It expands on previous water radiolysis models by incorporating planar or cylindrical solid–water interfaces and by explicitly including γ radiation in dose rate calculations. To illustrate our model's utility, we quantify radiolytic H2 and H2O2 production rates surrounding spent nuclear fuel under different conditions . These examples demonstrate the extent to which α, β and γ radiation contributes to total absorbed dose rate and radiolytic production rates. The different cases also illustrate how H2 and H2O2 yields depend on initial composition, shielding and age of the solid. In this way, the examples demonstrate the importance of including all three types of radiation in a general model of total radiolytic production rates.We present a mathematical model that quantifies the rate of water radiolysis near radionuclide-containing solids. Our model incorporates the radioactivity of the solid along with the energies and attenuation properties for alpha (α), beta (β), and gamma (γ) radiation to calculate volume normalized dose rate profiles. In the model, these dose rate profiles are then used to calculate radiolytic hydrogen (H We then calculate radiolytic production rates based on the dose rate profiles.stopR) is the maximum distance traveled by charged particles; it is determined by initial energy and matrix composition ] depends onP),power (I), andirradiance (a).attenuation  is the initial energy per unit time per solid angle associated with the radiation . It is determined by the decay energy and activity of each radionuclide. We calculate power for individual α and β particles by multiplying the radionuclide activity (A) for each radionuclide (i) by the initial radiation energy (E0) per decay (j) of the same radionuclide and dividing by 4π steradians P and the inverse square of the distance traveled (R) (The magnitude of radiance  for eacheled (R) (2)I at distance R and the attenuation formula from For both α and β particles, linear energy transfer (LET), energy attenuation per distance traveled, increases toward the end of the particle's path. We derive the energy remaining (a) isstopR are both determined using the experimental energy and range data from the ASTAR and ESTAR NIST databases . We sum over each radionuclide for all α and β decays with a unique E0 for the total radiant energy flux per radionuclide. We then sum the total radiant energy flux of all radionuclides to give the total radiant flux density,maxz depends on stopping distance and the distance of the radionuclide from the surface , based on F from Combining equations for μ).We use similar methods to calculate the flux and dose rate for γ radiation. Gamma ray absorption, however, obeys an exponential law (the Beer–Lambert law)  charactex½ = 0.693/μ, is the thickness of material required to reduce the initial energy flux by one half Fγ=∑P and I are the same as those for α and β particles, however the equation for γa is given by γF in known, we calculate the dose rate (γD) due to γ radiation.The equations for G-value (the number of molecules created per 100 eV of energy) (The production rate of radiolytic products is the absorbed dose rate from each type of radiation multiplied by its respective  energy) (8)ProduG-values  depend on radiation type and differ for each radiolytic product. We list the G-values that we use for this study in 2 fuel is directly related to α, β and γ dose rates near the fuel surface. If the integrity of the barrier around a fuel pellet is breached and the barrier is infiltrated by water, radiation from the fuel will dissociate the water. Production of H2O2 can increase dissolution of UO2 in fuel pellets, while H2 suppresses dissolution by consuming H2O2  is small compared to the curvature of the 1 cm diameter fuel rod. We test this assumption by comparing the area irradiated by α and β radiation using the curvature of the pellet to the area irradiated assuming a planar interface. We find that approximately 99% of the absorbed dose is accounted for when a planar surface is assumed.For our example, we chose to use 20-year-old UOon rates . WebKORIon rates . For thiG-values in our model and list the values we use in 2 yields, α radiation chemical yields have been shown to increase with the LET of the particle. However, for the energy range of α particles we are interested in this study (less than 5.8 MeV), LET appears to have a minimal effect on the G(H2) value. At high LET, G(H2) appears to plateau around 1.25 molecules/100 eV  values. Pastina and LaVerne conclude that for heavy ions, the same H2O2 yields, within ±20%, can be used for a wide range of LET. We did not find any work specifically on the effect of β-particle LET, therefore we assume the H2 and H2O2 yields to be constant for βG as well.We also assume constant s/100 eV . For H2O2O2 and H2 as a function of distance from the solid–water boundary. It is important to note that these are production rates, not concentrations. We apply our model to the fuel rod before and after barrier failure, to compare α, β and γ distribution patterns about the fuel rod, and radiolytic production rates, under both conditions. The profiles also differentiate between the contributions of α, β and γ radiation to the dose and H2O2 and H2 production rates. Below, we describe the results of our model with a spent nuclear fuel example.We use our model to quantify the total dose rate and the radiolytic production of HγD decreases away from the cladding–water interface. In βD is higher than the absorbed dose due to α particles. By 1000 µm, α and β particles only contribute about 17% to the average dose rate when barrier failure has occurred. Alpha particles emitted from the solid do not travel further than 50 µm away from the surface of the fuel, while γ radiation from the fuel pellet continues to be absorbed for tens of centimeters.We show the dose rate profiles from our 20-year-old spent nuclear fuel example in We also calculate the dose rate profile for 1000-year-old spent fuel with 55 MWd/kg burn-up after cladding failure . Similar2 production rate is 10 times higher after barrier failure, while radiolytic H2O2 production increases 8.5 times. After barrier failure, α and γ radiation together contribute over 80% of radiolytic H2O2 and H2 production near the pellet–water interface. Alpha radiation contributes to water radiolysis only within 50 µm of the pellet. Farther from the interface, radiolytic production due to β and γ radiation dominates.From the dose rates, we calculate radiolytic production rates . RadiolyDue to the high activity of young fuel assemblies and their close proximity to water in spent-fuel pools, application of our model to a 20-year-old fuel assembly nicely illustrates the model's capabilities. Typically, fuel assemblies are stored underwater in spent-fuel pools for up to 30 years to cool the fuel and provide shielding from radiation . The rad2 gas  and reductants  directly influences UO2 dissolution  spent fuel with breached cladding, β and γ radiation contributes greatly to chemical radiolytic production from young (20-year-old) spent fuel with breached cladding. In the young fuel, γ radiation dominates radiolytic chemical production adjacent to spent fuel with intact cladding.


However, for large scale applications single crystals are not suitable and it is essential to develop high-performance polycrystalline thermoelectric materials. In polycrystalline form, Ca3Co4O9 is known to exhibit much weaker thermoelectric response than in single crystal form. Here, we report the observation of enhanced thermoelectric response in polycrystalline Ca3Co4O9 on doping Tb ions in the material. Polycrystalline Ca3−xTbxCo4O9 (x = 0.0–0.7) samples were prepared by a solid-state reaction technique. Samples were thoroughly characterized using several state of the art techniques including XRD, TEM, SEM and XPS. Temperature dependent Seebeck coefficient, electrical resistivity and thermal conductivity measurements were performed. A record ZT of 0.74 at 800 K was observed for Tb doped Ca3Co4O9 which is the highest value observed till date in any polycrystalline sample of this system.The potential of thermoelectric materials to generate electricity from the waste heat can play a key role in achieving a global sustainable energy future. In order to proceed in this direction, it is essential to have thermoelectric materials that are environmentally friendly and exhibit high figure of merit, ZT. Oxide thermoelectric materials are considered ideal for such applications. High thermoelectric performance has been reported in single crystals of Ca One of the major challenges faced today by the thermoelectric community is to practically find stable TE material whose figure of merit is close to 1Thermoelectric (TE) power generation is considered as one of the most promising and emerging clean energy technologies for harvesting electricity from heat without producing any direct emission of greenhouse gases. Absence of moving parts makes TE devices attractive for many applications including automobiles. The thermal-to-electrical conversion efficiency of a TE material is quantified by a dimensionless quantity called figure of merit ZT which at a given temperature T is related to its Seebeck coefficient (S), electrical conductivity (σ), and thermal conductivity (κ) by the relation ZT = S2 layers sandwiching a block layer comprising of insulating rock salt structure2Co2O4, was reported by Teraski et al. in 1997 wherein they demonstrated that this material exhibits a high Seebeck coefficient of 100 μV K−1 with a low electrical resistivity of 200 μΩ cm at 300 K2Co2O4 for practical applications. Therefore, much research has lately focused on the development of other members of layered cobalt oxide familyet. al. demonstrated a high figure of merit of ~0.87 at 973 K in single crystal Ca3Co4O9 (CCO)3Co4O9 have limitations for practical applications. Over the last one and half decades or so, tremendous efforts have been put to realize high ZT values in polycrystalline Ca3Co4O9 samples. The enhancement in the thermoelectric performance of polycrystalline Ca3Co4O9 over the years has been summarized in Fig. 1(a)789101112131415161718In order to obtain a high figure of merit, the TE material should possess not only a high Seebeck coefficient but also high electrical conductivity and low thermal conductivity. A high electrical conductivity is needed to reduce the joule heating, and a low thermal conductivity is required to maintain a high temperature difference between the hot and cold ends of the TE device. These are the contradictory requirements and there are very few materials which satisfy the above conditions. Layered cobalt oxides are one of such materials. Crystal structure of these materials consists of two alternate CoO3Co4O9, as of yet, has been based on doping approach520212223242526272829302+ by heavy trivalent ions, especially rare-earth (RE) ions, has proved to be most effective in enhancing ZT values. The above substitution results in a decrease in the concentration of holes in the material which in turn increases the value of ‘S’. Heavy RE ions also cause scattering of phonons and hence the thermal conductivity of the material goes down. A brief survey of the ZT with the doping of different rare earth elements is shown in Fig. 1(b)52021222324252627282930Most of the research aimed on enhancing the ZT of polycrystalline Ca3Co4O9 sample. Specifically we are showing that when Ca in Ca3Co4O9 is replaced by Tb ions, its ZT value enhances significantly. Furthermore, we have shown that since the ionic size of Tb ions is quite close to that of Ca, a much higher level of dopants compared to other RE ions can be introduced in the lattice before causing the precipitation of the secondary phases. Tb doped Ca3Co4O9 showed a record ZT of 0.74 at 800 K which is the highest value observed till date in this system. Furthermore, compared to other RE elements, Tb is abundant in nature, economical and less toxic, which makes it a potential candidate for practical applicationsIn this paper, we are reporting the highest value of ZT observed so far for any rare earth doped polycrystalline Ca3−xTbxCo4O9 with x = 0–0.7 were prepared by solid state reaction technique. Stoichiometric amounts of Calcium oxide , Cobalt oxide  and Terbium oxide  were mixed together to obtain a homogenous powder. The resulting powder was then calcined twice at 973 K for 20 hours with intermediate grinding steps. The powder was then compacted into 0.5 inch diameter pellets under a pressure of 45 MPa using a uniaxial press. The compacted pellets were further compressed iso-statically under a pressure of 200 MPa to obtain high and uniform density pellets. Later, the pellets were sintered at a temperature of 1173 K for 20 hours, with re-pelletizing and sintering procedures repeated twice. Finally, the sintered pellets were shaped into rectangular shaped samples of dimensions 10 mm × 4 mm × 2 mm to perform various characterization experiments.Polycrystalline ceramic samples of Caα radiation. Micro-structural characterization was performed using an FEI Quanta 600 FEG scanning electron microscope (SEM). The SEM images obtained were used to determine the grains and their distribution. The transmission electron microscopy (TEM) including selected area electron diffraction pattern (SAED) and energy dispersive spectroscopy (EDS) was performed using a high-resolution JEOL 2800 S/TEM. The TEM specimens were prepared by ultra-sonication of the crushed powder in ethanol followed by drop casting on a Cu holey carbon grid. X-ray photoelectron spectroscopy (XPS) was performed using a Kratos Axis Ultra DLD system with a monochromatic Al Kα source (1486 eV). The component peaks were fitted using Casa XPS software to quantify the elemental composition of the samples. The electrical conductivity, Seebeck coefficient and thermal conductivity were measured over the temperature range 300 K to 800 K. For electrical conductivity measurements, 4-probe technique was employed in which a constant current was passed through the sample using a current source (KI 6220) and the resulting voltage was measured by a nano-voltmeter (KI 2182 A). For Seebeck coefficient measurements, a temperature difference was established between the two ends of the sample and resulting voltage was measured using high-purity platinum wire. Thermal diffusivity was measured by laser flash technique using a pulsed excimer laser (Compex Pro). The heat capacity was measured using a Netzch DSC 3500 Sirus differential scanning calorimeter. By using experimentally measured thermal diffusivity, specific heat and density, thermal conductivity of the samples was determined. The statistical analysis of the measurement result shows uncertainty about 1–2% for electrical conductivity, 2–5% for Seebeck coefficient and thermal conductivity, which is overall about 5–12% for power factor and 7–17% for figure of merit ZT.X-Ray diffraction (XRD) experiments were performed using Philips X’Pert diffractometer with Cu K3−xTbxCo4O9 samples are shown in 3Co4O9 sample shows all the characteristic peaks as reported in JCPDS database and previous reports332+ ion. On increasing the Tb content further, for x = 0.6 sample the intensity of the (002) and (003) peaks diminishes un-proportionally while for the x = 0.7 sample, few additional impurity peaks were detected (as marked in The XRD patterns of the CaSEM micrographs of x = 0 and 0.5 samples are shown in 2 and RS layers can clearly be seen in HRTEM images. The top-right panels show the selected area electron diffraction (SAED) patterns recorded from the respective samples. Since grain sizes in the samples, as seen in SEM micrographs, were of the order of several micrometers, the SAED covered the area containing just few grains. As a result SAED showed speckled ring kind of pattern mixed with diffraction spots of single crystals. The bottom-right panels show the EDS elemental mapping images of Ca, Tb, Co and O in the respective samples. EDS images show the distribution of elements over the whole area examined in the study.TEM results for x = 0 and 0.5 samples are shown in 3/2 and 2p1/2 peaks for the undoped sample lies at 779.93 eV and 795.33 eV, respectively. In literature, it has been reported that for Co3+ state, the 2p3/2 and 2p1/2 peaks occur at 779.6 eV and 794.8 eV, respectively, while for the Co4+ state, the 2p3/2 and 2p1/2 peaks occurs at 781.4 eV and 796.8 eV, respectively353+ and Co4+, it is inferred that the Co in the sample exists in the mixed Co3+/Co4+ valence state. In case of x = 0.5 sample, Co 2p3/2 and 2p1/2 peaks are observed at 779.80 eV and 794.87 eV, respectively. Though this position is still in between the peaks expected for Co3+ and Co4+ but is closer to the peak expected for Co3+. This indicates that in 0.5 Tb doped sample, the concentration of Co4+ ions (and hence of holes) is smaller than that in the undoped sample. In order to understand the observed change in the hole concentration on Tb doping, we recorded Tb 4d XPS spectrum  from 300 K to 800 K for the samples is shown in et al. and Bhaskar et al. where they observed an increase in the electrical conductivity on doping Bi3+ and Dy3+ in the Ca3Co4O9372) with temperature for all the samples. The x = 0.5 sample shows a PF of 1.15 mW m−1 K−2 is at 800 K. This value is about 4 times higher than the corresponding value for the undoped sample.The electrical resistivity vs temperature data for all the samples is shown in RT) was 2.41 Wm−1K−1. On increasing the Tb content, the value of κRT decreased to 1.34 Wm−1K−1 for x = 0.5. On further increasing the Tb content, κRT was found to increase slightly for the x = 0.6 and x = 0.7 samples. For all the samples, the value of κ decreases monotonously with increase in temperature. In order to understand the dominant mechanism responsible for the observed variation in the thermal conductivity, we separated out the electronic (κe) and lattice (κL) components of the thermal conductivities individually. e determined using the experimentally measured σ values in the Wiedemann-Franz law: κe = L × σ × T where, L is the Lorentz number (2.44 × 10−8 W.Ω.K−2)38e from κ, the value of κL was determined as per the relation: κL = κ − κe and is shown in e first increases for x = 0.01 and then starts decreasing while the value of κL shows a monotonous decrease. However, over the entire temperature range of our study the value of κL is much higher than that of κe. As a result total thermal conductivity shows a decrease with increase in x.Total thermal conductivity (κ) of the samples measured by laser flash technique is shown in −3 which increased to 0.12 at 800 K. On doping Tb in the system, ZT value first increased till x = 0.5 and then started decreasing. For all the samples, ZT showed an increase with increase in temperature. The x = 0.5 sample showed the highest value of ZT attaining a value of 0.74 at 800 K.Using the experimentally measured ‘S’, ‘σ’ and ‘κ’ values the figure of merit was calculated. The temperature dependence of ZT for all the samples is shown in 3Co4O9. A wide range of Tb doping was attempted and state-of-the-art structural characterization and electrical and thermal transport measurements were performed. We found that the maximum concentration of Tb that can be incorporated in the material without causing the precipitation of any impurity phase is x = 0.5. It was observed that Tb doping causes a reduction in the relative concentration of Co4+ ions which corresponds to a decrease in the concentration of holes in the material. Reduction in the hole concentration gives rise to a larger ‘S’ while enhanced phonon scattering caused by the heavier Tb ions results in lower ‘κ’. Because of high ‘S’ and low ‘κ’, a high figure of merit of 0.74 at 800 K was observed for the x = 0.5 sample. The enhanced figure of merit at high temperatures along with good chemical stability makes Tb doped Ca3Co4O9 a promising candidate for next generation thermoelectric module applications.In summary, we have explored the effect of Tb doping on the thermoelectric response of polycrystalline CaHow to cite this article: Saini, S. et al. Terbium Ion Doping in Ca3Co4O9: A Step towards High-Performance Thermoelectric Materials. Sci. Rep.7, 44621; doi: 10.1038/srep44621 (2017).Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.


Zizyphus lotus fruit. The trimethylsilyl (TMS) method was successfully used for derivatization of the monosaccharides units of extracted polysaccharides that were released by hydrolysis method. Sugars were identified based on their retention times compared with those of standards and the NIST MS Spectral Library. All sugars were quantified in TIC  mode using calibration curves. Data is related to “Optimization extraction of polysaccharide from Tunisian Zizyphus lotus fruit by response surface methodology: Composition and antioxidant activity”  Gas chromatography coupled to mass spectrometer (GC–MS) was used to identify and to quantify neutral sugars that constitute the water soluble polysaccharides from  Specifications tableValue of the data•The data can be used to identify neutral sugars of polysaccharides by GC–MS.•Method and data provide information for acidic hydrolysis conditions of polysaccharides.•The data provide information for trimethylsilyl (TMS) derivatization conditions of all neutral sugars for GC–MS analysis.1This data shows the GC-EI-MS total ion current profiles of the trimethylsilylated sugars obtained from ZLP and those of standards which injected under the same chromatographic conditions .Fig. 1GC20.45 mg of lyophilized ZLP was hydrolyzed by 1 mL of 2 M trifluoroacetic acid at 70 °C for 2 h in a sealed tube in nitrogen atmosphere. Then, 100 µL of D-myo-inositol  at a concentration of 1800 µg/mL in deionized water was added to the hydrolysate. The released monosaccharides were converted to their trimethylsilyl (TMS) derivatives by adding 100 µL of dry pyridine and 100 µL of N,O-bis(trimethylsilyl)-trifluoroacetamide (BSTFA) to the dried sample. The residue was dissolved in dichloromethane (100 µL) and 1 µL of silylated sample was injected to gas chromatography coupled to mass spectrometer analysis


The emission color can be tuned over the full visible spectral region making them promising for light‒emitting applications. Further control over the optical and magnetic properties of quantum dots (QDs) can be achieved through doping of transition metal (TM) ions such as Mn2+ or Co2+. Here we demonstrate how, following QD synthesis in the presence of a Mn‒precursor, dropwise addition of silicon tetrachloride (SiCl4) to the QDs in toluene results in the formation of Mn‒doped CsPbCl3 QDs showing bright orange Mn2+ emission around 600 nm. Evidence for successful doping is provided by excitation spectra of the Mn2+ emission, with all features of the CsPbCl3 QD absorption spectrum and a decrease of the 410 nm excitonic emission life time with increasing Mn‒concentration, giving evidence for enhanced exciton to Mn2+ energy transfer. As a doping mechanism we propose a combination of surface etching and reconstruction and diffusion doping. The presently reported approach provides a promising avenue for doping TM ions into perovskites QDs enabling a wider control over optical and magnetic properties for this new class of QDs.Nanocrystalline cesium lead halide perovskites (CsPbX Emission and quantum confinement of CsPbX3 QDs that form in Pb‒doped CsX scintillator crystals was reported in the late 1990’s2+ or Co2+ aimed at introducing new optical and/or magnetic functionalities1718192022232+ emission in ZnSe QDs is promising for application in solar concentrators as a result of the large spectral shift of the Mn2+ emission which prevents undesired reabsorptionThe all inorganic analogue perovskites CsPbX2+, but more often doping of QDs proved to be challenging2+ was considered impossible but was later realized using a single source precursor method. More recently with diffusion doping even concentrations up to 30% Mn2+ are now incorporated in CdSe28Doping can be straightforward by simply adding a dopant precursor in the reaction mixture, e.g. in ZnS:Mn2+ into perovskite CsPbCl3 QDs (CPC‒Mn) via a facile colloidal approach. A post synthesis treatment of CsPbCl3 QDs prepared in the presence of Mn‒stearate allows incorporation of Mn2+. Slowly dripping a small volume (5‒20 μl) of SiCl4 to the dispersion of QDs in toluene results in a slow incorporation of Mn2+ at room temperature. Evidence for successful doping is provided by a clear signature of efficient energy transfer from the CsPbCl3 QD exciton to Mn2+ ions. As Mn2+ incorporation proceeds, the excitonic emission intensity decreases and the bright orange Mn2+ emission increases. The excitation spectrum of the Mn2+ emission shows all the features of the CsPbCl3 QD exciton absorption. The life time of the exciton emission decreases as the Mn2+ concentration increases as a result of fast exciton to Mn2+ energy transfer. The proposed mechanism for incorporation involves surface etching and reconstruction and diffusion doping. Decomposition of SiCl4 by trace amounts of water in the toluene yields HCl. Surface etching by H+ and surface reconstruction enabled by chloride ions and Mn2+ form a surface layer and slowly drives Mn2+ into the QDs by ion diffusion. The decomposition of SiCl4 also gives rise to the formation of a protective SiO2 layer giving a relatively high stability to the QDs. Perovskites QDs are known to degrade and show a fast loss of emission intensity under ambient conditions. The PL of the CPC‒Mn QDs synthesized is relatively stable and retains over 25‒30% of the maximum intensity after stirring with SiCl4 for 40 days.Here we report doping of Mn3 QDs were prepared using a similar procedure as reported by Kovalenko et al.2 and the Mn‒precursor (Mn‒stearate) in 1‒octadecene (ODE) were dried at 120 °C after which dried oleylamine (OLAM) and oleic acid (OA) were added. The temperature was raised to 170 °C and a Cs‒oleate solution in ODE was injected. At this temperature the CsPbCl3 QDs form quickly. Next, the mixture was cooled and the NCs were extracted by centrifugation. To remove unreacted polar material the QD precipitate was redispersed in toluene and centrifuged. Precipitated particles were discarded. The final solution of CsPbCl3 QDs in toluene shows the characteristic violet QD emission but no Mn2+ emission, indicating that in this procedure Mn2+ is not incorporated. In the next step a small volume (5‒20 μl) of SiCl4 was dripped to 2 ml of the QD solution. During stirring a white precipitate forms. After different stirring times the precipitate was removed by centrifugation and a transparent QD solution was used for the spectroscopic measurements and characterization. As a control experiment SiCl4 was also added under vacuum conditions in which all water was removed from the toluene.CsPbCl3 QDs made without Mn2+. The XRD pattern of the QDs synthesized in the presence of the Mn‒stearate precursor (green line) is similar. For the product formed after addition of SiCl4 and stirring for 1 h (red line) the formation of amorphous SiO2 is evident from the broad diffraction peak of amorphous SiO2 (PDF Card No. 82‒1576), observed clearly from 15° to 30°. Stirring for 10 days, the pure CsPbCl3 XRD patterns (black line) was still found when the precipitate (SiO2) was removed by centrifuging. The lines in the XRD patterns of the QDs match the powder diffraction pattern of cubic phase CsPbCl3  induces faster decomposition of SiCl4 and silica formation2 covered CPC QD conglomerates all elemental signals of the CPC‒Mn QDs and SiO2  were observed, as shown in In a control experiment, SiCl 2 layer . The dif3 QDs: undoped QDs , CsPbCl3:5%Mn QDs  and CsPbCl3:5%Mn‒SiCl4 . All three QDs show exciton emission around 400~410 nm. The emission spectra of the CPC and CPC:5%Mn samples are almost identical. For the CPC:5%Mn‒SiCl4 sample an extra emission band is observed in the red/orange spectral region around 600 nm. The excitonic emission shows a small (5 nm) blue shift and is slightly broader .To monitor incorporation of luminescent dopants into QDs optical and time resolved spectroscopy are powerful tools. When a luminescent ion is inside of a QD, fast energy transfer from the QD exciton to the dopant is expected. Absorption of light by the QD is followed by exciton‒to‒dopant energy transfer and emission from the dopant ion. Monitoring emission and excitation spectra as well as changes in the time traces of the luminescence intensity of excitonic and dopant emission provides clear signatures for the successful incorporation of a dopant in a QD. In 4T1 → 6A1 emission from Mn2+. Mn2+ can emit in the green to deep red spectral region depending on the local coordination. The Tanabe‒Sugano diagram for ions with the 3d4T1 shifts to lower energies as the crystal field splitting increases363 the Mn2+ will substitute on the octahedrally coordinated Pb2+ site (the Cs+‒site is too large). Since Cl- ions are weak crystal field ligands in the spectrochemical series a green emission may be expected. However, octahedral coordination gives rise to a large crystal field splitting and even for weak ligands orange‒red emission is typically observed for Mn2+. For example, in CdCl2 and a variety of chloride perovskites where Mn2+ is in an octahedral Cl- coordination, the emission maximum of the 4T1‒6A1 emission of Mn2+ has been observed at wavelengths varying between 570 nm and 630 nm364T1‒6A1 emission of Mn2+ in an octahedral Cl- coordination.The 600 nm emission wavelength is typical of the 2+ emission for the CPC:5%Mn‒SiCl4 QDs (blue line) also shows the excitonic features of the QD absorption, indicating exciton‒to‒Mn2+ energy transfer, a signature of incorporation of Mn2+ in the CsPbCl3 QD. Careful inspection of the emission spectra for CPC:5%Mn before SiCl4 addition reveals a very weak emission feature in the orange spectral region  degradation of the CsPbCl3 QDs was observed. As the stability of the CPC QDs is compromised, also the Mn2+ emission intensity decreased for the highest volumes of SiCl4 added. For the reaction volume used (2 ml), the optimum SiCl4 amount equals 10 μl and this volume was used in subsequent experiments to follow the time evolution of Mn2+ incorporation.As shown above, the addition of a small volume of SiCl4 samples taken at different times after injection of SiCl4. In 2+ emission as well as the integrated emission intensity ratios of Mn/QD are shown for different times (1 hour to 10 days) after injection of 10 μl SiCl4 to a CPC:x%Mn QD dispersion prepared with 1, 3, 5, 7 and 10% of Mn‒precursor present in the initial reaction (concentrations in mole% relative to Pb2+). 2+ to QD emission intensity for samples taken after 1 h to 10 days. In 2+ ions may not be homogeneous, a higher concentration of Mn2+ ions in the surface layer can explain a reduced luminescence intensity by concentration quenching through energy migration in a Mn-enriched surface layer. Based on this, it is concluded that the reaction conditions leading to CPC‒Mn QDs with the highest Mn‒emission intensity is CPC:5%Mn with 10 μl SiCl4 for the reaction volume (2 ml) and QD concentrations (20 mg/ml) used here. The changes in luminescence properties over time were followed for the CPC:5%Mn‒SiCl4 (10 μl) by recording excitation and emission spectra as well as luminescence decay curves of the exciton and Mn2+ emission for samples taken after 1 h, and 1, 3, 10 days. The results are shown in To optimize the Mn‒concentration and to probe Mn‒incorporation over time, extensive spectroscopic characterization was done for CPC:x%Mn‒SiCl4 (10 μl) system, the results show a continued increase in the Mn2+ emission intensity relative to the excitonic emission  of CPC:5%Mn‒SiCl4 (10 μl) for various reaction times. Energy transfer from the excitonic state to Mn2+ results in a faster exciton emission decay. The Mn2+ emission decay can be different for surface absorbed Mn2+ than for Mn2+ incorporated in the QD as energy transfer to defects or surface quenching sites is more prominent for surface Mn2+ resulting in shorter decay times for emission from surface Mn2+ ions. The luminescence decay dynamics were recorded for samples taken at several times during the incorporation of Mn2+ for the CPC:5%Mn‒SiCl4 (10 μl) system as this was shown to give the most efficient Mn2+ emission. In i.e. for the decay curves recorded for samples taken after 1 h to 1 day, a lengthening of the decay time is observed. The decay curve for the exciton emission from the sample stirred for 1 day is close to single exponential with a ~16 ns decay time. The initial lengthening of the decay time can be explained by improved surface passivation as a result of the chloride formation following SiCl4 decomposition. Better surface passivation by the Cl- ions remove surface trap states resulting in a more efficient excitonic emission. The 16 ns decay time can be considered to be close to the pure radiative decay time of the exciton emission at room temperature. The results agree with the observation of higher absolute (exciton) emission intensities of CPC QDs in the initial reaction period (up to 1 day) as shown in To gain further insight in the energy transfer process and incorporation of Mn2+ emission intensity strongly increases, while the excitonic emission intensity decreases. Concurrently, the exciton emission decay becomes shorter and also non‒exponential  the decay is close to single exponential with a long 1.4 ms decay time suggesting that Mn2+ is incorporated inside the CPC host and decays with a ms radiative decay rate typical for the spin‒ and parity forbidden Mn2+ emission.4 to the CPC QDs in toluene in the presence of the Mn‒stearate precursor results in the slow incorporation of Mn2+ in the CsPbCl3 QDs. It is interesting to discuss potential incorporation mechanisms and the location of the Mn2+ ions in the CPC QDs. Control experiments help to provide insight in the mechanism. No incorporation of Mn2+ was observed after thorough washing of the CPC‒Mn QDs with acetone before the addition of SiCl4  and with the CPC QD surface to both etch the surface (H+) and reconstruct the QD surface (Cl-). The blue shift of the excitonic emission with reaction time and the TEM images showing a decrease in particle size are consistent with the occurrence of etching of the CPC QDs. The SiO2 formed during the decomposition reaction of SiCl4 and H2O gives rise to SiO2 spheres observed in TEM and possibly also form a thin protective SiO2 layer around the CPC QDs after prolonged reaction times. The increase of the QD exciton emission and lengthening of the exciton emission decay time in the initial hours after addition of SiCl4 signifies the positive role of the HCl in the passivation of surface defect states. Similar observations were made previously for hybrid perovskite thin films2+ and Cl- a Mn‒Cl surface layer can form on the CPC QD. Etching and reconstruction of the CsPbCl3 surface can occur with the incorporation of Mn2+ on Pb2+ sites in the surface layer of the CPC QDs. Subsequent diffusion of the Mn2+ into the CPC QD brings Mn2+ inside the CPC QDs, similar to the Mn2+ doping of CdSe QDs through diffusion doping2+ and SiO2 formation. The reaction mechanism is consistent with the long reaction times of hours to days observed for diffusion doping and cation exchange reactions at room temperature18Hydrolysis of SiCl2+ ions remain located in a surface layer of the CPC QDs. The observation of a transition from non‒exponential decay to a single exponential 1.4 ms decay of the Mn2+ emission is important as it gives evidence that the local surroundings for the Mn2+ ions changes during the reaction. Initially surface Mn2+ ions are present with a shorter life time because of the proximity of surface defect states. For longer reaction times the transition to a purely exponential decay indicates the transition to the same  coordination by Cl- for all Mn2+ ions. This observation however does not prove that the Mn2+ ions are homogeneously distributed through the CsPbCl3 QD. The continued observation of exciton emission indicates that the transfer efficiency is limited and that the Mn2+ ions may be preferentially situated in the surface layer where the coupling with the excitonic wavefunction is weaker. Further research on incorporation of Mn2+ in other perovskite QDs is also interesting. For example, for CsPbBr3 QDs the exciton shifts to lower energies (compared to CsPbCl3) while, because of a smaller crystal field splitting for the weaker Br- ligands, the Mn2+ emission is expected to shift to higher energies. This may lead to a cross-over between the exciton and Mn2+ emission depending on the quantum dot size, similar to what has been observed for Mn2+-doped CdSe quantum dots.24Further research is required to provide a more complete understanding of details on the incorporation mechanism. Based on the present experiments it cannot be excluded that the Mn4‒modified CPC‒Mn QDs solutions was followed over time. The emission intensity shows a slow decrease in intensity but retains 27% of the maximum intensity after 40 days in air, as shown in 4‒modified CPC‒Mn QDs (Mn2+ emission) after incorporation in a poly methyl methacrylate (PMMA) matrix. The decrease of emission intensity with temperature is shown in 4‒modified CPC‒Mn QDs maintained ~50% of their initial intensity.For practical application of perovskite quantum dots stability is an important issue. Thermal and photo stability are often limited for the Pb‒halide perovskites and also sensitivity to moisture is serious issue that hampers the introduction of this new class of materials in products for the solar cell and lighting market40412+ ions in perovskite CsPbCl3 QDs. Upon addition of SiCl4 to a CsPbCl3 QD/Mn‒precursor dispersion in toluene incorporation of Mn2+ occurs as demonstrated by the evolution of strong 600 nm Mn2+ emission (4T1 → 6A1) following excitation in the QD exciton states. Excitation, emission and luminescence life time measurements reveal the presence of energy transfer from the delocalized QD exciton state to the localized 3d2+. The proposed incorporation mechanism involves surface etching and reconstruction by HCl formed through decomposition of SiCl4 in the presence of trace amounts of water in toluene. The presently reported synthesis approach is promising for exploring other transition‒metal or lanthanide ion doped perovskite QDs in order to gain a wider control over the optical and magnetic properties of doped QDs for applications in the fields of lighting and spintronics.In conclusion, we have reported a new doping strategy for the successful incorporation of Mnet al.2CO3 , oleic acid , and 1‒octadecene  were loaded into a 50 ml 3‒neck flask and degassed under vacuum at 120 °C for 1 h. Subsequently, the solution was heated to 150 °C under N2 for all Cs2CO3 reacting with OA.The synthesis approach has been reported by Kovalenko 2 , C36H70MnO4 , and ODE (5 ml) were loaded into a 25 ml 3‒neck flask and degassed under vacuum at 120 °C for 1 h. Afterwards, dried oleylamine  and OA (0.5 ml) were injected into the solution at 120 °C under N2 atmosphere. The temperature was raised to 170 °C and Cs‒oleate solution (0.4 ml) was quickly injected into the mixture. After the desired different reaction time, the mixture was cooled by the ice‒water bath. The NCs were extracted from the crude solution through centrifuging at 3500 r.p.m. for 10 min. First, the supernatant was discarded and the precipitated particles were redispersed in toluene forming the stable colloidal solution. Second, the colloidal solution was centrifuged again, and the precipitated particles were discarded and the supernatant was redispersed in toluene forming the final solution.PbCl4  were introduced into a 20 ml sample bottle containing 2 ml of the colloidal CPC‒Mn QDs toluene solution (H2O content 0.0623%)2) and the transparent solutions were separated through centrifuging at 10000 r.p.m. for 10 min. In a control experiment, 10 μl SiCl4 was mixed with CPC‒Mn QDs (nearly water free) in a vacuum three‒necked flask.Various amounts of SiCl2 , and MnCl2  were mixed thoroughly and loaded into a quartz ampoule. The ampoule was evacuated and suspended with platinum wire in the high temperature region of a sintering furnace. The temperature was raised to 300 °C with a heating rate of 5 °C/min and kept for 2 h under continuous pumping to remove surface absorbed water. Then, the ampoule was sealed and the sample was sintered at 675 °C for 1 h, after which the sample was allowed to naturally cool to room temperature in the furnace.The CPC‒Mn microcrystalline powder was prepared by melting starting material in an evacuated and sealed quartz ampoule4 QDs (20 mg) were added into the PMMA‒toluene solution and stirred for 1 day. The transparent mixture was dripped into a glass container and a copper hold, as shown in the inset of PMMA (0.8 g) powder was dispersed homogeneously in toluene (2 ml) for stirring 1 day. Then CPC:5%Mn‒SiClXRD patterns were recorded using Cu Kα radiation (λ = 1.5418 Å) on a PW 1729 Philips diffractometer, operating at 40 kV and 20 mA. For XRD analysis, the samples were prepared by depositing the QDs colloidal solution on the silicon wafer and evaporating the solvent under vacuum atmosphere.Absorption spectra were obtained using a Perkin‒Elmer Lambda 950 UV/VIS/IR spectrophotometer. Excitation and emission spectra were measured by an Edinburgh Instruments FLS920 spectrofluorometer equipped with a 450 W Xenon lamp. Visible emission (400‒850 nm) was detected by a Hamamatsu R928 photomultiplier tube (PMT).3 QDs were recorded by using an Edinburgh Instruments diode laser with an excitation wavelength of 376.8 nm  and a Hamamatsu H74220‒02 PMT in combination with the Edinburgh spectrofluorometer. PL decay measurements of the Mn2+ dopants were done using an optical parametric oscillator (OPO) system (Opotek HE 355 II) pumped by the third harmonic of a Nd:YAG laser as excitation source and a Hamamatsu R928 PMT as light detection. The samples for optical analysis were prepared by dissolving the crude QDs production in toluene in a quartz cuvette.Photoluminescence (PL) decay curves of the CsPbClTM F200X transmission electron microscope (from FEI), equipped with a high‒brightness field emission gun (X‒FEG) and a Super‒X G2 EDX detector operated at 200 kV. TEM images were obtained using FEI TECNAI T12, operating at 120 kV. The samples for TEM imaging were prepared by depositing the QDs colloidal solution on the carbon‒coated copper TEM grid and evaporating the solvent under inert atmosphere.Scanning transmission electron microscopy (STEM) images and Energy dispersive X‒ray spectroscopy (EDX) measurements were performed using a TalosHow to cite this article: Lin, C.C. et al. Luminescent manganese-doped CsPbCl3 perovskite quantum dots. Sci. Rep.7, 45906; doi: 10.1038/srep45906 (2017).Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.


We also report the synthesis and reactivity of a fluorinated analogue and the structure of its covalently linked intermediate . We conclude that these inactivator reactions mainly involve motion of the pseudo‐anomeric carbon atom, knowledge that should stimulate the design of new transition‐state analogues for use as chemical biology tools.Glycoside hydrolases (GHs) have attracted considerable attention as targets for therapeutic agents, and thus mechanism‐based inhibitors are of great interest. We report the first structural analysis of a carbocyclic mechanism‐based GH inactivator, the results of which show that the two Michaelis complexes are in  We also present the de novo synthesis of a fluorinated inactivator (3) and structural characterization of its covalent adduct with TmGalA.Herein, we describe the first conformational itinerary for a non‐pyranosyl mechanism‐based inactivator of a glycoside hydrolase. That is, we have structurally characterized the Michaelis complexes of bicyclic carbasugar 1 upon binding to TmGalA, we undertook structural studies using X‐ray crystallography. Although a structure of TmGalA has been reported (PDB ID: 1ZY9), in our hands, TmGalA crystallized under different conditions and in a different space group to that reported (see the Supporting Information). The apo TmGalA structure was solved to 1.80 Å resolution (PDB ID: 5M0X) by using molecular replacement with the structure of TmGalA (PDB ID: 1ZY9) as the search model. Like other GH36 enzymes,TmGalA comprises three domains; an N‐terminal β‐sandwich fold domain, a canonical (β/α)8‐fold domain housing the active site, and a C‐terminal β‐sandwich domain with a Greek‐key motif  with 1 in an effort to observe the covalent complex formed through nucleophilic trapping of the bicyclobutenium ion intermediate  by Asp 327. Surprisingly, these crystals failed to yield a structure with any evidence of a small molecule bound in the active site of TmGalA. Crystals soaked for longer periods (1 h to 7 days) did, however, lead to structures with electron density present in the active site. Structures for TmGalA in complex with intact 1  for reaction of the enzyme–inhibitor complex (E:I complex) between  min for TmGalA in complex with 1 provides a rare glimpse of the interactions within a Michaelis complex, which is formed fleetingly upon substrate binding prior to catalysis. Here, the carbasugar binds in a half‐chair 2H3 conformation  and a water molecule (which was not present in the complex with 1) that partially fills the void created by the departure of the aromatic leaving group.The structure of 1 undergoes a pseudo‐glycosylation reaction within the crystalline phase via a transition state that has a different activation free energy barrier to that of the equivalent reaction in aqueous solution.TmGalA in the solid state. Toward this goal, we elected to adopt a common and very successful strategy used in the design of pyran‐based GH inactivators: improving the leaving‐group ability of the pseudo‐aglycone to ensure rapid formation of the covalent enzyme–inhibitor complex while replacing the C2 hydroxy group with a fluorine atom to slow the subsequent hydrolysis.3, in which the pseudo‐C2 hydroxy group is replaced with fluorine atom and the leaving group is a 2,4‐dinitrophenylate rather than a 3,5‐difluorophenoxide anion  as a single diastereomer in analytically pure form through precipitation and thus could be executed on a greater than 10 g scale.6 in excellent yield. A two‐step hydrolysis of the orthoester then gave the unstable epoxy acid 8, which was immediately coupled to TMS‐diazomethane to afford the diazoketone 9. As a key design element, we envisaged coincident formation of the bicyclo[4.1.0]heptane core and the all‐carbon quaternary center at C6 through an intramolecular rhodium‐carbenoid cycloaddition reaction.9→10), it was expected that the transition‐state structure leading to the l‐altro cyclopropane (not shown) would involve considerable strain and, consequently, that this process should favor formation of the desired d‐galacto‐configured 10. We were delighted then that our initial attempts to effect the intramolecular cycloaddition2(OAc)4 in anhydrous3 delivered a mixture of the desired cyclopropane 10 and the unexpected C−H insertion product 11 (10/11=2:1). While reaction in benzene afforded a 1:1 mixture of cyclopropane 10/11, in CH2Cl2, an acceptable 3:1 mixture of these compounds was produced in around 60 % yield.10 proved to be labile and partially decomposed during chromatographic purification. Therefore, it was optimal to treat the crude cyclopropanation reaction product directly with fluoride and isolate the stable α‐fluoroketone 12. Acetylation followed by reduction with l‐Selectride afforded the fluorohydrin 13, which underwent nucleophilic aromatic substitution with dinitrofluorobenzene to afford the arylether 14. Finally, sequential removal of the acetate and acetonide protecting groups delivered the second‐generation inactivator 3 in good overall yield.Our concise de novo synthesis of the densely functionalized fluorocarbasugar 3, this compound proved to be significantly less reactive with TmGalA than inactivator 1. Specifically, at both 25 and 37 °C, we noted little or no enzyme inactivation over 6 h. After raising the temperature to 50 °C and the concentration of 3 to 500 μm, TmGalA activity decreased with a first‐order rate constant of 2.86×10−5 s−1. We therefore measured the kinetic parameters for inactivation of this thermostable enzyme by 3 at 60 °C and found kinact/Ki=0.71±0.19 m−1 s−1 and kinact=(2.25±0.21)×10−4 s−1 ×10−2 s−1.TmGalA after removal of excess 3, however, after 5 days at 60 °C, we observed no increase in enzyme activity,Despite the enhanced leaving‐group ability of the dinitrophenolate in carbasugar TmGalA with 3, we incubated TmGalA with 3 at 60 °C overnight prior to crystallization, which was followed by buffer exchange to remove any unreacted 3 and the 2,4‐dinitrophenol produced during inactivation. New crystallization conditions were identified (see the Supporting Information), and we obtained structural data for this modified TmGalA to 1.55 Å resolution. As indicated in Figure TmGalA unambiguously confirms that the carbasugar skeleton was covalently linked to the nucleophilic acid (Asp327) of TmGalA.Using the information gleaned from the kinetic analysis of TmGalA alkylated by the carbocyclic fragment from 3 provides insight into the reactivity of the nonclassical carbocation intermediate formed within the constraints of the active site during catalysis. Interestingly, the trapped carbasugar intermediate maintains the bicyclo[4.1.0]heptane framework of 3, rather than one of the many possible rearranged skeletons,N2 displacement of the dinitrophenolate. The carbasugar binds in a non‐perfect half‐chair 2H3 conformation, where the C2 atom is also slightly above the plane containing the two bridgehead carbons and C1 and C4.1 and 2 (the C1 atom moves by 1.56 Å), the positions of the three unaltered hydroxy groups are essentially unchanged  should be addressed to the authors.SupplementaryClick here for additional data file.


The rate of redox reaction between KMnO4 and carbon was controlled by the concentration of KMnO4 in a neutral solution. The MnO2 nanoparticles grow along with (211) crystal faces when the redox reaction happens on the surface of fibers under 1D constraint, while the nanoparticles grow along with (200) crystal faces when the redox reaction happens in the bulk of fibers under 3D constraint. The composite, where MnO2 nanoparticles are formed in the bulk under a constraint, yields an electrode material for supercapacitors showing good electron transport, rapid ion penetration, fast and reversible Faradaic reaction, and excellent rate performance. The capacitance of the composite electrode could be 1282 F g−1 under a current density of 0.2 A g−1 in 1 M Na2SO4 electrolyte. A symmetric supercapacitor delivers energy density of 36 Wh kg−1 with power density of 39 W kg−1, and can maintain 7.5 Wh kg−1 at 10.3 kW kg−1. It exhibits an excellent electrochemical cycling stability with 101% initial capacitance and 95% columbic efficiency even after 1000 cycles of charge/discharge.Growing MnO Therefore various carbon materials have been employed, activated carbons78101113141516171820212224Depending on the specific charge storage mechanisms, capacitances are divided into double layer capacitance (EDLC), relying on the high surface area of electrode materials to adsorb charges at the electrode/electrolyte interfaces, and pseudo-capacitance, relying on Faradaic reversible redox reactions near the surface of the electrode12 nanoparticles with carbon materials is still a challenging task. Since reversible redox reaction of MnO2 occurs only near the surface of the electrode because of small diffusion length of the electrolyte across the MnO2 as less than 10 nm, it has been greatly desirable to prepare nanostructured amorphous MnO2 to rise the utilization of its electrochemical activity. There are 4 representative methodologies for MnO2 nanocomposite preparation, including electrodeposition, electrostatic interaction assembly, in-situ redox deposition and chemical co-precipitation2 nanocomposite electrode. Langet al. showed that a supercapacitors made of nanoporous gold and nanocrystalline MnO2 resulted in a specific capacitance of 1145 F g−1 at 50 mV s−1. Fei group2/3D RGO composites. The 3D RGO could provide channels for rapid ionic and electronic transport. The maximum specific capacitance of 709.8 F g−1 at 0.2 A g−1 of the MnO2 (66.4 wt%)/RGO composite. After 1000 cycles, 97.6% of the initial capacitance value can be maintained. Chenet al. prepared nanostructured MnO2/CNTs-sponge composite electrodes, which could give a high specific capacitance as 1230 F g−1 at 1 mV s−1. Based on the comparison in Homogeneous blending of MnOet al. and Maet al. prepared MnO2/C composite based on an in-situ reaction between MnO4− and carbon as the following reaction,To develop a flexible and cost-effective method, Wang2 obtained were submicrometers to micrometers. To improve the performance of MnO2/carbon fibers, Gao groupin-situ growth of MnO2 on the surface of activated carbon fibers. Gao−1. 116 F g−1 was kept after 3000 cycles. Quet al. reported that the MnO2/PCNFs exhibited a specific capacitance of 520 F g−1 at 0.5 A g−1 and 92.3% retention of the initial capacitance after 4000 cycles in a 6 M KOH aqueous solution. It should be indicated that although porous carbon materials were used, both of them pointed out the growth of MnO2 on the surface of the fiber.However, their attention was paid mainly on the influences of the length of reaction time17312/CNFs electrode, nano-MnO2 in the bulk of porous carbon nanofibers are needed. Amorphous carbon materials have strong reactivity with KMnO4. As the electrospun carbon fibers are amorphous3334354 diffusion into the bulk and reaction of KMnO4 with carbon on the surface of the carbon fibers. If the diffusion time is shorter than the reaction time, the KMnO4 can diffuse into the bulk. Otherwise, the MnO2 crystals cover the surface of the fibers and block the diffusion route of KMnO4 into the bulk. Developing MnO2 with much smaller sizes and high dispersion in carbon scaffold is urgently needed.To further improve the capacitance of the MnO2 in the bulk of porous CNFs by controlling the kinetics of an in-situ redox reaction of KMnO4 under nanopore-constraint of CNFs  is smooth and the diameters of fibers are in the range of 300–600 nm  observations were carried out by using thin slices of the fibers. 2 adsorption/desorption measurements were performed at 77 K to characterize the surface areas, pore volumes and the pore-size distributions of MC1, MC2, MC10 as well as pristine CNFs. All isotherms shown in 2 have only negligibly small amount of mesopores. When the concentration of KMnO4 increased to 0.12 g L−1 (MC10), pore-size distribution in micropore range becomes very similar to the pristine fibers (MC0), but the mesopores at 2–8 nm in MC0 disappear completely. On MC2 composite prepared by using the concentration of KMnO4 of 0.024 g L−1, very different distribution of micropore-size is observed.NBET and total pore volume Vtotal. With increasing KMnO4 concentration for the MnO2 loading onto MC0 fibers, both SBET and Vtotal increase slightly. For MC2 composite, however, surface area and pore volume contributed by the pores with the sizes less than 0.7 nm, S<0.7 nm and V<0.7 nm, decrease markedly and those by 0.7–2 nm increase markedly, although only slight increases in these parameters are detected for MC10.In 2 , the diffraction peaks of MnO2 at around 12, 17, 24, 37 and 66° in 2θ, of which index are 110, 200, 220, 211 and 002, respectively (JCPDS No. 44–0141).X-ray diffraction patterns are shown for the composites, including MC0, in D/IG was calculated to be 2.7, indicating the turbostratic structure of the CNFs. For the composites, one additional band from MnO2 is observed at 646 cm−1, which can be assigned to the symmetric stretching vibration of Mn-O in the MnO6 groups252 become more and more distinct with the increase of KMnO4 concentration.Raman spectra of the composites are shown in To investigate the chemical composition and chemical states of various elements in the composites, XPS analysis was performed. The spectrum in 422SO4 electrolyte with different scan rates from 2 to 100 mV s−1 are shown for the composites in Cyclic voltammetry (CV) curves in 1 M Na−1 are summarized for the composites in 2. The specific capacitance calculated from these curves with 0.2 A g−1 was 1282 F g−1 for MC2, while 151 and 515 F g−1 for MC0 and MC10, respectively. When increasing the charging/discharging current density to 10 A g−1, all curves present representative relatively linear voltage-time relationship with quasi-symmetric triangular shapes compare to the charging/discharging curve at 2 A g−1, revealing a remarkable reversibility, which indicate the capacity is mainly contributed by physically adsorption under larger current density.Galvanostatic charge/discharge curves measured with the current density of 0.2 and 10 A g−1 and then gradually. For MC2, the specific capacitance remains as high as 335 F g−1 at a current density of 50 A g−1, the capacity retention ratio being 26%. The capacitance retention with the current density of 1 A g−1 is shown for MC2 in 2 as the limitation of electrolyte diffusion into the bulk of the fibers. After 200 cycles, as the electrodes gradually soaked by electrolyte, more MnO2 are involved in the charge storage process. The retention ratio start to increase, then slowly exceed its initial value. Electrodes display an ideal cycling stability with full permeation of electrolyte after approximate 500 cycles, reaching a remarkable capacitance retention of 101%. 95% Coulombic efficiency remained after 1000 continuous cycles measured using the galvanostatic charge-discharge technique.The dependences of specific capacitance on current density are plotted for the composites in 2 and Warburg impedance Zw. It is seen that all composite electrodes exhibit the nearly vertical line along the imaginary axis in the low-frequency region, revealing an ideally capacitive behavior of the electrode materials due to the fast and reversible Faradic reaction of nano-MnO2. It should be noted that the slope is larger with MC10 than MC2. The Rs for composite electrodes remained almost the same because this parameter is insensitive to the electrode surface. In the high frequency region, the charge transfer resistance (Rct) was 11.5 and 2.5 Ω for MC2 and MC10, respectively, which seems to depend strongly on the concentration of KMnO4 during loading of MnO2.The electrochemical impedance spectroscopy (EIS) of the composite is displayed in 2/CNFs composites in −1 with power density of 39 W kg−1, and can maintain 7.5 Wh kg−1 at 10.3 kW kg−1. Comparing the present composite MC2 with those published on electrospun carbon fibers and activated carbons, the performance of MC2 surpasses most of MnO2/carbon composites reported12131420444546474849505152Ragone plot is presented to characterize the MnO2SO4 aqueous electrolyte, the potential window of the supercapacitor can be 1.6–2.2 V14492/C composite performance is about 292 F g−1 (2.5 A g−1). (2) The other is designing electrode with high specific capacitance. Many efforts have been done to design the host material micro-architecture for loading nanoscopic MnO2 to balance electron and ion transport inside the composites, as well as mass loading. Except the energy and power density shown in 4 and carbon in neutral solution can help to achieve thinner MnO2 with higher specific capacitance and high energy density12132 can be loaded2 nanocrystals in the bulk because the rapid intercalation/deintercalation of Na+ on the surface layer of MnO2 has been widely accepted as the charge-storage mechanism in mild electrolytes. The combination of MnO2 and porous CNFs can achieve excellent capacitive performance. More interface between the MnO2/conductive material can improve the MnO2 utilization mass ratio. In most of the cases shown in 2 crystal size is tens of nanometers. Although the MnO2 is exposed to the electrolyte, the interface between the MnO2 and carbon is limited. The interface between 3D contrained nano-MnO2 and porous CNFs has been enhanced dramatically while the resistance of the matrix increases with more loading of nano-MnO2 in the bulk of CNFs. So it is a trade-off between the energy density and power density. Furthermore, comparing CNFs with activated carbons and porous carbon spheres, the pores in CNFs are shallow ones since the fiber diameter is only sub-micros while the pores in activated carbons or carbon spheres are deep and narrow. The electrolyte penetration resistance of CNFs should be smaller than that of activated carbons. So, the electron transport loops and the diffusion loops of cations in most of the cases are longer than the structure shown in this work, which is related to the power performance of an electrode.Two strategies are commonly used to develop electrode with high energy density under high power density. (1) One is to extend the potential window in aqueous electrolytes via designing an asymmetric configuration50541449562SO4 electrolyte, two charge storage mechanisms of MnO2 are proposed, including surface adsorption/desorption and bulk insertion/extraction of Na+, which are showed belowAt the interface of electrode and Na2 crystals with an interlayer spacing of about 0.24 nm were composed of on the external surface of the fibers, corresponding to α-MnO2 (211) crystal faces, while MnO2 with an interlayer spacing of about 0.5 nm were deposited on the inner surface of the pores at the central part of fibers, exhibiting the character of α-MnO2 (200) crystal faces. The results are in good agreement with XRD information. The results indicate that MnO2 nanoparticles grow along with (211) crystal face when the redox reaction between KMnO4 and carbon happens on the surface of CNFs, in other words, under an un-limited 1D constraint; while MnO2 nanoparticles grow along with (200) crystal face when the redox reaction happens inside of mesopores, under a strong 3D constraint from the inner surface of pores.To get more insights into the advantages of the MC2 electrode, a high resolution TEM has been used. 4 and carbon to form MnO2, the electrospun CNFs prepared from PI without PVP were used as scaffold for MnO2 by the same condition as MC2 (KMnO4 concentration of 0.024 g L−1). Pore-size distributions for the pristine (MnO2-unloaded) and MnO2-loaded CNFs are shown in 2-loading is very similar to that happened in the composite MC2  crystal face when KMnO4 and carbon redox reaction happens with an un-limited 1D constrain; while the MnO2 growth along with (200) crystal face when KMnO4 and carbon redox reaction happens with a comparable 3D constrain. The optimized KMnO4 concentration in our case should be 0.024 g L−1. The good capacitance performance of the MnO2/CNFs composite electrode is clearly attributed to its unique nanostructure. The porous CNFs serves as a highly conductive matrix for fast ion and electron transport, while the ultrathin MnO2 nanocrystals on the inner surface of the pores enable a short diffusion path of electrolyte and provide a more electrochemically active surface area for pseudocapacitance through fast and reversible Faradic reaction.A 3D MnO−1, Sinopharm Chemical Reagent Co., Ltd), 4,4′-oxydianiline (ODA) , polyvinyl pyrrolidones (PVP) , N,N-dimethylformamide (DMF)  and potassium permanganate   were used as received.Pyromelliticdianhydride (PMDA) . The blended solution was stirred at 0 °C for 24 h. Electrospinning parameters were set as follows: applied voltage of 20 kV, tip-to-collector distance of 20 cm and flow rate of 0.8 ml h4 was dissolved into de-ionized water. The concentration of KMnO4 solution were prepared at 0.012 g L−1 (0.076*10−3 M), 0.024 g L−1 (0.152*10−3 M), and 0.12 g L−1 (0.76*10−3 M). Composites derived from these solutions were labeled as MC1, MC2 and MC10, respectively.KMnO4 solution and the weight ratio of CNFs to KMnO4 was set to 4:1. The CNFs was taken out until the purple color of KMnO4 solution had fade into golden brown. The mass loading remains to 10–11% for all samples. The MnO2/CNFs composites were rinsed by deionized water for several times, and finally dried at 80 °C for 10 h under vacuum condition.The carbon fiber electrode was soaked into the KMnO2-loaded, embedded in epoxy resin was cut into 20 nm slices using a microtome  with a diamond blade (DiATOME) and then placed onto a 300 mesh Cu grid for examination by TEM.The morphology of CNFs was investigated using a scanning electron microscope . A high-resolution transmission electron microscopy (HRTEM) was employed to characterize the distribution and crystallization of the composites. The CNFs, pristine and MnO−1 in a range of diffraction angle 2θ from 5° to 70°. Raman spectrum was recorded by HR800 (HORIBA spectrometer) with a wavelength of 633 nm. X-ray photoelectron spectroscopy (XPS) measurement was performed with a thermo ESCALAB 250 spectrometer. XPS data analysis was performed using Thermo Avantage software. Sorption/desorption isotherms of N2 at 77 K were measured by a Belsorp Max apparatus (Japan) and analyzed to evaluate specific surface areas of the composites by using the Brunauer-Emmett-Teller (BET) method, SBET, and pore-size distribution of the composites by density functional theory (DFT) method.X-ray diffraction (XRD) pattern of the composites was examined using D/MAX-RM 2000 at the scanning rate of 2° min2SO4 solution under vacuum for 24 hrs before being assembled. Cyclic voltammetry (CV) and galvanostatic charge/discharge analysis were carried out to evaluate electrochemical performance of the electrodes with a potential window ranging from 0 to 1 V in 1 M Na2SO4 electrolyte using an electrochemical workstation . The electrochemical impedance spectroscopy (EIS) was conducted in the frequency range of 0.01 Hz ~100 kHz with perturbation amplitude of 5 mV versus the open circuit potential. The average specific capacitance was calculated according to the charge/discharge tests based on the following equation:A symmetrical two-electrode supercapacitor was assembled using two pieces of the composite and a separator to investigate the electrochemical performance as an electrochemical capacitor. To achieve sufficient saturation of the electrodes and separator by the electrolyte, they were immersed in 1 M NaI, t, m and U are the charge/discharge current (A), the discharge time (s), the total mass of active materials in the two electrodes (g) and the potential after the deduction of IR drop (V), respectively.where How to cite this article: Le, T. et al. In-situ growth of MnO2 crystals under nanopore-constraint in carbon nanofibers and their electrochemical performance. Sci. Rep.6, 37368; doi: 10.1038/srep37368 (2016).Publisher’s note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.


We found that in contrast to naive expectations Fe is nearly 3+ in this material, which strongly affects its magnetic properties and makes it qualitatively different from well known sulfide analogue - FeS2. Doping, which is most likely to occur in the Earth’s mantle, makes FeO2 much more magnetic. In addition we show that unique electronic structure places FeO2 “in between” the usual dioxides and peroxides making this system interesting both for physics and solid state chemistry.Recent discovery of the pyrite FeO FeO2 is expected to appear in the Earth’s lower mantle below 1800 km, according to1 in the pyrite structure, and start to dominate over other Fe oxides at higher pressures. The composition of the mantle is extremely important for the seismology, since it determines convection processes. There were proposed a number of structural models based on different ratio of ferropericlase (a solid solution of FeO and MgO), bridgmanite O3, 2SiO4 olivine and other compounds5, but none of them took into account the existence of FeO2. Moreover, physical properties of this material are completely unexplored. One might expect that they can be highly unusual, since on one hand Fe ion in FeO2 formally should have exceptionally high oxidation state, 4+. Since the O-O distance in FeO2 is 1.89 Å it is not likely that there can be a strong bonding between the O ions, like in molecular oxygen (where the O-O bond distance is 1.21 Å) and one may indeed expect that Fe ions will adopt 4+ valence state and then FeO2 is in a negative charge transfer regime9. This may result in self-doping10 and also to bond or charge disproportionation12, inversion of the crystal field splitting13 or nontrivial magnetic structures. On the other hand, the presence of the ligand-ligand dimers may also strongly affect physical properties of FeO2 as it does in the actual pyrite FeS2 (“the fool’s gold”). However, O-O distance in FeO2 is 1.89 Å, much larger than in molecular oxygen (1.21 Å) or (O2)2− ion as in the usual peroxides like BaO2, MgO2 (1.49 Å).Recent discovery of a new iron oxide FeO21, see Fig. 2 as a complete analogue of FeS214. Since FeS2 is known to be a diamagnetic insulator with Fe ions adopting 2+ valence state17, one might expect that the same is true for FeO2. The first indication that such a picture is oversimplified follows from the recent theoretical study14, where FeO2 was found to be metallic at the pressures where it does exist.Iron peroxide was found to have the same pyrite crystal structure as FeS2. We show that FeO2 is completely different from FeS2, and so are the physical properties of these compounds. The oxidation state of Fe ion in FeO2 is not 2+, as in FeS2, but close to 3+. This strongly affects magnetic properties of FeO2, since having 3d5 electronic configuration, Fe3+ ions may have a magnetic moment. Our comprehensive theoretical calculations using combination of the density functional and dynamical mean-field theories (DFT + DMFT) demonstrate that there is indeed a highly nontrivial temperature dependence of the magnetic susceptibility in FeO2. We found out that the origin of the difference in magnetic properties between FeO2and FeS2 and of the metallic character of FeO2 is a much smaller bonding-antibonding splitting for ligand σ orbitals in the peroxide dimer O2 as compared with S2, and a total shift of oxygen 2p levels relative to 3p levels of sulfur. This feature of the electronic structure is rather general and important for other dioxides, which can exist in Earth’s mantle or in inner parts of exoplanets.In the present paper we describe electronic and magnetic properties of FeO2, electronic and magnetic properties of which are well understood. As discussed above, one might naively expect to have Fe4+ ions with 3d4 electronic configuration in FeS2, since usually sulfur has a valence 2-. This would shift Fe 3d band very low in energy, below S 3p, and would result in a self-doping and a metallic conductivity7, which strongly disagrees with the experimental fact that FeS2 is a semiconductor17.We start with FeS2 dimers. There are sulfurs 3p orbitals, which are directed exactly to each other in these dimers. They form such a strong bond that the antibonding eg orbitals, see Fig. 2)2− dimers), and Fe ions become 2+ with the 3d6 electronic configuration. Fe ions are in the ligand octahedra in pyrite structure. Strong crystal field splitting between the tg2 and eg bands  counteracts the Hund’s rule and stabilizes the low spin configuration with all six 3d electrons occupying tg2 sub-shell. This makes FeS2 diamagnetic and insulating19.The explanation of this contradiction lies in the specific features of its crystal structure, namely the presence of the S2 is rather different from a sulfide counterpart. We sketched how this difference appears in Fig. 18), starting from the hypothetical FeO2 having FCC lattice (like NaCl), where O ions do not form dimers and where there are basically three bands O p, Fe tg2, and Fe eg, see Fig. The electronic structure of FeOp levels are shifted down relative to the Fe 3d states, as compared with the 3p levels of sulfur. Besides, as was mentioned above, the presence of the ligand-ligand dimers in real FeO2 results in bonding-antibonding splitting, but since oxygen 2p orbitals are much less extended than sulfur 3p orbitals, this bonding-antibonding splitting in the O2 dimer is expected to be much smaller. As a result the antibonding O eg (like in FeS2), but exactly in the place, where Fe tg2 bands lie, see Fig. tg2 electrons would be transferred to oxygens, shifting Fe valence in the direction of 3+. Second, the hybridization between Fe 3d and O tg2 and O 2p , but a metal.This salient feature of FeOtg2 bands, each doubly degenerate with respect to spin, below the Fermi energy for the unit cell consisting of four formula units (f.u.), which are occupied by 4 electrons per each Fe ion . In addition there are four bonding tg2 states . On the other hand it is just on the border line between bands corresponding to localized tg2 electrons and antibonding molecular 2 as we will show latter.Second, the Fermi level appears to be in a very specific position. On one hand it is almost in the pseudogap, so that the Stoner criterion for ferromagnetism (FM) is formally not fulfilled, and this is the reason why magnetic solutions does not survive in the GGA . In spite of the fact that there are magnetic moments on Fe ions, they do not order, so that FeO2 stays paramagnetic down to 190 K (we checked FM and AFM-I). Even lower temperatures can be reached in our calculations by using a truncated Hamiltonian, which includes only Fe tg2 and O 2p states . In this case we were able to go down to 60 K, and again FeO2 does not order in our calculations even at these temperatures. This may seem somewhat surprising since having a rather large bandwidth (and hence hopping parameters) one might expect large superexchange interaction between Fe ions, if spins would have been localized.Correlation effects manifest themselves basically via the renormalization of the GGA DOS near the Fermi level, 23. The width of this correlator is inversely proportional to the lifetime of a magnetic moment. For example in a pure metallic iron, where 24. From the inset of Fig. 2 FWHM of the spin-spin correlator is ~3 eV, which demonstrates that the magnetic moments can hardly be considered as localized.In order to estimate the degree of the spin localization we calculated the analytical continuation on real frequency of the spin-spin correlator In DMFT one can calculate the uniform magnetic susceptibility m is the magnetization, 25. Detailed analysis of these data18 shows that such an unusual for 3D system behavior is due to a specific position of the Fermi level in between the localized tg2 and antibonding tg2 states and Here qualitatively different states . There are many different elements besides Fe (5.8%) and O (44.8%) in the Earth’s mantle, and one may expect that Mg (~22.8%), Si (~21.5%), Ca (~2.3%) or Na (0.3%)26 may dope FeO2 and change its properties dramatically. Indeed, the Fermi level in stoichiometric FeO2 is on the steep slope of a large peak in DOS, and changing its position we strongly affect both magnetic and electronic properties.This means that the electron and hole doping, which is likely to occur in Earth’s mantle, would result in a very different temperature dependences of magnetic susceptibility, since we shift the Fermi level to the peaks corresponding to 2 dimers, existing in the pyrite structure. Thus, at first sight the hole doping is expected to be much more effective for making FeO2 magnetic: the Fermi level would then be shifted to the large peak corresponding to localized Fe tg2 electrons.The electron doping will shift the Fermi level to antibonding molecular-like 2 group from 3- in FeO2 to 2- in MgO2 2− unit appears just above the Fermi level and does not provide any holes to the Fe ions. In NaO2 superoxide the O2 “molecule” is 1-, see Fig. 27, and hence by Na we depopulate O tg2 electrons. This leads to the shift of the Fermi level downwards, see Fig. 2:Si nonmagnetic, but only in unrelaxed crystal structure. After lattice optimization there appears two very different O2 dimers, which help to form magnetic moment ~0.4 We checked different types of hole and electron dopings by the GGA calculations (for ferromagnetic order) performing full structural optimization, starting from the pyrite structure and substituting 25% of Fe by different ions such ions as Mg, Si, and Na. Mg doping formally changes valence of the peroxide Osee Fig. , but it 2 to be magnetic due to different types of doping or because of non-stoichiometry. It is hard to expect, however, that FeO2 would order magnetically in the Earth’s mantle, because of very high temperatures, ~1000–2000 K, but even in a paramagnetic state it may still provide local magnetic moments. The direct DFT + DMFT calculations within the rigid-band shift approximation  from a ligand to a metal, 9.There is a well known concept in physics, introduced by Zaanen, Sawatzky and Allen2 part of electrons are transferred from sulfur to Fe. CoS2, NiS2, MgO2, KO2 and many other materials are just the same: ligand σ* and sometimes even 2 one returns to normal transition metal oxides, where oxygen’s valency is 2-, but there is still one step to make since O is 1.5- in FeO2. Thus, we see that FeO2 indeed lies “in between” oxides and peroxides/disulfides, which makes it an especially interesting material from physical point of view.In peroxides the situation with the CT energy is “inverted” from the beginning: as we have seen, in FeSsee Fig.  at least2, VO2, its “natural” state is O2−, and counting from that, we see that e.g. in FeO2 the CT energy would be negative, 2− to Fe . But in peroxides, as well as, e.g., in FeS2, the natural “structural unit” is the O2 or S2 dimer. Such dimer can be in different charge states: neutral O2 molecule, Fig. 2)− molecular ion , Fig. 2)2− ion as in MgO2, Fig. A simple qualitative difference between normal (di)oxides and peroxides is the following: On one hand, when the main “structural unit” in a system is a single O ion, like in dioxides of the type of TiO2 dimers. But till (O2)2− it is still reasonably harmless, we fill “weakly” antibonding states 3− dimers is 1.89 Å1 - much larger than 1.49 Å for (O2)2− dimer in MgO228 or 1.32 Å for (O2)− in NaO229. Already MgO2, having 4 electrons on antibonding 30. In FeO2 we lose even more energy occupying antibonding σ* orbital, see Fig. 2 even less stable in the pyrite structure, than MgO2, so that it can be stabilised only at a very high pressure.The more electrons we put on such a dimer, the more we fill antibonding states, which gradually destabilizes the very O21 is even more exotic than it was initially thought. Unexpected valence states, nontrivial magnetic properties, stabilization of local magnetic moments by non-stoichiometry or doping by such abundant constituents of Earth’s mantle such as Si (and Na) and finally its special place between (di)oxides and peroxides make FeO2 extremely interesting not only for geoscience, but also for the condensed matter physics and solid state chemistry.Summarising, we see that the recently discovered pyrite-like FeONo datasets were generated or analysed during the current study.Supplementary information


In addition, the S‐RGOH sensor is highly responsive to volatile organic compounds. More importantly, the characteristic patterns on the linearly fitted response–temperature curves are employed to distinguish various gases for the first time. The temperature of the sensor is elevated rapidly by an imbedded microheater with little power consumption. The 3D S‐RGOH is characterized and the sensing mechanisms are proposed. This work gains new insights into boosting the sensitivity of detecting various gases by combining chemical modification and 3D structural engineering of RGO, and improving the selectivity of gas sensing by employing temperature dependent response characteristics of RGO for different gases.Reduced graphene oxide (RGO) has proved to be a promising candidate in high‐performance gas sensing in ambient conditions. However, trace detection of different kinds of gases with simultaneously high sensitivity and selectivity is challenging. Here, a chemiresistor‐type sensor based on 3D sulfonated RGO hydrogel (S‐RGOH) is reported, which can detect a variety of important gases with high sensitivity, boosted selectivity, fast response, and good reversibility. The NaHSO Herein, we report a 3D sulfonated RGO hydrogel (S‐RGOH)‐based gas sensor with boosted performance by combining chemical functionalization and 3D structural modification of RGO with sulfonated groups  , but much smaller than that of GO (6.95 Å), demonstrating the existence of π–π stacking between RGO sheets in S‐RGOH.−1) and a G‐band (1593 cm−1) appear on the Raman spectra of GO and S‐RGOH. The D‐band implies the presence of the structural defects or attachments of functional groups on the carbon basal plane, which agrees well with previous study.ID/IG from GO to S‐RGOH suggests the formation of new sp2 clusters after NaHSO3 reduction.3 in the aqueous S‐RGOH solution by centrifugation and washing several times with deionized (DI) water, the elemental analysis of X‐ray photoelectron spectroscopy (XPS) was carried out , a small proportion of O atoms (23.7%), and some S atoms (3.6%) . The ratios of C/O and C/S were calculated to be 3.1 and 20.2, respectively, reflecting that twenty C atoms were modified with one HSO3− group on average.NaHSOerature <00 °C (Fierature <00 °C (Fierature <00 °C (FiI–V) curve of the S‐RGOH in Figure I–V lines demonstrate that the resistance of S‐RGOH (376 kΩ) is higher than that of unmodified RGOH (37 kΩ) at room temperature, which can be attributed to the increased defect sites introduced by NaHSO3 molecules.The synthesized 3D S‐RGOH was uniformly dispersed in water under ultrasonication. After depositing the aqueous dispersion of S‐RGOH on Au interdigital electrodes (IEs) by drop‐casting, the S‐RGOH bridged the gaps on the Au IEs. The Au IEs and microheater were fabricated by micromachining technologies. The current versus voltage (2.2R/R0 (%), in which R and R0 are the resistances after and before exposure of the sensor to a test gas, respectively. Furthermore, the response is defined as: (R0 − R)/R0 (%) = ΔR/R0 (%). It is clear that the normalized resistance of the S‐RGOH sensor decreases rapidly and significantly upon exposure to NO2 .The normalized resistance is defined as 2 Figure b. It is s Figure c,d. As s2 concentration relationship in the detection of NO2 with higher concentrations from 6 to 13 ppm . The LOD can be calculated from the linearly fitted response versus NO2 concentration line . The LOD of a sensor can be obtained when the noise level is lower than one third of the signal level.2 exposure, the sensitivity, noise level, and theoretical LOD can be extrapolated . As such, the sensitivity of this 3D S‐RGOH sensor was calculated to be 8.69 ppm−1, which was 65.3 times larger than that of RGOH sensor (0.133 ppm−1).2 sensors based on Gr/RGO materials (sensitivity: 0.001–0.443 ppm−1) .2 is recommended.2, which is the lowest NO2 concentration provided by our current experimental setup. The ability to detect NO2 with low concentration indicates the potential practical applicability of this gas sensor.Notice that the S‐RGOH sensor also displayed a linear response versus NO2 molecules. The HSO3− is an electron‐withdrawing group and has many lone pairs of electrons.2 tend to adsorb on electron‐rich sites of sulfonated groups.2 adsorption, the electron transfer from p‐type S‐RGOH to NO2 leads to a higher concentration of holes in S‐RGOH, reducing the resistance of S‐RGOH. Furthermore, Dai et al. predict that S‐doped Gr is able to chemically bind NO2 strongly based on first‐principle calculations.3− groups, gas molecules can be desorbed from the S‐RGOH surface facilely in the air purge process. In addition to functional HSO3− groups, the sensitivity is also enhanced by the pore‐filling of NO2 in the 3D porous structures of S‐RGOH, which provides increased adsorption site density.2 sensing performance of the sensor in high relative humidity (RH) condition is also studied . Due to the interference of water molecules, the response to 10 ppm NO2 deteriorated 30% from 0% to 70% RH. This is because the electron‐withdrawing NO2 and electron‐donating H2O molecules make the resistance of p‐type S‐RGOH decrease and increase, respectively. However, the prehumidification treatment of S‐RGOH can be employed to suppress the interference of humidity on NO2 detection.2 gas was introduced in the chamber at 70% RH. It was found that the response only reduced 7.1% in this case as compared with that at 0% RH. The prehumidification makes the response of the sensor to 70% RH reach a saturated state before NO2 exposure and therefore the sensor becomes less vulnerable to high RH.The improved response of the S‐RGOH sensor compared with its unmodified counterpart is ascribed to the strong interaction between sulfonated functional groups and NO2, this 3D S‐RGOH sensor is also highly responsive to NH3. The sensor displays fast and substantial increase in its resistance upon exposure to 20–1000 ppm NH3 . Although the lowest NH3 concentration limited by our current gas sensing setup is 20 ppm . Thus, the response time t50 of this sensor to 20 ppm NH3 is only 16 s, revealing the rapid response of the sensor . When the sensor was exposed to 200 ppm NH3 in three successive cycles to study the repeatability, a constant response of 13% with a small variation of 3.8% was observed . The sensor exhibited full recovery in each cycle, indicative of good reversibility. To investigate the effect of NH3 exposure time on the response, we found that the response increased monotonically with prolong NH3 exposure time until the sensor reached the stabilized/saturated response to 200 ppm NH3 at 800 s .In addition to NOH3 Figure4. The 2 gas, NH3 is an electron‐donating reducing gas.3 molecules on S‐RGOH surface, the electron transfer from NH3 to S‐RGOH reduced the carrier (hole) concentration in S‐RGOH. The hole‐depletion in p‐type S‐RGOH leads to an increased resistance. The boosted NH3 sensing response of S‐RGOH compared with unmodified RGOH is attributed to the chemical reaction between sulfonic groups and NH3, forming ammonium salts.3− groups and NH3. NH3 molecules can also fill in the pores of 3D S‐RGOH, increasing the interaction area. Furthermore, the hopping of charge carriers across the micro/nanoscale pores offers alternative paths for charge transport, leading to a greater signal level.3, the electron‐donating water molecules also cause the increase of resistance of S‐RGOH. Consequently, the response of the S‐RGOH sensor to 30 ppm NH3 at 70% RH was enhanced 118% due to the overlapped response from humidity . Nevertheless, the response of the sensor to 30 ppm NH3 only increased 6% from 0% to 70% RH after prehumidification treatment.Different from the electron‐withdrawing nature of oxidizing NO2 and NH3, the S‐RGOH sensor also exhibits high sensitivity to a variety of VOCs. For example, the responses of this sensor to saturated vapors of methanol, ethanol, acetone, toluene, and chloroform reached as high as 40%, 33%, 30%, 25%, and 16.2%, respectively .Ead) between S‐RGOH and gas molecules.2 molecules (−0.831 eV) than NH3 molecules (−0.003 eV) and H2O molecules (−0.019 eV).2 distance (1.47 ) is much smaller than that of S‐doped Gr—NH3 (4.03 ) and that of S‐doped Gr—H2O (3.88 ).2 compared with NH3 and H2O molecules. Therefore, the elevated temperature within proper range weakens the response of this Gr sensor to NH3 and H2O molecules, but not to NO2. There is a competition between adsorption and desorption of NO2 molecules on S‐RGOH surface during the NO2 sensing process.2 dominates, while desorption is negligible due to the strong bonding between S‐doped Gr and NO2.2 increased slightly from 22 to 45 °C. However, some reactions between electron‐withdrawing O2 and electron‐rich HSO3− groups may happen at further elevated temperature to 65 °C, leading to a reduced response. It demonstrates that an appropriate temperature is important to optimize the sensitivity of S‐RGOH sensor. To the best of knowledge, this is the first time that the locally elevated temperature is employed to enhance the selectivity of detecting various different gases. This method may shed light on the improvement of the selectivity of various chemical sensors by programing the temperature of sensing materials in future work.The varied response–temperature relationship for different gases should be attributed to different adsorption energy  was much larger than that of Pt lines (25 µm). The Au IEs were fabricated by another photolithography, sputtering of 5 nm Cr/100 nm Au, and lift‐off process. Both the width and gap of Au strips on the IEs were 20 µm. After drop casting of the aqueous dispersion of S‐RGOH on the Au IEs and drying, the 3D S‐RGOH bridged the gaps on the IEs and thus could be employed for gas sensing.Material Characterizations: FE‐SEM 7600 was utilized to characterize the morphology of 3D S‐RGOH, Au IEs, and microheater. X‐ray powder diffraction spectra of the prepared S‐RGOH were obtained by a Bruker D8 advance X‐ray diffractometer. Raman spectra of samples were acquired using a spectrophotometer  with a 514 nm laser. XPS spectra were obtained using a Kratos XSAM 800 spectrometer with a Mg Kα (1253.6 eV) X‐ray source.Gas Sensing Test: A fixed voltage (0.1 V) was applied on the sensor and the resistance change was monitored by a Keithley 2602 SourceMeter. Dry synthetic air was utilized to dilute the test gas to the desired concentration and also clean the gas chamber before and after test gas exposure. The VOCs were delivered by a room‐temperature saturated vapor stream method. The sensing test was performed on the same sample for five times to obtain the average values and the standard deviation of responses or recovery percent. The gas sensing properties were measured at 22 °C in ambient air, if without special notification.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


Their long phosphorescence lifetimes in living cancer cells give rise to high quantum yields for the generation of 1O2, with large 2‐photon absorption cross‐sections. 2 is nontoxic to cells, but potently cytotoxic to cancer cells upon brief irradiation with low doses of visible light, and potent at sub‐micromolar doses towards 3D multicellular tumor spheroids with 2‐photon red light. Photoactivation causes oxidative damage to specific histidine residues in the key proteins in aldose reductase and heat‐shock protein‐70 within living cancer cells. The oxidative stress induced by iridium photosensitizers during photoactivation can increase the levels of enzymes involved in the glycolytic pathway.Strongly luminescent iridium(III) complexes, [Ir Surface cancers are particularly accessible to PDT, but light can also be directed to a range of other tissues using, for example, fiber optics. Examples of clinical photosensitizers include hematoporphyrin derivatives (Photofrin)N,N‐chelated RuII complex,II bacteriochlorophyll derivative.III complexes, are often luminescent and might be useful in theranostic procedures in which optical imaging is used to monitor the uptake of the sensitizer before subsequent irradiation then kills the cancer cells.Selective activation of nontoxic photosensitizers in cancer cells by spatially‐directed light is an attractive regimen for therapy because of the minimal damage to normal cells, especially if the sensitizer is preferentially taken up by cancer cells.1O2 and/or other so‐called reactive oxygen species (ROS). The use of red and near IR light (λ=600–1100 nm) also enhances the clinical application of PDT because of its low energy compared to either UV or visible light.1O2 diffusion distance is short (≈1 μm),1O2,1O2 in living cells.A key advantage for metal complexes is the heavy‐atom effect, which favors fast singlet‐to‐triplet intersystem crossing (ISC), thus giving longer excited‐state lifetimest λ=600–10 nm also2]+ (1) and [Ir2] (2). Their absorption and emission properties depend on the chelated dithione  and diketonate  ligands, as well as their effectiveness in killing cancer cells in both two‐dimensional (2D) and three‐dimensional (3D) cell cultures, and the ability to generate 1O2 in cancer cells. Moreover, the photoactivation in cancer cells leads to specific oxidative damage in some key cellular proteins.Here we compare the photophysical, chemical, and biological properties of two highly luminescent organoiridium phenylquinoline  complexes,  by 1 and 2 upon irradiation with λ=465 nm light were determined as 0.73 and 0.81, respectively (see Table S7). These values are much higher than those for the well‐known [Ru(bpy)3]2+ [Φ(1O2)=0.22].III complexes can produce cellular 1O2 after irradiation with λ=465 nm light or λ=750 nm laser, 2D A549 monolayer lung cancer cells and 3D A549 tumor spheroids were incubated with the IrIII complexes and the fluorescence probe 2,7‐dichlorodihydro‐fluorescein diacetate (DCFH‐DA). Cells treated with only the DCFH‐DA control or the iridium complexes in the dark showed no enhancement of fluorescence. In contrast, a significant increase in fluorescence from DCFH‐DA was observed following light irradiation of the cells treated with either 1 or 2 . These findings suggest that 1 and 2 generate 1O2 efficiently in cancer cells upon light irradiation.To demonstrate that these Ir2 was nontoxic to both A549 lung cancer and MRC‐5 normal lung fibroblasts (IC50>100 μm), while 1 showed moderate toxicity to A549 cells . The potency of both 1 and 2 towards cancer cells increased markedly upon irradiation, notably sub‐micromolar for A549 cells. The phototoxicity index (PI) of 2 towards to A549 cancer cells is greater than 333. Under the same experimental conditions, 5‐ALA and cisplatin displayed much lower phototoxicity (IC50>100 μm).PDT in cell experiments was first tested by incubating monolayer cancer cells with the compounds at various concentrations for 2 hours. There was no loss of cell viability after irradiation in the absence of Ir complexes . In the dark, 1 and 2 towards 3D multicellular tumor spheroids (MCTSs) were further investigated. Untreated MCTSs were unaffected by 1P and 2P irradiation . The complex 2 showed no toxicity towards A549 MCTSs in the dark, and also no toxicity towards MRC‐5 MCTSs in either the dark or upon irradiation. However, the IC50 values for 2 towards A549 spheroids upon 1P and 2P irradiation were very low, 1.0 μm and 0.23 μm, respectively (see Table S8). Interestingly, 2P photocytotoxicity is higher than 1P photocytotoxicity. Although 1 showed good phototoxicity to A549 spheroids, it was not so selective, thus exhibiting some phototoxicity towards the MRC‐5 spheroids.The 2‐photon (2P) photocytotoxicities of 1 or 2 in 3D A549 spheroids was further characterized by 2‐photon confocal laser scanning microscopy. Strong red luminescence was observed on the surface of spheroids to a depth of about 130 μm with λ=458 nm excitation and down to about 212 μm with λ=750 nm excitation. The luminescence images were captured every 16.3 μm along the z‐axis. The spheroids exhibited much stronger luminescence in deeper layers of cells with 2P excitation, thus indicating deeper penetration of the 2P light  were treated with 2  either in the dark or with λ=465 nm light irradiation. Proteins from these cells were collected and digested as described in the Supporting Information. LC‐MS/MS data were searched for oxidative modifications of the side chains of cysteine, histidine, tyrosine, phenylalanine, methionine, and tryptophan, as well as asparagine and glutamine deamidation and serine, threonine, and tyrosine phosphorylation (see the Supporting Information). With 1 % false discovery rate (FDR) against the database, only proteins which presented three times or more (in triplicates) under all three conditions  were quantified. 212 proteins and 1500 peptides met these requirements and were identified. Two‐tail t‐tests ensured the consistency of the reference lysozyme peptide  throughout all samples (see Table S9). All p values were calculated to be larger than 0.05 . Hence, the oxidized peptide/lysozyme peptide area ratios were calculated using t‐tests to analyze the significance of changes in quantities of the oxidized peptides under different conditions. The levels of five oxidized peptides (detected three times or more out of the five replicates) were significantly different (see Table S10). These peptides were quantified against the standard lysozyme peptide as described in the Supporting Information. Two unique oxidized peptides have p values of less than 0.05  in the screening against 2D and 3D (spheroid) cancer cell models. In previous reports, the specific nature of the damage to proteins in the cell, induced by photodynamic therapy, has been little studied. We found that 1O2 generated by 2 can oxidize specific histidines in the proteins Hsp 70 and aldose reductase (AR), which have important functions in cancer cells. Hsp 70 is a molecular chaperone for nascent proteins, and aberrantly folded, damaged, or mutated proteins and AR is a monomeric reduced nicotinamide adenine dinucleotide phosphate (NADPH)‐dependent enzyme, a member of the aldo‐keto reductase superfamily. This work appears to be the first report showing that specific sites of cellular Hsp‐70 and AR can be oxidized during PDT. The combination of oxidative stress induced by the photoactivation of 2 together with the malfunction of mitochondria in cancer cells leads to the increased use of glucose to generate energy, and is consistent with the observed increase in the levels of all enzymes involved in the glycolytic pathway (by factors of about 2.1 to 5.3‐fold).In summary, we designed efficient new organoiridium photocatalytic sensitizers which were nontoxic in the dark and highly and selectively cytotoxic to cancer cells when irradiated by 1P and 2P irradiation  should be addressed to the authors.SupplementaryClick here for additional data file.