A single-phase CaCu3Ti4O12 with a unit cell ∼7.393 Å was detected in all of the studied ceramic samples. The grain sizes of sintered Cax1−CdxCu3Ti4Oy12−2Fy2 ceramics were significantly enlarged with increasing dopant levels. Liquid-phase sintering mechanisms could be well matched to explain the enlarged grain size in the doped ceramics. Interestingly, preserved high dielectric permittivities, ∼36 279–38 947, and significantly reduced loss tangents, ∼0.024–0.033, were achieved in CdF2 codoped CCTO ceramics. Density functional theory results disclosed that the Cu site is the most preferable location for the Cd dopant. Moreover, F atoms preferentially remained close to the Cd atoms in this structure. An enhanced grain boundary response might be a primary cause of the improved dielectric properties in Cax1−CdxCu3Ti4Oy12−2Fy2 ceramics. The internal barrier layer capacitor model could well describe the colossal dielectric response of all studied sintered ceramics.Ca 2 defect clusters result in enhancement of dielectric properties of the Cax1−CdxCu3Ti4Oy12−2Fy2 ceramics.CdF Focusing on capacitor applications, the colossal dielectric properties of polycrystalline ceramics, such as the codoped TiO2,1,2,17,20–22 codoped SnO2,3 and ACu3Ti4O12  and enhancing temperature/frequency stability. Reports related to CaCu3Ti4O12 (CCTO) based ceramics have been continuously published.4–9,26–29 Due to its high ε′, numerous researchers have proposed an internal barrier layer capacitor (IBLC) structure, consisting of semiconducting grains and insulating grain boundaries (GBs). This was the first model used to explain the colossal dielectric response in ACu3Ti4O12 ceramics.9–13,16,25–27 However, other factors, e.g., sample-electrode contact,27 charge defects or ionic defects,28 also have essential roles in the dielectric and electrical responses of ACu3Ti4O12 ceramics that cannot be neglected. High ε′ and non-ohmic properties measured in CCTO ceramics mainly originate from the insulating GB layers in IBLC structures.30Over the past decade, the survey of colossal dielectric permittivity ,27 ,8 and 7 codoped CCTO ceramics, it was found that the dielectric response is greatly enhanced in these ceramics as well. There are few reports related to Cd2+ doped CCTO ceramics.33,34 Additionally, experimental work and intensive discussions have not resulted in a clear understanding of these materials. Cd2+ and F− codoped CCTO ceramics are the focus of the current study. The roles of Cd2+ and F− in CCTO are systematically investigated. For high concentrations of Cd2+ ions doped in the CCTO, various reports showed that Cd2+ ions replace Ca2+ ions entirely.12,31 However, at low concentrations, Cd2+ may replace ions at other sites, such as those of Cu2+. Due to it charge state, Cd2+ can occupy ion sites with coordination numbers (CNs) of either 4 or 12,35 which are host sites of either Cu2+ or Ca2+, respectively. Therefore, a computational method, e.g., density functional theory (DFT) calculations, should be used to determine the most stable location of Cd in the CCTO host.The most severe problem of CCTO for capacitor applications is its rather high tan x1−CdxCu3Ti4Oy12−2Fy2 ceramics with x = y = 0, 0.10, and 0.15 were fabricated using a conventional solid-state reaction (SSR) method. We systematically examined both the microstructure and crystal structure of all studied sintered ceramics. First-principles calculation was employed to scrutinize the most stable sites of Cd and F in the CCTO structure. Experimental and computational details along with their results are discussed in subsequent sections.In this work, Ca2.2.1.2 , CuO , CaCO3 , CdF2 , and C2H5OH  which are used in a SSR method to synthesize Cax1−CdxCu3Ti4Oy12−2Fy2  ceramics. Stoichiometric ratios of all-ceramic conditions were weighed and ball-milled in C2H5OH at 150 rpm for 24 h. We heated these mixed materials in an oven at 80 °C for 24 h to vaporize all liquid from the sample materials. The resulting dried precursors were ground. We calcined the mixed powders in air at 850 °C for 12 h and then ground until their particle sizes were very fine. The Cax1−CdxCu3Ti4Oy12−2Fy2 powders were formed into pellet shaped samples. Pellets with diameters of ∼9.5 mm and thicknesses of ∼2 mm were sintered at 1050 °C for 6 h. Sintered Cax1−CdxCu3Ti4Oy12−2Fy2 ceramics with x = y = 0, 0.10, 0.15 are referenced as CCTO, CdF10, and CdF15 ceramics, respectively.We initially prepared the TiO2.2.x1−CdxCu3Ti4Oy12−2Fy2 ceramics were examined using an X-ray diffractometer . In our XRD measurements, we used 2θ ranging from 20 to 80°. All XRD spectra were analyzed using the Rietveld refinement method employing X'Pert HighScore Plus software, Version 3.0e. The adjusted parameters and coefficients in the Rietveld method are global parameters (zero shift and background parameters), scale factor, unit cell, atomic coordinates, and profile parameters. A desktop scanning electron microscope  was used to identify microstructural changes. An energy-dispersive X-ray spectrometer (EDS), included in the MiniSEM, was used to determine the chemical composition of the grains and GBs of sintered ceramics.The phase composition and crystal structures of all Ca2.3.x1−CdxCu3Ti4Oy12−2Fy2 ceramics were measured using a KEYSIGHT E4990A with an oscillation voltage (Vrms) of 0.5 V. The dielectric properties were studied over the frequency and temperature ranges of 40 to 107 Hz and −60 to 210 °C, respectively. The methods for calculation of dielectric and electrical parameters are given elsewhere.15Prior dielectric and electrical measurements, the surfaces of the ceramic samples were polished. These polished surfaces were coated with a conductive silver paint . Then all ceramic samples were fired in air at 680 °C for 0.5 h. Dielectric and electrical properties of Ca2.4.Ab initio Simulation Package (VASP)36 was used to calculate the total and formation energies of the Cd and F codoped CCTO ceramics. We used the Perdew–Burke–Ernzerhof (PBE)37 form of the exchange-correlation functional. For CCTO, the pseudopotentials of Ca, Cu, Ti, and O are given in ref. k-points samplings in the reciprocal space were employed. In the present study, a 5 × 2 × 1 CCTO supercell containing a total of 400 atoms was used.The Vienna 3.x1−CdxCu3Ti4Oy12−2Fy2 ceramics in the current study was completed using XRD. As revealed in Im3̄ no. 204 space group38 and presented a single-phase of CCTO.All structural parameters obtained from XRD measurements are summarized in R-factors, i.e., the profile R factor (Rp), weighted profile R factor (Rwp) and expected R factor (Rexp) of all Cax1−CdxCu3Ti4Oy12−2Fy2 ceramics were below 9.5%. As a result, low goodness-of-fit (GOF) values ∼1.12–1.27 were attained. As summarized in a) of sintered Cax1−CdxCu3Ti4Oy12−2Fy2 ceramics were similar, approximately ∼7.393 Å. The a values obtained in this work are comparable to previously published data.26–28,38 Co-substitution of Cd2+ into Ca2+ sites and F− into O2− sites should cause a unit cell of CCTO to decrease significantly. This is because the ionic radius of Cd2+ (r12 = 1.31 Å) for the CN = 12 is smaller than that of Ca2+ (r12 = 1.35 Å).35 Concurrently, the ionic radius of F− (r4 = 1.31 Å) is also smaller than that of O2− (r4 = 1.38 Å).35 However, the average size of the unit cell is constant. Although replacement of Cd2+ ions into all Ca2+ sites (CdCu3Ti4O12) can correctly form a body-centered cubic structure within an Im3̄ space group,10–12 substitution of a small amount of Cd2+ into Ca2+ sites might yield different results. It is reasonable to suggest that a small concentration of Cd2+ could replace Cu2+ in the CCTO structure. This is because the ionic radius of Cd2+ (r4 = 0.78 Å) for the CN = 4 is larger than that of Cu2+ (r4 = 0.57 Å). Consequently, if the larger Cd2+ ions replace Cu2+ at smaller sites and F− ions occupy O2− sites, this may allow the size of the unit cell to remain constant.No possible impurity phases were detected in any of the sintered ceramics. Rietveld refinement was used to calculate the vital parameters obtained from the XRD data of each ceramic sample. x1−CdxCu3Ti4Oy12−2Fy2 ceramics are shown in SEM and ImageJ software were used to collect data and analyse the microstructure of ceramic samples. Microstructural imagery and size distributions of sintered Ca6–8,27 This vastly increased grain size of the doped samples, compared to CCTO ceramic, might have been caused by liquid-phase sintering mechanisms. Generally, a liquid-phase can form at junctions between particles or GBs, resulting in a much increased diffusion rate. The liquid-phase might have originated from the direct melting of materials used in sample preparation or a reaction at a eutectic point between the raw metallic starting materials. As previously reported, a eutectic point between CuO–TiO2 could be formed at 950 °C.39 However, this reaction does have a great impact on the enormously expanded grain sizes of doped ceramics. Additionally, melting of CdF2 dopant during the sintering might not have been the cause of size expansion. According to previous research, the melting temperature of CdF2 is ∼1100 °C,40 much higher than the sintering temperature of CdF2 codoped CCTO ceramics (1050 °C). Hence, it is reasonable to propose that liquid-phase sintering mechanisms might be associated with replacing Cd2+ ions in the Cu2+ sites of CCTO structure during calcination at 850 °C for 12 h. Thus, a eutectic liquid of CuO–TiO2 is much more likely created due to surplus Cu ions, resulting in significantly enlarged grain sizes. This result is in agreement with our previously published work with  and  codoped CCTO.9,27The microstructural results reveal a significant increase in the grain sizes of the ceramic samples. A great increase in grain size due to doping was also reported in earlier published papers.2+ concentration detected inside the grains was much higher than that at the GBs, indicating the incorporation of Cd2+ into the CCTO lattice. An abundance of a Cu-rich phase was detected along the GB layers, as was seen in the EDS result. The presence of a Cu-rich phase directly relates to liquid-phase sintering mechanisms as a source of the increased grain sizes of the CdF2 codoped CCTO samples.6–8,27The chemical constitutions of the GBs and grains of the CdF15 ceramic were examined using EDS. As shown in 2 codoped CCTO ceramics were investigated under various conditions. The values of the parameters obtained are summarized in ε′ at 20 °C of all sintered ceramic samples is shown in 6 Hz. This result corresponds to its tan δ spectrum, as shown in the inset of Dielectric and electrical properties of CdFet al.,41 dielectric relaxation observed in the low-frequency range indicates the influence of DC conduction. In general, DC conduction in dielectrics can occur during the polarization process due to the long-range migration of charge carriers. It is notable that the frequency-independence of ε′ was achieved by doping with CdF2. The tan δ independence of frequency was observed in CdF10 and CdF15 ceramics over the frequency range of 40 to 105 Hz. The ε′ value of CdF10 and CdF15 ceramics slightly increased compared to CCTO ceramic. This small increase in ε′ is closely associated with an increase in grain size. Therefore, the dielectric response of CdF2 codoped CCTO ceramics can be explained with an IBLC model.6,26,27 At 1 kHz and 20 °C, the ε′ values of CCTO, CdF10, and CdF15 were 35 047, 36 279, and 38 947, respectively. In the current work, not only was the ε′ value highly conserved, but the tan δ of CdF2 codoped CCTO ceramics was also reduced, as revealed in the inset of δ values of CdF10 and CdF15 ceramics were respectively reduced to 0.033 and 0.024, compared to a value of 0.218 for the CCTO ceramic. The decrease in tan δ may have been induced by an enhanced GB response, as discussed in the next section. According to Maxwell–Wagner polarization, the ε′ value is reliant upon the GB capacitance (Cgb).8 Although the grain sizes of the CdF10 and CdF15 ceramics were much larger than that of the CCTO sample, their ε′ values were only increased by ∼3.5–11.1%. Hence, the ε′ value of Cax1−CdxCu3Ti4Oy12−2Fy2 ceramics might depend on both grain size and a change in Cgb due to CdF2 doping.8 Temperature dependences of ε′ and tan δ at 103 Hz are shown in ε′ of the CdF10 and CdF15 ceramics was found to be more stable than that of the CCTO ceramic in the temperature range below 120 °C. It is similar to that the influence of metastable-insulating layer at GB regions only affects the dielectric response of these codoped ceramics in the temperature range of −60 to 120 °C. Besides, the inset of δ of the CdF10 and CdF15 ceramics were much smaller than that of the CCTO sample through the temperature range of 60 to 120 °C. This dielectric result is likely to specifies the influence of co-substitution on improved dielectric response of CCTO ceramic. One possible origin of improved dielectric properties in the codoped CCTO ceramics might be the metastable-insulating phase along GB regions.9,27,42According to Li ε′ value is very high and rather constant over a wide frequency range. Also, the tan δ of the CdF10 sample is indeed low in comparison to the CCTO ceramics. Based on these results, CdF10 was the primary focus. The stability of the CdF10 structure was determined by means of density functional theory calculations. Initially, we needed to test the location stability of a Cd atom in a CCTO host. This led to calculation of the formation energies when a Cd atom is substituted into Ca, Cu and Ti sites, as displayed in the inset of As shown in 20Cd2Cu58Ti80O240 host. As clearly presented in 20Cd2Cu58Ti80O236F4 structure, O was replaced by F atoms at four sites in the Ca20Cd2Cu58Ti80O240 and then the total energy of these structures was evaluated. In the current work, four different positions of F atoms in the Ca20Cd2Cu58Ti80O236F4 structure were considered, i.e., structures I–IV in 20Cd2Cu58Ti80O236F4 is structure I of Based on our DFT results presented in 2+ and F− can enlarge grain sizes and enhance dielectric properties. Although the grain sizes of CdF10 and CdF15 ceramics were greatly increased in comparison with the CCTO sample, the ε′ values of these two samples are only ∼3.5–11.1% larger. In general, the ε′ value of CCTO should increase with its grain size and more than double.9,27,42 However, this case is different. In structure I of 2+–F− atoms can form a small insulating barrier inside the grain and/or near GB regions. This insulating barrier is graphically presented in ε′ value. According to Maxwell–Wagner polarization, the total capacitance of a ceramic formed in an IBLC structure should be consistent with the total polarized charges or range of charge displacement under an external electric field. Therefore, limitation of the region for polarization due to a small insulating barrier might cause the ε′ values to decrease or remain constant. This result is similar to a cases of Li+, F− codoped CCTO ceramics.8In the experimental and DFT results, it was found that co-substitution of Cd2 codoped CCTO ceramics to study the ceramic samples' electrical responses. An impedance complex (Z* = Z′ − jZ′′) plot can be made using the following relationship:ω = 2πf is the angular frequency. C0 = ε0A/d is the capacitance of free space. ε* is the complex dielectric permittivity consisting of a real part (ε′) or dielectric permittivity and an imaginary part (ε′′) or total loss factor. Resistances of grains (Rg) and GBs (Rgb) can be approximated from the nonzero intercept at high-frequency and a diameter of a semicircular arc of the Z* plot, respectively. However, for highly resistive ceramics, a semicircular arc is not formed at room temperature (RT). Thus, we considered Rgb at a relatively high temperature. The Z* plots at 100 °C for CCTO, CdF10, and CdF15 ceramics are given in Rgb values of the doped ceramics are much larger than that of CCTO. Additionally, Rgb significantly increased with dopant concentration. The Rgb values of CCTO, CdF10, and CdF15 ceramics are listed in 4, 2.86 × 105, and 4.32 × 105 Ω cm, respectively. As shown in the inset of δ at low-frequency might be caused by an enhanced Rgb. This result closely agrees with previously published results, which suggest that the trend of Rgb and low-frequency tan δ is opposite.4,6–11,29 According to Mao et al.,43 a Cu-rich phase found at GB layers can reduce tan δ values to relatively low values, ∼0.048–0.083. This reveals that the enhanced GB response is due to a Cu-rich phase observed at GBs. Other researchers also reported a similar tendency.6–8,27Impedance spectroscopy was performed to estimate the resistance of grains and GBs of CdF2+–F− clustering ]. Simultaneously, by considering the high-frequency and relatively low-temperature, nonzero intercept indicates the electrical response inside the grains shift to a lower Z′-axis as soon as the temperature increases [the inset (2)]. This shifting spectrum shows that Rg variations have the same tendency as Rgb in a low-temperature range. This electrical response is usually observed in CCTO and related ceramics, indicating an IBLC structure.4–8,10,11Clearly, the diameter of the semicircular arc of Rg and Rgb can be used to estimate the conduction activation energies of grains (Eg) and GBs (Egb) through a calculation using the Arrhenius law as follows:Temperature-dependent variations of R0 is a pre-exponential term. Also, kB and T are the Boltzmann constant and absolute temperature, respectively. Temperature dependencies of Rgb (open symbols) and Rg (solid symbols) for CCTO, CdF10, and CdF15 ceramics are illustrated in Rg and Rgb follow an Arrhenius relationship well. The slopes of fitted lines were used to estimate Eg and Egb. As shown in Eg of CCTO, CdF10, and CdF15 ceramics were 0.078, 0.105, and 0.102 eV, respectively. Egb of these three samples were also considered. These values for the CCTO, CdF10, and CdF15 ceramics were 0.406, 0.592, and 0.521 eV, respectively.Eg and Egb show the dominance of the IBLC effect in CdF2 codoped CCTO ceramics.43 Interestingly, the Egb of the doped samples increased compared to the CCTO ceramic. This result depicts an increase in the potential barrier height (ΦB) at GBs. Generally, a ΦB located in the non-ohmic region possibly originates from either oxygen enrichment at the GBs and/or decomposition of metal oxide phases or precipitated phases during the sintering process.27The quite divergent values of Rgb observed in the CdF10 and CdF15 ceramics might be primarily caused by two significant factors. The first factor is precipitation of highly resistive insulating phases at GBs.27,42,43 A second factor arises from strong bonding between Cd2+ and F− (structure I of δ might be suppressed to very small values.From the EDS, DFT, and impedance results, it is reasonable to propose that the enhanced 4.2 codoped CCTO ceramics were systematically investigated. Crystal structure results indicate a single-pattern of the CCTO phase. CdF2 addition induces liquid-phase sintering that decomposes along grain boundary layers, resulting in a very great enlargement of grain sizes. DFT calculations indicate that a possible cause for liquid-phase sintering mechanisms might be the substitution of Cd2+ into Cu2+ sites, inducing an enhanced CuO–TiO2 eutectic liquid owing to excess Cu ions. The colossal dielectric properties of Cax1−CdxCu3Ti4Oy12−2Fy2 ceramics were enhanced. Very-high ε′ ∼ 36 279–38 947 and low tan δ ∼ 0.024–0.033 values were accomplished in the doped ceramics. According to an impedance spectroscopy result, the largely reduced tan δ in the codoped CCTO ceramics could be induced by strongly increased grain boundary resistance and conduction activation energy of grain boundary. Therefore, improved dielectric properties in CdF2 codoped CCTO ceramics were created by enhancing grain boundary responses.The structural formations, dielectric properties and electrical responses of CdFJ. B., P. T. and P. S. designed this project. J. B. and P. S. performed experiments and computations. J. B., J. P., J. M., S. K., P. T and P. S. analysed the computational and experimental data. J. B. and P. S. write the manuscript.The authors declare no competing financial interests.


However, industrially discharged H2S has a concentration range of several tens of ppb to several partsper million (ppm), which conventional methods are unable to process.Therefore, advanced and sustainable methods for treating very lowconcentrations of H2S are urgently needed. TiO2-based photocatalysts are eco-friendly and have the ability to treatenvironmental pollutants, such as low-concentration gases, using lightenergy. However, there are no reports on H2S decompositionor oxidation at concentrations below several ppb. Therefore, in thisstudy, we employed anatase/TiO2(B) nanotubes, which havea high specific surface area and an efficient charge-transfer interface,to treat H2S. We successfully reduced 10 ppm of H2S to 1 ppb or less at a kinetic rate of 75 μmol h–1 g–1. The suitability of our method was furtherdemonstrated by the generation of sulfate ions and sulfur (as detectedby X-ray photoelectron spectroscopy and ion chromatography), whichare industrially useful as oxidation products, whereas sulfur dioxide,a harmful substance, was not produced. This is the first study toreport H2S decomposition down to the ppb level, providingmeaningful solutions for malodor problems and potential health hazardsassociated with H2S.Hydrogen sulfide(H For example, the Yanaizu-Nishiyama geothermalpower station in Fukushima Prefecture, Japan, releases up to 60 ppbof H2S,7 and the Nesjavellirand Hellisheidi geothermal power plants in Reykjavik, Iceland, produce153 ppb of H2S.16 However, asset by the World Health Organization, the concentration at which theodor of H2S becomes a nuisance is 4.9 ppb over a 30 minaverage.5 In addition, Collins and Lewisreported that 83% of humans can smell the odor of H2S ata concentration of 30 ppb, 50% at a concentration of 8 ppb, and 6%at a concentration of 1 ppb.17 The odorthreshold of H2S in Japan is 0.5 ppb, that in Netherlandsis 0.3 ppb,18 and that in the United Statesis 0.5 ppb, as set by the Agency for Toxic Substances and DiseaseRegistry.5 As the human sense of smellis extremely sensitive, even if the concentration of the releasedH2S is below the emission control concentration and thereis no health hazard, the emission concentration may be consideredodorous. Therefore, it is necessary to decompose H2S toless than 0.3–1.0 ppb to avoid odor problems.Furthermore, sulfur compounds such as H2) is widely used to remove environmentalpollutants because it is highly stable, inexpensive, and nontoxic.TiO2 can also generate both oxidizing and reducing speciesor directly oxidize and/or reduce contaminants adsorbed on its surfaceunder UV light.21 After Canela et al. reported the first photocatalyticdecomposition of H2S using TiO2,22 several studies reported the gas-phase decompositionof high-concentration H2S (approximately 30–2000ppm) using TiO2-based photocatalysts.24 However, onlya few studies have determined the efficiency of TiO2 photocatalystsfor the low concentrations of H2S (below the emission controlconcentration) discharged from factories without processing. In addition,to the best of our knowledge, there are no reports on decomposingH2S at the ppb level because of the low surface area ofthe catalyst and inactivation due to charge recombination.Photocatalyticdecomposition is an efficient technique. In particular,titanium dioxide , which accumulates and deactivates the catalyst.26 However, Alonso-Tellez et al. suggested that SO42– reacts with photogenerated holes to generate thesulfate radical, resulting in SO42– beingcompletely converted into SO2 gas.27 Grześkowiak et al. showed that H2S is oxidizedto elemental sulfur and SO42– under dryconditions.28 The experimental conditionsin these reports, such as H2S concentration, humidity,and light intensity, were different, yielding different reaction products.While these previous studies focused on analyzing reaction products,research on the reaction mechanism by controlling the reaction conditionsat low H2S concentrations is lacking.29Toelucidate the reaction mechanism of H2 powder, whichis conventionally regarded as highly active TiO2, to achievehighly efficient H2S decomposition. However, H2S at a low concentration of approximately 10 ppm was not completelyprocessed. In addition, silver particles or copper particles depositedon P25 as a cocatalyst were used to decompose low concentrations ofH2S. Although improved activity was observed, completetreatment was not achieved.30 In our previouswork, we focused on anatase/TiO2(B) nanotubes which werefound to be more photocatalytically active than P25 owing to a largesurface area and efficient interfacial charge transfer.32 TiO2(B) is a widely studied35 monoclinicmetastable phase that has attracted attention as a photocatalyticallyactive TiO2 phase. Therefore, in this study, we appliedanatase/TiO2(B) nanotubes to the gas-phase decompositionof H2S and succeeded in decomposing H2S to 1ppb or less, which are concentrations that do not cause odor problems.We also investigated the effects of oxygen, humidity, and sulfur specieson the photocatalyst surface on the reaction mechanism. The resultsshowed that anatase/TiO2(B) nanotubes are a useful purificationmaterial for air pollutants. This photocatalytic method can be proposedas a new approach for decomposing H2S to a level of severalppb, which has not be achieved previously, and solving odor problemscaused by H2S.We previously used the Degussa P25 TiO22.12-350 °C, TiO2-450 °C, and TiO2-700 °C samples. The characteristicpeaks at 24.2, 29.6, 43.3, 48.5, and 67.2° in the diffractionpattern of TiO2-350 °C can be indexed to the (110),(002), (003), (020), and (023) crystal planes of the metastable monoclinicTiO2(B) phase (JCPDS no. 00-046-1237), respectively. Themain peaks observed for TiO2-700 °C at 25.3, 37.8,48.0, 53.9, 55.1, and 62.7° can be indexed to the (101), (004),(200), (105), (211), and (204) crystal planes of the anatase phase(JCPDS no. 00-021-1272), respectively, while the other peaks at 24.1,29.3, and 43.3° are attributable to the TiO2(B) phase.Based on relative peak intensities, TiO2-450 °C showsmixed TiO2(B) and anatase phases and TiO2-700°C is mainly composed of anatase, and anatase and TiO2(B) are present at a suitable composite ratio in TiO2-450°C. The characteristic anatase peaks intensified as the annealingtemperature increased from 450 to 700 °C while those correspondingto TiO2(B) decreased. These results indicate that the anatase/TiO2(B) composites are formed by the gradual transformation ofthe TiO2(B) phase to the anatase phase, as clarified bythe presence of both anatase and TiO2(B) in TiO2-450 °C.2-350 °C  images ofthe prepared samplesare shown in 2-350 °C 2a has a 2-450 °C 2b owing 2-700 °C2c has a 2-350 °C, TiO2-450 °C, and TiO2-700 °C were 3.29, 3.21, and 3.20 eV, respectively, from whichadsorption edges of 377, 387, and 388 nm, respectively, were obtained.The diffuse reflectance spectroscopy (DRS) results for thepreparedsamples are shown in 2.22S decomposition afterUV irradiation, with the outlet concentration of H2S shownas a function of time. TiO2-450 °C exhibited the bestperformance during gas-phase H2S decomposition. Using thisphotocatalyst, 10 ppm H2S was decomposed to 0.1 ppm orless within 3 h. However, the detection limit of the H2S concentration detector tube used in this experiment was 0.1 ppm;therefore, it was necessary to determine whether the method yieldedH2S concentrations at the ppb level. When the measurementwas performed using gas chromatography, the H2S concentrationafter 3 h of light irradiation was below the detection limit of gaschromatography (≤1.0 ppb). As 1.0 ppb is the concentrationat which 6% of humans detect the odor of H2S, the photocatalyticreaction performed in this work can decompose H2S to aconcentration that is sufficiently low to avoid bad odors. In addition,this is the first report of a ppb level, which is significantly lowerthan parts per million (ppm) levels previously reported for the photocatalyticdecomposition of H2S (Table S1).2(B) heterojunction. In our previouswork,32 the electron energy structure ofanatase/TiO2(B) nanotubes was suggested using the energy-resolveddistribution of electron traps (ERDT) of TiO2(B), anatase,and a 1:1 mixture of anatase and TiO2(B) measured by reverseddouble-beam photoacoustic spectroscopy.38 These results revealedthat electron-transfer excitation to the electron traps of TiO2(B) occurs based on the high density of states in the valenceband of anatase. Furthermore, the valence band top (VBT) of TiO2(B) is located at a deeper energy level (approximately 0.7eV) than the VBT of anatase in the mixture, and the ERDT patternsof the anatase/TiO2(B) nanotubes in the mixture match.Using photoluminescence emission spectroscopy and a transient photovoltagetechnique with a photocatalyst synthesized by the same preparationprocedure as in this study, Wang et al. revealed that the electronlifetime of the anatase/TiO2(B) composite is longer thanthat of either anatase or TiO2(B).31 These results indicate that a type II band structure40 is constructed between anatase and TiO2(B) and that interfacialcharge-transfer excitation occurs in the composites. TEM imaging alsoconfirmed the interface between anatase and TiO2(B).31The superior performance of our method can be attributed to thelarge surface area and enhanced separation of photogenerated chargesinduced by the anatase/TiO2-450 °C  nanotubes synthesized bya two-step hydrothermal method is more than 3 times that of P25 oranatase synthesized by two-step hydrothermal methods. The photocatalyticreaction proceeds by adsorbing the decomposed substrate on the surfaceof the photocatalyst; hence, the specific surface area of the catalystis important from the viewpoint of adsorption.42 Therefore, TiO2-450 °C, which has a high specificsurface area and more reaction sites for H2S, exhibitsa superior H2S-decomposing ability. TiO2-350°C, which has a pure TiO2(B) phase, exhibits low photocatalyticactivity that is ascribable to charge recombination and the low crystallinityof TiO2(B). TiO2-700 °C has mixed anataseand TiO2(B) phases that promote charge separation; however,it exhibits lower photocatalytic activity owing to the lower specificsurface area produced at the higher annealing temperature. The phasecompositions, band gaps, specific surface areas, and kinetic ratesof TiO2-350 °C, TiO2-450 °C, and TiO2-700 °C are summarized in Based on these results, we conclude thatcharge recombination issuppressed by electron transfer in TiO2-450 °C 5. Furthe2S decomposition withN2 as the carrier gas. Under these conditions, all thephotocatalysts decomposed less H2S than when air was usedas the carrier gas (2S decomposition underhigh-humidity conditions (RH = 76–95%), which reveals thatthe amount of decomposed H2S decreases with increasinghumidity for all the photocatalysts. Because TiO2 is hydrophilic,water molecules are easily adsorbed onto its surface,20 which makes it difficult for H2S to approachthe TiO2 surface, thereby inhibiting H2S adsorptionand decomposition. Under low-humidity conditions (RH = 2–4.6%), whereas TiO2-350 °C and TiO2-700 °C decomposed moreH2S under low-humidity conditions than under high-humidity(RH = 76–95%) or normal conditions (RH = 18–24%). Previousreports23 have mainly considered H2S tobe decomposed by the hydroxyl radicals generated by oxidizing watermolecules with holes. However, in this study, the decomposition reactionproceeded even under low-humidity conditions; therefore, the mainreaction is considered to be the direct reaction of H2Swith holes. We conclude that a H2S molecule reacts directlywith a hole at the oxidation side and an oxygen molecule reacts withan electron to generate a superoxide anion radical at the reductionside 4a, the aionside 5. Sincer 9 days 4e. While4+ incorporated into TiO2 or oxidizedsulfur accumulated on the TiO2 surface.46 The peak at 167.6 eV can be assignedto either the S–O bonds in SO42– or S6+ in the TiO2 lattice.45 The O 1s XPS spectra are shown in 47 and the shoulder peak at 532.2 eV might be dueto the O–H bonds of chemisorbed water molecules on the TiO2 surface.52 After light irradiation, the peak at 529.2 eV can be assigned toO–Ti–O and Ti–O–S bonds, and the broadpeak at 532.2 eV is due to the S–O bonds in SO42– and the O–H bonds of chemisorbed water molecules,52 which was expected because SO42– accumulateson the photocatalyst surface during H2S decomposition.X-ray photoelectron spectroscopy (XPS) was usedto investigatethe factors affecting the color and photocatalytic activity. The bindingenergies due to relative surface charging were corrected using theC 1s peak at 285 eV. The S 2p spectra are shown in 2-450 °C was washed three times with 5 mL of deionizedwaterafter light irradiation, and the washing solution was analyzed usinghigh-performance ion chromatography (HPIC). Only SO42– (7.42 mg) was detected in the washing solution. Becauseno SO32– was detected, we conclude thatthe peak at 168.2 eV is not due to S4+ in SO32–. In addition, the intensities of the peaks at167.6 and 532.2 eV decreased after washing, suggesting that thesepeaks correspond to S6+ in free SO42– and SO42– adsorbed on the surface ofTiO2-450 °C. However, a weak peak corresponding toS4+ was observed after washing, whereas the peak of elementalsulfur was not detected. Based on these results, we conclude thattwo factors are responsible for the reduction in photocatalytic activityupon long-term light irradiation. The first is SO42– deposition on the TiO2 surface by oxidizingH2S to reactive oxygen species and the second is sulfurdeposition on the TiO2 surface by the direct oxidationof H2S with holes ,and the photocatalytic ability was restored by simply washing withwater. When the same experiment was conducted under visible-lightirradiation (λ > 420 nm), TiO2-450 °C didnotdecompose H2S during the first cycle, whereas approximately2.0 ppm of H2S was decomposed by TiO2-450 °Cduring the second cycle . Thisbehavior implies that sulfur deposited on the TiO2 surfacefunctions as a photocatalyst and forms a new band structure with TiO2, which exhibits visible-light responsiveness.55To determine the regenerative ability and visible-light responsivenessof the photocatalyst, TiO32S was successfully decomposedto 1.0 ppb or less within 3 h using an anatase/TiO2(B)composite photocatalyst owing to its high specific surface area andefficient charge-transfer interface. This study is the first reportedexample of H2S treatment to 1.0 ppb or less using a photocatalyst.We also clarified the effects of changing the reaction conditions,such as the carrier gas and humidity. H2S decompositionproceeded best under low-humidity conditions and in the presence ofoxygen. Furthermore, XPS and HPIC analyses of the TiO2 surfaceafter long-term light irradiation revealed that the reaction mechanismfor H2S decomposition involved oxygen consumption on thereducing side and holes directly oxidizing H2S on the oxidizingside, eventually generating SO42– andsulfur. The photocatalyst after washing was found to decompose 2.0ppm of H2S under visible-light irradiation because sulfuracts as a photocatalyst. In addition, the catalytic activity couldbe successfully restored by simply washing with water. The findingsof this study show that photocatalysis is a useful approach for solvingodor problems and provide useful information for designing catalyststo treat H2S.In this study, 10 ppm ofH44.12 powder (Degussa P25) was obtained from Nippon Aerosil Co.,Ltd. All chemicals were used as obtained without further purification.Sodium hydroxide , potassium hydroxide , and hydrochloric acid were purchased from Kanto Chemical Co., Ltd. TiO4.22 nanotubes were synthesized using a previously reported method.31 Degussa P25 (8.46 g) was mixed with 80 mL of10 M NaOH solution. The mixture was then transferred to a Teflon-linedstainless steel autoclave and heated at 130 °C for 36 h. Afterthis treatment, the autoclave was allowed to cool to 20–25°C. The intermediate product was then collected and washed severaltimes with deionized water and 0.1 M HCl solution until the pH ofthe solution was approximately 7. Subsequently, 80 mL of 15 M KOHsolution was added. The mixture was transferred to an autoclave, andthe temperature was increased to 200 °C for 24 h. The systemwas cooled to 150 °C at a rate of 1 °C min–1 and subsequently maintained at 150 °C for 12 h. The autoclavewas then quenched to 20–25 °C. The white product was thencollected and washed several times with deionized water and neutralizedwith the appropriate amount of 0.1 M HCl solution. It was again washedwith deionized water to remove any remaining traces of NaCl and KCl.The white product was then filtered and dried at 70 °C for 4h in air and gently ground in a mortar. Finally, the obtained powderwas calcined at 350, 450, and 700 °C for 2 h in air. These sampleswere labeled as TiO2-350 °C, TiO2-450 °C,and TiO2-700 °C, respectively.The TiO4.32 was identified using a powder X-raydiffractometer  at 40 mA and 40 kV using Cu Kα1radiation. The morphologies of the TiO2 samples were observedusing a transmission electron microscope .DRS was performed using a UV–vis–NIR spectrometer withan integrating sphere . The chemical states of the photocatalysts were investigatedusing an X-ray photoelectron spectrometer .The washing solution was analyzed using a high-performance ion chromatograph.The crystal structureof the synthesized TiO4.42S was conducted in a plug flowreactor at room temperature (25 °C). The vessel was made of quartzglass . The synthesized photocatalyst (0.1g) was placed in a glass reaction tube. Compressed air was then circulatedto a permeator , followed by heating a permeation tubecontaining liquid H2S  to 35 °Cin the permeator to generate H2S gas. A mixture of H2S gas (10 ppm) and compressed air  was ultimately produced with a relative humidityof 18–24%; this mixture was then introduced into the glassreaction tube under single-pass conditions. After adsorption equilibrationin the dark for 1 h, the reactor was placed below a spiral UV lamp, which emitted UV lightwith an irradiance of 18 mW cm–2. The experimentalsetup is shown in 2S concentration was monitoredevery 60 min for 3 h using a gas detector tube  and a gas chromatograph . The flow rate and concentration of H2S wereadjusted using a permeator. Experiments were conducted under bothhigh- and low-humidity conditions. The high-humidity experiments wereperformed by bubbling the H2S/air gas mixture adjustedwith a permeator into deionized water and then circulating this gas,which had a relative humidity of 76–95%, through the reactionvessel. The low-humidity experiments were conducted by passing thegas mixture through a calcium chloride tube, which imparted a relativehumidity of 2.0–4.6%, and then passing it through the reactionvessel. The humidity of the generated gas was measured using a digitalthermo-hygrometer . To determinethe reaction product, the photocatalyst was washed three times with15 mL of deionized water after light irradiation, and the washingsolution was analyzed using HPIC. In addition, the products on thephotocatalyst surface after the reaction were analyzed using XPS.Thedecomposition of gaseous H


In this study, a novel graphitic carbon nitride/NiCo-layered double hydroxide (CNLDH), a two-dimensional nanohybrid, is prepared by a simple hydrothermal synthesis. The homogeneous interpolation of g-C3N4 nanosheets into NiCo LDH stacked nanosheets effectively increases the overall performances of the g-C3N4/NiCo LDH nanohybrid. The improved morphology of the nanohybrid electrode upon the addition of g-C3N4 to the NiCo LDH yields a specific capacity of 183.43 mA h g−1 in 6 M KOH at 1 A g−1, higher than those of bare g-C3N4 (20.89 mA h g−1) and NiCo LDH (95.92 mA h g−1) electrodes. The excellent supercapacitive performance of the CNLDH nanohybrid is complemented by its low internal resistance, excellent rate capability, and large cycling lifetime. Furthermore, the hybrid supercapacitor is assembled using CNLDH 0.1 as a positive electrode and activated carbon (AC) as a negative electrode. The hybrid supercapacitor device of CNLDH 0.1//AC shows the maximum specific capacity of 37.44 mA h g−1 at 1 A g−1 with remarkable energy density, power density and good cycling performance. This confirms that the CNLDH 0.1 nanohybrid is an excellent electrode material for supercapacitor applications.The limited number of edge nitrogen atoms and low intrinsic electrical conductivity hinder the supercapacitive energy storage applications of the nitrogen-rich graphitic carbon nitride (g-C 3N4 and NiCO LDH shows the maximum specific capacity of 183.43 mA h g−1 with remarkable electrochemical performance.A two dimensional CNLDH 0.1 nanohybrid supercapacitor electrode prepared by simple hydrothermal hybridization of g-C The combination of two different nanostructures with distinct physical and chemical properties into a new hybrid with a unique structure can usually inherit cumulative advantages of the individual materials or even lead to the formation of a new heterostructured material with better properties. In recent years, the syntheses of different two-dimensional (2D) layered materials, such as graphene, hexagonal boron nitride, layered double hydroxides (LDHs), carbon nitride, MXenes, and transition metal dichalcogenides, have attracted considerable attention for the development of energy materials having unique properties.6–9 The 2D layered structures of these materials provide metallic or semiconducting properties.10 The 2D layered structures have advantages over bulk materials, such as the facilitated intercalation/deintercalation of ions and volume expansion,11 and thus have been employed in various applications, such as sensing,12 electrochemical energy storage and conversion,13 catalysis,14 and transistors.15 The increased interlayer distance in 2D layered materials is advantageous for energy storage as it can facilitate the ion transport and provide a better tolerance to the volume change, unlike in bulk materials.16 Additionally, the large surface area and number of active sites in 2D layered materials make them desirable for use as energy storage devices.Energy storage technologies involving supercapacitors require high energy and power densities.3N4) and LDH of nickel–cobalt (NiCo). g-C3N4, a nanocarbon material consisting of two earth-abundant elements (carbon and nitrogen), is a novel 2D layered organic semiconductor17 which has attracted increasing interest owing to its layered structure, simple synthesis, and low synthesis cost. The LDHs, such as the NiCo LDHs, are a class of 2D inorganic layered matrices consisting of positively charged layers, with anions and water molecules intercalated in the interlayer region. The 2D NiCo LDH provides a short ion diffusion path, abundant active sites, and different valence states, resulting in an increased electrochemical performance.18,19 However, the limited electroactive sites, low conductivity, low cycling stability of the NiCo LDH, and low ion transport capability of g-C3N4 hinder the applications of both materials. To overcome the disadvantages of these materials, we develop integrated heterostructures of g-C3N4 with NiCo LDH by successfully coupling them. This concept elucidates the specific challenges that need to be addressed for the introduction of 2D heterostructured electrodes into next-generation energy storage devices.Despite the extensive studies on 2D layered materials, their functionality should be improved. A novel approach for this purpose is the development of heterostructures of organic–inorganic 2D layered materials. This approach can pave the way for the development of novel electrodes with increased electrochemical performances by combining the advantages of the individual building blocks while eliminating their shortcomings. To demonstrate our concept, we hybridize 2D graphitic carbon nitride (g-C3), absolute ethanol (99%), nickel nitrate (Ni(NO3)2·6H2O), cobalt nitrate (Co(NO3)2·6H2O), and potassium hydroxide (KOH) were purchased from Sigma-Aldrich. None of the materials used in the experiments were further purified, unless otherwise stated.Analytically pure urea, nitric acid (HNO3N4 was fabricated by the simple pyrolysis of urea.20 The urea precursor (5 g) was kept in an alumina crucible with a cover. The crucible was heated for 2 h at 80 °C to remove moisture and obtain an oxygen-free product. The crucible was then heated at 550 °C in the ambient air for 3 h at a rate of 5 °C min−1. The resulting product was washed with 0.1 M of nitric acid and distilled water several times and dried at 80 °C for 24 h. A pale-yellow g-C3N4 powder was obtained after ultrasonication for 2 h.g-C3)2·6H2O and 0.874 g of Co(NO3)2·6H2O were dissolved in 70% ethanol. The prepared solution was transferred in a 200 mL Teflon autoclave and heated at 160 °C for 3 h. During the cooling of the reaction mixture to room temperature, the resultant solution was washed with copious amounts of water and ethanol. The washed powder was kept in the vacuum oven at 60 °C for 12 h. Different nanostructures of NiCo LDH were synthesized by varying the synthesis time .The NiCo LDH powder was synthesized by a hydrothermal method. In the typical synthesis, 0.582 g of Ni. The specific surface area was calculated using the Brunauer–Emmett–Teller (BET) method, while the pore size distribution was obtained using the desorption branch data by the Barrett–Joyner–Halenda (BJH) method. Fourier-transform infrared (FTIR) spectroscopy was performed using a Nanofinder 3.0 Raman spectrometer at an excitation wavelength of 488 nm. X-ray photoelectron spectroscopy (XPS) was performed using a Kratos-Axis spectrometer with monochromatic Al Kα (1486.7 eV) X-ray radiation . Electrochemical measurements of the g-C3N4, NiCo LDH, and CNLDH heterostructures were carried out in a three-electrode system, in a 6 M KOH aqueous solution, where Hg/HgO and Pt coil were used as the reference and counter electrodes, respectively. For the fabrication of the working electrode, the active material were mixed with poly (vinylidene fluoride)  and carbon black  with a weight ratio of 8 : 1 : 1 in N-methyl-2-pyrrolidone (NMP) to form a slurry. The slurry was coated onto a carbon paper current collector (1 × 1 cm2) using a spatula. The prepared electrodes were dried at 110 °C for 12 h under vacuum. The same process followed for the preparation of activated carbon (AC) as a negative electrode for hybrid supercapacitor. The total mass of CNLDH 0.1 and AC was about 3.5 mg cm−2. Cyclic voltammetry (CV) and galvanostatic charge–discharge (GCD) measurements were carried out using an electrochemical workstation  in the range of 0 to +0.5 V/Hg/HgO for half-cell test and 0 to +1.2 V for hybrid supercapacitor at various scanning rates and current densities. Prior to the CV and GCD measurements, the electrodes were activated by cycling the operation voltage for 25 cycles at 50 mV s−1. The ion kinetics within the electrode material were investigated by electrochemical impedance spectroscopy (EIS) in a frequency range of 0.01 to 105 Hz at the open-circuit potential. The specific capacity , energy density  and power density  were calculated using the GCD curve,I (mA), t (h), V (V) and m (g) are the discharging current, time, voltage and mass of active material, respectively.The morphologies and microstructures of the as-obtained products were characterized by field-emission scanning electron microscopy , transmission electron microscopy , and high-angle annular-dark-field scanning transmission electron microscopy . X-ray diffraction  patterns were recorded using Cu KThe morphologies of the as-prepared nanostructures were characterized by FESEM and TEM. The SEM images of the NiCo LDHs in Fig. S13N4, NiCo LDH, CNLDH 0.1, CNLDH 0.3, and CNLDH 0.5 nanohybrids are presented in 3N4.21 As shown in 3N4 exhibits a lattice spacing of 0.32 nm corresponding to the (002) reflection of the graphitic structure of carbon nitride, which corroborates the XRD results. 3N4 nanosheets in the hydrothermal autoclave, the CNLDH hybrid nanosheets could be synthesized in the form of “sheet-on-sheet” structure. In CNLDH 0.1 . The firmly immobilized NiCo LDH nanosheets obtained during the TEM specimen preparation remained unseparated for days, indicative of a strong interaction between g-C3N4 and NiCo LDH. The high-resolution TEM image in 3N4. Energy-dispersive X-ray spectroscopy (EDS) mapping was carried out to determine the compositional and elemental distributions; the results are shown in 3N4 nanosheets between the layers of the NiCo LDH. Such unique 2D heterostructure can provide more electrochemically active sites for the exposure to the electrolyte, favoring the ion transfer and diffusion and accelerating the surface redox reaction.24CNLDH 0.1 was investigated by HAADF-STEM. The elemental maps  show the25 As Ni and Co hydroxides have similar structures, the differentiation of the two phases was not simple. The diffraction peaks around 11°, 23°, 33°, 34°, and 61° are attributed to the (003), (006), (220), (012), and (113) planes of α-Co(OH)2 and α-Ni(OH)2 (Joint Committee on Powder Diffraction Standards (JCPDS) 460605 and 380717, respectively). The 2θ diffraction peak at 11.6° (0.76 nm) corresponds to both α-Co(OH)2 and α-Ni(OH)2. The stacking of the layers along the c-axis indicates the difference between the α- and β-phases. The Co(OH)2 and Ni(OH)2 layers are randomly oriented and separated by intercalated water molecules bonded to the hydroxyl groups by hydrogen bonds in the α-phase with an interlamellar distance of approximately 0.8 nm. In the β-phase, Co(OH)2 and Ni(OH)2 layers are perfectly stacked along the c-axis with an interlamellar distance of 0.4 nm without intercalated species. α-phase compounds are theoretically expected to exhibit superior electrochemical activities to those of the β-phase compounds.26 The peak at 19.38° observed for the sample synthesized for 12 h corresponds to the (001) plane and confirms the layered structure of the synthesized NiCo hydroxide.The sharp peaks in the XRD patterns of the NiCo LDHs obtained using different hydrothermal synthesis times Fig. S2 indicate3N4, and CNLDH nanohybrids. The peaks at 12.74° and 27.43° observed for g-C3N4 can be indexed to the (210) and (002) planes (JCPDS 000660813), corresponding to the in-plane structural packing of tri-s-triazine units with a distance of 0.68 nm and periodic interlayer stacking of aromatic segments with a distance of 0.32 nm in the conjugated aromatic system, respectively. The peaks corresponding to the (002) plane are usually attributed to the distance between the layers of the graphitic material; however, for g-C3N4, the peaks corresponding to the (210) plane are attributed to the intralayer d-spacing.27 In addition, the d-spacing of the (002) plane of CNLDH 0.1 suggests that g-C3N4 is intercalated between the LDH layers. Further, the high-angle (2θ) shifts of the (002) XRD peaks of the LDH composites  diffraction peak  are analyzed by FTIR spectroscopy  scale" fill="currentColor" stroke="none"><path d="M0 440 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z M0 280 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z"/></g></svg>N–C, tertiary nitrogen (N–(C3)), and amino functional group having a hydrogen atom, respectively. The synthesis of a defect-free g-C3N4 with the perfect stoichiometric ratio of 0.75 is often challenging. The obtained C/N ratio of 0.76 of our g-C3N4 is attributed to the various optimizations carried out along with the systematic reaction design 3, and amino N peaks of CNLDH 0.1 confirm the intercalation between NiCo LDH and g-C3N4 sheets , which indicates that the introduction of the small amount of g-C3N4 relieves the agglomeration of NiCo LDH layers. Fig. S4b, d, and f3N4, NiCo LDH, and CNLDH 0.1, respectively, showing the average pore sizes of 5 nm.The as-fabricated g-C3N4, NiCo LDH, and CNLDH heterostructures were evaluated by electrochemical measurements in a three-electrode system. 3N4, NiCo LDH, and CNLDH 0.1 at a scan rate of 2 mV s−1. The rectangular CV curves of g-C3N4 2/NiOOH and CoOOH/CoO2 in the NiCo LDH, respectively.48 With the increase in scan rate, the anodic and cathodic peaks shift toward positive and negative potentials, respectively, owing to the polarization of the electrode.49 Charge–discharge curves of CNLDH 0.1 at different current densities are shown in 3N4 in the composite is most suitable for the full utilization of the layered LDH  and Rct (3.72 Ω) compared to pure g-C3N4 (Rs = 3.71 Ω and Rct = 4.63 Ω) and NiCo LDH (Rs = 4.02 Ω and Rct = 4.67 Ω). The superior electronic transport in CNLDH 0.1, attributed to the enhanced interfacial contact by the anchoring of g-C3N4 on the layered material (NiCo LDH). The improved interfacial contact promotes a fast and smooth ion diffusion through the layered structure and decreases its transfer resistance. To evaluate the application feasibility of these materials, the electrochemical cycling stabilities of g-C3N4, NiCo LDH, and CNLDH 0.1 were evaluated. 3N4, NiCo LDH, and CNLDH 0.1 at a current density of 10 A g−1. CNLDH 0.1 retains 79% of its initial capacity after 5000 cycles showing its long-term stability. The corresponding GCD curves are shown in Fig. S7.3N4 and NiCo LDH 2D nanostructures.However, at a high current density, only the external active surface can be utilized for charge storage.um inset , where R−1 is obtained at a current density of 1 A g−1. The specific capacity at current densities of 2, 3, 5, 7 and 10 A g−1 are 30.81, 26.96, 22.51, 18.96 and 14.55 mA h g−1, respectively  as a negative electrode was assembled. ectively . The capone plot  of CNLDH3N4 and 2D NiCo layered materials for supercapacitor electrodes. A series of CNLDH nanohybrids with different mass ratios of g-C3N4 and NiCo LDH was prepared by the hydrothermal method. The structures and morphologies of the nanohybrids suggested that the NiCo LDH nanosheets were in situ-anchored on the surfaces of the g-C3N4 nanosheets or vice versa. CNLDH 0.1 exhibited excellent electrochemical performances, higher than those of CNLDH 0.3 and CNLDH 0.5, owing to the synergetic effect of g-C3N4 and NiCo LDH. In addition, CNLDH 0.1 exhibited a high specific capacity (182.43 mA h g−1 at 1 A g−1) and long-term cycling stability. The assembled CNLDH 0.1//AC hybrid supercapacitor presented the maximum energy density of 22.46 W h kg−1 at a power density of 600 W kg−1 and maintained an energy density of 8.73 W h kg−1 at a high power density of 6000 W kg−1. These results indicate that the CNLDH nanohybrid is a promising material for use as a supercapacitor electrode.We presented a simple strategy for the development of CNLDH nanohybrids of g-CThere are no conflicts to declare.RA-009-C9RA06068E-s001


Compared with the control mice, the propofol-treated wild-type mice exhibited lower levels of IL-6 (0.57 ± 0.23 vs 1.00 ± 0.39 at 24 h), and smaller numbers of TLR4-expressing microglia in the penumbra (11.7 ± 3.1 vs 25.1 ± 4.7 cells/0.1 mm2). In conclusion, propofol administration prior to ischemic brain insult attenuated brain injury by blocking the TLR4-dependent pathway and suppressing pro-inflammatory cytokine production.Ischemic brain injury is one of the most serious perioperative complications. However, effective preventative methods have not yet been established. This study aimed to investigate whether propofol has neuroprotective effects against ischemic brain injury, with a specific focus on Toll-like receptor 4 (TLR4). Focal brain ischemia was induced via a combination of left common carotid artery occlusion and distal left middle cerebral artery coagulation in mice. Either propofol (10 mg/kg) or vehicle was intravenously injected 10 min prior to the induction of brain ischemia in wild-type and TLR4 knockout mice. Infarct volume, pro-inflammatory cytokine expression, inflammatory cell infiltration, and neurobehavioral function were assessed. Propofol administration significantly reduced infarct volume in wild-type mice (26.9 ± 2.7 vs 15.7 ± 2.0 mm Consequently, only a limited number of patients receive reperfusion therapy during the acute phase. Therefore, preventative strategies for minimizing ischemic brain injury are vital. In addition, the effects of commonly used anesthetic drugs on the ischemic brain and its underlying mechanisms are relevant to research targets. Given these considerations, continuous efforts have been made in recent decades to elucidate the effects of anesthetic drugs on ischemic brain injury, and volatile anesthetic drugs have been shown to exert preconditioning and neuroprotective effect via activation of mitochondrial adenosine triphosphate-sensitive potassium channels ,7. Some Thus, we tested the hypothesis that propofol attenuates ischemic brain injury via inhibition of the TLR4-dependent pathway and suppression of consequent inflammatory cytokine production. The primary outcome of this study was infarct volume; the secondary outcomes were pro-inflammatory cytokine mRNA expression levels, the number of TLR4-expressing microglia in the penumbra, and neurological function.22.1Male C57BL/6 mice (8 weeks old) were purchased from Japan SLC . Two breeding pairs of TLR4 knockout mice with a C57BL/6 genetic background were purchased from Oriental Bio Service , and their male offspring mice (10–14 weeks old) were used in the study. All mice were group-housed at 23°C ± 2°C with free access to standard food and water and a 12-h light/dark cycle. All experiments were performed between 09:00 and 17:00 under normal room light and temperature (23°C ± 2°C) conditions.Ethical approval: All experiments were conducted in accordance with the National Institutes of Health guidelines for the care and use of laboratory animals. The experimental protocol was reviewed and approved by the University of Yamanashi Animal Care Committee.2.2Focal brain ischemia was induced via a combination of permanent left common carotid artery occlusion and distal left middle cerebral artery (MCA) coagulation . Briefly2.3Propofol  diluted with fat emulsion by 10 times  was administered via the tail vein 10 min prior to MCA occlusion. In control mice, an equal volume of the fat emulsion alone was administered. The dose was chosen based on an earlier study, in which propofol produced neuroprotection against ischemic brain injury in mice .2.4Seven days after MCA occlusion and after hemodynamic measurements, mice were deeply anesthetized with 5% isoflurane and euthanized via cervical dislocation. Brains were removed and cut into 1-mm-thick coronal sections. The brain slices were immersed in 2% 2,3,5-triphenyltetrazolium chloride  at 37°C for 15 min in a dark room. The infarct area was traced and measured using image analysis software  by an individual who was blinded to the grouping and study design. To correct for the contribution of edema, the infarct area was calculated as follows: total ipsilateral hemisphere – infarct region . Total i2.5Heart rate and arterial blood pressure were measured non-invasively using a tail-cuff monitor  to evaluate the effects of propofol on hemodynamics. Values were recorded 1 h before MCA occlusion (baseline), 10 min after the injection of either propofol or the control solution, 1 h after MCA occlusion, and 7 days after MCA occlusion.2.62, PO2, and glucose levels . Blood samples were collected before induction of MCA occlusion (baseline), 1 h after the injection of either propofol or the control solution, and 7 days after MCA occlusion.Blood gas analysis was performed to evaluate pH, PCO2.7TM real-time PCR system  using the PowerSYBR® Green PCR Master Mix and corresponding primers to quantify target genes (Table S1). The relative changes were expressed as a ratio to glyceraldehyde-3-phosphate dehydrogenase (GAPDH) mRNAs of the same sample. The data were analyzed using the 2−ΔΔCT method. The 2−ΔΔCT value of the target transcript from each mouse was normalized with those of the control group mice as 1.0. The normalized 2−ΔΔCT values derived from the two groups were then compared.Real-time PCR was used to measure mRNA expression levels of interleukin (IL)-6, IL-1β, and tumor necrosis factor α (TNF-α). The brain was the target mRNAs from the brain slices of propofol-treated wild-type mice that were compared with those of control wild-type mice. The RNeasy Mini Kit  was used for the extraction of total mRNA from brain slices. About 1µg of mRNA was reverse transcribed into complementary DNA with a QuantiTect Reverse Transcription Kit . PCR was performed on a StepOne2.8The mouse brains were fixed with 4% paraformaldehyde in phosphate-buffered saline (PBS) for 6 days and cut into 1mm-thick sections using a slicer . Prior to immunofluorescence, CUBIC tissue clearing  was perf2 areas within the cortex were randomly chosen, and the number of TLR4 + CD11b + double-positive cells was automatically counted using Image J software. Ten readouts per animal were averaged.TLR4 + CD11b + double-positive cells in the penumbra, defined as the region immediately adjacent to the infarct area, were counted in 400× magnification microscopic fields. For each animal, ten 0.1 mm2.9Neurological function was assessed using the neurological deficit scores , lateral push test, and body asymmetry test at baseline and 7 days after the induction of ischemic brain injury. The lateral push test and body asymmetry test were performed as previously described ,22.2.10F test, respectively. Infarct volumes, numbers of TLR4 + CD11b + double-positive cells, and body asymmetry test results were analyzed using two-tailed Student’s t test; cytokine expression levels and neurological deficit scores were analyzed using Mann–Whitney U test, based on the results of the Shapiro–Wilk and F tests. Heart rate, blood pressure, and blood gas parameters  were analyzed using the two-way analysis of variance for repeated measures; the lateral push test results were analyzed using Chi-square test. For infarct volumes and cytokine expression levels, the sample size was calculated to detect significance with 95% confidence, assuming alpha of 0.05 and power of 0.8 (G*Power 3.1.9.3). Values are presented as mean ± standard error of mean for the infarct volumes and numbers of cells, as median ± quantile for cytokine mRNA expression levels, and as mean ± standard deviation for heart rate, blood pressure, and blood gas parameters. p-values of less than 0.05 were considered statistically significant.Statistical analysis was performed using Prism 8 software . The collected data were assessed for the normal distribution and equal variance using the Shapiro–Wilk and 33.13, n  =  10 each, p  <0.05, p  > 0.99; lateral push test, baseline and day 7: 0.0 ± 0.0 vs 0.0 ± 0.0, p  > 0.99; body asymmetry test, baseline: 6.0 ± 5.4 vs 1.0 ± 6.4, p  = 0.56, day 7: 9.0 ± 12.5 vs 26.0 ± 9.1, p  = 0.29, all n  =  10 each, To investigate the effects of propofol on ischemic brain injury, 8- to 9-week-old wild-type mice were treated with either propofol or 10% fat emulsion (control) 10 min before the induction of ischemic brain injury. Propofol-treated mice exhibited significantly smaller infarct volumes than control mice 7 days after ischemic brain injury  compared with the control mice.To investigate the possible involvement of anti-inflammatory action of propofol in the neuroprotective effect observed, mRNA expression of pro-inflammatory cytokines was measured 24 h after ischemic brain injury. As shown in 3.32, n = 10 each, p  <  0.05).Immunofluorescence targeting TLR4 and microglia was performed to assess whether the propofol treatment influenced TLR4 expression or infiltration of microglia in the penumbra. As shown in 3.43, n = 10 each, p = 0.91, p  > 0.99; lateral push test, baseline and day 7: 0.0 ± 0.0 vs 0.0 ± 0.0, p  > 0.99; body asymmetry test, baseline: 7.5 ± 3.1 vs 11.1 ± 8.9, p  = 0.72, 7 day after: −10.0 ± 6.5 vs −1.1 ± 2.0, p  = 0.19, all n = 10 each, In experiments testing the effects of propofol treatment against ischemic brain injury in TLR4 knockout mice, there was no significant difference in the infarct volume between the propofol-treated TLR4 knockout mice and control-treated TLR4 knockout mice (25.6 ± 3.8 vs 26.5 ± 3.1 mm4In the present study, we found that propofol significantly reduced infarct volume after ischemic brain insult in wild-type mice but not in TLR4-knockout mice. The reduced infarct volume was associated with reduced pro-inflammatory cytokine expressions and TLR4-expresing microglia in the penumbra. These results indicated that propofol exerts neuroprotective effects via its anti-inflammatory capacity, specifically the inhibitory effect on TLR4. Since the hemodynamic and blood gas parameters were not affected by the propofol administration, the effects of propofol on cardiovascular and respiratory function and blood sugar levels were not likely to be involved in the neuroprotective effects of propofol observed in the present study.The immune system is promptly activated once ischemic brain injury occurs . Among vin vivo studies in which propofol suppressed inflammatory cytokine production via inhibition of TLR4-dependent pathways in various disease models, including lung [in vitro studies using lipopolysaccharide-treated microglia [In the current study, the neuroprotective effect of propofol against ischemic brain injury was abrogated by depletion of TLR4, indicating that the TLR4-dependent pathway was substantially involved in the mechanism underlying that neuroprotective effect. The results of the current study are concordant with recent ing lung , liver [ing lung , and gasing lung , and asting lung . Severalicroglia , macrophicroglia , spinal icroglia , and alvicroglia  also sugOur result showed significantly lower IL-6 mRNA expression levels in the brains of propofol-treated wild-type mice than in the brains of control mice, which is consistent with an earlier study in which there was a correlation between infarct volume and IL-6 mRNA expression in the brain  Among vaOur immunofluorescence study demonstrated that the pre-treatment with propofol reduced the number of TLR4-expressing microglia in the penumbra. Recent studies have reported that microglia are involved in the exacerbation of cerebral infarction ,38,39. TIf ischemic brain injury occurs during general anesthesia, the consequences are tragic. The clinical importance of the present study is that propofol administration prior to ischemic insult may have the potential to protect against ischemic brain injury, presumably by blocking the TLR4-dependent pathway in microglia in the penumbra. This understanding of the mechanism underlying the neuroprotective effect of propofol against ischemic brain injury may lead to a novel strategy to prevent exacerbation of ischemic brain injury during general anesthesia.This study has some limitations. First, we did not find neurological improvement in the propofol-treated mice, possibly because of the relatively small infarct volumes and low sensitivities of the neurological assays we used in the current study. Although infarct volume is clinically relevant and was the primary outcome of this study, further studies with more sensitive neurobehavioral evaluation are needed to elucidate the effects of propofol on neurobehavioral function after ischemic brain injury. Second, we only used relatively young male mice. The neuroprotective effects of propofol should be tested in aged mice and female mice at different menopausal stages because those biological variables can affect the outcomes of ischemic brain injury. Third, we used only GAPDH as a reference gene for the real-time PCR analysis. Although this gene is widely used as a reference gene for evaluating inflammatory responses after ischemic stroke in mice, it is indicated that hypoxia upregulates GAPDH mRNA expression , which c5Propofol administration prior to ischemic brain insult attenuated brain injury by blocking the TRL4-dependent pathway and suppressing pro-inflammatory cytokine production. This understanding of the mechanism underlying the neuroprotective effect of propofol against ischemic brain injury may lead to a new strategy to prevent exacerbation of cerebral infarction during general anesthesia.IL-1βinterleukin-1βIL-6interleukin-6MCAmiddle cerebral arteryPCRpolymerase chain reactionTLR4Toll-like receptor 4TNF-αtumor necrosis factor-α


The results indicated that the adsorption amount was highly dependent on the pH of the solution, and reached the highest value at pH = 3. The kinetic behavior of defluoridation on LMF11 NFs was dominated by the PSO model with the highest fitted determination coefficients of 0.9999. Compared with the other three isotherm models, the Langmuir model described defluoridation characteristics well with larger correlation coefficients of 0.9997, 0.9990, 0.9987 and 0.9976 at 15 °C, 25 °C, 35 °C and 45 °C, respectively. The optimized LMF11 NFs exhibited superior monolayer defluoridation capacities for 173.30–199.60 mg F−/g at pH 3 at 15–45 °C according to the Langmuir isotherm model. A thermodynamic study proved that the defluoridation by LMF11 NFs is a spontaneous, endothermic along with entropy increase process. In addition, the LMF11 NFs still showed high defluoridation performance after three reused cycles. These findings unveil that the synthesized LMF11 NFs adsorbent is a good adsorbent for fluoride remediation from wastewater owing to its low cost, high defluoridation performance and easy operation.The occurrence of fluoride contamination in drinking water has gained substantial concern owing to its serious threat to human health. Traditional adsorbents have shortcomings such as low adsorption capacity and poor selectivity, so it is urgent to develop new adsorbents with high adsorption capacity, renewable and no secondary pollution. In this work, magnetic electrospun La-Mn-Fe tri-metal oxide nanofibers (LMF NFs) for fluoride recovery were developed via electrospinning and heat treatment, and its defluoridation property was evaluated in batch trials. Modern analytical tools  were adopted to characterize the properties of the optimized adsorbent, i.e., LMF11 NFs with a La:Mn molar ratio of 1:1. The surface area calculated via BET method and pH Surplus fluoride in drinking water may lead to irremediable diseases in humans, namely fluorosis. The tolerance level of fluoride in drinking water recommended by the World Health Organization is no more than 1.5 mg/L [−, namely carbon-based materials [Fluoride contamination in groundwater caused by the processes of natural geochemical and human production activities is a burning issue, which has attracted substantial attention since it is closely related to the safety and health issue for humans ,3,4,5,6.1.5 mg/L ,8. Howev1.5 mg/L ,12,13,141.5 mg/L . A varieaterials , activataterials , carbonaaterials , polymeraterials , etc. Wh− to form colored complexes with high selectivity, respectively. It provided a new way for scientists to remove fluoride from wastewater. Since then, various materials, such as compounds of multivalent metal elements such as La(III), Ce(III) and Ce(IV) in the form of oxides, hydrous oxides and basic carbonate [−/g for fluoride. Cotton-like Ca-Al-La adsorbent with optimal Ca/Al/La molar ratio of 1/4/2 exhibited the largest fluoride binding amount of 29.30 mg F−/g [−/g at 25 °C and pH 8.25 [As far as we know, Belcher R. ,21 in 19arbonate , lanthanarbonate , La(III)arbonate , La-chitarbonate , Mn-La marbonate , Fe-La-Carbonate , La-MOFsarbonate , La-MOF@arbonate , had beearbonate ,32,33,34 mg F−/g . Similar pH 8.25 . Althoug3O4 [3O4 nanoparticles. To overcome the limit of Fe3O4 nanoparticles, numerous magnetic adsorbents were fabricated by coating metal oxides on Fe3O4 nanoparticles to form core-shell composites or by loading Fe3O4 and other metal oxides on the carrier [−/g for fluoride. The use of magnetic La2O3-CeO2-Fe3O4 nanofibers for fluoride removal from aqueous solution has been investigated by our group [−/g at 15–35 °C, and the adsorbent can be easily and quickly separated from the solution by an external magnet. The obstacles to the industrial application of La2O3-CeO2-Fe3O4 nanofibers are the scarcity and relatively high cost of rare earth elements. There is no literature about the defluoridation by magnetically separable electrospun La-Mn-Fe tri-metal oxide nanofibers in wastewater. In view of the high affinity of La towards fluoride ions and the low cost of metals Mn and Fe, this research is to fabricate a new magnetic tri-metallic adsorbent  with less expenditure, which can exhibit structural stability, high defluoridation efficiency, excellent recyclability and reusability.Ease in localization and separation can promote the recovery and reuse of materials for successive treatment cycles. Due to the good magnetic responsiveness of Fe3O4 ,38, its 3O4 ,40. Wher carrier . For exa carrier  showed aur group , the max3O4 NPs are evenly dispersed along the nanofiber axis without obvious agglomeration. Moreover, the magnetic fibrous LMF NFs can effectively prevent the agglomeration of adsorbent during the defluoridation process and can be quickly separated from the solution using an external magnet after adsorption. The effects of the molar ratio of La/Mn, initial F− concentration, pH, LMF11 NFs dosage, shaking time, interfering anions, isotherm models and adsorption kinetics were investigated in detail.In this study, magnetically separable electrospun La-Mn-Fe tri-metal oxide nanofibers (LMF NFs) were prepared by electrospinning and heat treatment. The LMF NFs were used as a new type of adsorbent for fluoride remediation. The magnetic Fe3)3·6H2O, Mn(CH3COO)2·4H2O, NaCl, DMF, Na2CO3, NaNO3, Na3PO4, CH3COOH, Na3C6H5O7·2H2O, Na2SO4, HCl, NaOH, FeCl3·6H2O, HOCH2CH2OH and CH3COONa were acquired from Aladdin Chemical Reagent Co., Ltd.  and used as received. Polyacrylonitrile  was purchased from Sigma-Aldrich .NaF, La(NO3)3·6H2O and Mn(CH3COO)2·4H2O were mixed in a beaker with La:Mn molar ratios of 2:1, 1:1 and 1:2, then completely dissolved in a small amount of DMF. Afterward, 1.8 g PAN powder, 10 mL DMF and 0.075 g Fe3O4 nanoparticles 3/Mn(CH3COO)2/Fe3O4/PAN/DMF). The electrospinning processes were conducted to produce precursor fibers [−1 and a flow rate of 1.2 mL h−1 at room temperature, to generate precursor nanofibers. After electrospinning, the resulting precursor nanofibers were dehydrated for 8 h and then put into a tube furnace for heat treatment at 500 °C for 2 h to prepare La-Mn-Fe tri-metal oxide nanofibers under an air atmosphere. La-Mn-Fe tri-metal oxide nanofibers with molar ratios of La/Mn being 2/1, 1/1 and ½ were noted as LMF21 NFs, LMF11 NFs and LMF12 NFs, respectively. The graphic plan of the fabrication process of electrospun LMF NFs and the adsorption of fluoride is portrayed in La(NOl method ) were inr fibers ,43,45,46− solution with initial F− concentrations of 20 and 50 mg F−/L at 25 °C and pH 3. Adsorption isotherm studies were carried out at varying original F− concentrations of 10–60 mg F−/L at a temperature range of 15–45 °C with LMF11 NFs dosage of 0.2 g/L in 50 mL of respective solution. Various concentrations  of five commonly occurring anions, including nitrate (NO3−), sulfate (SO42−), chloride (Cl−), carbonate (CO32−), phosphate (PO43−) were studied under initial F− concentration of 20 mg F−/L with an adsorbent dosage of 0.2 g/L at 25 °C and pH 3 for 12 h. The effect of pH (2–9) along with LMF11 NFs dosage (0.2–0.6 g/L) on adsorption property was assessed under 50 mL initial F− concentration of 20 mg F−/L at 25 °C for 12 h.The defluoridation performance of LMF NFs was evaluated using batch experiments . Adsorpt− adsorbed magnetic LMF11 NFs was determined to quantify the cost-effectiveness of the adsorbent for F− remediation in the successive adsorption–desorption cycles. The original concentration of F− was 20 mg F−/L at pH 3 and 25 °C for 8 h for each regeneration cycle. The F− adsorbed magnetic LMF11 NFs were recovered by an external magnet, and then eluted with 100 mL of 0.1 M NaOH to regeneration [The reusability of the Fneration ,47,48, vThe crystal structures of LMF11 NFs before and after adsorption were established by X-ray diffraction , The BET specific surface area of the LMF11 NFs adsorbent was examined using a Micrometrics ASAP 2020 analyzer . The TEM and SEM images of the magnetic electrospun LMF11 NFs were acquired from TEM JEM 2100  and SEM 300U . The point of zero charge of LMF11 NFs was assessed via a pH drift method .3)3/Mn(CH3COO)2/Fe3O4/PAN and LMF11 NFs were determined by SEM, as portrayed in 3)3/Mn(CH3COO)2/Fe3O4/PAN fibers are continuous and smooth with an average diameter of 1.28 ± 0.10 μm. Fe3O4 NPs are evenly distributed on the surface of the La(NO3)3/Mn(CH3COO)2/Fe3O4/PAN and LMF11 NFs. After heat treatment, the volume and diameter of fibers decline due to the decomposition of Mn(CH3COO)2, La(NO3)3 and PAN. The continuous LMF11 NFs are hollow and rougher compared to that before calcination, which possess an average diameter of 626 ± 57 nm. 3O4 NPs are uniformly dispersed along the fiber axis without obvious agglomeration. The chemical composition of LMF11 NFs is verified by EDS analysis, as shown in 2 adsorption–desorption isotherm (3 g−1 and 55.81 m2 g−1. The inset of pzc) of LMF11 NFs calculated according to the literature [− can be could easily absorbed into the positively charged LMF11 NFs owing to their columbic attractions at a pH range below pHpzc [The morphological details of La and removal efficiency (95.58%) for LMF11 NFs occurs at a pH of 3 owing to the columbic attractions (Equation (1)). LMF11 NFs show a dramatic decrease in defluoridation property at pH > 3 and pH < 3. Most fluoride ions exist in the form of electrically neutral HF (pKaHF = 2.95) [− and negatively charged LMF11 NFs along with the competition between F− and OH− for active sites [− solution changed from 2 to 2.18, 3−5.49, 4–6.59, 5–6.75, 6–6.59, 7–6.34, 8–6.83, 9–7.14 after adsorption, respectively. The improvement of pH after adsorption in an initial pH range of 2–6 suggests the existence of an ion exchange of hydroxyl groups bonded on the surface of the adsorbent with fluoride ions (Equations (2) and (3)). While the decline of pH after adsorption in an initial pH range from 7 to 9 (>pHpzc) is mainly ascribed to the competitive adsorption between hydroxyl groups and adsorbate (F−) (Equation (4)). A pH of 3 was chosen to utilize for subsequent studies.The effect of pH on the defluoridation performance of LMF11 NFs was evaluated at various pH values from 2 to 9 . The max = 2.95)  when pH  = 2.95) . Meanwhive sites . Additio−/L fluoride solution at 25 °C. Evidently, the adsorption uptakes at an adsorbent dosage of 0.2, 0.3, 0.4, 0.5 and 0.6 g L−1 are 188.21, 153.44, 123.81, 99.33 and 82.77 mg F−/g, the fluoride removal percentages are 75.29%, 92.07%, 99.05%, 99.32% and 99.33%, respectively  of LMF11 NFs was poured into a 50 mL of 50 mg Fectively . It can uts down . The ads−. As shown in C0 (10 mg F−/L to 45 mg F−/L). The reason is that with the increase in F− concentration, the number of fluoride ions near the surface of the LMF11 NFs increases apparently, and the binding sites on the surface of LMF11 NFs are more fully surrounded by fluoride ions, thus, more fluoride ions are adsorbed by the adsorbent, resulting in the increasing of fluoride binding amount [C0 due to the saturation of active sites. While the fluoride removal efficiency declines with the rise of fluoride concentration. It is because the LMF11 NFs adsorbent have a limited adsorption uptake. When LMF11 NFs adsorbent achieve adsorption saturation, the removal percentage will decline with the rise of the initial fluoride concentration. In addition, at C0 = 60 mg F−/L, the binding amount leaps from 175.82 mg F−/g at 15 °C to 193.53 mg F−/g at 45 °C. The higher the temperature of the adsorption system, the larger the binding amount, validating the endothermic nature of the defluoridation process [The fluoride uptake significantly depends on the initial concentration of Fg amount . Then it process . − removal on LMF11 NFs by varying the time intervals in the range of 0–60 min using an adsorbent dosage of 0.2 g/L with given concentrations (C0 = 20 and 50 mg F−/L) at 25 °C and pH = 3. As depicted in − could be removed within the first 20 min at C0 = 20 and 50 mg F−/L, showing an excellent defluoridation property. In this stage, numerous binding sites in LMF11 NFs and the high concentration of fluoride ions in the solution decide the rapid adsorption rate. Hence the larger adsorption rate in the incipient step is ascribed to the strengthening of the diffusion rate of F− provided by concentration gradient along with the existence of numerous available binding sites on the LMF11 NFs [−/L solution and 30 min in 50 mg F−/L solution. Clearly, at higher C0, it takes longer shaking time to attain equilibrium. Based on the aforementioned analysis, the effect of pH (A), initial concentration of F− (B) and LMF11 NFs dosage (C)   was detesage (C) .3−), sulfate (SO42−), chloride (Cl−), carbonate (CO32−) and phosphate (PO43−) on defluoridation by LMF11 NFs were investigated. The competitive experiments were conducted by weighing 0.01 g LMF11 NFs into a binary system that contained 50 mL 10 mg F−/L F− paired with different concentrations  of the interfering anions, respectively. The result is depicted in 3−, Cl−, SO42− and CO32− anions at all concentrations of 10–100 mg/L almost do not prevent the F− removal, illustrating the excellent selectivity and stability of LMF11 NFs adsorbent. Whereas the binding capacity drops quickly from 98.20 to 52.16 mg F−/g with the rise of PO43− concentration from 0 to 100 mg/L, manifesting a remarkable competitive effect of PO43− on defluoridation. This may be explained by the Ksp of LaPO4 (3.7 × 10−23) [43− to displace the adsorbed F− on LMF11 NFs compared with that of LaF3, thereby deteriorating the fluoride binding capacity. The same result has also been reported by several studies using rare earth element-based materials for fluoride remediation [One of the major barriers that limit the widespread application of adsorption technology in practical water treatment is the selectivity of adsorbent. The effects of five commonly occurring anions including nitrate (NO× 10−23) , which fediation ,57.R2) at 15 to 45 °C is Langmuir  > Temkin  > D–R  > Freundlich , illustrating that the homogeneous and single-layer adsorbed on LMF11 NFs dominates the defluoridation process [Qm values range from 173.30–199.60 mg F−/g at 15 to 45 °C, exceedingly higher than many reported adsorbents including MOFs such as La–BTC (105.2 mg/g), La–BPDC (125.9 mg/g), La–BHTA (145.5 mg/g), La–PMA (158.9 mg/g) and La–BDC (171.7 mg/g) [RL) > 1 means unfavorable adsorption, while RL < 1 indicates favorable adsorption processes [RL at evaluated temperatures lie between 0 and 1  , Uio–66 .7 mg/g) ). Table rocesses . All the 0 and 1 , reflectn (0.1961–0.2543) established from the Freundlich model lie in the range of 0 < 1/n < 1 at 15 to 45 °C, signifying beneficial adsorption by LMF11 NFs under studied conditions [E value derived from the D-R model represents free energy. E values in ranges of 1–8 kJ mol−1, 8–16 kJ mol−1 and greater than 16 kJ mol−1 are indicative of electrostatic physical adsorption, electrostatic interaction or synergy and chemical adsorption, respectively [E value calculated from the D-R model gradually elevates from 14.06 to 17.96 kJ/mol with ascending the temperature from 15 to 45 °C, which denotes that the defluoridation on LMF11 NFs proceeds from ion exchange to chemisorption [The values of 1/nditions . The E vectively . The E vsorption . C0 = 20 and 50 mg F−/L and pH = 3. Correlative parameters are listed in −/L and 0.9999 at 50 mg F−/L) derived from the PSO model are larger than those of PFO (0.7212 and 0.9019) and Elovich (0.8714 and 0.9698) models, elucidating the chemisorption behavior plays a critical role in defluoridation process, which is consistent with the reported results of fluoride removal on other La-based adsorbents [Qe, cal, 98.23 and 184.50 mg F−/g at 20 and 50 mg F−/L) derived from the PSO model are in agreement with measured adsorption amounts . The second-order rate constants (k2) are 3.7012 × 10−2 and 0.5236 × 10−2 g mg−1 min−1 at C0 = 20 and 50 mg F−/L and 25 °C, respectively. This indicates that LMF11 NFs exhibit a higher fluoride removal rate at a lower F− concentration, resulting in a shorter time needed to attain the adsorption equilibrium. Similar results were observed in other La-modified adsorbents [C) are not zero  derived from the plot of lnKD versus 103/T  and ΔS0 (69.12 J mol−1 k−1) confirm the defluoridation by LMF11 NFs is endothermic along with an entropy increase process [The thermodynamic factors , (110), (111), (200), (211), (220) and (310), respectively, portray the cubic perovskite phase of LaMnO3 (JCPDS NO. 74-0440) [3O4 has been generated [2O3. In summary, the LMF11 NFs are composed of La2O3, LaMnO3 and Fe3O4. After defluoridation, the disappearance of the diffraction peak ascribed to La2O3 and the presence of new reflections belonged to the hexagonal structure of LaF3 (PDF No.32-0483) [− is the dominating defluoridation mechanism of hydrous La2O3 in adsorption system. Moreover, the reflection intensity of Fe3O4 and LaMnO3 turns slightly weak after fluoride uptake owing to the decrease in crystallinity caused by the entry of F− into the crystal lattice. The defluoridation mechanism of LMF11 NFs was analyzed by PXRD and FTIR. 74-0440) ,77. The enerated  assigns 32-0483)  state th−1 assign to the surface hydroxyl groups stretching mode, chemisorbed water bending mode, CO32− stretching mode, La-O stretching mode, and Mn-O along with Fe-O stretching mode [−1 are detected. After fluoride uptake, the bands at 3430 and 1632 cm−1 turn obviously weak and two peaks at 1127 and 1067 cm−1 are not detected, validating that the bonding of fluoride through electrostatic interaction as well as the replacement of -OH by F− via ion exchange [As seen from the FTIR spectrum of the neat LMF11 NFs b, the baing mode ,77,79. Ming mode ,80 at 11exchange . Moreoveexchange .− loaded LMF11 adsorbent was further verified by SEM and EDS 3O4 NPs were uniformly dispersed along the fiber axis without obvious agglomeration, and the total pore volume and surface area of LMF11 NFs were 0.276 cm3 g−1 and 55.81 m2 g−1, respectively. The isoelectric point (pHpzc) of LMF11 NFs was 6.47.The LMF11 NFs presented fiber morphology with an average diameter of 626 ± 57 nm. The results showed that the spherical Fe(2)−/g, and the adsorption amount of LMF11 NFs increased with the ascend of initial fluoride and then reached equilibrium. The fluoride adsorption by the LMF11 NFs does not affect by the interfering anions, such as NO3−, Cl−, SO42− and CO32−, with the exception of PO43− anions.The responsible factors such as initial pH values, shaking time, interfering ions, temperature, initial concentration of adsorbate and adsorbent for fluoride retention on LMF11 NFs were systematically conducted. The maximum adsorption amount of LMF11 NFs achieved at a pH of 3. With ascending LMF11 NFs dosage, the adsorption uptake descended from 188.21 to 82.77 mg F(3)− adsorption on LMF11 NFs was seen as a single-layer chemisorption. The maximum binding amount calculated from the Langmuir isotherm model was as high as 173.30–199.60 mg F−/g at pH = 3 at 15–45 °C. A thermodynamic study proved that the defluoridation by LMF11 NFs is a spontaneous, endothermic as well as entropy increase process. The regeneration of the F− loaded LMF11 NFs exhibited a high defluoridation efficiency of 94.76 mg F−/g at C0 = 20 mg F−/L up to three repetitions.The Langmuir isotherm and PSO model were favored divulging that the F(4)On the basis of the isoelectric point, FTIR and PXRD analysis, it was confirmed that LMF11 NFs worked with the defluoridation mechanisms including hydrogen bonding, electrostatic attraction and ion exchange.In this study, magnetic electrospun La-Mn-Fe tri-metal oxide nanofibers (LMF NFs) were developed through electrospinning and heat treatment and used for fluoride recovery. The prepared LMF11 NFs with a La:Mn molar ratio of 1:1 gained the highest fluoride adsorption amount. The results of the investigation were summarized as follows:In summary, the fabricated LMF11 NFs possessed good defluoridation performance and high selective adsorption ability, indicating that the material has great potential as a promising adsorbent for fluoride decontamination. Further study will be conducted on the removal of phosphate, arsenate and other pollutants in the future.