Why are we reducing graphene oxide
Restoration of thermally reduced graphene oxide through selenium doping at the atomic level
- Chemical engineering
- Electronic properties and devices
The use of reduced graphene oxide (rGO) suffers irreparable damage due to topological defects and residual heteroatoms that degrade the inherent properties of graphene. In order to restore its electrical transport properties, chemical doping by charge transfer with d-electron-rich heteroatoms has been proposed. Here we report on the effects of selenium doping at the atomic level in rGO. With the help of First Principles calculations, we found out that selenium atoms can be selectively bound at certain points, e.g. B. at the pseudo edge locations of hole cluster defects in the ground plane and at the edge defects of graphs. However, we found that the intrinsic topological defects of the basal plane were unfavorable for binding. Numerous selenium atoms were introduced into the completely amorphized rGO surface, which caused a dramatic change in their electrical transport properties through electron doping. The large metal regions formed by the selenium atoms on rGOs increased the electrical conductivity by 210 S cm –1 at 300 K. In addition, the temperature-dependent conductivities (σ) / σ 20K of selenium-doped rGOs (Se-rGOs)) were almost constant in the temperature range of 20–300 K, which indicates that the charge carrier mobility of Se-rGOs becomes independent of temperature after selenium doping, similar to pure graphene.
Due to its unique properties, such as its high electrical conductivity (~ 10 6 S cm −1 ) and the high intrinsic mobility (200,000 cm 2 V. −1 s −1 ) has attracted great attention in physics, chemistry and materials science modulus of elasticity (~ 1.0 TPa) and exceptional thermal conductivity (~ 5000 W m −1 K −1 ). 1, 2, 3 However, the practical application of graphene is limited due to the difficulty of making it on a large scale. Therefore, a chemical way of producing the so-called graphene oxide (GO) was determined. 3, 4, 5 GO has numerous topological defects and oxygen heteroatoms that distinguish it from pure graphene, and it cannot be restored to its original state even by healing through various reduction processes. The large number of sp 3 - Locations and residual oxygen groups of reduced GO (rGO) negatively influence the intrinsic properties of graphene. 6, 7, 8 It is therefore essential to give rGO properties similar to those of graphene.
It is expected that chemical doping with heteroatoms or molecules that can increase the density of free charge carriers in rGO increases its electrical conductivity. 9, 10, 11, 12, 13 the chemical charge transfer doping of graphene can be divided into surface transfer doping and substitution doping. 11 In general, the surface transfer doping does not destroy the graphene structure and is therefore reversible. In contrast, substituting doping refers to the permanent substitution of carbon atoms in the aromatic hexagonal lattice of graphene. The surface transfer doping takes place via the charge transfer from the adsorbed dopant to graphene and is closely related to its density of states (DOS). Therefore, adsorption of electron-withdrawing molecules leads to p-type doping, while adsorption of electron-donating molecules induces n-type doping. Similarly, the p- or n-type doping can be induced by substitution with dopants each having fewer or more valence electrons. In a large number of studies, substitutional doping of carbon-based materials with nitrogen, 9, 10, 14, 15 Boron, 14, 15, 16 Sulfur, 17, 18, 19, 20 and phosphorus 20, 21, 22, 23 reported . These types of dopants affect the carbon crystal structure by acting as defects while providing the free charge carrier density and surface electroactive properties. In addition, surface transfer doping using p-type dopants such as NO 2 -, Br 2 - and I. 2 -Steam, 24, 25, 26, 27, 28 and n-type dopants such as K and Rb, 27, 28 performed. 29, 30, 31, which leads to greatly improved electrical properties. However, since these dopants have air instability and toxicity, their use in practical applications is limited.
Selenium is a polyatomic, non-metallic chalcogen with a large atomic size (in period 4 together with potassium and bromine) and high polarizability and is d-electron-rich (in group 16). 32 Because of these properties, selenium is a promising potential dopant for improving the electrical transport properties of rGOs. A number of studies have recently been published on the catalytic effects of selenium-doped carbon materials in reducing oxygen 33, 34 as well as the electro-oxidation of methanol 35 and ethanol reported. 36 Although these results indicated the presence of a direct chemical bond between selenium and carbon, the precise configuration of selenium and its effects on the intrinsic properties of carbon materials are still obscured. Further applications of selenium-doped carbon materials are only possible after a deeper understanding of the role of selenium in carbon host structures.
In this study we report on selenium doping at the atomic level and its effects on rGOs. Selenium-doped rGOs (Se-rGOs) were produced by heating GO with elemental selenium powder above the sublimation temperature (> 685 ° C) of selenium. The d-electron-rich selenium atoms were introduced homogeneously over the entire rGO surface, which led to greatly improved electrical transport properties caused by the n-type doping. As a result, the carrier mobility of rGOs was dramatically restored by the introduction of selenium, reaching a level comparable to that of graphene. This study not only provides a basic understanding of the selenium doping of rGOs, but also shows their potential for energy storage applications.
Preparation of Se-rGO
GO was made from natural graphite (Sigma-Aldrich, St. Louis, MO, USA) using the Hummers method. 18 Aqueous GO suspensions were frozen in liquid nitrogen and freeze-dried for 72 h at −50 ° C. and 0.045 mbar using a lyophilizer (LP3, Jouan, France). After freeze-drying, low-density, loose-pack GO powders were obtained. A mixture of the lyophilized GO powder obtained (100 mg) and selenium powder (300 mg) was in a tube furnace from room temperature to 800 ° C with a heating rate of 10 ° C min –1 heated and the temperature was maintained for 2 h. An Ar flow rate of 200 ml min –1 was used. The obtained product (Se-rGO) was stored in a vacuum oven at 30 ° C. rGO was made using the same procedure without the elemental selenium powder.
The morphologies of Se-rGO were observed by atomic force microscopy (AFM, NT-MDT, Russia) with an NSG-10 cantilever (NT-MDT, Russia) in semi-contact mode and FE transmission electron microscopy (TEM; JEM) -ARM-200F , JEOL, USA). The chemical states of Se-rGO were determined by X-ray photoelectron spectroscopy (XPS; PHI 5700 ESCA) using monochromated Al K α Radiation (h ν = 1486, 6 eV) was examined. Elemental analysis was carried out using an elemental analysis 1112 instrument (CE instrument, Italy). Raman measurements were carried out using an NTEGRA spectrometer (NT-MDT, Russia) equipped with a 473 nm (2.62 eV) laser in backscattering configuration. The spectral resolution was ~ 2 cm –1 for a grid with 600 grooves per mm. A × 100 objective (NA = 0.9) provided a laser spot with a diameter of approx. 330 nm. The laser power was kept well below 0.3 mW for non-destructive Raman measurements.
All geometric optimizations and DOS calculations were made with the Vienna ab initio simulation package software 37 of the framework for density functional theory. We used a plane wave basis set with a kinetic energy cutoff of 500 eV and a projector-enhanced wave pseudopotential 38, which was provided by the Vienna ab initio simulation package. One from Perdew-Burke-Ernzerhof 39 proposed generalized gradient approximation was adopted to correct the exchange correlation energy. To calculate the adsorption energies, we constructed a model of a graphene structure with 72 and 192 carbon atoms in a single layer. A zigzag graphene nanoribbon structure was used to describe the edge adsorption, which contained the same amount of carbon atoms as the graphene model. A 15 Å vacuum plate was inserted to avoid interaction between the graphene layers. The adsorption energy E ad was defined as
where E C + Se, E C., E Se are the free energies of Se-adsorbed graphene, perfect graphene and the chemical potential of Se. For the DOS calculation, 18 k-points were evenly distributed along the graphene nanoribbon, the structures of the latter being completely relaxed until the system energies converged to <0.05 eV.
Characterization of electronic transport properties
Se-rGO and rGO were dispersed in DMF and placed on a 300 nm SiO 2 / highly p-doped Si substrate deposited. Electrode fabrication was performed by standard electron beam lithography, evaporation of Cr / Au (5/120 nm), and lift-off methods. The temperature dependence of the I - V characteristics was measured using four probe configurations in a Janis cryosystem with a power source (6221, Keithley) and a nanovoltmeter (2182A, Keithley). The temperature dependence of the conductivity of Se-rGO and rGO was investigated between 20 and 300 K by four-probe measurements with a Keithley 6221 / 2182A Delta Mode System. For Se-rGO and rGO, currents of 100 and 10 nA were applied.
The electrochemical properties of Se-rGO and rGO as anode materials for Li-ion batteries were evaluated using a WonATech automatic battery cycler and CR2032 coin cells. The working electrodes were made by mixing the active material (80% by weight) with conductive carbon (10% by weight) and polyvinylidene fluoride (10% by weight) in N-methyl-2-pyrrolidone. The resulting slurries were evenly applied to Cu foil. The electrodes were dried at 120 ° C. for 2 hours and roller-pressed. The button cells were assembled in a glove box filled with argon, using a composite electrode with metallic lithium foil and 1 M LiPF 6 (Aldrich, 99, 99%) dissolved in a solution of ethylene carbonate / dimethyl carbonate (1: 1 v / v) was used. as an electrolyte. The cells were galvanostatically between 0.01 and 3.0 V against Li / Li + clocked at different current densities.
The air electrode of the Li-air battery consisted of catalyst materials (Se-rGO or rGO) and the polyvinylidene fluoride binder in a ratio of 8: 2 (w / w). The slurry was prepared using N-methyl-2-pyrrolidone and coated on the Ni mesh substrate. Li metal (3/8 inch diameter), a glass fiber separator (Whatman GF / D microfiber filter paper, 2.7 µm pore size), and the fabricated air electrodes were assembled in a Swagelok cell. 1 M lithium bis (trifluoromethylsulfonyl) imide dissolved in tetraethylene glycol dimethyl ether was used as the electrolyte. The cells were operated at a constant external pressure of pure oxygen at 770 torr.
Results and discussion
Se-rGO samples have a thickness of ~ 15 nm, a lateral size of several micrometers, and numerous waves on their surface (Figure 1a). This indicates morphologies that resemble those of rGOs reduced at 800 ° C in the absence of elemental selenium (Supplement, Figure S1). TEM and electron energy loss spectroscopy mapping images of Se-rGOs shows that a large amount of selenium atoms was doped over the entire surface of the carbon host structure of rGOs (Figure 1b-d). In addition, Scan-TEM images (STEM) of Se-rGOs and rGOs show clear differences, namely that Se-rGOs have a relatively rough surface, which is indicated by the presence of numerous bright spots (Figure 1e and Figure S2). The presence of selenium atoms on the rGO surface (with an atomic selenium content of ~ 13%) is also confirmed by energy dispersive X-ray spectroscopy (EDS; supplementary Figure S3). However, the high-resolution TEM image (Figure 1f) of Se-rGO shows no aggregates or nanoparticles on its surface, suggesting selenium doping of rGO at the atomic level. Notably, the carbon structures of Se-rGOs and rGOs were strongly amorphized (Figure 1f and supplementary Figure S4). Enlarged high-resolution TEM images of Se-rGOs also show selenium atoms (Figure 1g). The large d-electron-rich selenium atoms (radius: 120 pm) are shown as bright spots that are observed on the strongly amorphized Se-rGO surfaces. The chemical states of Se-rGOs were characterized by XPS (Figure 2a - c). The C 1 spectrum shows several distinct peaks due to CO and CSe bonds with a centering of 285.8 eV and those due to CO bonds with a centering of 290.5 eV. In addition, the spectrum shows the main CC binding peak at 284.5 eV (Figure 2a), which is similar to the value observed for rGOs (Figure S5a). Oxygenated functional groups were also found in the O 1 spectrum, which is indicated by the peaks due to CO and CO bonds at 533.3 and 530.5 eV, respectively (Figure 2b). The C / O atomic ratio of Se-rGOs was ~ 10.9, which was higher than the corresponding value of 9.2 for rGOs, indicating the reduced nature of Se-rGOs. Selenium atoms were mainly in the form of C-Se-C bonds, which is indicated by the peak centered at 56.3 eV, with the peak at 59.3 eV due to the slight Se-O bonds (Figure 2c) chemically doped on the rGO -surface. The C / Se atomic ratio was calculated to be ~ 40.7, which is much lower than the value suggested by the energy dispersive X-ray spectroscopic data. To determine the selenium content more precisely, an elemental analysis was performed and we obtained ~ 15 wt% selenium content for Se-rGOs in accordance with the XPS results. Fourier transform infrared spectroscopy data also confirm the presence of the C-Se-C bonds (Figure S5b). It is noteworthy, however, that the Raman spectra have similar D / G intensity ratios (I. D. / I G ) as well as integrated sp 3 - and sp 2 -Area ratios (A sp3 / A sp2 ) for Se-rGO and rGO (Figure) 2d, supplementary figures S6 and S7). The expanded Raman spectra of Se-rGO and rGO and their integrated area ratios of sp 3 and sp 2 (A. sp3 / A sp2 ) are shown in the supplementary figures S6 and S7. 40 The D band corresponds to the disturbance in A 1g- Breathing mode of the six-membered aromatic ring near the basal edge, while the G band reflects the hexagonal structure associated with the E 2g- Vibration mode of the sp 2 -hybridized carbon atoms. Hence the mean the similar ID / IG- Values that the selenium doping slightly influences the hexagonal carbon structure, which suggests that the doping could be close to the surface transfer doping.
Direct observation of selenium atoms doped on rGO. Topographic images of Se-rGO captured using ( a ) Atomic force microscopy and ( b ) FE-TEM. Se-rGO has a thickness of ~ 15 nm and a lateral size of several micrometers and has a high aspect ratio. Electron Energy Loss Spectroscopy (EELS) for mapping images of ( c ) Carbon and ( d ) Selenium. Selenium atoms were doped homogeneously over the entire surface of rGO. ( e ) Scanning electron micrograph of Se-rGO. Numerous bright spots are observed on the entire surface of Se-rGO. ( f ) High resolution transmission electron microscopy (HR-TEM) image of Se-rGO showing its highly amorphous surface. ( G ) HR-TEM images with higher magnification, obtained from f . The atoms observed as bright spots with a larger size than the others are considered to be selenium atoms.
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Chemical configuration and microstructure of Se-rGO. ( a ) C 1 s, ( b ) O 1 s, ( c ) Se 3 d X-ray photoelectron spectra of Se-rGO. Se-rGO has chalcogen heteroatoms that are involved in CO, CO, CSe and SO bonds. Most Se atoms are in the form of C-Se bonds. ( d ) Raman spectra of rGO and Se-rGO.The two Raman spectra are similar, as are their I. D. / I G- Intensity ratios, indicating that the carbon microstructure was not changed by selenium doping.
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More specific information about the chemical bond was obtained from density functional theory calculations (Figure 3). Various defect sites of rGOs, including intrinsic topological defects, pseudo edge sites of hole cluster defects in the ground plane of graphene, and graph edge defects, are considered to be possible adsorption sites for selenium atoms (Figure 3a). As shown in Figure 3a, the configurations 1–5 represent the selenium adsorption over the carbon atoms (1), CC bonds (2) and carbon rings (3) at the pseudo-edge locations of hole cluster defects in the graphene ground plane (4) and on the graphene Edge defects (5; additional figure S8). The calculated selenium adsorption energies for each adsorption site are given in Figure 3b. Note that several points are described for configurations 1–3, as there are several different local environments for selenium adsorption, which are composed of 5-, 6-, and 7-membered carbon rings (see the Supporting Information for details). The most stable adsorption sites in configurations 1–3, indicated by red circles in Figure 3b, were carbon atoms 577 of configuration 1, 77 of configuration 2, and the five-membered carbon ring of configuration 3. However, all selenium adsorption energies were in configurations 1–3 it positive, indicating that the binding of Se atoms to intrinsic topological defect sites such as monovacancer, divacancer and Stone-Wales defect is unfavorable. These results agree with the previous calculation reports on the substituting sulfur doping of graphs 41 and the surface sulfur doping of defective graphene 18 agree, both of which show a positive binding energy for sulfur adsorption at the intrinsic defects of graphene. On the other hand, the binding energies of selenium atoms are negative both at pseudo-edge hole cluster defects in the ground plane and at edge defects of graphene, which indicates that Se atoms can only be bound at the pseudo-edge and spontaneously at edge defects, which are relatively unstable are. In addition, the negative energy of selenium adsorption implies a chemical bond with graphene at the atomic level, not just physical adsorption. We also considered the possibility of selenium adsorption in the presence of oxygen-containing functional groups (Figure S9). We find that selenium atoms are unlikely to bind to oxygen directly. However, their individual adsorption is possible in almost oxygen-containing functional groups, in agreement with the XPS result, which showed an SeÀO bond (Figure 2b). In addition, we calculated the DOS for H- and Se-terminated graphene nanoribbons to investigate the effect of Se doping on the electronic properties (Figure 3c and d). The detailed model structures used for DOS calculations are shown in the supplementary figure S8e. As shown in Figure 3c and d, the Fermi content of Se-terminated graphene nanoribbons is 0.5 eV higher than that of H-terminated ones. Since selenium has more valence electrons than carbon atoms, they are naturally transferred from selenium to rGO, hybridize with the carbon atoms of rGO and assume high-energy states, which leads to a positive shift in the Fermi level and makes Se-rGO n conductive material . This effect is also used for doping with K 42 and NH reports. 43 Since n-doped graphene typically has an increased electrical conductivity due to the introduction of charge carriers 43 can Se-rGOs have improved electrical conductivity compared to rGOs.
Density functional theoretical calculations for the configuration of selenium on rGO. ( a ) Various possible adsorption sites on defective graphene (1 - atom, 2 - bridge, 3 - ring center, 4 - hole defect and 5 - edge) and ( b ) Se - adsorption energies on it. ( c ) DOS of Se-terminated GNR (scarlet) and PDOS of Se (navy) and ( d ) DOS of H-terminated GNR. The dashed lines in c, d indicate the Fermi level, while the occupied energy levels are filled in in color.
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A temperature-dependent current-voltage characterization (I-V) of rGO and Se-rGO was carried out (Figure 4a and b). For both samples, the IV curves show a highly symmetrical and linear behavior in all temperature ranges, the slope of the IV curves increasing with temperature. The curves show an increase in conductivity with temperature. Interestingly, the current of Se-rGO is two orders of magnitude higher than that of rGO, and the slope of the IV curve for Se-rGO changes little despite the temperature fluctuation from 20 to 300 K. The dependent conductivities (σ (T)) of rGO and Se-rGO were obtained when the temperature was varied continuously from 20 to 300 K (Figure 4c). The σ (T) of Se-rGO is 1–2 orders of magnitude larger than that of rGO. The defects and functional groups cause an increase in dipole polarization and thus reduce the conduction path length in GO. 44 For rGO, the dipole polarization is decreased due to the decreased amount of defects and oxygenated functional groups, which increases the probability of electron jump. Therefore, the electronic transport mechanism of rGO has been attributed to variable jumping. 45 The rGO follows the Efros-Shklovskii jump with a variable range (σ = σ 0 exp (- T 0 / T) 1/2, where σ 0 and T 0 87.8 S cm –1 or 508.2 K) Known for the transport properties of rGO with different proportions of sp 2 -Carbon (blue color in Figure 4c). 46 The electrical transport properties of Se-rGO were not adjusted by the Efros-Shklovskii jump model with variable range, but by the fluctuation-induced tunneling mechanism 47 (red color in Figure 4c) well explained. The & sgr; (T) in the sway induced tunnel model is expressed as & sgr; =? exp (- T 1 / (T + T 0 )), where T 0 and T 1 200 and 96 K, respectively. T 0 is the starting temperature of the thermally activated line and T 1 is a measure of the energy that an electron needs to cross the insulator gap between conductive areas. These values indicate that Se-rGO has large metallic areas that are separated by a relatively small isolation gap. These metallic areas produced by selenium doping increase the conductivity at 300 K to 210 S cm –1. This is a very large conductivity value compared to rGO (obtained with N 2 H 4, ~ 2 S cm –1 ) 45 and is comparable to The ultra-high volume conductivity of GO papers is reduced by metal iodides. 48 Note that the variation of σ (T) in Se-rGO is quite small (Figure 4d). The σ / σ 20 K- The value of rGO increases by 41.5 times when the temperature rises from 20 to 300 K, while that of Se-rGO remains almost constant (1.3 times increase from 20 to 300 K). This means that the charge carrier mobility becomes independent of temperature due to the selenium doping; The mobility of the carriers is similar to that of pure graphene. 49
Temperature-dependent electron transport properties of rGO and Se-rGO. Temperature-dependent electrical transport properties of ( a ) rGO and ( b ) Se-rGO, measured between 20 and 300 K in steps of 20 K with inserted optical images of the configurations with four electrodes. The slope of the IV curves increased with temperature, indicating an increase in conductivity with temperature. ( c ) Temperature dependence of rGO and Se-rGO conductivities obtained from IV curves and VRH and FIT models (fluctuation-induced tunneling). rGO and Se-rGO were well matched by VRH and FIT models, respectively. ( d ) Temperature versus σ / σ 20 K- Diagrams for rGO and Se-rGO. The above curve for Se-rGO is almost flat in the temperature range from 20 to 300 K (1.3-fold increase), which indicates that the carrier mobility becomes temperature-independent due to the selenium doping, which implies a restoration of mobility.
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Furthermore, the effects of selenium doping on the electrical transport properties of monolayer rGO were investigated directly by in situ measurements at 300 K (Figure S10). The sample was loaded into a doping tube with selenium powder placed on the bottom. The tube with the sample was evacuated with a rotary pump (<3 × 10 –2 Torr; Additional figure S10a). When the tube sheet was heated, the conductivity of rGO rose abruptly from 0.27 to 0.58 ½ cm –1 and then became saturated (Figure S10b). In addition, we examined the voltage generated by the gate (V G ) dependent conductivity at 300 K and observed the behavior of a graph-like field effect. The minimum conductivity (& sgr; min ) from rGO was close to V G = 60 V (Supplementary figure S10c). However, this value shifted in the direction of negative V after the selenium doping G which implies that this doping was n-type. Based on the electrical transport properties, we conclude that the chemisorbed selenium atoms on rGO lead to electron doping of the latter and improve its electrical conductivity.
The superior electrical transport properties of Se-rGOs lead to significantly improved electrochemical properties of energy storage systems. As an anode material for Li-ion batteries, Se-rGOs have a much higher power and capacity compared to rGO (Figure 5a - c). The reversible capacity of Se-rGO is 570 mAh g –1 at a current density of 1 C (372 mAh g –1 ; Figure 5a), which is much higher than the 310 mAh g observed for rGO –1 (Figure 5a). The increased capacity of Se-rGO suggests that selenium atoms contribute to the reversible storage of Li ions, which is supported by the change in the galvanostatic discharge / charge profile. The same test carried out at low temperature (0 ° C) shows a more pronounced capacity gap between Se-rGO and rGO, as shown in the supplementary figure S11. In the low temperature cell test, the reversible capacity of Se-rGO is 360 mAh g –1which corresponds to ~ 63% of the capacity at 25 ° C. In contrast, rGO has a reversible capacity of ~ 140 mAh g -1 onwhich corresponds to ~ 45% of the capacity at 25 ° C. These results suggest that the improved electrical properties of Se-rGO greatly affect the electrochemical properties at low temperatures. Highly stable capacities of 165 mAh g –1 could be achieved at a high current rate of 50 C, which is approximately seven times the value of rGO anode materials (Figure 5b). In addition, the high speed performance of Se-rGO is significant compared to graphene-based electrode materials and nitrogen-doped graphene nanosheets for Li-ion batteries. 50, 51 After 200 cycles with a successive increase in current density of 1 C (372 mA g –1 ) to 50 C (18,600 mA g –1 ) the specific capacity of Se-rGO returns to its initial value of approx. 50. 460 mAh g –1, which indicates good reversibility. In addition, it was found that the cycle stability is maintained over 1000 repetitive cycles, which shows a Coulomb efficiency of almost 100% (Figure 5c).
Electrochemical properties of rGO and Se-rGO as anode and catalyst for Li-ion and Li-air batteries. ( a ) Galvanostatic discharge / charge profiles of Se-rGO and rGO at a current density of 1 C (372 mA g –1 ), ( b ) Performance of Se-rGO and rGO with continuously increasing current densities of 1 C (372) mA g –1 ) up to 50 C (18 600 mA g –1 ) with subsequent reversal to 1 C (372 mA g –1 ) and ( c ) Cycle performance of Se-rGO as a Li-ion battery anode for 1000 repetitive cycles at a current density of 3 C (1116 mA g –1 ). ( d ) Discharge / charge profiles of rGO- and Se-rGO-based air electrodes in a potential range of 2.0–4.5 V at a constant current rate of 200 mA g –1 for Li - O 2 Cells (inset of a modified cell) Li-O 2 Swagelok-type cell). ( e ) Comparison of the electrochemical properties of rGO- and Se-rGO-based Li-O 2 - Cell air electrodes with and without LiI catalyst. The capacity was up to 500 mAh g -1 used . The dotted line gives the theoretical voltage of formation of Li 2 O 2 at 2.96 V versus Li / Li + on .
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We thought that Se-rGO was a potential catalyst for Li-O because of its high energy density 2 -Batteries could be. However, we found that such batteries were highly polarized while charging, resulting in poor cycling stability (Figure 5d and e). The above behavior was due to the formation of non-conductive Li 2 O 2 due during discharge. If a conventional ether-based electrolyte (i.e., tetraethylene glycol dimethyl ether) is used for Li-O 2 Batteries were used, amorphous Li 2 O 2, that entangled on the carbon surface, and crystalline Li 2 O 2 observed with toroidal morphology. The preference between amorphous and crystalline Li 2 O 2 is significantly influenced by the properties of the electrolyte, the current rates and the depth of discharge. 52, 53 In the fully discharged state, as shown in FIG. 5d, nucleated Li 2 O 2 enough time to grow into large toroidal particles with high crystallinity. Based on theoretical studies of the charge transport properties of Li 2 O 2 its intrinsic ionic and electronic conductivities became ~ 10 –19 S cm –1 calculates what is unacceptable to charge conduction. 54 Therefore, the charge profile of Se-rGO showed a large polarization despite the increased electrical conductivity, that of rGO due to the low conductivity of the large toroidal Li 2 O 2 -Particle is similar. In this regard, a specific electrochemical protocol should be used to assess the effect of Se-rGO in Li-O 2 -Cells clearly to be detected. So we used a capacity limited discharge to discharge surface-based Li 2 O 2 to create. Since the surface-derived Li 2 O 2 comes in direct contact with the Se-rGO surface, the improved electrical properties of Se-rGO could reduce the charge polarization of Li-O 2 - strongly affect cells. As shown in Figure 5e, the air cells tested under a protocol in which the discharge capacity was set to 500 mAh g -1 limited was, an obvious difference in their charge profiles. Overall, the Se-rGO-based electrode showed a lower charge potential than rGO; Especially the beginning part (~ 200 mA h g –1 ) had a charge potential below 4.0 V. Based on this result, we were able to demonstrate the effect of selenium doping in rGO. When a soluble LiI catalyst was used, the Se-rGO-based electrode showed an abrupt decrease in charge polarization near 0.5 V (Figure 5e), which was superior to previous reports on rGO electrodes in terms of energy efficiency. 55, 56, 57 These results thus show the possibility of using Se-rGO-based electrodes for Li-O 2 -Batteries to use.
In this work we investigated the effects of selenium doping in rGO. Numerous selenium atoms were homogeneously doped onto the surface of rGOs. it caused little structural change, indicating surface transfer doping. High resolution TEM images showed no aggregates or nanoparticles on the surface of rGOs, indicating that selenium atoms were introduced at the atomic level. On the basis of first principles calculations, we found that the selenium atoms are mainly bound at the surface edge defects of highly amorphized rGO, which leads to an improvement in electrical conductivity due to the n-type doping. In contrast, all of the intrinsic topological defects of the graphene ground plane showed an unfavorable selenium atom binding energy. The doped selenium atoms changed the electrical transport properties of rGO. The properties were well matched by the fluctuation-induced tunneling mechanism, and the results indicated the presence of a large metallic area separated by a relatively small isolation gap. The σ value of Se-rGO reached 210 S cm –1 at 300 K, while the σ / σ 20 K- Parameter was nearly constant (1.3-fold increase from 20 to 300 K), suggesting that Se-rGO is mobile was similar to graph. In addition, due to the effects of selenium doping in rGO, Se-rGO showed superior electrochemical performance as an anode for Li-ion batteries and as a catalyst for Li-O 2 -Batteries.
Additional information supplements the paper on the NPG Asia Materials website (//www.nature.com/am).
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