Our findings demonstrated that the hepatic epigenetic landscape can be affected by both nutritional status and dietary carbohydrates in trout. The mechanism underlying the setting up of these epigenetic modifications has now to be explored in order to improve understanding of its impact on the glucose intolerant phenotype in carnivorous teleosts.The rainbow trout, a carnivorous fish, displays a ‘glucose-intolerant’ phenotype revealed by persistent hyperglycaemia when fed a high carbohydrate diet (HighCHO). Epigenetics refers to heritable changes in gene activity and is closely related to environmental changes and thus to metabolism adjustments governed by nutrition. In this study we first assessed in the trout liver whether and how nutritional status affects global epigenome modifications by targeting DNA methylation and histone marks previously reported to be affected in metabolic diseases. We then examined whether dietary carbohydrates could affect the epigenetic landscape of duplicated gluconeogenic genes previously reported to display changes in mRNA levels in trout fed a high carbohydrate diet. We specifically highlighted global hypomethylation of DNA and hypoacetylation of H3K9 in trout fed a HighCHO diet, a well-described phenotype in diabetes.  Oncorhynchus mykiss) is considered to be a glucose-intolerant species, displaying persistent hyperglycaemia after intake of a high-carbohydrate diet12456et al.in vivo9in vitrog6pcb2 ohnologs; ohnologs are paralogs formed by a whole genome duplication event) by dietary carbohydrates or insulin in vitro was hypothesized to be involved in the establishment of the glucose-intolerant phenotype in trout fed a high carbohydrate diet.The rainbow trout  in rainbow trout. Indeed, only one study has reported global DNA hypomethylation in the livers of trout fed a high carbohydrate diet without prospecting for histone mark changes at glucose metabolism-related gene locii.e, fasted vs fed fish with or without dietary carbohydrates) affected global epigenome modifications by targeting DNA methylation and histone marks previously reported to be affected in metabolic diseases linked to disturbance of glucose metabolism . As changes in the global epigenome induced by environmental modifications may have a potential epigenetic impact at the gene regulation level, we then examined whether dietary carbohydrates could affect the epigenetic landscape at duplicated gluconeogenic gene loci previously reported to display mRNA level changes in trout fed a high carbohydrate diet . This last step would also provide new information to improve understanding of how duplicated genes encoding the same enzyme, but which are regulated differently by dietary carbohydrates, are regulated.There is little information concerning epigenetic modifications at global and target gene levels mediated by nutritional status or by dietary carbohydrates and their physiological consequence  in fasted animals and trout fed the NoCHO diet than in trout fed the HighCHO diet.Classical ChIP -qPCR analyses were performed to profile the patterns of H3K4me3, H3K9me3, H3K9ac and H3K36me3 in relation to the nutritional status in trout along /g6pcb2b , pck1 an1/fbp1b2  5′-upstr1/fbp1b2  revealedet al.910g6pc paralogous genes, fbp1b ohnologs and pck1). DNA methylation levels of g6pca, fbp1b1, fbp1b2 and pck1, remained stable whatever the nutritional condition  or the CpG site analysed  and in trout fed the NoCHO diet compared to fasted trout or trout fed the HighCHO diet .Using targeted Next Generation Bisulfite Sequencing (t-NGBS), we analysed DNA methylation levels at specific CpG sites along the 5′-upstream region of gluconeogenic genes whose mRNA levels had previously been shown to be affected by nutritional status and/or dietary carbohydrates by Marandel analysed . A signii.e. the feeding status  and then also according to the content of the diet in dietary carbohydrates (0% or 30%). We also examined the data in relation to the resulting hyperglycaemia observed in trout fed the HighCHO diet. Indeed, increase in blood glycaemia and its persistence can be considered an environmental factor and thus likely to induce epigenetic modifications.Epigenetics refers to heritable changes in gene activity that are not caused by changes in the DNA sequence. Since one of their first definitions by WaddingtonWe initially considered epigenetic modifications induced by the feeding status . The global analysis of the hepatic epigenome first revealed that H3K9me3 and H3K9ac levels were affected. Indeed, these two histone modifications were increased in trout fed the NoCHO diet compared to fasted trout, whereas H3K4me3, H3K36me3 and the DNA methylation level remained stable. To our knowledge, global epigenome changes in the livers of fasted and refed healthy animals have rarely been studied. These findings together suggested that the global hepatic epigenome was remodelled by the feeding status, and demonstrated that the global epigenome could be modified in the short term as fish were fed for only four days before sampling.fbp1ba/fbp1b2) this modification was enriched around the TSS  in mammals and in diabetes at gluconeogenic gene loci22242528293033 the TSS . These fin utero nutrition in mammals22g6pca, fbp1b1, fbp1b2 and pck1 whatever the feeding status of the animal. On the other hand, the methylation level was affected at specific CpG sites along g6pcb1b, g6pcb2a and g6pcb2b by the feeding status , all of them displaying higher methylation levels in fed trout than in fasted trout. These dynamic changes in DNA methylation at specific loci could not be correlated with changes in mRNA level described on the same samples by Marandel et al.g6pcb1b) or up-regulated (for g6pcb2a and g6pcb2b) in trout fed the NoCHO diet compared to fasted troutg6pcb2 ohnologs displayed higher DNA methylation levels than g6pb1 ohnologs, and that g6pca was at an intermediate level between both pairs of ohnologs. This strongly confirms that these genes suffered a sub- or neo-functionalisation after the Ss4R at the epigenetic regulatory level which may have led to or participated in the previously described9104041Finally, DNA methylation at specific gluconeogenic gene loci has previously been shown to be involved in the regulation of their expression mainly in cancer3536i.e. fasted vs fed state, affected the hepatic epigenetic landscape at the global level as well as at the level of some gluconeogenic genes even in the short term (only 4 days of starvation and feeding).These findings together showed for the first time that feeding status in trout, in vitro4620We also looked at the effects of dietary carbohydrates on the hepatic epigenetic landscape. Our results demonstrated that global DNA methylation and H3K9ac levels decreased in trout fed the HighCHO diet compared to trout fed the NoCHO diet, suggesting that dietary carbohydrates may directly or indirectly affect these two epigenetic marks. Hypomethylation of DNA and hypoacetylation of H3K9 have already been fully documented in metabolic diseases, particularly in diabetes . These findings demonstrated for the first time nutrient-induced promoter-specific methylation at g6pc loci in trout. In addition, when CpG methylation was affected in trout fed the HighCHO diet it was also in favour of a hypomethylation compared to trout fed the NoCHO diet and interestingly displayed the same level of methylation as in fasted trout. This pattern reflected what happened at the global epigenome level at these specific loci (see above). As explained above, DNA methylation acts in general to repress gene transcription. g6pcb2 ohnologs have previously been shown to be up-regulated in trout fed the HighCHO diet and this was believed to contribute to the glucose intolerant phenotypevia the de-repression of g6pcb2 ohnologs. Functional analysis of affected CpG sites must now be tested to explore the relationship between their methylation levels and the expression levels of genes.Finally, dynamic changes in DNA methylation at specific CpG sites were monitored at g6pcb2 ohnologs previously believed to contribute to the initialisation of the glucose intolerant phenotypeTaken together, our findings showed for the first time that changing dietary carbohydrate content in the trout diet has an effect on DNA methylation level at specific gluconeogenic gene loci and may be involved in the regulation of We reported here for the first time remodelling of the hepatic epigenetic landscape by nutritional status and dietary carbohydrate content of the trout diet both at global and target gene levels. We demonstrated that global hypomethylation occurred in trout fed the HighCHO diet which in particular mimicked a phenotype described in diabetes. DNA hypomethylation can lead to dramatic consequences such as induction of expression of oncogenes or miRNAInvestigations were conducted according to the guiding principles for the use and care of laboratory animals and in compliance with French and European regulations on animal welfare . This protocol and the project as a whole were approved by the French National Consultative Ethics Committee.et al.−1 for fasted trout and those fed the NoCHO diet and around 1.8 g.L−1 for trout fed the HighCHO diet, data previously published by Marandel et al.et al.Data published in this manuscript were obtained from the analysis of biological material used in an article previously published by Marandel Gut content of each fish was systematically checked to confirm that the fish sampled had effectively consumed the diet. The livers were dissected and immediately frozen in liquid nitrogen, and stored at −80C until further analysis.Livers of three fish per condition were analyzed. DNA extraction was performed on 10 mg of tissue as previously describedThe livers of three fish per condition were analysed . One hundred mg of tissue was homogenized using Precellys24  in 2 ml tubes containing 1 ml of TEB ) and four 2.8 millimeter ceramic beads, 2 × 10 s–15 s off at 5,000 rpm. Samples were then left to stand at 4 °C for 20 min before centrifugation at 2,000 rpm and 4 °C for 10 min. Pellets were then resuspended in 0.5 N HCl containing 10% glycerol  and incubated on ice for 30 min (with vortexing every 10 min). Samples were centrifuged at 12,000 rpm at 4 °C for 5 min. Three volumes of iced acetone were added to the supernatant and precipitation was performed at 20°C overnight. Samples were centrifuged at 12,000 rpm at 4 °C for 5 min and iced cold acetone was used to wash the pellets 3 times. Pellets were resuspended in distilled water at 60 °C for 1h. Total protein concentration was determined using the Bio-Rad protein assay kit . Lysates  were subjected to SDS-PAGE and western blotting using the appropriate antibody on 15% gel  for 40 min at 100 V and then 56 min at 150 V.The same antibodies were used for both western blotting and ChIP analysis. Anti-IgG (C15410206), Anti-H3K9me3 (C15410056), anti-H3K4me3 (C15410003), anti-H3K36me3 (C15410192) and anti-H3K9ac (C15410177) were purchased from Diagenode (Belgium) and anti-H3 (ab1791) from Abcam (UK).http://www.urogene.org/cgi-bin/methprimer/methprimer.cgi)in silico bisulfite converted as a reference. Alignment BAM files were then sorted by target and condition using BAMtoolsDNA extracted from the three livers per condition for global DNA analysis was used for gene-specific analysis of DNA methylation. Bisulfite conversion was performed using EZ DNA Methylation-Gold™ Kit  on each DNA sample according to the manufacturer’s instructions. PCR primers  for the g6pcb2a and g6pcb2b), the same primer set amplified both genes as we were not able to design specific primers because of the high percentage of identity between the ohnologous sequences. The assays were performed using a reaction mix of 6 μl per sample, each of which contained 2 μl of diluted DNA template, 0.3 μl of each primer (5 μM), 3 μl Light Cycler 480 SYBR® Green I Master mix and 0.7 μl DNAse/RNAse free water . The PCR protocol was initiated at 95 °C for 10 min for initial denaturation of the DNA and hot-start Taq-polymerase activation, followed by 45 cycles of a two-step amplification programme . Each PCR assay was performed in duplicate. Results were expressed as % input and fold enrichment over IgG  on a rotating wheel. The cross-linking reaction was quenched by adding 0.12 M of glycine 5 min at RT. Samples were washed twice in 8 ml PBS and centrifuged for 10 min at 4 °C and 2000 g. Pellets were resuspended in 1 ml iced-cold PBS. Samples were homogenised using Precellys® 24  fitted with Cryolys® in 2 mL tubes containing six zirconium beads (2.8 millimiter), for 2 × 10 seconds, separated by 1min off, at 5,000 rpm. Homogenized samples were centrifuged for 10 min at 4 °C and 1300 g. Pellets were resuspended in 10ml classical lysis buffer  and 20mM NaBu) for 15 min at 4 °C on a rotating wheel and then centrifuged for 5 min at 4 °C and 1300 g. Extracted chromatin was resuspended in 130 μl of shearing buffer , left to stand on ice for 10 min and transferred to 1.5 ml Bioruptor® Plus TPX microtubes . Shearing was performed in Bioruptor® Plus : 7 cycles, 30 sec ON-30 sec OFF, High power. Sheared chromatin was centrifuged for 10 min at 4 °C and 14000 g and 10 μl was used to check the efficiency of the shearing (smear size expected between 150 bp and 500 bp). The remaining chromatin was diluted at 1:8 in immunoprecipitation (IP) buffer . Ten μl of Dynabeads® Protein A  were coated with 3 μg of antibody  diluted in 90 μl IP buffer for 2 h at 4 °C. Coated beads were resuspended in 100 μl of diluted chromatin overnight at 4°C. Beads were washed four times in IP buffer, and once in 100 μl TE buffer . Immunoprecipitated chromatin was finally eluted in 150 μl of ChIP elution buffer , decrosslinked for 2 h at 68 °C, 1300 rpm, and resuspended in a 900 μl final volume of ChIP elution buffer. Ten μl of diluted chromatin was also decrosslinked to assess input. DNA from immunoprecipitated chromatin and input were isolated by classical phenol-chloroform extraction and resuspended in a 15 μl final volume of DNAse-free water. The Roche Lightcycler 480 system was used for qPCR . The primer sequences used are listed in  IgG see . NormaliHow to cite this article: Marandel, L. et al. Remodelling of the hepatic epigenetic landscape of glucose-intolerant rainbow trout (Oncorhynchus mykiss) by nutritional status and dietary carbohydrates. Sci. Rep.6, 32187; doi: 10.1038/srep32187 (2016).


Noble metals can promote the direct participation of lattice oxygen of very stable oxide materials such as aluminum oxide, to oxidize reactant molecules, while the fundamental mechanism of noble metal catalysis is elusive. Here we report that a single atom of rhodium, a powerful noble metal catalyst, can promote the transfer of five oxygen atoms to oxidize carbon monoxide from a nine-atom rhodium–aluminum oxide cluster. This is a sharp improvement in the field of cluster science where the transfer of at most two oxygen atoms from a doped cluster is more commonly observed. Rhodium functions not only as the preferred trapping site to anchor and oxidize carbon monoxide by the oxygen atoms in direct connection with rhodium but also the primarily oxidative centre to accumulate the large amounts of electrons and the polarity of rhodium is ultimately transformed from positive to negative. Noble metals can promote the stable lattice oxygen of metal oxides to take part in oxidation reactions. Here, the authors report the preparation and reactivity of rhodium-aluminum oxide cluster ions, in which the rhodium ion promotes the transfer of five oxygen atoms from a nine-atom parent cluster to oxidize carbon monoxide. The oxygen atom transfer (OAT) from metal oxide clusters to small molecules is one type of extensively studied reactions1218222324253O5+ and PtAl3O7− can transfer at most two oxygen atoms to oxidize CO 262O3+ and Al4O7− are all oxygen-rich species. Identification of multiple OAT from a single Rh-atom-doped cluster to reactant molecules is an important step to understand the participation of lattice oxygen promoted by noble metals. This gas-phase study that a nine-atom rhodium–aluminum oxide cluster oxidizes five CO molecules is a sharp improvement in the field of cluster science and provides a strictly molecular level understanding of the fundamental mechanism of noble metal catalysis in the related condensed phase.Atomic clusters are considered as the intermediate matter to bridge atoms and their bulk counterpart111213141516 CO refs . Here we2Om+ (m=2–6) cluster ions were generated by laser ablation of a mixed-metal disk compressed with Rh and Al powders. The generated RhAl2Om+ cluster ions were mass-selected, cooled and then interacted with N2 and CO in an ion trap reactor, as shown in 2O6+ with 150 mPa N2 (2 adsorption (RhAl2O6N2+) and N2/O2 exchange (RhAl2O4N2+) products were generated. Generation of RhAl2O4N2+ suggests the possible presence of superoxide (O2−·) or peroxide (O22−) unit in RhAl2O6+ (RhAl2O6+ + N2 → RhAl2O4N2+ + O2). In sharp contrast, on the interaction of RhAl2O6+ with CO 2+  clusters could also react with CO to generate products, from RhAl2Om−1+ to RhAl2O+  on the interaction of RhAl2Om+ (m=6–2) cluster ions with CO can be well fitted (2Om+ (m=6–2) with CO are 4.9±1.5 (m=6), 6.2±1.9 (m=5), 1.6±0.5 (m=4), 6.9±2.0 (m=3) and 2.4±0.7 (m=2), which correspond to the reaction efficiencies2O5+ are more reactive towards CO oxidation than clusters with even number of oxygen atoms such as RhAl2O6+.The strong signals that can be assigned as RhAlO4(CO)2+  on the ilacement , and fur RhAl2O+  and 2b,dl fitted  and the 2Om+ (m=2–6) are shown in 2O6+ + 5CO → RhAl2O+ + 5CO2) is highly exothermic (−9.00 eV). The low-lying energy isomers of clusters RhAl2Om+ (m=6–1) are provided in 2O6+ is in the triplet spin state 2+) on the interaction of RhAl2O6+ with N2 and CO 2+) than RhAl2O4N2+ observed in the experiment  can trap CO tightly at the first step  for CO oxidation by PRhAl2O5+. This step is to subject to a barrier of 1.26 eV. The subsequent steps follow a nearly downhill pathway characterized by small barriers to yield RhAl2O4+ and CO2  is often considered to be the crucial step in oxidation reactions3O8−22 is not enough to promote the dissociation of the O2 2− unit, while the Au dimer in Au2VO4− can promote O2 2− unit dissociation or direct participation in CO oxidationPRhAl2O5+ can promote the dissociation of the O2−· unit and the process is much more favourable than CO2 desorption . Direct oxidation of CO by the O2−· unit in PRhAl2O5+  is of great importance in CO oxidation by Au-doped clusters222324M bond. In sharp contrast, natural charge analysis demonstrates that after the transfer of four oxygen atoms from RhAl2O6+ to CO, the Rh atom is still positively charged  after the oxidation of the second CO due to the dissociation of the O2−· unit  in RhAl2O2–6+ but at the same time is relatively weak to deliver oxygen atoms to oxidize CO (O–CO: 5.52 eV)in situ X-ray diffraction and transmission electron microscopy, Stierle and colleaguesThe positively charged metal centre is crucial to provide not only the characteristic site for the adsorption of CO ref.  but also +0.53e,  in produ2O+ cluster is a linear structure (Al–O–Al–Rh) and the single oxygen atom is sandwiched between two Al atoms. This structure with triplet spin state has been confirmed to be the lowest energy isomer of RhAl2O+ by more accurate CCSD(T) calculation . The electron configuration calculations (4d8.755s0.645p0.03) provide solid evidence for the negatively charged Rh in RhAl2O+. Thus, the oxidation state of Rh changes from +2.5 in RhAl2O6+ to −1 in RhAl2O+ , covering from cationic to anionic, and it is the driving force to accumulate the electrons that are stored originally in the released oxygen atoms and promotes the unique oxidative reactions to proceed.The negatively charged Rh has been theoretically predictedculation . The twotion μB, , indicat2O6+. This study leads a leap ahead towards OAT reactions in the field of cluster science and represents an important step to understand the participation of lattice oxygen promoted by noble metals. The preferable Rh–O rather than the reductive Rh–Al bond formation together with the capability of Rh to accumulate the large amounts of electrons are crucial factors to drive the unique reactions. This gas-phase study reveals the molecular-level origin for the puzzling experimental observation that trace amounts of Rh can promote the reactivity of lattice oxygen of Al2O3  in the presence of O2 (0.4%) seeded in a He carrier gas with a backing pressure of 6.0 standard atmospheres. The cluster ions of interest were mass selected using a quadrupole mass filter and then entered into a linear ion trap (LIT) reactor, where they were cooled by collisions with a pulse of He gas and then interacted with a pulse of 5% (for reactions RhAl2Om+ (m=2–5) + CO) or 2% (for reaction RhAl2O6+ + CO) CO seeded in He for around 0.6∼1.1 ms. The temperature of cooling gas (He), reactant gases (CO or N2) and the LIT reactor was around 298 K. The cluster ions ejected from the LIT were detected by a reflector time-of-flight mass spectrometer. The details of running the time-of-flight mass spectrometerThe Rhk1) of cluster reactions in an ion trap reactorIR is the signal intensity of the reactant cluster ions, IT is the total ion intensity including product ion contribution, kB is the Boltzmann constant, T is the temperature (∼298 K), tR is the reaction time and P is the effective pressure of the reactant gas in the ion trap reactor.2O5+ + CO, the relative ion intensity of RhAl2O5+ is the reactive component generated in the experiment and the unreactive component is not included. The unreactive component of RhAl2O5+ could be well-fitted by equation (4)To calculate the reaction efficiencies (the possibilities of reaction on each collision), the collision rate constants were calculated on the basis of the surface charge capture model developed in the literaturexinert is the relative intensity of the unreactive component of RhAl2O5+ and k1 is the pseudo-first-order rate constant of the reactive component of RhAl2O5+. The xinert was determined to be about 12%, indicating that the experimentally generated RhAl2O5+ may have isomers that are not or less reactive with CO.in which 2O6+). To find an appropriate functional for the Rh–Al–O system, the bond dissociation energies of Rh–O, Rh–C, Al–O, O–O, Rh–Al and O–CO were computed by various functionals and compared with available experimental data .Supplementary Figures 1-13 and Supplementary Tables 1-2


The composition, bond state and coverage fraction of the boron carbide coating on the diamond particles were investigated. The boron carbide coating prefers to grow on the diamond (100) surface than on the diamond (111) surface. A stoichiometric B4C coating completely covered the diamond particle after maintaining the raw mixture at 1200 °C for 2 h. The contribution of the boron carbide coating to the oxidation resistance enhancement of the diamond particles was investigated. During annealing of the coated diamond in air, the priory formed B2O3, which exhibits a self-healing property, as an oxygen barrier layer, which protected the diamond from oxidation. The formation temperature of B2O3 is dependent on the amorphous boron carbide content. The coating on the diamond provided effective protection of the diamond against oxidation by heating in air at 1000 °C for 1 h. Furthermore, the presence of the boron carbide coating also contributed to the maintenance of the static compressive strength during the annealing of diamond in air.A boron carbide coating was applied to diamond particles by heating the particles in a powder mixture consisting of H Diamond has the highest hardness and excellent thermal conductivity, making it useful for reinforced composite materials, such as diamond compacts, which are potential candidates for use in a variety of applications, such as drill bits, saw blade segments, grinding wheels, cutting and polishing tools, and heat sinks for electronic devices12342O3) has many advantageous properties that make it useful for oxidation protection applications. B2O3 possesses low oxygen permeability, higher fluidity and good wettability on carbon materials below 1000 °C, which results in self-healing coatings on carbon materials6789102O3 causes the glass to swell and crumble, which can result in coating spallation at room temperature due to glass swelling or spallation during heating due to moisture release2B4O7•10H2O) is a common substitute. However, the usefulness of borate glasses is also limited because hydrated borates are highly volatile. Therefore, glass depletion may occur at relatively low temperatures in moist environments1415et al.Boron oxide (B4C) coatings are of interest due to their potential use for improving the oxidation resistance of diamond. The oxidation of B4C occurs at approximately 700 °C and forms a B2O3 oxygen barrier4C is a refractory hard material that is insoluble in water and chemically inert below 700 °C 4C requires a high synthesis temperature and long holding time19204Cet al.4C coating on diamond particles using a mixture of B and H3BO3 as the boron sources. The nucleation of B4C was obtained after 2 hours of holding at 1050 °C, and complete coverage of the diamond particles with B4C was achieved after 6 hours of holding at 1150 °C. However, the product quality of the coating  has yet to be report. Hu and Kong4C coating on diamond particles using the same method that was reported in ref. 4C coating on different surfaces of a single crystal diamond has not been previously studied. Moreover, the influence of the boron carbide coating on the enhancement of the oxidation resistance of diamond particles has not been previously investigated.Boron carbide . Therefore, the coating can impede corrosion of diamond from the oxygen in air at a high temperature.In this study, the ability of a boron carbide coating to protect diamond and improve its oxidation resistance has been studied. During annealing of the coated diamond in air, a B4C, B13C2 or B10C structure. The following XPS and Raman results suggest that the phase structure of the coating is B4C. In addition, as shown in 4C phase. The XPS C1s spectra for the boron carbide coatings on the (100) surface of the diamond particles synthesized at different temperatures are shown in 151526The chemical bonding and composition of the boron carbide coatings on the (100) surface of diamond particles were studied by XPS . The com−1. A high-intensity peak located at 1333 cm−1, which was observed in all the spectra, corresponded to diamond−1 were assigned to vibrations of the principal structural elements, icosahedra and three-atom linear chains in the B4C crystal29303132−1 decreased as the synthesis temperature increased and completely disappeared when the synthesis temperature reached 1200 °C. These peaks are related to chain-icosahedral linkages of amorphous boron carbide−1 and D peak located at 1350 cm−1Raman analysis was performed using a confocal system. The Raman laser was forced on the (100) surface of the diamond particles to characterize the binding states of the boron carbide coating. 343534352O3 structure, and the high-intensity peak corresponding to the single crystal diamond substrate  was partially truncated. In addition, the grain size calculated using the Williamson-Hall plot in All the samples were annealed at 1000 °C in an air atmosphere at a heating rate of 5 °C/min to characterize the oxidation resistance. The XRD patterns for the uncoated (A0) and boron carbide coated (A1-3) diamond particles after annealing are shown in −1, the Raman spectra for all the coatings after annealing in air contain two sharp peaks at 1360 and 1400 cm−1, which are assigned to B-O bonds4C crystal located in a range from 400 to 1200 cm−1 disappeared, which suggests that B4C was oxidized during the annealing process. This result confirms the XRD (−1) rather than amorphous boron carbide peaks (269 and 325 cm−1) were observed in the Raman spectra for samples A1 and A2. The use of a higher synthesis temperature to prepare the boron carbide coating results in a lower contribution from the D and G peaks after annealing. The D and G peaks were not observed in the Raman spectrum for sample A3, which indicates that amorphous boron carbide was also oxidized during the annealing process, and a larger amount of amorphous boron carbide contributed to the formation of amorphous carbon. Furthermore, the formation of amorphous carbon restricts grain growth of the B2O3 crystal during annealing, resulting in a refined B2O3 crystal grains in A1, as shown in To further investigate the binding states of the coatings on the diamond particles after annealing in air, Raman spectroscopy was employed see . In addi the XRD  and XPS  the XRD  results.2O3. During holding at 1000 °C, the coated diamonds exhibited a slow decrease in their weight due to evaporation of B2O3. When the protective B2O3 completely evaporates, the samples exhibited a significant decrease in weight due to the oxidation of diamond. As shown in Detailed thermal analyses of all the samples were performed using thermal gravimetric analysis (TGA). The temperature for the TGA was programmed to increase from room temperature (20 °C) to 1000 °C at a rate of 5 °C/min, and then, the temperature was maintained at 1000 °C for 4 h. The results of the TGA measurement for the uncoated (D0) and boron carbide coated diamond (D1-3) particles are shown in –1 in an air atmosphere. As shown in 2O3 coating, and no oxygen corrosion pit was observed.The static compressive strengths of one diamond particle from each of the samples are shown in 3BO3), and magnesium (Mg). After holding the raw mix at 1200 °C for 2 hours (Sample D3), the coating on the diamond exhibited a stoichiometric B4C crystal structure without any impurities or amorphous structures. The overall reaction can be expressed as394041In combination with the XRD , XPS Fi and RamaThis reaction proceeds in the following three steps.2O3 is reduced by Mg  surface increased from 1 to 3 μm as the synthesis temperature increased from 800 to 1200 °C. In addition, the B4C grain size increased as the synthesis temperature increased. Furthermore, the amorphous structure content decreased. For the boron carbide coating synthesized at a lower temperature , a large amorphous content (29%) was observed based on the XPS  is much lower than the oxidation temperature of diamond (720 °C), indicating that the B2O3 formed during annealing exhibits fluidity. During annealing, the fluidity of B2O3 is beneficial for self-healing and densification of the protective coating on the diamond surface prior to oxidation, which improves the oxidation resistance of the diamond particles. Therefore, the slight oxygen corrosion observed on sample A1 (2O3 coating that completely covers the diamond (111) surface during annealing is difficult to achieve. Furthermore, sample A2 was completely covered by the B2O3 coating  surface rather than on the diamond (111) surface because atomic diffusion on the diamond (111) surface is more difficult. The diamond particle was completely covered by a stoichiometric B4C coating after holding the raw mixture at 1200 °C for 2 hours. The boron carbide coating on the diamond particles contributed to the enhancement of the oxidation resistance due to formation of a B2O3 oxygen barrier layer. In this study, all the boron carbide coatings were oxidized prior to the oxidation of diamond. Moreover, the B2O3 exhibits fluidity, which gives rise to secondary coating growth on the diamond particles , resulting in repair of defects . The coating on diamond can provide effective protection against oxidation by heating in air at 1000 °C for 1 h. In comparison to the uncoated diamond particles, the initial time prior to oxidation of diamond was delayed for 2 h. Furthermore, the boron carbide coating on the diamond particles also contributed to the maintenance of the static compressive strength during annealing in air.To summarize, the boron carbide coating on the diamond particles was synthesized by heating the diamond particles in a powder mixture consisting of H3BO3) powders. The diamond-powder mixture was mixed using vigorously mechanical stirring at room temperature for two hours. This mixture was placed on an alumina boat and placed into a tube furnace. The chamber was evacuated using a turbomolecular pump to 5×10−3 Pa prior to sample heating. The diamond-powder mixture was heated to 800, 1000 and 1200 °C for 2 hours to synthesize the boron carbide coating on the diamond particles. The vacuum was only broken when the furnace temperature reached room temperature. After cooling of the furnace, the samples were treated with dilute nitric acid to remove the soluble phases (including B2O3 and MgO) from the product. The coated diamonds were separated from the excess B powders using a sieve. The uncoated diamond particles are referred to as D0. In addition, the coated diamond particles synthesized at 800, 1000 and 1200 °C is referred to as D1, D2 and D3, respectively.Synthetic HPHT diamond particles , Henan huanghe whirlwind international Co., Ltd., China) were used. 20 g of the diamond particles were immersed in a mixture consisting of 10 g of boron (B), 5 g of magnesium (Mg) and 40 g of boric acid . All the samples were annealed at 1000 °C with a heating rate of 5 °C min–1 in an air atmosphere to investigate the oxidation resistance. After annealing, the uncoated diamond particles (D0) are referred to as A0. After annealing, the corresponding samples of D1, D2 and D3 are referred to as A1, A2 and A3, respectively. Thermal gravimetric analysis (TGA) was performed using a Netzsch STA449F3. The temperature of the TGA was increased from room temperature (20 °C) to 1000 °C at a rate of 5 °C/min, and then, the temperature was maintained at 1000 °C for 4 h. An air atmosphere was flowed through the chamber during the measurement. The static compressive strength of one diamond particle was measured by continuously adding load until the particle broken. At least forty particles were measured for each sample to obtain higher measurement accuracy.The microstructure of each sample was characterized by X-ray diffraction (XRD) using a Bruker D8 with a Cu Kα source. The chemical composition of the samples was determined using X-ray photoelectron spectroscopy (XPS) on a Physical Systems Quantum 2000 spectrometer with monochromatic Al Kα radiation. The samples used for XPS analysis were cleaned using ArHow to cite this article: Sun, Y. et al. Enhancement of oxidation resistance via a self-healing boron carbide coating on diamond particles. Sci. Rep.6, 20198; doi: 10.1038/srep20198 (2016).


When the Mg is impregnated into the support before the Co and Ru addition, some Mg is incorporated into the support in the form of MgxAl2Ox3+ if the material is calcined at 550°C or 800°C after the impregnation, while the remainder is present as amorphous MgO/MgCO3 phases. After subsequent Co-Ru impregnation MgxCox3−O4 is formed which decomposes on reduction, leading to Co(0) particles intimately mixed with Mg, as shown by high-resolution transmission electron microscopy. The process of impregnating Co into an Mg-modified support results in dissolution of the amorphous Mg, and it is this Mg which is then incorporated into MgxCox3−O4. Acid washing or higher temperature calcination after Mg impregnation can remove most of this amorphous Mg, resulting in lower values of x in MgxCox3−O4. Catalytic testing of these materials reveals that Mg incorporation into the Co oxide phase is severely detrimental to the site-time yield, while Mg incorporation into the support may provide some enhancement of activity at high temperature.The effectiveness of Mg as a promoter of Co-Ru/γ-Al FTS can be used to convert cheap gas sources, such as shale gas, to exploit stranded gas reserves, to prevent gas flaring during oil extraction or to convert feedstocks such as coal and biomass to higher value liquid hydrocarbons –5. The r2in situ to reduce the Co oxides to Co metal.Cobalt catalysts for FTS generally consist of the active metal dispersed as an oxide over an oxidic support with various quantities of metal oxide or platinum group metal promoters to improve the reducibility of the active phase. As it is widely accepted that FTS proceeds over Co metal sites , prior t2/kieselguhr catalyst [et al. [2O3 suppressed the formation of difficult to reduce Co-surface phases. The effect was thought to be due to the formation of MgAl2O4 spinel on the surface of the γ-Al2O3, which reduced the contact between Co and γ-Al2O3 and therefore the amount of cobalt aluminate spinel, which is not FT-active. Hence, Mg at low loadings could increase reducibility and activity. The same study found that more than 1 wt% MgO decreased the reducibility due to the formation of an MgO–CoO solid solution, which required high-temperature reduction. Chernavskii et al. [2O3 support. A study by Niemelä et al. [2-supported catalysts reduced the X-ray diffraction (XRD) particle size of Co metal particles from 14 nm to around 10 nm and decreased reducibility presumably due to the smaller particle size. It should be noted that Co/SiO2 catalysts tend to be more reducible than Co/γ-Al2O3 catalysts and therefore a decrease in reducibility of MgO-promoted SiO2 may be expected. The study by Niemelä also found that activity was higher when Mg was impregnated after Co onto the support rather than before. A more recent study by Holmen’s group observed a decrease in activity for Mg-promoted Co/γ-Al2O3 catalysts which could not be explained by particle size or dispersion effects [Magnesium has been used numerous times as an additive to FTS catalysts. Indeed, some early work used Mg as a promoter in a Co/MgO/ThOcatalyst . Most stcatalyst . However [et al.  found thi et al.  also fouä et al.  found th effects . The aut effects . Mg in gIn this work, we aim to identify the origin of the effects on particle size, activity and stability. We conduct detailed characterization of Mg-promoted catalysts, and a series of experiments aimed at controlling the location of Mg in the catalysts to better understand the effects.2.2O3 , was dried at 120°C overnight. The impregnation was carried out in plastic bags . The impregnation was a multi-step process, the first step being the impregnation of magnesium nitrate hexahydrate  followed by heat treatment at either 550°C or 800°C, then the impregnation of the cobalt nitrate hexahydrate  and ruthenium (III) nitrosyl nitrate , followed by a final heat treatment at 250°C or 400°C. To the γ-Al2O3 was added the appropriate amount of hot aqueous nitrate solution (ca 60°C) to achieve the desired loading. The impregnated solid was then kneaded in the plastic bag until homogeneous. For thermal treatments, the damp solid was spread out on a stainless steel tray and treated in static air in a muffle furnace. After thermal treatment, subsequent impregnations were carried out using the same procedure. The volume used for the initial impregnation was 1 ml g−1 for unmodified Puralox TH100/150, and the volume required was recalculated after each impregnation.Catalysts were synthesized by an incipient wetness impregnation process. The support, γ-Al2O3 materials—acetic acid  and an acetic acid–ammonium acetate buffer . Acetic acid was chosen as it has a similar pH to Co(NO3)2 (ranges from approx. 3 at 1 M concentration to approx. 2 at 4 M) and the buffer solution was chosen to determine whether a milder treatment could produce the same results. Additionally, residual species left after the washes could be removed by appropriate heat treatments, unlike for instance acetic acid–sodium acetate buffers. The washing procedure was thus:Some of the Mg-modified supports were subjected to an acid wash. Two different solutions with different pH were used to wash the Mg/γ-Al2O3 sample (1 g) contained in a centrifuge tube (50 ml) was added the weak acid solution (40 ml) followed by shaking for 30 s. Tubes were then centrifuged for 30 min at 4000 r.p.m. before removing the supernatant with a pipette. In some cases, the Mg and Al content of the supernatant was analysed by inductively coupled plasma (ICP) analysis. The wet solid was washed with water (40 ml) via shaking by hand for 30 s before returning to the centrifuge for 30 min at 4000 r.p.m. The supernatant was removed by pipette before drying the wet solid for 18 h at 80°C. The dry solid was calcined at 550°C for 2 h in static air to remove residual acetate or ammonium ions.To each Mg/γ-Alet al. [α1 radiation (1.7890 Å). Variable divergence slits were used with an illuminated length of 15 mm. The measuring time for XRD scans in the 2θ range 18°–82° was 1 h. Samples were spiked with KCl in order to more accurately calculate the lattice parameters of the Co3O4 phase. Two methods were used to extract quantitative information from XRD patterns: empirical peak fitting and Le Bail whole pattern fitting. Empirical peak fitting was used to calculate particle size from the Scherrer equation while Le Bail fitting was used to calculate the lattice parameters of Co3O4 phases. TOPAS 4.2 was used to refine the patterns of the Co3O4/γ-Al2O3 samples. In addition to the background polynomial, the lattice parameters and effective particle size parameters were refined, although determination of instrumental broadening was not performed so these particle size measurements are not discussed as they would lead to a false impression of their accuracy. Where particle size measurements are quoted, these are calculated from single peaks using the Scherrer equation.XRD data on samples from Boldrin et al.   radiation, Fe filter, 5 mm mask and automatic divergence slits with an illuminated length of 3 mm. The Soller slits used were 0.02° on the incident side and 0.04° on the reflection side. The measuring time for an XRD scan in the 2θ range 38°–57° was 13 min.In addition to this, the acid-washed (AW) supports were characterized using the high-throughput XRD technique described by Boldrin et al. . Samples2O3. Spectra were collected from 600 to 4000 cm−1 with a resolution of 4 cm−1 using a Perkin Elmer Spectrum 100 spectrometer fitted with the Spectrum 100 Universal Diamond/ZnSe ATR.Fourier transform infrared (FTIR) spectra were recorded for Mg-modified γ-Al−1 before the gas was switched to 10% H2/Ar at 30 ml min−1. Samples were ramped from room temperature to 1100°C at 10°C min−1, holding at 1100°C for 10 min. Five injections were performed at the end of the experiment, which allowed the peak area to be related to H2 consumption.Temperature-programmed reduction (TPR) was carried out on an Altamira AMI-200 unit. Samples (70 mg) were purged with flowing argon at 30 ml min2/N2 as the sample gas flowing at 25 ml min−1. Samples (ca 5 mg) were placed in aluminium pans. The dynamic high-resolution mode was used with a ramp rate of 50°C min−1, sensitivity of 1.0 and a resolution of 4.0. Samples were heated from room temperature to 600°C, where the temperature was held for 5 min. The maximum temperature was limited to 600°C due to the use of aluminium pans.Thermogravimetric analysis (TGA) scans were carried out on a TA Q5000IR thermal analyser using 5% H2 temperature-programmed desorption (TPD) was carried out on an Altamira AMI-200 unit. Samples (250 mg) were first heated to 250°C at 10°C min−1 and purged with flowing He at 30 ml min−1 for 10 min before being cooled to 35°C at 10°C min−1. Subsequently, the gas was switched to 10% CO2/He at 30 ml min−1 for 1 h at 35°C followed by 30 ml min−1 He for 30 min. Finally, the sample was heated to 800°C at 10°C min−1 in He flowing at 10 ml min−1, holding at 800°C for 10 min. Five injections were performed at the end of the experiment, which allowed the peak area to be related to CO2 desorption.CO2 chemisorption at 150°C in a Micromeritics ASAP 2020C by extrapolating the total gas uptakes in the H2 adsorption isotherms at zero pressure. Prior to adsorption, the samples (ca 0.5 g) were pre-treated in flowing He at 120°C for 1 h. Afterwards, the samples were reduced in situ by flowing pure H2 and raising the temperature to 425°C and maintaining this temperature for 6 h. After reduction, the samples were degassed, the temperature lowered to 150°C and H2 dosed over the sample in the pressure range of 100–760 mm Hg.Co surface areas were determined by HTransmission electron microscopy was performed to calculate Co particle size distributions and to produce elemental maps. Samples were first reduced and passivated. Powder samples were embedded in resin, cured, microtomed and placed on holey carbon-coated Cu grids. The samples were examined in the probe-corrected JEOL-2100F transmission electron microscope (TEM) at the University of Birmingham, UK, using the scanning transmission mode with high-angle annular dark-field and bright-field detectors and a voltage of 200 kV. Parallel electron energy loss spectroscopy (EELS) using a GATAN Enfina spectrometer was used to investigate the composition.ICP analysis was performed to determine the compositions of solid samples. Following complete dissolution of the samples using microwave digestion in HCl, analysis was performed on a Perkin Elmer Optima ICP-OES2O3 solids by washing with acetic acid or ammonium acetate on a Ciros CCD optical emission spectrometer.ICP analysis was performed to determine the amount of Mg and Al ions removed from Mg-modified γ-Al2 prior to testing. After reduction, the reactor was cooled to 160°C and H2 was replaced with syngas (H2:CO:Ar=2:1:0.1) and pressure was raised to 20 bar. Conversion was measured by the consumption of the syngas mix (CO+H2) relative to Ar. Wax and water were collected from the gas–liquid separators and analysed using an offline gas chromatogram with a flame ionization detector on a SimDist column in order to calculate the chain growth probability, α, from the C22–C40 fraction.The FTS testing was carried out in a stainless steel isothermal fixed bed six-way micro-reactor connected to an online Varian CP-3800 gas chromatograph with three detectors. A Valco-valve was used to select between the gas streams for online analysis. Catalysts were reduced at 425°C for 9 h in pure H3.2O3. The Mg was found to have several benefits: reducing the Co particle size, improving the Co surface area and increasing the stability of the catalysts during FTS. There were also detrimental effects on the reducibility of the catalysts and the initial activity. In order to investigate the origin of these effects, detailed characterization was carried out on both the initial catalysts and the modified supports.In our previous work , we desc2O3 support and one assigned to Co3O4. Le Bail refinements were undertaken with peaks inserted empirically for the alumina phase due to the difficulty in assigning a structural model to this phase [3O4 phase for the 0% Mg sample was slightly smaller than expected for bulk Co3O4. This slight contraction of the unit cell could be the result of the small size (14.7 nm) of the Co3O4 particles causing deviations from bulk properties, for instance, due to the presence of vacancies or other crystal imperfections [3O4 lattice parameter increased. By comparison with Mg0.93Co2.07O4, which is reported to have a lattice parameter of 8.1390 Å[2+ has a larger ionic radius than Co2+ 2 incipient wetness impregnation, there are Mg species which are sufficiently reactive to form a mixed oxide with Co ions. By comparison with the literature data on the lattice parameter of MgxCox3−O4 compounds [xCox3−O4 spinel relative to the total amount of Mg can be calculated. For the 6% Mg sample, the composition of the spinel was Mg0.76Co2.24O4, which means that 62% of the total magnesium in the sample is incorporated in this phase. The particle sizes calculated from the broadening of the Co3O4 peaks indicate that catalysts containing Mg had smaller Co3O4 particles than the catalyst without Mg, which agrees with our previous work [XRD analysis of the calcined catalysts  showed the presence of two phases, one assigned to the γ-Alis phase ,15, withfections . As the  8.1390 Å, and conal sites ) this caompounds , an esti2O3 phase, information relating to the phase can be gathered by simply looking at the shifts in the peaks. Although somewhat crude, by assuming that γ-Al2O3 has the defect spinel structure, a lattice parameter was calculated by measuring the peak position of a single reflection at approximately 53.6° 2θ . The modified supports were also characterized by XRD to examine the crystalline phases present before Co(NO3)2 impregnation. This analysis confirmed that the shift in lattice parameter of the γ-Al2O3 phase observed in the Co-loaded catalysts was primarily due to Mg incorporation . All XRD patterns only contained peaks due to the MgyAl2Oy3+ phase except for the sample with 6% Mg calcined at 550°C, which also contained peaks due to MgO. The emergence of the MgO peaks only at a loading of 6% Mg indicates that below 6% Mg the MgO particles were too small or insufficiently crystalline to be detected by XRD under the experimental set-up. Because peaks due to MgO were seen in this sample before Co(NO3)2 impregnation and calcination but not seen after, this indicates that MgxCox3−O4 formation in the catalyst is probably due to a dissolution or solid-state reaction causing Mg ions to incorporate into the Co3O4 structure. The lattice parameter for the γ-Al2O3 phase is similar before and after Co(NO3)2 impregnation, suggesting that the Mg in MgxCox3−O4 does not come from MgyAl2Oy3+ and is therefore likely to be from an amorphous or highly dispersed Mg-containing phase.Although a Le Bail refinement could not be performed on the γ-Al53.6° 2θ . It can 3O4and the γ-Al2O3 structures. From the measurement of x in MgxCox3−O4, it is possible to calculate the proportion of Mg which is present in the MgxCox3−O4 phase, and, by subtraction, the remaining Mg which is not accounted for by this  indicates that any Mg which was not incorporated into MgxCox3−O4 or MgyAl2Oy3+ was highly dispersed or poorly crystalline. The proportion of the total Mg which is incorporated into the MgxCox3−O4 phase increases with increasing Mg loading.It is evident from the XRD data that, in the pre-activated catalyst, Mg is incorporated into the Co2O3 samples were calcined at a higher temperature of 800°C to provide insight into the behaviour of the samples calcined at the lower temperature of 550°C. The XRD patterns showed a single MgyAl2Oy3+ phase for all compositions and, again, an increase in the lattice parameter with Mg loading was observed . However, for these samples, the increase in the lattice parameter was greater than that at the lower calcination temperature. This indicates that more Mg was incorporated into the γ-Al2O3 structure, presumably due to the higher temperature facilitating more diffusion into the bulk. As the lattice parameter increased more at 800°C, this implies that after calcination at 550°C there are Mg species which are XRD-silent and not incorporated into the MgyAl2Oy3+ phase. As this available Mg was not seen by XRD for samples other than 6% Mg calcined at 550°C this means that it is either poorly crystalline or highly dispersed, perhaps as a monolayer as suggested previously [xCox3−O4.Mg/γ-Aleviously . It is tyAl2Oy3+ phase. As expected, every sample had major bands due to O−H (ca 3500 cm−1), which is mainly due to adsorbed water in the pore network but may also have a contribution from hydroxide species. The presence of a peak at ca 1385 cm−1 in the samples calcined at 550°C is indicative of carbonate ions. The intensity of this peak increased with Mg loading, indicating that MgCO3 was present. This peak is absent in the samples calcined at 800°C. CHN analysis confirms the presence of carbon in all the modified supports calcined at 550°C, and the two highest loadings of Mg in the supports calcined at 800°C . The amount of carbon increases with Mg loading, reinforcing the hypothesis that the peaks in the FTIR are due to MgCO3. CHN analyses conducted over time show an increasing carbon content with time since synthesis , suggesting that initially the XRD-silent Mg species are amorphous oxides, which convert to carbonates over time in air.FTIR spectroscopy on the modified supports  can giveyAl2Oy3+ phase. For the samples calcined at 550°C, there is a peak at ca 824 cm−1 which is attributed to tetrahedral Al−O bonds [b). Furthermore, for the samples calcined at 800°C, an additional band at ca 683 cm−1, attributed to octahedral Al−O bonds [3+ ions in tetrahedral sites and more in octahedral sites. In a normal spinel such as MgAl2O4, the divalent cation resides in tetrahedral sites and the trivalent cation is in octahedral sites. Although the structure of γ-Al2O3 is similar to spinel, it is often reported as having many defects, non-spinel sites and with a proportion of the Al3+ cations in tetrahedral sites rather than octahedral [3+ cations’ occupation of crystallographic sites resembles more closely that of bulk MgAl2O4 spinel. Presumably, as the Mg loading increases, more Mg2+ is incorporated into the tetrahedral sites of the γ-Al2O3, which results in fewer Al3+ cations in the tetrahedral sites and more in the octahedral sites. This move towards a spinel-type MgyAl2Oy3+ structure is consistent with the XRD data.Examination of the low-energy region of the FTIR spectrum reveals peaks typically due to metal–oxygen bonds and gives information about the Mg−O bonds . As the −O bonds , increastahedral ,21. The 2 TPD . The unmodified support possesses a single CO2 desorption peak at around 100°C, indicating the presence of a weakly basic site, as is expected from γ-Al2O3 [2O3 crystal structure does not increase the basicity of the support’s surface. The total amount of CO2 desorbed does not change with Mg loading in these samples, which reinforces this hypothesis. However, modified supports with Mg loadings of 2% and above calcined at 550°C also possessed a more strongly basic site reflected in an additional desorption peak at 300°C. This is likely to be from the XRD-silent Mg species identified previously, especially as, in these samples, the total amount of CO2 does increase with Mg loading.Further information about the nature of the supports can be deduced from CO γ-Al2O3 ,19. This2 desorption in the TPD with the amount of Mg present in each phase in the final catalyst, we can see that the correlation is better with the amount of Mg in the MgxCox3−O4 phase than with the amount of ‘other’ Mg. Because the total amount of CO2 desorbed is mainly influenced by the XRD-silent Mg, this suggests that the XRD-silent Mg is ending up in the MgxCox3−O4 phase rather than as ‘other’ Mg in the final catalyst, so the ‘other’ Mg is almost entirely the MgyAl2Oy3+ phase (rather than being a combination of Mg in MgyAl2Oy3+ and XRD-silent Mg species). This is in agreement with the XRD data, which show a good correlation between the MgyAl2Oy3+ lattice parameter and the ‘other Mg’ content, and gives an indication that the cobalt impregnation solution is dissolving the majority of the amorphous Mg, which is then incorporated into the MgxCox3−O4 phase.Comparing the total CO2O3 with Mg has two effects: (i) the creation of an MgyAl2Oy3+ phase that is intermediate between γ-Al2O3 and MgAl2O4 in crystal structure, which has a similar basicity to the unmodified support and is unchanged by the impregnation of cobalt, and (ii) the deposition of highly basic XRD-silent MgO or MgCO3 species which are later incorporated into the cobalt oxide spinel forming MgxCox3−O4. We hypothesized that it is the MgO/MgCO3 species that is dissolved by the acidic Co(NO3)2 impregnation solution and thus forms the MgxCox3−O4 phase. We therefore attempted to remove this second species by acid washing the modified supports before Co(NO3)2 impregnation, allowing us to separately examine the two effects. The washing solutions chosen were acetic acid (pH 2.4), which has a similar pH to concentrated cobalt nitrate solutions (pH 2–3 in the range of 3–4 M concentrations), and ammonium acetate buffer (pH 4.7) to assess if a milder treatment could achieve the same results.The evidence detailed above strongly suggests that the modification of γ-Ala,b (buffer wash) and the electronic supplementary material, figure S6 (acetic acid wash). Washing with either acetic acid (pH 2.4) or ammonium acetate buffer (pH 4.7) resulted in almost identical amounts of Mg being dissolved from the sample for both the supports calcined at 550°C and 800°C, indicating that the two washing solutions were equally effective in dissolving Mg. As would be expected, much more Mg was dissolved from samples calcined at 550°C than from those calcined at 800°C. Interestingly, some Al was also dissolved by both ammonium acetate buffer and acetic acid , including, in the unmodified support, raising the prospect that some Al may also be incorporated into the Co3O4 spinel, an effect which to our knowledge has not been discussed previously in the literature on γ-Al2O3-supported catalysts. The amount of Al dissolved increases with Mg loading, suggesting that Al in mixed Mg-Al species is more easily dissolved. Less Al is dissolved by the higher pH buffer.ICP analysis of the supernatant allowed us to characterize the amount of Mg removed during washing, while the amount left behind is calculated by subtraction from the ICP results of the materials before washing. The results are shown in figure 2yAl2Oy3+ phase from XRD of the washed and unwashed samples  shows little change on washing for all samples, indicating that this phase is largely stable towards washing with acid. There is a small decrease in the lattice parameters for the samples calcined at 550°C, and this is reflective of the ICP results given above, which suggest that, although the majority of the dissolved Mg comes from highly dispersed or amorphous Mg-containing phases, some mixed Mg-Al species are also dissolved.Comparison of the lattice parameters of the Mg samples a,b shows3O4 crystallite size was measured using XRD . In samples where the support has been calcined at 550°C, Mg loading still has an effect in reducing the particle size, but the effect is reduced compared with unwashed samples. In addition, higher loadings of Mg show a bigger difference in Co particle size between washed and unwashed samples, suggesting that a large part of the particle size reduction is due to the Mg species which are removed by washing. The washed and unwashed supports calcined at 800°C give very similar particle sizes to those seen in the washed samples calcined at 550°C, which again indicates that it is the free Mg which largely causes the decrease in particle size.A series of Co-Ru catalysts were made using the washed and unwashed supports calcined at 550°C and 800°C, and the Cosing XRD c,d. In s2 in N2 was carried out on these catalysts. Samples where the supports were calcined at 550°C show clearly that washing with either acid or buffer improves reducibility, as measured by the temperature maxima for the second reduction peak (Co(II)–Co(0); electronic supplementary material, figure S8). This is in line with the hypothesis that it is the free Mg which is detrimental to reducibility rather than total Mg.TGA in 5% HSeveral samples washed with ammonium acetate buffer solution (henceforth AW) were scaled up for characterization and testing based on these results. Characterization of these samples is shown in xCox3−O4, indicating that the washing is not fully effective in preventing this phase from forming. It is not clear why the impregnation solution is able to dissolve Mg from the washed supports, possibly because the impregnation solution is hot, or there is some other chemical effect than simply the pH. The washes were effective in reducing the Mg in MgxCox3−O4 by 60% in the 3% Mg samples, and by 80% in the 6% Mg samples. The washes also removed an appreciable fraction of the Mg not incorporated into MgxCox3−O4—around 17% in the 3% Mg samples and 33% in the 6% Mg samples. The calcination at 800°C was more effective at reducing the formation of MgxCox3−O4—in the 3% Mg sample less than 1% of the total Mg was found in the MgxCox3−O4 phase, while in the 6% Mg sample it was 4% of the total Mg.XRD analysis showed that the washed catalysts still contained some MgxCox3−O4 spinel phase can be calculated from XRD, and it can be seen that acid washing has a strong effect in reducing the particle size. The reason for this is unclear, and it occurs even in samples without Mg . The relative effects of the Mg in MgxCox3−O4 and Mg in MgxAl2Ox3+ on particle size are difficult to deconvolute, because they are heavily autocorrelated , and the calcination at 800°C which would provide a comparator also has a separate effect on particle size. In general, the results agree with earlier results that higher Mg contents reduce particle size.The crystallite size of the MgxCox3−O4 rather than the total amount of Mg which determines the reduction temperature. Calcining the support at 800°C also lowers the reduction temperature relative to both washed and unwashed samples with the same initial Mg loading, and this also reinforces the hypothesis. This is complicated slightly by the 3% Mg samples calcined at 550°C, which show no change in peak position between washed and unwashed samples despite a large difference in the level of Mg in MgxCox3−O4. However, taking these results together with the results of TGA in 5% H2 carried out on the earlier samples , it becomes clear that it is the Mg in MgxCox3−O4 which is detrimental to reducibility.The effect of the different Mg levels on the reducibility, as seen in the TPR, is somewhat clearer . The was2 chemisorption followed the same trend for the unwashed samples as the Mg-promoted samples scaled up previously but were slightly lower, which could be due to the higher calcination temperature. The Co surface area increased from 57.4 m2 g−1Co for 0% Mg to 63.9 m2 g−1Co for 3% Mg before decreasing to 59.2 m2 g−1Co for 6% Mg. Again this is attributed to decreasing particle size coupled with decreasing reducibility as the Mg loading was increased.The cobalt surface area measurements as determined by H2/gcat compared with 11.2 m2/gcat. The testing procedure involved examining the activity and selectivity at 210°C, 230°C and 240°C, before returning to 210°C over a total of 245 h time online at testing temperature.AW2–AW6 were tested in the Fischer–Tropsch reaction to examine the effect of acid washing and high-temperature calcination prior to Co addition, along with a control sample, which was the 0% Mg sample discussed at the beginning of the Results section and in our previous work. To avoid confusion, we will refer to this sample as 0% Mg. The main difference between this sample and AW1 is that AW1 had a final heat treatment at 450°C rather than at 250°C, and a slightly lower Co surface area of 9.8 mxCox3−O4 and reducibility discussed above, it appears that the trends in activity cannot be explained simply by reducibility; for example, AW5 has a second reduction peak 35°C higher than any other catalyst but still has higher activity than AW2 or AW6.The trends in activity at 210°C shown in  cannot bex situ) of each catalyst at 210°C and plotting them against the lattice parameter of the MgxCox3−O4 phase measured in the as-calcined state provides insight into the trends in activity just discussed. Figure 6a shows the site-time yields with respect to the lattice parameter for all six catalysts. It is clear that the level of Mg substitution into the Co3O4 spinel has a marked effect on the intrinsic activity of the catalysts—as the MgxCox3−O4 phase becomes more Mg-rich, the site-time yield decreases approximately linearly with increasing lattice parameter.However, calculating the site-time yields  and Co(II) ions are reduced to Co(0), destroying the MgxCox3−O4 structure but remaining in close proximity to the Mg ions which they were previously intimately mixed with.Because a single-mixed Co-Mg phase is formed after calcination, it is highly likely that, on reduction in the reactor, Mg which was mixed intimately with Co remains in close proximity to the reduced Co metal atoms. Hence the more Mg-rich the mixed spinel, the greater the amount of Mg in close proximity to the Co metal active sites in the activated catalyst. In order to investigate this, EELS was applied to a reduced and passivated 3% Mg catalyst . The cat3O4 appears not to be simply due to the total amount of Mg present in the sample, exemplified by the fact that AW6 has almost twice the total amount of Mg as AW2 but approximately 150% of the activity at 210°C.Hence, it is proposed that Mg ions have a poisoning effect on the intrinsic activity of Co metal sites and the degree of poisoning depends on the concentration of Mg ions close to the Co sites. The decrease in activity with increasing Mg incorporation into CoPoisoning by Ca has been observed previously with the reduction in activity being attributed to electronic effects causing a decrease in surface hydrogen concentrations and increased CO adsorption and dissociation . Anotheractivated them. The percentage increase in activity is strongly related to the level of Mg in the MgxCox3−O4 (c), suggesting that the higher temperatures are causing Mg to migrate away from Co, causing a reduction in the poisoning effect of Mg.After approximately 65 h at high temperature, the catalysts were returned to 210°C. Interestingly, all the samples except the 0% Mg standard had higher activity after the high-temperature run than before, as shown in xCo3−xO4 c, suggesb). Again, the relative increase in site-time yield is somewhat correlated to the level of Mg in the MgxCox3−O4, although as the yields fall back below those of the standard when returning to 210°C, this effect cannot be entirely due to a high-temperature activation process and must be at least somewhat an intrinsic property of Mg promotion. Arrhenius plots for the 0% Mg standard, AW3 and AW4  show no deviation from a linear trend at the highest temperature, which also indicates that there is no large change in the number of active sites. The fact that it is AW4, in which x in MgxCox3−O4 is only 0.01, which shows the largest increase in site-time yield with temperature raises the interesting prospect that the improved high-temperature performance is due to Mg incorporated into the alumina phase.During the high-temperature phase of the test at 230°C and then 240°C, the site-time yields of all the catalysts except AW5 increase relative to the 0% Mg standard sample, to the extent that, by 240°C, AW4 has a higher site-time yield than the standard, while AW3 is the same within error b. Again,d), there is a consistent correlation between selectivity for all products and conversion among all samples at all temperatures, strongly suggesting that Mg in any form is not having a detectable influence on the selectivity or on the relative rates of possible competing reactions such as water gas shift.Looking at the selectivity of the samples d, there 4.The fact that magnesium is incorporated into the cobalt oxide spinel phase as well as the alumina phase despite the high-temperature treatment before impregnation has important implications for the use of structural promoters. We have shown that, while incorporation into the alumina phase may have beneficial effects on the high-temperature performance, incorporation into the cobalt phase has severely deleterious effects on the site-time yield, which may hold for other structural promoters. We have also shown that the impregnation solution has leaching effects greater than its pH implies, and also that the solution is potentially capable of dissolving measurable amounts of aluminium from the support, which presumably is also incorporated into the cobalt spinel phase, with unknown effects.Quantification of the extent and effect of this doping of promoters and elements from the support into different structures of a catalyst is a difficult problem requiring detailed investigation with XRD, high-resolution TEM and elemental analysis, as well as a carefully designed experimental strategy, but should be a fruitful avenue of research in assessing new materials and reassessing older materials.


Utilizing the bandgap modulation property, a tunable bandgap transistor, which can be in general made of a two-dimensional semiconductor, is proposed.Semiconductors with a moderate bandgap have enabled modern electronic device technology, and the current scaling trends down to nanometer scale have introduced two-dimensional (2D) semiconductors. The bandgap of a semiconductor has been an intrinsic property independent of the environments and determined fundamental semiconductor device characteristics. In contrast to bulk semiconductors, we demonstrate that an atomically thin two-dimensional semiconductor has a bandgap with strong dependence on dielectric environments. Specifically, monolayer MoS Monolayer (ML) molybdenum disulfide , the strong unscreened Coulomb interaction  makes the quasiparticle (QP) renormalization of electrons huge. Within the GW approximation, the electronic bandgap of a freestanding ML MoS2 has been predicted to be about 2.8 eV 2 on a Au(111) substrate, the ARPES bandgap of about 1.39 eV has been measured, which is very small2 on graphite substrate has been measured to be 2.15 eV  as a fundamental material property may not be applicable to 2D semiconductors, and that the electronic bandgap may have strong dependence on the environments. Considering the fundamental role of the bandgap size in electronic device applications, it is critical to develop a fundamental and quantitative understanding on how the environmental effects change the bandgap sizes of 2D semiconductors.A moderate bandgap size is a determining characteristic property of a semiconductor. Nonetheless, an accurate evaluation of the bandgap in low dimensional semiconductors has not been as simple as in conventional bulk semiconductors. It is well established that the bandgap size of MoS eV refs , 11, 12. eV ref.  and that eV ref. , which a2 with various environments based on GW calculations, and predict a wide range of bandgap size determined by the strong effects on the ML MoS2 embedded in a device structure. The bandgap of ML MoS2 is found to change by the surrounding medium according to the dielectric constant (κE) of the environment. Specifically, the GW bandgap changes from 2.8 eV of freestanding MoS2 down to about 1.9 eV for MoS2 in a sandwich structure between two high-k dielectrics. On the other hand, the bandgap changes down to 2.2 eV for a supported structure on ultrahigh-κE dielectrics. These GW bandgap changes are continuous functions of dielectric constant of the surrounding medium. Based on this finding, it is suggested that there should be transport barriers to electrons and holes in the ML MoS2 channel between near the metallic contacts and near the gate dielectric in a device structure, because of the different screening environments (and correspondingly different bandgap sizes) surrounding the ML MoS2. When the barriers are controlled by an external source, a tunable bandgap transistor can be made possible utilizing the environment-dependent property of the 2D semiconductors.In this study, as a representative 2D semiconductor, we investigate the bandgaps of a ML MoS2 model systems with different dielectric environments for full DFT and GW calculations: (A) a freestanding ML MoS2 surrounded by vacuum, (B) ML MoS2 on a HfO2 substrate, (C) ML MoS2 sandwiched by HfO2, (D) ML MoS2 on a Au substrate, and (E) ML MoS2 sandwiched by Au. We construct the model atomic structures for the five systems as shown in 2’s have a direct bandgap of 1.8 eV at the K valley irrespective of the environments. The electronic structure analysis shows that the orbitals of the valence band maximum (EV) and the conduction band minimum (EC) at the K valley are characterized as the Mo 4d atomic orbitals. These atomic orbitals are located in the middle of the three atomic layers of ML MoS2 layer, and they have negligible hybridization with the orbitals of the nearby surrounding materials that interacts with the ML MoS2 through the van der Waals gap. The LDA bandgap at the K valley is about 1.8 eV agreeing with previous DFT calculations and found to remain unchanged with different environments. For the metallic Au environments (s) are found inside the bandgap of the ML MoS2, but they can be clearly distinguished from the ML MoS2 states (shown as blue dots). Within the LDA, the Au 5d states are found inside the valence bands of the ML MoS2, and the Fermi level (EF) is found to be located at EV + 0.8 eV or EC − 1.0 eV.First, we consider five ML MoSronments , the Au GW band structures of the ML MoS2 with the same environments are plotted in 2 has a direct bandgap at the K valley irrespective of the environments, as in the case of the LDA. The direct bandgap at the K valley of the freestanding ML MoS2 . Since the electronic orbital hybridization between the MoS2 and the nearby HfO2 dielectric is negligible, as shown in the LDA results  into the dielectric matrix of the screened Coulomb (W) interaction in GW calculations. Details of the procedure are described in the GW, there are several advantages besides making it possible to study separately the EDS effect: reducing the computational costs of GW calculations with environments, and making it possible to include additional polarizability into the screened Coulomb (W) interaction. Note that the dielectric effect of liquid medium on ML MoS2 bandgap can be modeled by the effective dielectric medium. The quasiparticle renormalization in the GW approximation (W) interaction. Since high-κE dielectrics such as HfO2 (ε0 ≅ 26) usually have large ionic contribution (ε0−ε∝ ≅ 21) to the dielectric screening, the renormalization of electrons in the ML MoS2 would be further modified by the ionic screening of the surrounding materials. However, this ionic screening effect is neglected in the GW calculations of the band structures shown in GW plus the lattice polarization effect (LPE) can be applied to the calculation of the bandgap renormalization. In some ionic solids, the inclusion of the LPE has been reported to lead to a large shrinkage of the bandgap20GW + LPE, effective dielectric medium method provide an alternative efficient approach to include full dielectric effects on the MoS2 bandgap.In order to investigate the primary effect of EDS on the bandgap of ML MoSximation  includes ε0−ε∝ ≅  to the dGW bandgaps of ML MoS2 with an effective dielectric constant (κE) of the environments are plotted as a function of the κE in E = 1, the ML MoS2 represents the freestanding isolated one in vacuum, and the GW bandgap is found to be 2.8 eV at the K valley. With increasing the dielectric constant, κE, the GW bandgap is found to drop rapidly down. With the one-side dielectric, in such a case of the supported ML MoS2 on a substrate, the GW bandgap is found to reduce down to about 2.2 eV with a ultrahigh-κE dielectric (at κE = 30). With the both-side dielectric, as in a typical top-gate FET structure, the GW bandgap of the ML MoS2 is found to be smaller down to about 1.9 eV with a ultrahigh-κE dielectric of κE = 30. It is notable that the bandgap reduction is very rapid in the range 1 < κE < 5, and most of the bandgap reduction, about 80%, occurs with κE = 5. Thus, the presence of a moderate dielectric material in vicinity of a ML MoS2, can strongly affect the bandgap renormalization in the ML MoS2, even though it is not an ultrahigh-κE dielectric.The calculated 2 on graphite substrate is close to the obtained asymptotic value of 2.2 eV for the one-side dielectric, and the STS bandgap of 2.16 eV for the ML MoS2 on a bilayer graphene substrate2 (ref. 2 embedded in high-κE dielectric (1.9 eV). The measured ARPES bandgap of 1.39 eV for the ML MoS2 on Au substrate2 interface study2 in the Al2O3/MoS2/SiO2 stacked top-gate FETn-type ML MoS2 FET.The experimentally measured bandgap of 2.15 eV in STSte2 ref.  is close2 with including the EDS effect. 2 in LDA and those in GW with various κE. The EV and EC in LDA are found to be −6.16 and −4.31 eV, respectively, which are close to the previous LDA calculations (−5.98 and −4.29 eV)GW band edge levels are obtained using the bandgap center alignment schemeEV and EC in GW with κE = 1 are −6.64 and −3.83 eV, respectively, in good agreement with the previous GW calculations GW band edge levels with κE plotted in EV increases up and the EC decreases down monotonically with increasing κE, approaching the LDA values of EV and EC with ultrahigh-κE.We now investigate the absolute band edge levels (relative to vacuum level) of ML MoSEV + 0.8 eV or EC − 1.1 eV with the LDA band edge levels. They agree with those obtained in our atomistic GW calculations (EV + 0.8 eV or EC − 1.0 eV) , the Au work function level is located at EV + 1.2 eV or EC − 1.6 eV, which is far from the atomistic GW calculations , the Au work function level of EV + 0.8 eV or EC − 1.1 eV is close to the atomistic GW results (EV + 0.9 eV or EC − 1.2 eV)  Au metal, the work function (5.4 eV) level is located at  1.0 eV) . In the tructure . If we u 1.2 eV) . Note th ML MoS2 , which i ML MoS2 . AlthougoS2 refs ,25. Au hoS2 refs ,27, whicoS2 refs , 29.2. In order to match the lattice constant of Au with the MoS2 within a few % of strain, a larger supercell, for example 2 × 2 Au(111) slab and √3 × √3R30° MoS2 with 5.0% of compressive strain, is required with the number of atoms exceeding 33, with which the GW calculation is computationally demanding. One effect of the tensile strain on the Au(111) slab is lowering the Fermi level with smaller band dispersions, and the work function of the strained Au(111) is 5.4 eV, while that of the unstrained Au(111) is 5.1 eV, as indicated in 2 bandgap. Both are metallic with the 6s Fermi electrons, having the infinite static dielectric constants. Since the quasi-particle bandgap dependence on the dielectric constant is very weak with the large static dielectric constant, less than 0.1 eV when κE > 10 . We check another metallic slab, Ag(111), to test the quasi-particle bandgap of ML MoS2 with metallic screening. The Ag(111) slab is strained by 9.4% (tensile) to match with the 1 × 1 ML MoS2. The calculated band structures of the ML MoS2 with Ag(111) in the sandwich structure are shown in GW). The obtained GW bandgap of the ML MoS2 is 2.11 eV with the metallic Ag(111) environment, which is nearly the same to the GW bandgap (2.09 eV) of the ML MoS2 with the metallic Au(111).In our calculations, the Au(111) slab is strained by +9.6% (tensile) to match in lattice to the 1 × 1 ML MoS> 10 see , the qua2 in various environments. In an isolated freestanding ML MoS2, the strong Coulomb interaction (through the free space) between electrons makes the QP renormalization of electrons huge leading to a large bandgap . In a typical top-gate FET device structure, a channel is located between a substrate and a gate dielectric. The channel is also connected with metallic contacts in the source and drain regions. Such typical device structure is shown in 2) in the central region. For this structure, the EDS effect surrounding the ML MoS2 should be different for the source/drain and channel region. The different EDS strengths of metal and gate dielectric would result in different bandgaps in different regions. Under the metallic source/drain contact regions, the electronic bandgap should be smaller than in the channel region between the gate dielectric and the substrate. According to our atomistic GW calculations, the bandgaps in the source and drain regions with the both-side Au contacts are 2.1 eV, while the bandgap in the channel region with the both-side HfO2 dielectrics is 2.4 eV. The EV and EC offsets (ΔEC and ΔEV) in the ML MoS2 channel are 0.1 and 0.2 eV, respectively. It indicates that there are an electronic transport barrier (ΔEC) of 0.2 eV from the source to the channel, and a hole transport barrier (ΔEV) of 0.1 eV, even though the ML MoS2 channel itself is homogeneous atomically, due to the different EDS effects  externally. Recently, the electrically controlled dielectric materials utilizing ferroelectric properties have been suggested323334If the κ2, depends on the nearby dielectric environments, through the screened QP renormalization of electrons in the ML MoS2. The bandgap tends to reduce with increasing the environmental dielectric constant. In a ML MoS2 FET, the vicinity of metallic contacts gives smaller bandgap than that of the gate dielectric, and there should be valence and conduction band offsets between the regions. The band offsets can play a role of barriers to electron and hole transports through the channel. Utilizing the environment-dependent property of the bandgap, a tunable bandgap FET is suggested, which operates with the bandgap controlled by an external source to control the electron and hole transport barriers.Electronic bandgap of a 2D semiconductor, ML MoSk-point sampling in the hexagonal Brillouin zone (BZ) of the 1 × 1 ML MoS2 was used. The GW calculations were performed with the BerkeleyGW code382. For the 1 × 1 ML MoS2 with the HfO2 and Au environments, the number of conduction bands used was around 700. The limited number of conduction bands can affect the GW eigenvalues at the M point in the hexagonal BZK point converge fast with respect to the number of conduction bands. When we used more number of conduction bands, and the GW bandgaps at K were not significantly affected. The generalized plasmon pole (GPP) approximation was used for the frequency dependence of the dielectric matrix. We applied slab truncation scheme for the Coulomb interaction to minimize the supercell interaction for the free-standing ML MoS2 and the ML MoS2 with the HfO2 environments.The mean-field density-functional theory (DFT) calculations were performed with the Quantum-Espresso code with the local density approximation (LDA)2 was used. The lattice constant was fixed to 3.16 Å, which is the LDA optimized value for the free-standing ML MoS2. With the 1 × 1 in-plane periodicity of the ML MoS2, the HfO2 was modeled by the O-terminated 1 × 1 HfO2(111) slab, and the Au was modeled by the 1 × 1 Au(111) slab. With these interface structures, the HfO2 and Au slabs are hydrostatically strained by −8.5% (compressive) and +9.6% (tensile), respectively, to match their lattice constants to that of the ML MoS2 (3.16 Å). The atomistic environments are only model systems that represent dielectric and metallic environments. Six Hf atomic layers were used for the HfO2 dielectric slab, and six Au atomic layers were used for the Au metallic slab, as shown in 2 layers in bulk 2H-MoS2. The vacuum thickness of about 12 Å in the supercell was used. The ideal (as-cleaved) atomic structures for the HfO2 and Au slabs were used to see only the electronic EDS effect.The 1 × 1 hexagonal unit cell for the ML MoSHow to cite this article: Ryou, J. et al. Monolayer MoS2 Bandgap Modulation by Dielectric Environments and Tunable Bandgap Transistors. Sci. Rep.6, 29184; doi: 10.1038/srep29184 (2016).


There is an error in the correction published on April 24, 2015. All instances of IFN-β in the published article should be replaced by IFN-α.