data suggest that TrkB-ICD is mainly located on nucleus.Calpain-mediated cleavage of several proteins, namely of p35, β-catenin or mGluR1α proteins, leads to the formation of intracellular fragments known to be involved in neuronal death, hyperphosphorylation of Tau protein, excitotoxicity, gene transcription modifications  and by the Ethical Comission of Centro Hospitalar de Lisboa Norte and Centro Académico de Medicina de Lisboa. In addition, this study was carried out in accordance with the recommendations of “Directive 2010/63/EU”.orcid.org/0000-0001-7915-3517) performed the experiments, analyzed the data and wrote the manuscript. AJ-S (orcid.org/0000-0002-9599-6597) performed the experiments and reviewed the manuscript. AL (orcid.org/0000-0003-4163-0044) and PC (orcid.org/0000-0002-1090-4631) collaborated in the experimental work. EC (orcid.org/0000-0002-1402-2791) and AS (orcid.org/0000-0001-9030-6115) discussed and reviewed the manuscript. MD (orcid.org/0000-0001-5486-6246) designed the experimental work, discussed the data and reviewed the manuscript.JF-G (The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.


TM, was purified obtaining PVAm films with two degrees of purification: Low Grade (PVAm-LG) and High Grade (PVAm-HG). These two-grade purified PVAm were loaded with 3 wt% of graphene and graphene oxide to improve mechanical stability: indeed, pristine tested materials proved to be brittle when dry, while highly susceptible to swelling in humid conditions. Purification performances were assessed through FTIR-ATR spectroscopy, DSC and TGA analysis, which were carried out to characterize the pristine polymer and its nanocomposites. In addition, the membranes′ fracture surfaces were observed through SEM analysis to evaluate the degree of dispersion. Water sorption and gas permeation tests were performed at 35 °C at different relative humidity (RH), ranging from 50% to 95%. Overall, composite membranes showed improved mechanical stability at high humidity, and higher glass transition temperature (Tg) with respect to neat PVAm. Ideal CO2/N2 selectivity up to 80 was measured, paired with a CO2 permeability of 70 Barrer. The membranes’ increased mechanical stability against swelling, even at high RH, without the need of any crosslinking, represents an interesting result in view of possible further development of new types of facilitated transport composite membranes.In the present study, the separation performance of new self-standing polyvinylamine (PVAm) membranes loaded with few-layer graphene (G) and graphene oxide (GO) was evaluated, in view of their use in carbon capture applications. PVAm, provided by BASF as commercial product named Lupamin T. T−1 ref−1, which is quite reasonably associated to the neutralization of amine group from the R–NH3+ to the R–NH2 form, as also confirmed by the results obtained by Annenkov et al. [Concerning the polymer peaks, then, an interesting change in the different spectra is the shift of the peak related to N–H stretching before and after the resin treatment (PVAm-HG). This one moves from 3100 to 3300 cmv et al. . This neThermal and thermo-oxidative stability of Lupamin 9095 and purified PVAm was assessed through TGA analysis . The oveUntil 180 °C, TGA curves display weight loss due to absorbed water. The polymer degradation occurs in two main steps with onset temperature at 190–220 and 305–320 °C with only very minor differences among different materials, as visible in Sample weight is stable in the range 500–600 °C, indicating the end of the thermal degradation. Organic residues were oxidized by switching in air atmosphere at 600 °C and keeping the sample at that temperature for 30 min. After this treatment, purified films showed a significantly lower residue with respect to untreated Lupamin 9095. In particular, PVAm-LG and PVAm-HG showed residues in the order of 15% and 8% of the initial weight, respectively, confirming the better purification obtained when the ion exchange step was used after ethanol washing.No major modifications of the TGA curves of the polymer were observed upon graphene or graphene oxide addition; these nanofiller, therefore, did not seem to affect the thermo-oxydative stability of the matrix.g) of the films, as shown by DSC analyses reported in Interestingly, the degree of purification and nanofiller loading affected the glass transition temperature , which seem to act as a nanometric filler able to efficiently rigidify and reinforce the matrix through polar interactions with the polymeric chains. This hypothesis can explain the higher Tg in High Grade films (+20 °C), as in this case no rigid inclusions are present in the highly purified polymer, thus the nanofiller can properly act as reinforcement, hindering macromolecular motion.On the other hand, 3% in weight addition of both G and GO to PVAm-HG determined a sharp increase of TThrough SEM analysis, the fractured surfaces of produced films were examined to investigate the dispersion and the interaction of G and GO with the polymeric matrix.As shown in Differences in smoothness between the fracture surfaces may be considered as further proof of the achieved purification: the most purified materials presented fewer salt impurities on the surface.In the chart, it can be seen how Lupamin 9095, the commercial form, can absorb the largest amount of water, reaching 90 wt% of water mass gained at 60% water activity, as also shown in a previous work . It is fThe addition of a graphene-based phase presents different effects on the water sorption of the different materials. Indeed, the reduction of water uptake upon filler loading is negligible for PVAm-LG composites, while it results significantly higher for PVAm-HG based materials. In the latter case, even at relatively small amounts, the graphene-based nanofillers significantly decreases the water uptake of PVAm-HG based composites with respect to that of the non-loaded corresponding polymer. For example, pure PVAm-HG presents an uptake of 21 wt% at around 60% activity, while the addition of 3 wt% graphene and 3 wt% graphene oxide lower that value to, respectively, 15.1 and 7.8 wt%. This decrement can be interpreted as a positive interaction between the matrix and the filler, which is capable of stiffening the overall material structure, preventing an excessive swelling at high relative humidity. This result is more evident in the case of GO rather than G, likely due to the larger aspect ratio of the sheets and to the possible presence of electrostatic interaction among the carboxylic groups of GO and the primary amine groups of PVAm.Concerning the PVAm-LG + 3% GO behavior, it can be noted instead that it is in accordance with that discussed above regarding the materials glass transition temperature variation: the limited effect of the filler seems indeed to be related to the presence of finer dispersion of sodium formate crystals, which somewhat interact with the GO decreasing its ability to impact the properties of the polymeric matrix, and also to the cationic nature of low-grade samples, which impairs the possible interactions among the amine groups of the polymer and the GO’s carboxylic acid groups.2/N2 selectivity are also reported, when available. As quite common for hydrophilic membranes, permeability increased exponentially with the degree of humidification [The single gas permeation results for carbon dioxide for the different membranes tested are presented in fication  due to t2; in the case of nitrogen tests, no reliable data could be obtained, likely due to the longer experimental time required, which resulted in an excessive stress on the film. On the other hand, all films prepared using a graphene-based filler resulted mechanically stable and allowed extensive tests to be performed without any loss in permselective properties for several days. Apart from the general behavior and considering the different materials, as a first observation, it can be noted that for PVAm-HG a single point at intermediate humidity was obtained. This is due to the fact that PVAm in its pure form appeared to be quite unstable as a self-standing film, especially when purified. This led to the impossibility of running tests at higher humidity, due to the rupture of the film, verified on several specimens. For the case of PVAm-LG, the presence of residual sodium formate appeared to slightly increase the mechanical stability, allowing a full permeation curve to be acquired, at least for CO2 varies from 16.5 Barrer at 63% RH, up to 73.8 Barrer at 93% RH. For the same polymer, filled with 3 wt% of graphene oxide, values range from 1.7 Barrer at 53% RH to 71.0 Barrer at high humidity. PVAm-HG, loaded with the same quantity of graphene oxide, presented instead a permeability for carbon dioxide of 1.6 and 25.1 Barrer at 75% and 93% RH, respectively. When loaded with few-layer graphene, the same matrix presented values from 2.0 Barrer at 77% RH to 23.1 Barrer at 92% RH. From the experimental data, it can be seen how PVAm-HG presents a permeability of 4.2 Barrer at 56% RH, but no particular trend can be obviously inferred. For PVAm-LG, permeability of COThe non-loaded materials therefore present a higher permeability at the same relative humidity with respect to their counterparts containing a graphene-based nanophase. This is a somewhat expected result, since the platelets of graphene, due to their high aspect ratio, can easily increase the tortuosity of the diffusion pathway of the gases in the matrix. This effect appeared to be less pronounced in the case of PVAm-LG, possibly due to the polar interaction between graphene oxide and salt nanocrystal in these materials as already discussed by considering DSC and water sorption experimental results.2 permeability, as also shown by previous works [Regarding PVAm-HG, it is quite interesting to notice how no particular differences in gas permeability can be discerned between the materials obtained using two different fillers. This could be contradictory considering that GO usually presents a higher aspect ratio than few-layers graphene and that PVAm-HG + GO showed a lower water uptake than PVAm-HG + G. Usually, indeed, when no other factors are in play, a lower concentration of humidity in the matrix resulted in a lower COus works ,37,39,402 solubilization and facilitated transport execution, as previously shown in Equations (1)–(5) and As previously mentioned, nitrogen permeation tests were only possible on the nanocomposite membranes, due to their enhanced stability. For this reason, the ideal selectivity could be calculated only for these materials, using an exponential interpolation to report data from different gases at the same RH values. In the case of PVAm-LG+GO in particular, selectivity was shown to vary from 3.1 to 59.2 when humidity was raised from 60% to 95%. For PVAm-HG, instead, this value ranges from 3.0 at 75% RH to 80.7 at 95% RH for the graphene oxide sample and from 1.1 at 82% RH to 45.2 at 92% RH for the graphene one.2 permeability (≈70 Barrer), with lower CO2/N2 selectivity (≈60) compared to the high-grade one, which has higher selectivity (≈80) and lower permeability (≈35 Barrer). These differences are probably due to the swelling difference of the two composite materials, as already discussed, and shows a PVAm-LG + 3% GO more hydrophilic than the PVAm-HG-based composite. These results fit also with the chemical structure differences between the two materials: PVAm-HG has an almost totally neutral charge upon the polymeric chain, leading to a tighter packing of the chain itself, which cause the overall lower permeability, while the PVAm-LG material has much higher ionic repulsion [2/N2 selectivity.Analyzing PVAm-LG and -HG reinforced with graphene oxide, therefore, it resulted that the low-grade polymer has higher maximum COepulsion . Neverth2/N2 Robeson’s upper bound [Despite these differences, the two materials show a very similar behavior in terms of overall permselective performances as they lie very close, but are not able to overcome the COer bound , even atFrom this point of view, therefore, contrary to what expected, the further purification step seems not to give substantial advantages in terms of membrane performance. 2 permeabilities in the order of 50 Barrer and CO2/N2 selectivity lower than 10 [, as also reported in Among the many papers found in the literature about PVAm based system, it is very difficult to find references for a consistent comparison of the present results in order to give a more general picture of the overall potential of such membranes. Available data on FTM are usually measured at higher temperatures and/or different upstream pressures, and often refer to very complex systems also including mobile carriers. Among others, Zhao and coworkers, studying PVAm-based mixed-matrix membranes, tested PVAm thin films at 22 °C and 100% RH and reported CO than 10 . On the  than 10 , at 20%  than 10 , as also, however, is accompanied by a remarkable improvement of mechanical strength of the material: pristine PVAm sample indeed cannot be tested as pure polymer self-standing membrane and also literature data [This reduction of the separation performance with respect to the supported pure polymer from Kim et al. , howeverure data  were obt2 capture and purification. Graphene-based PVAm composite membranes were fabricated and tested to assess their potential in the field of COThe materials were produced by mixing polyvinylamine (PVAm) at different degrees of purification with few-layer graphene and graphene oxide, in an attempt to increase the mechanical stability of the water-soluble polymer without excessively reducing its transport properties.2 and N2 permeability at different relative humidity.Commercial PVAm, Lupamin 9095, was purified via a multistep procedure to obtain two grades of pure materials, PVAm-LG and PVAm-HG, which have, respectively, low and high purification grade. These samples were then loaded with 3 wt% of few-layer graphene or graphene oxide and were characterized through FTIR, SEM, TGA, and DSC techniques as well as by measuring water vapor sorption and COThe results obtained confirm the sodium formate removal during the purification steps and also the neutralization of amine groups on the polymeric chain in the purest product (PVAm-HG), which should improve the facilitated transport performance of the membrane. In addition, PVAm-HG also seems to be able to better disperse the nanofiller with respect to PVAm-LG; both DSC and water sorption results were greatly affected by the addition of the nanofiller in the highly purified polymer while remained substantially unchanged in the partially purified material. SEM images in this concern showed in the PVAm-LG materials the presence of sodium formate nanocrystal, which may interact with the polar graphenic nanofiller affecting its dispersion in the matrix. 2 permeability of almost 70 Barrer (PVAm-LG + 3% GO) and maximum CO2/N2 selectivity of about 81 (PVAm-HG + 3% GO). In general, both permeability and selectivity increased with the increasing of relative humidity and water content in the polymer. Moreover, PVAm-HG resulted to be more selective and less permeable than PVAm-LG-based materials and, among different nanofiller GO, resulted to give similar permeability but slightly higher selectivity with respect to graphene. Permeation tests, conducted at 35 °C, gave some interesting results at high relative humidity (95% RH), with a maximum COTherefore, the further purification step did not seem to remarkably modify the materials separation performances, which in any case resulted to be below the Robeson’s upper bound, for all the materials considered. This is likely because of the barrier effect performed by the high quantity of filler added to the system. Graphene-based nano-sheets, however, allowed achieving complete permeation tests on self-standing membranes, which was not possible in polymer without filler, due to film rupture at high relative humidity. In this way, graphene addition has confirmed the ability to act as a reinforcer for this type of polymers containing in side-chain amino groups, thus presenting some potential as an additive for FTM membranes applications. Further research, however, has to be performed to prevent excessive reduction of the pristine PVAm separation performances upon nanofiller addition.


The electronic band structure of bilayer graphene is studied systematically in the presence of substitutional B and/or N doping, using density functional theory with van der Waals correction. We show that introduction of B-N pairs into bilayer graphene can be used to create a substantial band gap, stable against thermal fluctuations at room temperature, but otherwise leaves the electronic band structure in the vicinity of the Fermi energy largely unaffected. Introduction of B-N pairs into B and/or N doped bilayer graphene likewise hardly modifies the band dispersions. In semiconducting systems (same amount of B and N dopants), however, the size of the band gap is effectively tuned in the presence of B-N pairs. Substitutional doping with foreign atoms is one of the main approaches to intrinsically alter the electronic properties5. In particular, graphene is often doped with B6 and N7, as these atoms have radii  similar to that of C (67 pm) and the planar geometry therefore is preserved9. Gradual transition from n-type to p-type character in contact with air can be neutralized by intentional overdoping15. Bilayer graphene early has been proposed as viable alternative to the monolayer16. In the case of AA stacking the material consists of two graphene layers located directly on top of each other, whereas in the case of AB stacking half of the C atoms from one layer sits on top of C atoms of the other layer and the other half sits on top of hollow sites. Both stackings have been observed experimentally18, AB stacking being 5 meV per C atom lower in energy than AA stacking19 and nowadays can be grown in high quality and with large area20. It has been demonstrated that both B and N doping requires less energy than in the monolayer21.Absence of a band gap limits applications of graphene in electronics, calling for band gap engineering22 both for the top and hollow sites (which refers here and in the following to the position of the dopant atom relative to the other layer) and similar electronic band structures have been found for the two sites in each case. B and N favor hollow sites over top sites by 11 meV and 16 meV, respectively21. In AB stacked bilayer graphene with one N atom in each layer the total energy declines significantly for increasing distance between the two dopant atoms, while for corresponding doping with B atoms hardly any dependence of the total energy on the distance is found23. Codoping of B and N in graphene (monolayer) is energetically favorable in the form of B-N pairs in ortho position, for example, by 20 meV as compared to next nearest neighbours24. Reference25 has studied bilayer graphene with one B atom in one layer and one N atom in the other, claiming that the electronic properties depend on the occupied sublattice rather than on the mutual distance. In the present work, we investigate the electronic properties of bilayer graphene in the more general case that B and N also can be codoped in the same layer. Interestingly, it turns out that the presence of such B-N pairs has only minor effects on the electronic band structure except for the position of the band edges in the case that the system is semiconducting. For this reason, our results indicate that intentional doping of B-N pairs is an effective way to control the size of the band gap.Introduction of one B or one N atom into AB stacked bilayer graphene has been studied in ref.26 and the generalized gradient approximation for the exchange-correlation potential (Perdew-Burke-Ernzerhof parameterization), taking into account a van der Waals correction with Becke-Johnson damping27. The plane wave energy cutoff is set to 500 eV. We employ a Γ-centered 2 × 2 × 1 k-mesh for the structure optimization and a Γ-centered 8 × 8 × 1 k-mesh for evaluating the electronic properties. Both for AA and AB stacking, a 5 × 5 supercell of bilayer graphene is constructed with a total of 100 C atoms (to consider doping effects) to which a 16 Å thick vacuum layer is attached along the c-axis in order to avoid artificial interaction between periodic images. The interlayer distance after structure optimization turns out to be 3.6 Å for AA stacking and 3.4 Å for AB stacking, in agreement with the experimental observations in refs29. These values are not modified by the presence of dopant atoms for the concentrations considered in this study. We obtain for both stackings an in-plane lattice parameter of 2.46 Å and a C-C bond length of 1.42 Å. As in no case atoms develop significant magnetic moments (≤0.01 μB), we present in the following spin-degenerate band structures.Spin-polarized first-principles calculations are performed using the Vienna ab-initio simulation package30. We next introduce one B atom into one layer and one N atom into the other, representing a dopant concentration of 2%, and examine different distances between them as well as doping at both the top and hollow sites. In each case we obtain minimal energy for minimal distance between the dopant atoms. The energetically favorable configuration with AB stacking has 467 meV (~5 meV per C atom) lower energy than that with AA stacking. In the case of AB stacking, when B sits at a top site and N at a hollow site the energy is 14 meV lower than when both sit at hollow sites, 75 meV lower than when N sits at a top site and B at a hollow site, and 65 meV lower than when both sit at top sites. The electronic band structure of the minimum energy configuration, see Fig. 25. There are two buried Dirac cones at the K point, which can be attributed to the layers doped with B (unoccupied cone) and N (occupied cone), and we obtain a band gap of 298 meV.For reference, we show the electronic band structure of AB stacked bilayer graphene in Fig. 31. Concerning the position of the B-N pair relative to the other layer we find that N occupying the hollow site gives an energy advantage of 22 meV, which agrees with the finding of ref.21 that N favors the hollow site in bilayer graphene more strongly than B (by 22 meV). It turns out that the total energy is 590 meV (~6 meV per C atom) lower for AB stacking than for AA stacking. The electronic band structure of the minimum energy configuration is shown in Fig. 32), with a band gap of 70 meV due to the presence of the B-N pair. As a remark, we note that the electronic properties depend on the sites  occupied by the dopant atoms rather than on the occupied sublattice as claimed in ref.25.Turning to the situation that both the B and N atoms are located in the same layer while the other layer is undoped, we confirm that for both AA and AB stacking the formation of a B-N pair is energetically favorable, as a strong covalent bond is formedWe next study, for different relative positions, one B-N pair in one layer and one N atom in the other. It turns out that the N atom of the B-N pair still favors the hollow site, as the coordination geometry is not modified. The numbers in Fig. Keeping one layer undoped, we now add one B atom to the layer containing the B-N pair. The numbers in Fig. We finally study, for different relative positions, one N atom in one layer and one B-N-B triple in the other. The numbers in Fig. In conclusion, the effects of B and/or N doping on bilayer graphene have been studied by first-principles calculations based on density functional theory. Special attention has been given to the modifications of the electronic band structure due to the presence of B-N pairs. We have demonstrated that such B-N pairs hardly modify the band dispersions in bilayer graphene, while opening a substantial band gap . Similarly, introduction of B-N pairs into N and/or B doped bilayer graphene largely maintains the electronic band structure (for any relative position between the dopant species). Interestingly, for semiconducting systems (same amount of B and N dopants) our results show that the presence of B-N pairs shifts the band edges  towards each other, i.e., the band gap shrinks. For this reason, introduction of B-N pairs into bilayer graphene can be used to tune the size of the band gap without critically affecting the band dispersions.


Background and objectives: Idiopathic pulmonary fibrosis (IPF) has a particularly poor prognosis, and most IPF-related deaths are due to acute exacerbation (AE) of this condition. Few reports about biomarkers to predict prognosis of AE-IPF have been published since the release of the new AE-IPF criteria in 2016. The present study investigated relationships between serological markers and in-hospital mortality after the onset of AE-IPF. Methods: Demographic, serological, and imaging data from patients hospitalized at the Maebashi Red Cross Hospital  between 1 January 2013, and 31 December 2017, were retrospectively reviewed. Subjects fulfilling the diagnostic criteria for AE-IPF were divided into those who survived or died; statistical analysis of risk factors was performed using data from these two groups. Results: Diagnostic criteria for AE-IPF were fulfilled by 84 patients ), with a median age of 78 years . IPF was diagnosed before hospitalization in 50 (59.5%) patients and 38 (45.2%) died in hospital. Among the serological markers at hospitalization in the deceased group, C-reactive protein (CRP) was significantly higher than in the survivor group (p = 0.002), while total serum protein (p = 0.031), albumin (p = 0.047) and total cholesterol (p = 0.039) were significantly lower. Cox hazard analysis of factors predicting mortality, corrected for age, sex and BMI, revealed the following: CRP (hazard ratio (HR) 1.080  1.022–1.141); p = 0.006), LDH (HR 1.003 (95% CI 1.000–1.006); p = 0.037), and total cholesterol (HR 0.985 (95% CI 0.972–0.997); p = 0.018). Conclusions: Our data suggest that CRP, LDH, and total cholesterol may be biomarkers predicting mortality in patients with AE-IPF. However, only prospective controlled studies can confirm or not our observation as a generalizable one. Acute exacerbation (AE) is the cause of most IPF-associated deaths: the mortality rate among patients with IPF increases to 40–80% once AE has occurred ,2,3. Ris2/FiO2) , and bei2/FiO2) . There h2/FiO2) , and rec2/FiO2) ; howeverThere have been numerous studies investigating the relationships between imaging and physiology tests and AE-IPF; however, reports describing the relationships between serological markers and AE-IPF have addressed only KL-6 and a small number of other markers. In addition, the number of published reports relating to the most recent AE-IPF diagnostic criteria, released in 2016 , is limiThe present study was a retrospective investigation of the relationships between serological/physiological markers and the onset of AE-IPF defined by the most recent diagnostic criteria , and an The present investigation was a single-center, retrospective study. Patients who were hospitalized at Maebashi Red Cross Hospital  between 1 January 2013, and 31 December 2017, were evaluated; those who fulfilled the diagnostic criteria for AE-IPF were included in the study. All participants in this study were confirmed to have met the 2016 AE-IPF diagnostic criteria at the enrollment of this study. We diagnosed IPF with a multidisciplinary approach. At least two respiratory specialists and a radiology specialist discussed and reviewed subjects. A histologist was also consulted when histological data were obtained.The AE-IPF diagnostic criteria, as described by Collard et al. in 2016 , were inThis study was reviewed and approved by the Ethics Committee of Maebashi Red Cross Hospital on 12 July 2018 (acceptance no. 30-19).2/FiO2 and alveolar-arterial oxygen difference (AaDO2) were calculated from arterial blood gas analysis results. In addition, the duration of hospitalization, and whether death occurred while in hospital, were obtained from the medical records.The following data were obtained from medical records: demographic information ; sociological information ; IPF treatment before admission; treatment in hospital; serological test results in hospitalization; arterial blood gas analysis results; imaging results; and treatment progression. Body mass index (BMI) and body surface area were calculated from data. PaOp < 0.05 was considered to be statistically significant.Statistical analysis was performed using R version 3.4.1 and EZR version 1.37. Continuous variables were analyzed and compared using the Mann–Whitney-U test and expressed as median (maximum and minimum). Nominal variables were analyzed and compared using the Fisher’s exact test and expressed as numbers and ratios; Cox hazard analysis was performed, and corresponding hazard ratio (HR) and 95% confidence interval (CI) were calculated. For parameters that demonstrated significant differences, factors previously demonstrated to be correlated with AE-IPF mortality were included. Receiver operating characteristic analysis was performed using continuous variables that were judged to be significant factors, and cut-off values were calculated.Of the patients who were hospitalized at the Maebashi Red Cross Hospital between 1 January 2013, and 31 December 2017, consecutive 84 patients met the diagnostic criteria for AE-IPF and were included in the present study . The medp = 0.071). No differences in the history of diabetes, dyslipidemia, myocardial infarction, cerebral infarction, malignant tumors, or bronchial asthma were found between the deceased and survivor groups. Regarding serological markers at hospitalization, c-reactive protein  was significantly higher in the deceased group (11.57 (95% CI 0.77 to 27.73) versus 6.69 (95% CI 0.1–21.84); p = 0.002), and serum total protein , albumin , and total cholesterol  were significantly lower (6.5 (5.2–8.4) versus 6.8 (5.5–9.4); p = 0.031), 2.75 (1.7–3.5) versus 2.95 (2–4.1); p = 0.047), and 133 (82–477) versus 159 (66–222); p = 0.039), respectively. Surfactant protein-A  tended to be higher in the deceased group (124.2 (30.4–251) versus 91.2 (27–270.7)); however, the difference was not significant (p = 0.067). No significant differences in KL-6 or surfactant protein D (SP-D) were found. Arterial blood gas analysis revealed no significant differences in PaO2/FiO2 ratio or AaDO2.The subjects were divided into two groups: those who died (deceased); and those who survived (survivor). Statistical analysis was performed using data from these groups. As shown in p = 0.059). No significant differences were found in the number of subjects who underwent steroid pulse therapy, cyclophosphamide pulse therapy, or recombinant thrombomodulin administration. The number of subjects administered sivelestat was significantly higher in the deceased group (7 (18.4%) versus 1 (2.2%); p = 0.02). We confirmed the presence or absence of malignant tumor from medical records. Fifteen and nine patients had malignancies as comorbidities and past tumors, respectively. The most frequent comorbid/past tumor was lung cancer, which was observed in 11 patients.No significant differences between the groups were found in the number of subjects who were admitted to the intensive care unit, or in whom tracheal intubation or nasal high-flow therapy were performed. The number of subjects in whom non-invasive positive-pressure ventilation was performed tended to be higher in the deceased group (8 (21.1%) versus 3 (6.5%)); however, the difference was not significant , and KL-6 and lactate dehydrogenase (LDH), which were already known to be correlated with mortality, were corrected for age, sex, and BMI, and Cox hazard analysis was performed (p = 0.006), LDH (HR 1.003 (95% CI 1.000–1.006); p = 0.037), and T-chol (HR 0.985 (95% CI 0.972–0.997); p = 0.018) were predictive factors for time until death. Receiver operating characteristic analysis was performed, and the cut-off values of the mortality-predicting continuous variables were calculated  has been low; new alveolar lesions have developed; and other causative diseases, such as infection and heart failure, could be ruled out.AE-IPF accounts for 15% to 40% of IPF-related deaths ,12. In tSingle-center studies by Kataoka et al. and Sakamoto et al. reported 44 cases and 80 cases of AE-IPF in five years and 10 years, respectively ,14. ThusThe new diagnostic criteria proposed in 2016 widened the range of AE-IPF , in thatIn addition, among the patients who had already been diagnosed, only four patients used antifibrotic drugs, and many patients remained untreated. It is plausible that primary care physicians did not introduce antifibrotic drugs into the treatment of elderly patients in consideration of the drug’s side effects and that non-specialists administered corticosteroids according to past recommendations, despite not corticosteroids but antifibrotic drugs having been recommended for IPF treatment. Moreover, while anti-fibrotic drugs suppress the incidence of AE, the corticosteroid administration group exhibited an increase in the incidence of AE; this may account for the limited prescription of antifibrotic drugs in this study.Previously, low forced vital capacity , elevateThe relationships between serological markers and AE-IPF were investigated in this study, and findings suggested that CRP, LDH, and T-chol might be predictive factors for in-hospital mortality after AE-IPF. KL-6, LDH, and CRP have been known to be mortality-predicting factors at AE-IPF onset ,3,14.CRP is a protein that is produced by hepatocytes in response to stimulation by cytokines such as tumor necrosis factor-alpha, interleukin-1, and interleukin-6. During the chronic progression of interstitial pneumonia, it has been reported that a baseline CRP level associated with scleroderma was a long-term prognostic factor , and thaIt has been reported that elevation of serum LDH level in interstitial lung disease was a predictive factor for the onset of AE  and elevThere have been virtually no reports addressing the relationship between serum cholesterol level and AE-IPF. Correlations between low serum cholesterol level and mortality rate are known in several diseases, and it has been reported that hypocholesterolemia was an independent mortality-predicting factor with malignant tumors  and thatIt may be difficult to predict the prognosis of AE-IPF by using a single biomarker because various causes contribute to the development of AE-IPF. It has previously been reported that hypolipidemia was a risk factor for respiratory tract infection; it impairs immune responses to lung infection, resulting in the development of respiratory tract infection . Low totConsistent with this supposition, experimental and clinical investigations have revealed that high levels of cytokines, such as CRP, decreased circulating cholesterol levels during severe infection . Sepsis In addition to the presence of untriggered AE, this study included triggered AE caused by various factors, including infection and operation. Due to the heterogeneity among the pathologies of patients with AE-IPF, a single biomarker cannot reflect the various factors underlying the development of AE-IPF; hence, the AUC and specificity were relatively low.There are a few limitations in this study. Because this is a small, single-center retrospective study, the study population may not represent the entire AE-IPF patient population. In particular, a few patients in this study were treated with antifibrotic drugs. The administration of antifibrotic drugs is currently standard for the treatment of IPF; future analyses should thus be performed with populations featuring a greater proportion of patients treated with antifibrotic drugs. In addition, we were able to analyze serological markers only at admission and could not, therefore, record continuous changes in the levels of serological markers. This study cannot confirm serial changes of serological markers during AE. Therefore, the results in this study should be confirmed in a large size and in a multi-center study. In this study, we obtained spirometry data of only part of IPF patients before the onset of AE-IPF. Many of the 50 cases diagnosed with IPF before hospitalization were followed up at the local clinic when the IPF was at a stable condition and hence we could not obtain the data regarding spirometry performed at the clinic; alternatively in some other cases, a follow-up spirometry was not adequate at that time due to AE-IPF onset. Follow-up of spirometry in IPF patients is crucial as low forced vital capacity is associated with the onset of AE-IPF , and lowOur data suggest that CRP, LDH, and total cholesterol may be biomarkers predicting mortality in patients with AE-IPF. However, only prospective controlled studies can confirm or not our observation as a generalizable one.