In comparison, the anodic electrochemically exfoliated graphene exhibited a capacitance of approximately 44 F g−1, the highest of the ‘pure’ graphene materials, which all exhibited superior capacitance retention, owing to their higher conductivity. The cyclability of all of the materials, with the exception of reduced graphene oxide (70 %), was found to be greater than 95 % after 10 000 cycles. These results highlight the importance of matching the graphene production method with a specific application; for example, graphene oxide and anodic electrochemically exfoliated graphene would be best suited for high energy and power applications, respectively.A comparison of the performance of graphene‐based supercapacitors is difficult, owing to the variety of production methods used to prepare the materials. To the best of our knowledge, there has been no systematic investigation into the effect of the graphene production method on the supercapacitor performance. In this work, we compare graphene produced through several routes. This includes anodic and cathodic electrochemically exfoliated graphene, liquid phase exfoliated graphene, graphene oxide, reduced graphene oxide, and graphene nanoribbons. Graphene oxide exhibited the highest capacitance of approximately 154 F g Production matters: Comparing the performance of graphene‐based supercapacitors in the literature is difficult, owing to the variety of testing regimes and production methods used to prepare the materials. To the best of our knowledge, there has been no systematic investigation into the effect of the graphene production method on the supercapacitor performance. In this work, graphene produced through several “standard” routes is compared. This includes anodic and cathodic electrochemically exfoliated graphene, liquid phase exfoliated graphene, graphene oxide, reduced graphene oxide, and graphene nanoribbons. A further attraction is the long cycle life, typically 100–1000 times the value for batteries, but the devices tend to have comparatively low energy density. Therefore a key focus for the rapidly growing research into supercapacitors is to combine increased energy density−1).2 g−1),−1).Porous carbons are currently used as an electrode material, with the most commonly used being activated carbon due to its moderate conductivity, high surface area, low density and high specific capacitance , anodic (AEEG) and cathodic (CEEG) electrochemically exfoliated graphene, graphene oxide (GO), reduced graphene oxide (rGO) and graphene nanoribbons (GNR).In this work we have produced five types of graphene: liquid phase exfoliated graphene (LEG), cathodic electrochemically exfoliated graphene (CEEG), anodic electrochemically exfoliated graphene (AEEG), graphene oxide (GO) using a modified Hummers method and chemically‐reduced graphene oxide (rGO) using ascorbic acid as the reductant for GO. We have also acquired graphene nanoribbons (GNR) from a commercial source, which were fabricated using oxidative unzipping of carbon nanotubes. The graphene materials were produced by “standard” conditions of the production method to give typical materials e. g. the rGO was fully reduced. These materials were first characterised by a variety of techniques including Raman spectroscopy, X‐ray photoelectron spectroscopy (XPS), nitrogen gas adsorption and scanning electron microscopy (SEM) to determine the quality of the electrode materials.Electrodes were formed through the filtration of dispersions of the materials onto polyvinylidene fluoride (PVDF) filters to make thin membranes. This made it possible to stack the membranes back to back using the PVDF filter as the porous separator. These were then characterised electrochemically in a symmetrical two electrode CR2032 coin cell configuration using an aqueous electrolyte. This approach to electrode preparation removes the need for binding agents or conductive additives, thereby simplifying the device design and enabling the intrinsic properties of the graphene materials to be assessed. It has been shown that the use of a three electrode system can exaggerate the performance of an electrode, more than doubling the measured capacitance.22.1Figure −1  to ∼2690 cm−1, and is approximately four times more intense than the G peak.−1 and has multiple contributions indicating the presence of multi layered flakes (>5). As both production methods used have been shown to produce few layer (1–3) flakes in dispersion, it means that restacking of the materials upon filtration has occurred. The D and D’ band in graphene is caused by the presence of edges and defects within the sample. Therefore, the LEG, AEEG and GNR spectra indicate that the flakes contain many defects due to either the ultrasonic exfoliation or oxidation. However, in the case of CEEG there is no D’ band present and the D band has a very small intensity indicating that the CEEG material is much less defective than the LEG, AEEG and GNR.Typically, for monolayer graphene, the 2D band contains one peak contribution, which is blue shifted from ∼2750 cm3 hybridised carbon bound to the oxygen‐containing groups. Upon reduction of the GO, a majority of the oxygen‐containing groups are replaced with hydrogen. This also results in a large D band in the rGO spectrum as the sp3 hybridised carbon defect is still present.In the case of the GO and rGO, the D and G bands are shifted to higher wavenumbers due to the disorder introduced by defects and surface functionalisation. The GO spectrum contains a large D band again due to defects, which in this case are the spTo investigate the chemical state of the graphene materials further, XPS was performed on the graphene material membranes Figure b. XPS caThe only elements detected within the materials were carbon and oxygen as expected. In the case of CEEG and LEG the oxygen concentration is low (<5 %), this is expected as the two production methods are not oxidative.To further investigate the morphology of the graphene membranes, SEM analysis was performed. Typical SEM images of each membrane are displayed in Figure Figure Finally the surface area of the electrodes was investigated using nitrogen gas adsorption. Table 2 g−1). This is due to the restacking of the materials during electrode fabrication (filtration) which reduces the surface area. Restacking during electrode formation has been demonstrated previously for a wide range of 2D materials.−1). However, it can also be seen within table vide infra). Therefore, there is a trade‐off between increasing the surface area of the graphene material through the introduction of defects and increasing the conductivity through the removal of defects. This demonstrates the care needed in ensuring that the graphene material employed is suitable for the specific application.It can be seen from table 2.24) in acetonitrile, allows a larger operating potential to be used, of ∼2.7 V. However, organic electrolytes have some limitations such as low conductivity and lower solubility, as well as their sensitivity to moisture. Figure After physical characterisation of the membranes was complete, electrochemical testing was carried out in a symmetrical coin cell. The electrolyte used was 6 M KOH: this limits the operating potential of the supercapacitor to 1 V due to the electrolysis of water at this pH.It can be seen that at low scan rates and current densities all the graphene materials exhibit the response typical of an EDLC. Treating Faradaic processes as pseudocapacitive, the area enclosed by the CV curve is proportional to the capacitance. Therefore, when all the materials are compared, it can be seen that GO, AEEG and rGO show significantly increased performance over the other materials. This large difference between the performance of the AEEG, GO and rGO and the other graphene materials is also seen in the galvanostatic charge‐discharge (GCD) response, as the gradient of the slope after the initial decay due to iR drop is lower for the AEEG, GO and rGO materials. However, it can be seen that when the scan rate and current density were increased both rGO and GO showed a large drop in performance (∼50–60 %). This loss in performance is attributed to the high resistance of the GO and rGO caused by the large number of defects and the oxygen functionalities present. Furthermore, it was noted that the coulombic efficiency of the rGO material was lower (60 %) when compared to the graphene materials (>95 %). The ESR of the materials was measured using electrochemical impedance spectroscopy (EIS), which was fitted using a standard Randles circuit to obtain the ESR of the materials and the results are displayed in Table 2 or MnO2,This high resistance limitation has been encountered for many different materials, such as MoSFigure The low frequency values of C’ corresponds to the capacitance of the system during discharge, while C’’ corresponds to the irreversible energy losses of the system. Figure Figure The cyclability of the supercapacitor electrodes is one of the most important metrics for comparison between materials. An ideal EDLC should have an extremely high cyclability with minimal degradation in performance after repeated GCD cycles. This high cyclability is due to the energy storage mechanism being purely physical with virtually no surface reactions occurring. For the GCD approach to cycling of graphene‐based electrodes, there are several factors that contribute to the slow degradation in capacitance. These include increased resistance between the flakes, collapse of the porous structure, desorption of active material from the current collector and for AEEG, GO and rGO surface reactions between the functional groups and the electrolyte. Therefore, the cyclability of the material is a key factor in determining its suitability for commercial supercapacitor applications. The capacitance loss for graphene including rGO and GO over 10,000 cycles can range from 30 % to no loss depending on the testing conditions.−1. A pre‐treatment regime of ∼100 cycles at 1 A g−1 was applied to all cells to ensure that the electrolyte had permeated the porous structure of the electrode.Figure In the case of the LEG, CEEG, AEEG and GNR there was no observable degradation in performance of the material over the 10,000 cycles. This cyclability behaviour is indicative of EDLCs suggesting that this is the primary energy storage mechanism for these materials even after cycling. For GO there was a slight increase in capacitance by ∼4–5 % after 10,000 cycles. This slight increase in capacitance has been observed before and is ascribed to the gradual exfoliation of the GO electrode by the water during cycling.−1.Figure It can be seen in Figure The normalised real and imaginary components of the capacitance were also calculated, as described above, to give further insight into the processes occurring at the electrode‐electrolyte interface. These normalised values of the capacitance are shown in Figure After cycling, the normalised C’ value for the LEG, CEEG, AEEG, GNR and GO materials did not show any noticeable change, indicating that the capacitor response with discharge time had not changed dramatically. In the case of the rGO a sharp drop off in the C’ value was observed corresponding to the material no longer acting as an ideal capacitor for any of the measured frequencies. This is consistent with the CV and GCD results which also show a deviation from EDLC behaviour for rGO after cycling. Figure For rGO no peak was present in Figure The gravimetric capacitances calculated from the CVs and GCDs after cycling for all graphene materials at the high and low scan rates and current densities are shown in Table To investigate the large loss of capacitive performance of the rGO material, the cycled coin cells were opened to allow inspection of the rGO electrodes. In all cases it was found that the rGO electrodes had mechanically failed, resulting in delamination from the current collector and PVDF support. This explains the large increase in resistance as the internal wiring of the rGO electrodes had been compromised, shown in figure S1. It also explains the loss of capacitance, asymmetry in the CV and poor coulombic efficiency shown by the rGO as the amount of active material and surface area present on each electrode has been reduced. This would result in the potentials of each electrode no longer being symmetrical during operation due to the loss of mass on the positive electrode. This means that the positive electrode would reach higher potentials to equalise the charge between the electrodes resulting in degradation reactions occurring at higher voltages due to the positive electrode potential exiting the electrochemically stable potential window of the system.Another important metric when discussing the application of supercapacitor electrodes is the self‐discharge of the active material. The self‐discharge severely limits its ability to be applied to real world situations. There are three mechanisms that cause the spontaneous decrease in voltage within a supercapacitor: Faradaic processes, leakage current and charge redistribution.For supercapacitors the voltage decreases rapidly from the initial value and the rate decreases after that with time.Figure It can be seen that the GO and rGO electrodes exhibit the best charge retention with the pure graphene materials  showing progressively worse retention. This is believed to be due to the pseudocapacitive functional groups present on the surface allowing for reduced Faradaic leakage and charge redistribution across the membrane.3In summary, the electrochemical performances of six different graphene‐derived materials were compared as electrodes for supercapacitors. The materials that were tested included liquid exfoliated graphene, cathodic and anodic electrochemically exfoliated graphene, commercial graphene nanoribbons, graphene oxide and reduced graphene oxide. It was observed that graphene oxide and reduced graphene oxide showed the highest specific capacitances of ∼153 F/g and ∼119 F/g, respectively. It was also observed that the anodic electrochemically exfoliated graphene showed the highest capacitance, ∼44 F/g, of the pure graphene samples. The increased capacitance in these materials is due to the oxygen containing functional groups present on the materials imparting an additional pseudo‐capacitive contribution. However, the reduced graphene oxide electrodes underwent a severe degradation in performance and mechanically failed during galvanostatic charge discharge cycling. The other graphene materials showed no loss in performance after cycling. However, due to the decreased conductivity caused by the presence of the oxygen groups, the graphene oxide electrodes showed a large loss in performance for faster discharge rates due to resistive losses. This can be overcome through the addition of conductive additives such as graphene or CNTs. For the other graphene materials, very little loss in capacitance was observed for increased discharge rates. For high energy applications involving graphene materials, graphene oxide would be the material of choice whereas for high power applications involving graphene, anodic electrochemically exfoliated graphene would be the best choice. These electrochemical results suggest that the graphene morphology and hence production method plays a large role in its supercapacitive performance. Therefore, the possible supercapacitor application should dictate the type of graphene used.−1) and ultrasonically processed in a bath sonicator (Elmasonic P70H). The bath sonicator was operated at 37 kHz and 40 % amplitude for 12 hours with cooling to maintain a stable temperature of 20 °C. The sonicated dispersions were centrifuged at 6000 rpm (3139 g) for 30 minutes to remove any unexfoliated material. The supernatant was decanted and fresh solvent was added before repeating the centrifugation so that a narrow distribution of flake dimensions and thicknesses was obtained. These graphene dispersions were stable in solution for several months with virtually no detectable sedimentation. The yield of few‐layer LEG flakes for this technique is typically ∼2 %.Graphene dispersions were created by a liquid phase exfoliation method that has been reported previously.Graphene dispersions were created using a modified electrochemical exfoliation method that has been reported previously.4)2SO4 in water and a potential of 10 V was applied for 2 hours. After the graphite exfoliation was completed, the product was washed several times with deionised water by vacuum filtration. The resultant graphene was then dispersed in N,N’‐dimethylformamide (DMF) by sonication (Elmasonic P70H) at 30 % amplitude for 30 minutes. The dispersion was left for 48 hours to precipitate any un‐exfoliated graphite flakes or particles. The top part of the dispersion was filtered and washed twice with water. The filter cake was then freeze‐dried to obtain electrochemically exfoliated graphene. The yield of few layer AEEG flakes from this method is typically ∼30–40 % and the oxygen content for these flakes was ∼9 % determined by XPS.Graphene dispersions were created using an anodic electrochemical exfoliation method.Graphene oxide was produced using a modified Hummers method.The chemical reduction of GO was done using a method that had been previously reported.Commercially sourced GNR powders were dispersed in NMP (10 mg/ml) before sonicating (Elmasonic P70H) at 37 kHz and 40 % amplitude for 30 minutes to ensure a homogenous suspension. This GNR dispersion was put through the same centrifugation treatment as the liquid phase exfoliated graphene material to ensure that no exfoliated material remained. These GNR dispersions were stable in solution for several months with virtually no detectable sedimentation.Membranes of the graphene materials were created by first diluting the NMP solutions (15‐fold dilution) with isopropanol (IPA). These diluted solutions were filtered through preweighed PVDF filters (Millipore) with a 0.1 μm pore size. The filtered membranes were then washed with IPA to remove any residual NMP and dried in an oven overnight at 80 °C to remove any residual solvent. The mass of the graphene materials was measured by weighing the membrane after drying and was typically between ∼4–5 mg. The graphene membranes were analysed by Raman spectroscopy, SEM, XPS and then used as the electrodes within coin cells.−1 was used, which gave a spectral resolution of ∼1 cm−1.Raman spectroscopy was performed to give an initial characterisation of the filtered membranes. This technique can enable the determination of the layer number and defect density of the graphene materials after filtration. To prepare the samples for Raman analysis the filtered membranes were placed in a vacuum oven at 100 °C to ensure that all solvent had been removed. Raman measurements were performed using a Renishaw inVia Microscope with 532 nm (2.33 eV) laser excitation at a power of 1 mW with a 100× objective. A grating of 1800 mmScanning electron microscopy (SEM) analysis was performed on the filtered membranes supported by the PVDF with a FEI/Phillips XL30 E‐SEM under high vacuum conditions with an accelerating voltage of 15 kV. All images were taken using the secondary electron detector.α X‐ray source , a hemispherical electron energy analyser and a multichannel plate and delay line detector (DLD). The analysis area of the membranes was 300 μm×700 μm. Survey spectra were collected to determine the elemental composition. Charge neutralisation was used for all samples.X‐ray photoelectron spectroscopy (XPS) was performed using a Kratos Axis Ultra DLD spectrometer with a monochromated Al KNitrogen gas adsorption measurements were performed on the graphene material membrane electrodes using a Quadrasorb‐evo Gas Sorption Surface Area and Pore Size Analyser (Quantachrome). To avoid the contribution of the highly porous PVDF support, the graphene material was removed from the PVDF before testing. To ensure that the measurement could be completed in the 36 hours of testing time available, approximately 30–40 mg of sample was tested during each run. The specific surface area of the materials was calculated using non localised density functional theory with a slit pore model.Cyclic voltammetry (CV), electrochemical impedance spectroscopy (EIS), galvanostatic charge‐discharge (GCD) and self‐discharge (SD) were performed using a PGSTAT302 N potentiostat (Metrohm Autolab). All electrochemical measurements were performed using a sealed symmetrical coin cell (CR2032) system; for each material a minimum of four coin cells were measured. The membranes were stacked back‐to‐back within the coin cell with the active material making direct contact with the stainless steel spacer, which acted as a current collector; a steel wave spring was used to ensure a good contact was kept during the testing. The electrolyte used was aqueous 6 M KOH. CV was performed at scan rates ranging from 10 to 500 mV/s. EIS was performed at a frequency range of 0.1 Hz to 100 kHz with a 10 mV (rms) perturbation and 0 V dc bias. GCD of the cells was performed at current densities ranging from 0.5 A/g to 4 A/g. The specific capacitance was calculated using the best practice methods established by Ruoff.The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


Here, we report a novel method for low-temperature synthesis of monolayer graphene at 450 °C on a polycrystalline bimetal Ni-Au catalyst. In this study, low-temperature chemical vapor deposition synthesis of graphene was performed at 450 °C on codeposited Ni-Au which shows successful monolayer graphene formation without an extra annealing process. The experimental results suggest that electron beam codeposition of bimetal catalyst is the key procedure that enables the elimination of the pre-growth high-temperature annealing of the catalyst prior to graphene synthesis, an indispensable process, used in previous reports. The formation was further improved by plasma-assisted growth in which the inductively coupled plasma ionizes the carbon precursors that interact with codeposited Ni-Au catalyst of 50 nm in thickness at 450 °C. These combined growth conditions drastically increase the graphene’s sheet uniformity and area connectivity from 11.6% to 99%. These fabrication parameters enable the graphene formation that shifts from a bulk diffusion-based growth model towards a surface based reaction. The technique reported here opens the opportunity for the low-temperature growth of graphene for potential use in future CMOS applications. While direct growth of graphene on Si or SiO2 has not yet been achieved, graphene synthesis on thinner catalysts represents a substantial step towards this goal. On a thinner catalyst, the resultant graphene layer is closer and closer to the target substrate. This creates the potential for developing a transfer process that is less harmful to the graphene by minimizing the amount of manual handling of the graphene due to its close proximity to the target substrate. While graphene is commonly synthesized using chemical vapor deposition (CVD) techniques, the addition of a remote plasma can help to reduce synthesis temperatures. Plasma growth energizes the precursor gases via ionization, overcoming the thermal energy that is lost when growing in the 400 °C ~ 500 °C range compared to the 800 °C ~ 1000 °C range. The advantages of a remote inductively coupled plasma are twofold: plasma is created away from the growth catalyst which reduces damage from ion bombardment on the synthesis surface, and that the plasma is produced via induction coils which are outside of the graphene growth chamber where the catalyst substrate is located. In a typical capacitive plasma system, the synthesis stage is between two metal plates inside the vacuum chamber with an electric potential between them which exposes the growth surface to any foreign material from the plasma source. Using exterior induction coil wrapped around the gas flow tube, we ionize the precursor gases by creating an alternating electromagnetic field inside the tube. This completely removes the plasma source from the chamber which does not allow for foreign material from the plasma source to potentially contaminate the growth catalyst. Here we report the synthesis of monolayer graphene (MLG) at 450 °C using inductively coupled plasma chemical vapor deposition (ICPCVD) on thin (50 nm) Ni-Au catalyst codeposited by the electron beam evaporation technique.It has been more than 10 years since the isolation of graphene , 3. As a2/Si wafers to the desired thickness (50 nm). For Ni-Au catalyst preparation, Au and Ni pellets were first mixed by electron beam heating, where the electron beam system directs a beam of electrons to a crucible containing metal pellets of both Au and Ni. Electron beam deposition was preferred here due to the fact that it allows for accurate control of the weight percentage of Au in the mixture. The electron beam heats and mixes the pellets, and the resulting mixture is then codeposited via evaporation onto SiO2/Si wafers as schematically illustrated in Fig. Two methods of catalyst preparation were used for this study to compare the effects of Ni-Au versus pure Ni, and all catalyst depositions were performed in a separate Kurt J. Lesker Physical Vapor Deposition (PVD) tool. For pure Ni graphene growth experiments, the Ni catalyst was prepared via magnetron sputtering onto SiO−6 Torr after which H2 was introduced at 15 sccm as the sample was heated to the process temperature of 450 °C. Once the samples reached the process temperature, the H2 flow was ceased, and the chamber was purged using Ar. After base pressure was reestablished, Ar was ceased and C2H2 was flown at 0.1 sccm bringing the chamber pressure to 6 × 10−6 Torr. For plasma-enhanced growth, RF power varying from 0 to 10 W was introduced for different growth samples. Ten watts were the highest RF power that could be supplied while maintaining a stable, remote inductively coupled plasma (ICP) since at such a low C2H2 pressure the gas resistance is very high and the induced current begins flowing elsewhere at higher plasma powers. Once the desired growth duration was reached (7 min for CVD growths and 30 s for ICPCVD growths), C2H2 flow was stopped, followed by another Ar purge to flush the remaining process gases out of the chamber. The sample heater was then shut off allowing the sample to cool to room temperature.Our graphene growth regimes were performed in a custom-built cold-wall ICPCVD chamber. Temperatures were measured via a thermocouple gauge affixed to the sample surface in order to detect accurate in situ surface temperatures. Gases were flown into the chamber via a showerhead. The samples were placed on a radiantly heated stage roughly 25 cm below the gas showerhead. The prepared catalysts were entered into the ICPCVD chamber and pumped to a base pressure of 1 × 103 to etch the Ni, followed by iodine/potassium iodide (40 mL H2O/4 g KI/1 g I2) solution to etch the remaining Au. The resulting film was then transferred to clean SiO2/Si wafers, and the PMMA was etched away with acetone. Graphene spectroscopic analysis was performed using a Horiba Jobin Yvon HR800 UV Raman spectrometer with a 532 nm laser to identify the key spectral peaks expected for graphene thin films. The D:G and D:D’ peak intensity ratios  provide information about the defect density and defect type in the graphene respectively. The 2D:G peak intensity (I2D:G) as well as the 2D peak full width at half maximum (FWHM) provide information on the number of total graphene layers. Defect-free, monolayer graphene displays ID:G of approximately 0, however, when defects are present an I2D:G > 1.0 and FWHM2D < 100 cm−1 are indicative of monolayer graphene [I2D:G = 1 which suggests the percentage of monolayer graphene produced, and Raman peaks were fitted and analyzed using a program written in R to identify peak ratios and FWHM.Following synthesis, graphene samples were spin-coated with PMMA. The wafers were then placed in FeCl2H2 exposure on (a) pure Ni and (b) the codeposited Ni-Au catalyst without annealing in contrast to previous literature [I2D:G = 1.2 and FWHM2D = 48.5 cm−1 which suggests monolayer graphene formation with defects. There is a relatively large ID:G = 0.68 and a corresponding ID:D’ = 5.0 which suggest that vacancy and lattice mismatch defect types are present [Here, it is shown that the use of codeposition of Ni-Au as a catalyst preparation technique eradicates the previously required process of 600 °C pre-annealing to alloy the Ni-Au catalyst by comparing to a baseline Ni only control catalyst and to previous reports . To compterature . The accterature  suggeststerature . Figure  present , however present , and thi present  if locat present . Most imI2D:G map of codeposited Ni-Au grown graphene via ICPCVD  with a 30s C2H2 exposure time. Figure I2D:G, represented by large areas of uniform color, with increasing RF power. This shows that plasma addition assists in creating a larger, more uniform graphene sheet which is further supported by the data chart in Fig. I2D:G, a decrease in ID:G, and a decrease in FWHM2D which all are significant of monolayer graphene. To visually explore graphene thin film continuity, Fig. While codeposited Ni-Au does produce graphene films, it is crucial to be able to produce large areas of continuous and uniform graphene for practical applications. To address this issue, three growths on codeposited Ni-Au were performed to compare the effects of incorporating a remote plasma with three different plasma powers into the growth. Figure I2D:G > 1.0 while the 30s ICPCVD growth on the other hand displays a 99% coverage of I2D:G > 1.0. This suggests the plasma plays an important role in graphene connectivity and uniformity on a thin catalyst and prevents the absorption of large amounts of C into the catalyst as opposed to the CVD growth in which over absorption of C leads to overproduction of graphene and thus less uniformity.The benefits of ICPCVD over CVD can also be seen when comparing multilayer (MLG) coverage. A set of samples from both ICPCVD and CVD syntheses was compared and the results are shown in Fig. We have demonstrated the validity of codeposition as a unique catalyst preparation technique which effectively removes the necessity of annealing for Ni-Au catalysts by remotely pre-alloying the catalyst during e-beam evaporation and producing an immediately capable catalyst for 450 °C graphene growth. The addition of using an inductively coupled plasma during the growth serves to increase graphene thin film area and layer uniformity by shifting the synthesis process to a surface dominated mechanism which is beneficial when thin catalysts are used for growth. The study presented here demonstrated the significant progress of using plasma-enhanced CVD and codeposited Ni-Au thin catalyst to grow graphene with improved quality at low temperature. However, the growth parameters need to be tailored and optimized with respect to the specific applications. For example, the catalyst design and optimization to further increase the graphene grain size under the low-temperature growth conditions and the direct growth of graphene on desired substrates. These are the issues that will be addressed in the ongoing investigations.


Numerous viruses can cause upper respiratory tract infections. They often precede serious lower respiratory tract infections. Each virus has a seasonal pattern, with peaks in activity in different seasons. We examined the effects of daily local meteorological data  from Edinburgh, Scotland on the seasonal variations in viral transmission. We identified the seasonality of rhinovirus, adenovirus, influenza A and B viruses, human parainfluenza viruses 1–3 (HPIV), respiratory syncytial virus (RSV) and human metapneumovirus (HMPV) from the 52060 respiratory samples tested between 2009 and 2015 and then confirmed the same by a generalised linear model. We also investigated the relationship between meteorological factors and viral seasonality. Non-enveloped viruses were present throughout the year. Following logistic regression adenovirus, influenza viruses A, B, RSV and HMPV preferred low temperatures; RSV and influenza A virus preferred a narrow “humidity-range” and HPIV type 3 preferred the season with lower humidity. A change (i.e. increase or decrease) in specific meteorological factors is associated with an increase in activity of specific viruses at certain times of the year. URTI are mostly self-limiting but can progress to lower respiratory tract infections (LRTI) particularly in those with underlying conditions4. Adults suffer from a cold two to five times a year, and children can have 7–10 episodes annually5. Numerous viruses can cause URTI6. Rhinoviruses , are responsible for around 30% to 70% of all respiratory infections6. Since there are >100 different rhinovirus types, reinfections are very common7. Coronavirus is the second most common cause of URTI, causing 7% to 18% of cases8. Other viruses including respiratory syncytial viruses (RSV) types A and B, human parainfluenza viruses (HPIV) types 1–4, adenoviruses, human metapneumovirus (HMPV) and influenza viruses A-C  also cause URTI9.The common cold is typically a mild upper respiratory tract infection (URTI), with symptoms such as nasal stuffiness and discharge, sore throat, coughing and sneezing8. Three theories have been put forward to explain viral seasonality. 1: The effect of climatic conditions on host resistance to infection (low vitamin D levels following lack of sun exposure can affect our ability to fight infection)10. 2: The effect of meteorological factors  on virus survival and hence on infection rates11. 3: The effect of behavioural changes on transmission (e.g. spending more time indoors in close proximity to others or aggregation of susceptible children at schools during the colder months)9. There are other hypothesis such as diminished immune responses in a chilled host, or the reactivation of dormant viruses by chilling for the increase in URTI during the colder months12.Most respiratory viral infections have seasonality. In temperate regions, URTI increase in frequency in autumn and spring, and remain raised through winter13. Integrity of IAV envelopes are better at lower temperatures, whilst in warm temperatures, the envelope becomes disordered and the virus susceptible to damage14. Reports suggest that IAV survival is better associated with absolute humidity (which measures the amount of water vapour in air regardless of the temperature) rather than relative humidity16. An association of absolute humidity with hospitalisation due to influenza LRTI (but not with LRTI due to Rhinovirus or RSV) has recently been reported17.Temperature has an important effect on viral activity, particularly in the case of enveloped viruses. In a guinea pig model, IAV transmission is more effective in cold and dry conditionsMost reports on viral seasonality concentrate on viruses that appear in the winter months, particularly influenza viruses and RSV. Here we aim to determine the seasonality of all the respiratory viruses detected over 6.5 years, and identify any associations between changes in meteorological factors (such as temperature or humidity) and the appearance of different viruses across a whole year.rd of the year. HPIV-3 was present in 24.6% of days, followed by HPIV-1 (15.5%) and IBV (15.2%). HPIV-2 was present for the fewest number of days (6.5%) , followed by adenovirus which was present on 52.3% of the days. RSV (36.4%), HMPV (32.8%) and IAV (32.7%) were present in around 1/3Figures Many viruses have a definite seasonality Figs –3. RSV aTemperature appears to have the greatest effect on viral seasonality Fig. . The aveThe average relative humidity was 81% over the study period Fig. . RSV 84. The aveTable th December, 12th January and 8th February respectively. Even though rhinovirus and adenovirus are present throughout the year, their GLM’s point to periods of increased activity (around the 5th of March for Adenovirus and around the 6th November for rhinovirus). The GLM corroborates HPIV-3’s preference for warmer temperatures and its seasonal peak around the 4th of May. Finally, the GLM suggests that HMPV peaks around the 11th of March, HPIV-1 peaks around the 31st October, and HPIV-2 peaks around the 15th of November.Figure From Fig. et al. have reported that air temperature and relative humidity were important factors affecting viral seasonality in Germany over a 6-year period18. As in our study, air pressure did not have a significant correlation. However, to our knowledge, this is the first study to report the association between meteorological conditions and the sequential appearance of viruses through the year and the effect of daily fluctuations in humidity (i.e. “humidity-range”) on virus seasonality.Du Prel Our data shows that all the viral agents of URTI show degrees of seasonal variation corroborated by the GLM. Based on similarity in structures and seasonality, there appear to be three groups. 1: Non-enveloped viruses (rhinovirus and adenovirus) which are present throughout the year. 2: Enveloped viruses with winter preponderance  and 3: Parainfluenza viruses 1–3, enveloped viruses with a preference for warmer temperatures.19. Since this is very close to the mean humidity for the study period (81%), this could explain the presence of adenovirus throughout the year. Davis GW showed that adenovirus survival is better at high humidity levels (89%) than at lower humidity (50%)20. The peak in March found in our GLM for adenoviruses is different to the December peaks observed for enteric adenoviruses in Japan, though a peak in the winter months is seen when GLM was drawn for 2011, a year when visual peaks for all viruses were present 22, but the November peak found using the GLM is later than the September-October peaks found in other research. However, as these two peaks correspond to when children go back to school, it is possible that meteorological factors do not solely account for these peaks.The year-round presence of rhinoviruses is similar to other reports18. Whether this reflects the climatic conditions of the UK (being an island) compared to Germany needs further evaluation.As rhinovirus is present throughout the year, there appears to be no significant difference in meteorological factors on days rhinovirus is present or absent. This is unlike the negative correlation with temperature and a positive correlation with relative humidity reported from Germany24, IAV25 and IBV25. RSV and IAV appear to be associated with lower temperatures, dew point, “humidity-range” and high humidity, however, by logistic regression significant association is only seen with lower temperatures and lower “humidity-ranges”. IBV and HMPV which appear to be associated with lower temperatures, dew point, and “humidity-range”, are only associated with lower temperature by logistic regression. This illustrates the complex interaction between different meteorological factors and hence the difficulty in understanding their specific effects on viral prevalence. IBV has a biennial pattern in Edinburgh unlike the bimodal pattern from Okinawa, a subtropical region in Japan26. This suggests that local meteorological variation and factors affecting transmission and herd immunity might have a role to play in viral seasonality. Other studies have also reported early spring peaks of HMPV, association with low temperature28 and raised relative humidity29. The biennial pattern for HMPV reported from Germany30, Austria30 and Sweden31, is not seen in our study.In our study, enveloped viruses with peaks in the winter months, had an association with lower temperatures, dew points and humidity-ranges, and preference for higher humidity. A significant relationship with low temperature, dew point and high relative humidity have been reported for RSV18, and others reporting that enveloped viruses are better transmitted in low humidity32. For example, Iha Y et al. reported an inverse correlation with relative humidity for influenza A and a direct association with relative humidity for influenza B from Japan26. Since relative humidity is significantly associated only with HPIV-3 on logistic regression, the absolute value of humidity may be less important than fluctuation in humidity to the survival of the virus, particularly with viruses predominantly seen in the Autumn and Winter. It is possible that temperature and “humidity-range” could both have a combined impact on viral stability. In laboratory conditions, RSV survives for longer at lower temperatures, with infectivity falling rapidly as temperature increases. It is known that freezing and thawing leads to marked reduction in infectivity of RSV33. The preference for a narrow “humidity-range” and the change in relative humidity associated with changes in temperature34 could explain why viruses are not stable in frost-free freezers which have to warm up regularly to prevent accumulation of ice in the freezer compartment. This needs further investigation. HPIV 1-3, despite being enveloped RNA viruses, seems to prefer warmer temperatures (8.2 °C-9.4 °C) than the other enveloped RNA viruses. Due to its low incidence over our study period, the seasonality of HPIV-1 and HPIV-2 remain unclear though the GLM suggests it is around October-November. There is a possibility of a biennial seasonality for HPIV-1, similar to that reported in the US35, but the trend remains unclear due to the low prevalence in our study. Unlike HPIV-1 and 2, HPIV-3 increases when temperature is highest (9.4 °C) and when relative humidity is lowest (79.9%). HPIV-3 also appears to tolerate large variations in “humidity-range” (29.7%). Following logistic regression, HPIV-3 is associated with lower relative humidity. The seasonal peak for HPIV-3 in May35, its stability at lower relative humidity36 and higher temperatures36 have been reported previously. Our results indicate that most viruses prefer stable humidity (i.e. a lower “humidity-range”) through the day. This requires confirmation from further studies.Both RSV and IAV appear when the relative humidity is high (83% to 84%) in Edinburgh. Others have reported variable findings, with some reporting a positive association37. We hence attempted to tease this out by performing logistic regression on variables which showed an association in the univariate analysis. Since dew point was removed to prevent collinearity affecting the results of logistic regression, a new regression could in the future be performed to analyse the impact of dew point by removing temperature. There is a preference for stable “humidity-range” in most enveloped viruses except for HPIV-3. Following a visual inspection of the minimum and maximum for each variable and comparison with virus seasonality patterns for different viruses, there appeared to be a relationship between “humidity-range” and viral seasonality. Hence, we investigated the relationship between “humidity-range” and viral seasonality. It would be interesting to assess if a similar relationship is seen with other ranges such as “temperature-range”, “dew point-range”, etc in the future. Relationship with other variables appears to be virus specific. Therefore, we need models which can investigate multiple variables for each virus. Changes in weather, human behaviour and the immune system can all affect an agent’s seasonality and are all interlinked, as seen in the decrease in vitamin D in temperate climates due to the decrease in daylight hours (shorter/cloudier days). This also makes it difficult to describe the effect of weather on viral seasonality in isolation. Finally, the GLM used in this study has some disadvantages. The GLM used does not account for the lag period between people becoming infected and then seeking medical help, and it also cannot model the multiple peaks seen for rhinoviruses. Furthermore, average data was used to calculate the GLM. This has led to some of the curves not being as smooth as when yearly data is analysed , examining the effects of individual meteorological factors on viral seasonality can lead to an incomplete picture38. The further you are from the equator, winter seasonality and or association with monsoon rains is notable40. Our findings on “humidity-ranges” need confirmation in a tropical setting. The order of occurrence of different viruses in tropical and temperate regions needs to be investigated in relation to the parameters highlighted. Additionally, whether these peaks in viral activity change depending on the age of the population analysed could also be a topic for future research.Nonetheless, our findings on temperature, humidity and “humidity-range” could potentially help explain the variation in seasonality between temperate and tropical regions of the world. IAV is known to occur throughout the year in South East Asian countries closer to the equator In conclusion, non-enveloped viruses were present throughout the year. Enveloped viruses were more seasonal starting from IAV in January to RSV in December with other viruses in between. RSV, HPIV-3 and IAV and IBV showing the clearest patterns. RSV, IAV and IBV show a clear preference for colder temperature and dew point, as well as consistent humidity. In contrast, HPIV showed a preference for higher temperatures. Additionally, HPIV-3 demonstrated a significant tolerance to low humidity. Hence, meteorological factors like temperature, humidity and “humidity-range” have a significant effect on the incidence of the causative agents of the common cold and changes in meteorological factors could potentially predict the decline of one virus and the emergence of the next.st April 2009 to the 30th November 2015 were available. Small periods of weather data were missing in October 2011 and June 2015 (reasons unknown). We calculated means and ranges for each variable for each day. We also created weekly and monthly aggregations of the data. As rhinoviruses and HMPV testing only started from the 4th of September 2009, results for these two agents were only available after this date.Following discussion with the Met Office, Edinburgh Gogarbank weather station, which captures data on an hourly basis, was identified as the most representative weather station for the catchment areas examined in this study. We downloaded meteorological data collected by the Edinburgh Gogarbank weather station from the Met Office Weather Observation Website. Hourly air temperature (°C), air pressure (hPa), relative humidity (%) dew point (°C), wind speed and daylight length measurements from 1th-24th February 2011, samples were not tested for rhinovirus due to high work load. Between September 15th and December 9th 2010, IBV was not tested due to technical reasons. We removed repeat positive samples from the same individual collected within 14 days from analysis, resulting in a total of 52060 samples for analysis. We calculated the proportion of positive results (as %) for individual viruses for each day, to help compare results over time. We used daily data for statistical analysis to avoid loss of accuracy from averaging data. To identify seasonal trends, we plotted the proportion of positive results for each day against individual meteorological factors.We collated anonymised virological results on all forms of respiratory samples  tested at the Department of Laboratory Medicine, Royal Infirmary of Edinburgh. These included all samples from patients attending 8 hospitals around Edinburgh , and 126 primary care centres across in Edinburgh and Lothian, tested at the Department of Laboratory Medicine, Royal Infirmary of Edinburgh. We had results from 36,191 individual patients over the study period. Of these patients, 17,936 were male and 18,241 were female (information on gender was not available for 14). The age groups of the cohort are in Table Univariate analysis of continuous variables  were compared between days where one or more samples were positive for a given virus, and days where no samples were positive for that virus using two-sample t-tests. To determine the impact of multiple factors, multivariable logistic regression was conducted using those variables with a univariate p-value of <0.1 in the t-test and models were built using the “Enter (forced entry)” method on SPSS. Due to the likelihood of collinearity affecting results because of the close relationship between dew point and temperature, dew point was removed from the logistic regression.An analysis of variance (ANOVA) in conjunction with Tukey’s range test were done, to determine if the mean temperature, dew point, pressure, relative humidity, and “humidity-range” at which individual viruses were present were different between viruses. This was performed in conjunction with post-hoc analysis.et al.41. The GLM was calculated using daily incidence of each virus, averaged over the entire study period.Finally, seasonality of individual viruses identified in our study was confirmed with a generalised linear model (GLM) based on temperature following the method of Naumova Supplementary Figure 1


The microstructure of compression molded Estane 5703 films exposed to 11%, 45%, and 80% relative humidity and 70 °C for 1 and 2 months has been studied by small-angle neutron scattering (SANS), Fourier transform infrared spectroscopy (FTIR), gel permeation chromatography (GPC), and differential scanning calorimetry (DSC). Scattering data indicated increase of the interdomain distance and domain size with a higher humidity and longer aging time. GPC data showed a progressive shortening of polyurethane chains with increasing humidity and aging time. The shortening of the polyurethane chains caused a drop of the glass transition temperature of soft segments, and promoted crystallization of the soft segments during long-time storage of the aged samples at room temperature. FTIR showed a substantial increase in the number of inter-urethane H-bonds in the aged samples. This correlates with the increase of the hard domain size and the degree of phase separation as measured by SANS. The data collected reveals that the reduced steric hindrance caused by hydrolysis of ester links in polybutylene adipate residues promotes the organization of hard segments into domains, leading to the increase of domain size and distance, as well as phase segregation in aged Estane. These findings provide insight into the effects of humidity and thermal aging on the microstructure of aged polyester urethane from molecular to nanoscale level. Thermoplastic polyurethanes (PUs) are a very attractive type of segmented polymers that bridge the gap between rubbers and plastics . PUs havhard domains), greatly increasing the modulus and strength by acting as rigid fillers. Early investigations on aged Estane mainly focused on chemical degradation and cross-linking [et al. [et al. [The system of interest in the present work is a kind of polyester based PU, Estane 5703 , which has been widely used in adhesive for inks and lacquers, fabric coating, as well as magnetic media. Previous studies on the long term stability of Estane in thermal and humid environments revealed the destroying effect of humidity on the microstructure. As a common characteristic for PUs, the performance of Estane is not only determined by the chemical properties, but also strongly inﬂuenced by the microphase structure, in which the hard segments (HSs) segregate into domains  how the hard domain morphology and the degree of phase separation are affected by humid aging at elevated temperature; (2) what processes take place in the material on the molecular scale.In this work, the microstructural changes in Estane aged in humid air with low, median, and high RH at elevated temperature is studied by various techniques. The domain morphology has been analyzed by SANS. In order to increase the contrast and data quality, the soft domains of Estane were swollen by deuterated solvent following a method proposed by Mang . We haveEstane 5703 was obtained as pellets from the Lubrizol Advanced Materials, Inc., Westerlo-Ovel, Belgium. It is a segmented polyurethane copolymer, containing about 23 wt % HSs, prepared by the reaction of 4,4’-diphenylmethane diisocyanate (MDI) and poly (butylene adipate) (PBA) with 1,4-butanediol (BDO) as the chain extender . It cons3)2, and LiCl solutions produced relative humidity (RH) of 80%, 45% and 11%, respectively. The containers were placed in convection ovens operating at 70 °C for 1 and 2 months. After treatment, the films were kept at RT in sample storage box with a constant humidity of 30% for not less than 6 months before the microstructural characterizations. During this time, the excess absorbed water was largely removed from the samples aged at medium and high humidity, and the phase separation of hard and soft domains took place. Estane films stored in the same storage box at RT without thermal and humid aging were taken as the reference samples.The films (15 mm × 8 mm × 1 mm) were sealed in containers to maintain the specified temperature and humidity environment. Aqueous saturated KCl, Mg(NOYellow Submarine at the Budapest Research Reactor. The scattering intensity I(q) is measured as a function of scattering vector q = 4πsinθ/λ, where λ is the wavelength of the incident neutrons, and θ is half of the scattering angle. A mean neutron wavelength of 0.47 nm and 1.175 nm with 0.2 FWHM was produced by a multidisk mechanical velocity selector. Sample-detector distances 1.57 m and 5.50 m were used. By changing the wavelength and sample-detector distance, a q range of 0.08–2.5 nm−1 has been covered. In order to enhance the scattering contrast between hard and soft domains, all the samples were swollen by deuterated toluene for 48 h before measurements, and were measured in the swollen state. Immediately before the measurements, the swollen samples were transferred into quartz cells of path length 2 mm and ﬁlled with deuterated toluene. The schematic diagram of the setup can be found in the reference [q-resolution has been taken into account by convoluting the theoretical curve with the instrument resolution function.The SANS measurements were performed with the small-angle neutron scattering diffractometer eference . The acqeference , and the−1 were summed to achieve a good signal-to-noise ratio.FTIR was performed using NICOLET 6700 FTIR spectrometer  in single attenuated total reflection mode on ZnSe crystal. Sixteen scans with spectral resolution 1 cmw/v)) was used as internal standard. The data from a refractive index detector were collected and treated by using CSW 1.7 software . For the determination of molar masses, a universal calibration equation calculated from the data on polystyrene standards  was used.The GPC measurements were performed using a Biospher GMB 100 column  filled with 10 μm sorbent particles. Tetrahydrofuran , distilled and dried over molecular sieves (4 Å), was used as a mobile phase at a flow rate 1 mL/min. Small amount of toluene . Purge nitrogen was let through the DSC cell with a flow rate of 20 mL/min. The temperature of the equipment was calibrated with mercury, water and indium. The melting heat of indium was used for calibrating the heat flow. The samples were subjected to a cooling scan from 20 to −60 °C, holding at −60 °C for 5 min, followed by heating scan to 200 °C. Both scans were carried out at a rate of 10 °C/min. Glass transition temperatures were determined by Gaussian peak fitting to the numerical first derivative of the first heating thermogram.r/acor), where acor is the characteristic length over which the structural correlation decays [acor can be considered as the hard domain size, since in Estane 5703 the HS content is low (23 wt %) and consequently the hard domains are much smaller than the chords crossing the soft and hard domains [A and G are algebraic functions of the RHS and v, the hard sphere interaction radius and volume fraction, respectively. The first strong peak in the PY structure factor is related to the average distance between adjacent hard domains. Herein, we take the DAB combined with PY as an approximate empirical model to fit the scattering intensities by the equationK is a scale factor, and IB is a constant which accounts for the incoherent scattering. The model performs well over a q range from 0.1 to 2 nm−1. Form factors of compact particle models, spherical and ellipsoidal, even taking into account the size distribution, resulted in worse fitting in the high q range, while required more free fitting parameters. Therefore, we consider that the present model is reasonable for describing the scattering data and following the development of the relevant structural parameters [I(q) curves of Estane aged in 11%, 45%, and 80% RH air at 70 °C for 1 and 2 months are shown in q = 0.25 nm−1, which shifts towards smaller q with increasing aging time, indicating increase of the interdomain distance. Results of model fitting are shown in acor and RHS. After aging in 11% RH air for 1 month, the hard domain distance (2RHS) increases slightly and acor increases from 3.1 to 3.5 nm. After aging in 80% RH air for 1 month, 2RHS increases from 14.9 to 24.9 nm and acor increases from 3.1 to 3.9 nm. The development of the hard domain distance with the aging time and humidity is presented in q values indicates the appearance of a new population of scatters or phase segregation on the length scale above 30 nm (estimated by π/qmin). For this sample, the model used herein is no longer accurate due to the strong forward scattering, which masks the details of the domain structure.Polyurethanes containing 40 wt % or less HSs are known to have a discrete hard microdomain structure , which sn decays . The DAB domains ,20. In trameters ,21. The i.e., the inter-urethane and urethane-ester H–bonds. The absorption at about 3440 cm−1 corresponds to free N–H group, and the absorption band around 3340 cm−1 is assigned to H–bonded N–H group. After 1 month of aging, the IR data show no obvious variation on the N–H bands, only a slight shift to low frequency for the H–bonded N–H band . The absorbance of the shoulder increases with the humidity. For the sample aged in 80% RH for 2 months the C=O band splits into two peaks, located at 1726 and 1700 cm−1. The FTIR data of the aged samples reflect the packing changes of the polymer chains in hard and soft domains.The corresponding variation of the C=O stretching vibration band at 1760–1660 cmca. 1500 g·mol−1. A probable reason of this effect is presence of a layer of more degraded polymer at the surface of specimen used for GPC analysis. The number- and weight-average molecular weight of the polymer (Mn and Mw) aged in 80% RH air for 2 months decreased by 20 times as compared with that of the original sample.The original Estane sample exhibits a single broad GPC peak . Aging aTg,s) associated with the SSs is observed around −37.6 °C, followed by an endotherm with onset temperature (T2) at 33 °C. The Tg,s and T2 of the aged samples are lowered by a few degrees with increasing the RH and aging time relative to the original sample, except for the sample aged at the highest RH . The specific melting enthalpy (Δh2) grows gradually, which means that the degree of crystallinity of the molecular chains is increased due to thermal and humid aging. Another wide and weak endotherm is present at 70–100 °C for all the samples, and the corresponding onset temperature (T3) shifts to higher temperature with aging. In the sample aged at the highest RH for 2 months two melting endothermic peaks emerge, centered at about 42 and 11 °C, and an exothermic peak centered at about −18 °C. This phenomenon is the same as the samples aged for 1 month at the highest RH, however, the changes are much less pronounced. The detailed parameters are listed in DSC thermograms for the original and aged Estane are shown in et al. [q. In this experiment, we measured the structure change induced by aging, biased by the subsequent swelling of the material. Nevertheless, the trends of the domain size and interdomain distance variations are informative for the morphology development of the non-swollen aged materials.The studied polyurethane, Estane 5703, has a weak scattering power for neutrons because of the modest contrast between the hard and soft domains. Therefore, we employed swelling by a selective solvent, deuterated toluene, to increase the contrast. This way, the coherent scattering signal increased by one order of magnitude and allowed the use of accurate model fitting to reveal details of the microphase morphology. The early work by Nierzwick et al.  proved tRHS remained unchanged for Estane aged at RT in 100% RH for 2 months. The present study provides additional information on the behavior of Estane treated at 70 °C at different humidities. The structural changes can be explained in the following way. The chemical structure of Estane has a compositional heterogeneity of the HSs length distribution. Upon aging, HSs of shorter lengths prefer to dissolve into the soft domains [v/RHS3, is decreased for the aged Estane, while the domain distance increased, as shown in acor) and their distance (2RHS), it follows that a certain number of smaller hard domains disappear, being dissolved in the SSs matrix, while some of the larger hard domains grow even larger during the aging, as shown in In our previous study , the RHSDecrease of hard domains in annealed waterborne polyurethanes has been also seen in TEM investigations employing ruthenium staining . The domTg,s  for the HSs (MDI+BDO) is around 75 °C [Tg,h, and the thermal and humid aging results in the increase of the hard domain size as observed by SANS, hence the T3 increases with aging. During the second heating, there is no endothermic peak at around 40 °C, indicating that longer time is necessary for ordering of the molecular chains; previously the samples were stored at RT for about one year. This strongest endotherm can be assigned to the melting of crystallites of well-ordered SSs. The rate of crystallization during the RT storage is strongly dependent on the mobility of the polymer repeating units. Decrease of the molecular weight due to hydrolysis and the subsequent increase of SSs mobility during the aging lead to faster crystallization during storage. Therefore, the degree of crystallinity of SSs is increased upon aging.The glass transition temperature (nd 75 °C , therefoIn summary, the above findings prove that, with decreased molecular weight, organization of HSs into domains is promoted by reduced steric hindrance in aged Estane, which appears to be the main mechanism of the morphological changes observed by SANS. While it has been reported earlier that Estane containing deuterated plasticizer underoges stronger phase segregation upon aging ,9, new iTg,s, leading to higher degree of crystallinity of the SSs after long-time storage of the samples at RT following the aging. The FTIR, GPC, and DSC results all indicated, from different perspectives, that the hydrolysis of ester links on polymer backbone at 70 °C is the main reason for the structural changes during aging at humid conditions.The microphase morphology and chemical modifications of Estane 5703 aged at three different humidities at 70 °C for 1 and 2 months have been studied by multiple techniques. The scattering data showed a significant increase of the domain distances with humidity and aging time. For the sample aged in high humidity for 2 months, the number of inter-urethane H-bonds increased significantly, accompanied by the decrease of the urethane-ester H-bonds. It is consistent with the increase of hard domain sizes and the degree of phase separation as measured by SANS. GPC results showed progressive shortening of the polyurethane chains with increasing humidity and aging time. The decrease of molecular weight of the polyurethane chains caused a decrease of the


In this work, Bi2Se3 nanoplates were deposited under the same experimental conditions on different types of graphene substrates . Dimensions of the nanoplates deposited on graphene substrates were compared with the dimensions of the nanoplates deposited on mechanically exfoliated mica and highly ordered pyrolytic graphite flakes used as reference substrates. The influence of different graphene substrates on nucleation and further lateral and vertical growth of the Bi2Se3 nanoplates is analysed. Possibility to obtain ultrathin Bi2Se3 thin films on these substrates is evaluated. Between the substrates considered in this work, graphene grown on silicon carbide is found to be the most promising substrate for obtaining of 1–5 nm thick Bi2Se3 films.Knowledge of nucleation and further growth of Bi For instance, these include thermoelectric devices3 with Bi2Se3 being one of the best thermoelectric materials for near-room temperature applications. Other important applications are flexible composites for the use as anodes in lithium-ion batteries4 and active elements in high-performance optoelectronic devices5.Recently, the integration of graphene with other Dirac materials  is formed by five covalently bonded atomic sheets (Se-Bi-Se-Bi-Se). Bi2Se3 can be deposited on solid substrates by variety of methods, including molecular beam epitaxy (MBE)7, physical and chemical vapour deposition (PVD and CVD)8, thermal evaporation9.Bi10. Deep knowledge of Bi2Se3 nucleation and growth on various substrates is crucial for obtaining ultrathin continuous films of this material.Recent theoretical calculations performed for thermoelectrical properties of thin bismuth chalcogenide films predict that the reduction of the thickness of thin film below 10 nm may result in an increase of the existing thermoelectric figure of merit of this thin film by an order of magnitude2Se3 nanostructures or continuous films by PVD deposition on different surfaces remains challenging. Hampering of lateral growth of the nuclei by physical and chemical defects of the substrate surface prevents layer-by-layer growth of Bi2Se3 films11. For obtaining of ultrathin nanostructures or films, atomically-flat, chemically-inert substrates are favourable.The PVD method remains one of the most promising to be used at industrial scale as it is simple, cost-effective and allows for large-scale production. However, unlike the highly controllable MBE method, growth of ultrathin Bi2Se3 growth. Mica is chemically inert substrate with atomically flat surface ideal for Van der Waals epitaxy of Bi2Se316. Despite a significant lattice mismatch between the mica and Bi2Se3 (~25%), the reported thicknesses for Bi2Se3 nanoplates deposited by PVD technique on synthetic fluorophlogopite mica surface are as small as 2–3 nm15. PVD synthesis on natural muscovite mica, which is cheaper alternative to synthetic mica, results in formation of Bi2Se3 nanoplates with thicknesses down to approx.10 nm. Continuous Bi2Se3 thin films grown on mica have thicknesses starting from 50 nm16. The thermoelectric power factor of these thin films is comparable with the Bi2Se3 thin films obtained by MBE techniques12. As for graphene, a very small lattice mismatch between graphene and Bi2Se3 (~2.9%)1 makes it very attractive substrate to promote epitaxial growth of Bi2Se3 for the applications, where electrically conductive substrate is required.Mica and graphene are recognized as excellent substrates for Bi17. All these methods have advantages and drawbacks in terms of controllability, graphene quality, and costs. It is important to clarify how graphene substrates prepared by different techniques may differently affect the nucleation, further growth and properties of deposited on these substrates Bi2Se3.Graphene substrates can be prepared by different methods, including CVD synthesis of graphene followed by its polymer-assisted transfer onto the desired substrate, mechanical exfoliation of highly ordered pyrolytic graphite (HOPG), thermal decomposition of SiC substrates (G/SiC), and others2Se3 nanoplates deposited by catalyst-free PVD technique on prepared by different methods  graphene substrates. The dimensions of deposited under the same experimental conditions onto graphene substrates nanoplates are compared with the dimensions of the Bi2Se3 nanoplates deposited on mica and exfoliated multilayer HOPG flakes used as reference substrates. Factors affecting nucleation and growth of the nanoplates are discussed, and the possibilities for further obtaining of ultrathin Bi2Se3 films on graphene substrates are evaluated.This work is focused on a study of growth dynamics of Bi2Se3 nanoplates deposited on as-transferred and post-annealed CVD graphene, G/SiC, exfoliated multilayer HOPG flake and mica substrates.Figure 18, normally the most energetically favourable orientation of Bi2Se3 nanoplates is with c-axis oriented perpendicularly  to the commonly used for Bi2Se3 deposition substrates as graphene9, SiO213 and mica15. Diffraction peaks of XRD patterns obtained for the continuous Bi2Se3 thin films formed from the coalesced nanoplates on as-transferred CVD graphene surface , and are affiliated to a group (0 0 3n). This proves the growth of the Bi2Se3 layers in a single-crystalline phase with c-axis oriented perpendicularly to the substrate.Due to the intrinsically anisotropic bonding natureace Fig.  can be i2Se3 nanoplates are Se-passivated20, resulting in much faster lateral growth of nanoplates than vertical due to diffusion of Bi and Se adatoms to the edges of the nanoplates. However, it is found in this study that the ranges of thicknesses and surface areas of the deposited under the same synthesis conditions nanoplates differ greatly for different graphene substrates. For relative comparison nanoplates dimensions, a parameter “surface-to-thickness ratio” (S/T) is used. For the S/T calculation, the surface area of the upper facet of a nanoplate is divided by the averaged thickness of this nanoplate. Both surface area and thickness are determined by the analysis of AFM scans of the nanoplates. The domination of the lateral growth direction over the vertical is indicated by higher values of S/T parameter. Data on the ranges of nanoplates thicknesses, surface areas and S/T ratios are summarised in Table Generally, the top surfaces of the Bi2Se3 nanoplates deposited on mica (considered as a reference substrate) are below 25 nm, and 17% have thicknesses between 6 and 10 nm . Most probably, this effect can be attributed to the influence of the large domain boundaries, formed during the graphene growth on the copper substrates and remaining on the CVD graphene surface after its post-annealing. It should be noted that the nanoplates with thicknesses below 10 nm and large S/T parameters are crucial for formation of ultrathin continuous Bi2Se3 thin films. The conclusion is that both as-transferred and post-annealed CVD graphene substrates may be useful for obtaining of thick (~150–300 nm) nanostructured Bi2Se3 films. For obtaining of ultrathin Bi2Se3 films, alternative to CVD graphene substrates as, for instance, G/SiC should be considered.Post-annealing of as-transferred CVD graphene leads to the partial healing of the chemical surface defects on the graphene surface2Se3 nanoplates.Using of G/SiC substrates allows to avoid the influence of such surface defects as domain boundaries or transfer-caused defects (polymer residues etc), increasing the number of nucleation centres and hampering the lateral growth of the Bi2Se3 islands formed on the surface of G/SiC are found to be below 5 nm  and post-annealed CVD graphene (81% of S/T is below 30), the S/T parameters for the Bi2Se3 islands formed on the G/SiC substrates are starting from 30 and are spread up to 2500  higher S/T values. 45% of the Bi2Se3 islands on G/SiC substrates are having higher S/T parameters than the largest S/T of mica-supported nanoplates  surface areas, thicknesses up to 20 nm  and continues epitaxially. Lateral growth of Bi2Se3 islands is limited by the edges of the terraces. Similar growth mechanism is observed on the mechanically exfoliated multilayer HOPG flakes, where the deposited Bi2Se3 nanoplates were lined-up along the graphite terraces. However, the average thickness of the nanoplates deposited on HOPG flakes (~21 nm) is by at least 5 times larger than the average thickness of the nanoplates deposited on G/SiC (~4 nm). Possibly, this may be explained by specific issues of the G/SiC substrates as double-layer graphene with higher surface energy along the edges of the terraces and strain-less graphene surface. Thus, the G/SiC substrates are found to be the most perspective for obtaining of Bi2Se3 thin films with thicknesses below 5 nm. However, the drawback of using the G/SiC substrates is limitation by the wafer sizes (currently 6” available) and costs.Graphene grown directly on the SiC substrate by its thermal decomposition method is found to be a very good substrate for 2D growth of Bi2Se3 nanostructures were synthesized by modified catalyst-free vapour-solid deposition method reported elsewhere35 using a quartz tube furnace . The source material  was placed in a quartz boat in the centre of the furnace tube, while the substrate was located downstream relative to the source material.Bi2, pumped down to the base pressure 5 Torr and sealed. The centre of the quartz tube was heated at a rate of 13 °C/min to 585 °C. The substrate temperature for Bi2Se3 deposition was 410–450 °C. The temperature was held constant (at 585 °C) for 15 minutes for the initial formation of the nanoplates. Than the furnace was naturally cooled down to 465 °C. During this cooling process, the further growth of the nanoplates together with their enrichment with Se occurs. At furnace temperature of 465 °C, the furnace tube was filled with N2 to atmospheric pressure, sealed and cooled down to room temperature. For the study of the initial formation stages of the nanoplates, the substrate was taken out from the furnace tube at 585 °C and rapidly cooled down under the N2 gas flow.For the synthesis, the furnace tube was flushed with N2 flow. Muscovite mica sheets (Agar Scientific) were mechanically cleaved immediately before the synthesis process.Glass substrates were washed in 2-propanol and distilled water, followed by drying under N20.Monolayer graphene was prepared by chemical vapour deposition method on copper foil and transferred on the quartz/glass substrates using polymer-assisted technique described elsewhere2/Ar flow .Glass substrates with transferred by polymer-assisted technique graphene was annealed consequently in air  and HTM 8805).HOP graphite flakes were mechanically exfoliated from the macro piece of HOPG and transferred onto the pre-cleaned glass substrate using adhesive tape (3M2. Chips are then cleaned using RCA 1 and 2 processes and dipped in HF (3%) for 30 seconds. The samples are loaded in a reactor for epitaxial graphene growth equipped with an inductive heater. Graphene growth takes place in argon atmosphere (P = 600 Torr) at a temperature of 1700 °C for 5 min. After growth, optical microscopy reveals a typical surface coverage of >90% monolayer graphene.Monolayer graphene was grown on the Si-face of SiC using thermal decomposition of 7 × 7 mm SiC substrates . The 3” SiC wafer is diced into 7 × 7 mm chips, which are cleaned using solvents  and blow-dried with N2Se3 nanostructures and nanostructured layers were inspected using field emission scanning electron microscope (SEM) Hitachi S-4800 equipped with an energy-dispersive X-ray (EDX) analyser Bruker XFLASH 5010 and atomic force microscopes Asylum Research MFP-3D and Bruker Dimension ICON.Morphology, thickness, structure and composition of obtained BiFor the statistical analysis, 5 to 9 AFM scans of size 20 × 20 μm each were obtained at different locations on the sample. The number of measured nanoplates was 110–150 for each type of the substrate with an exclusion of exfoliated graphene substrate. Due to the technical difficulties, 6 nanoplates were measured on the surface of exfoliated graphene.XRD characterization of the thin films was performed by a Philips X’Pert MRD with a Cu Kα radiation source.