The nanoparticles are characterized by means of energy‐dispersive X‐ray spectroscopy (EDX), powder X‐ray diffraction (PXRD), scanning electron microscopy (SEM), and transmission electron microscopy. In addition, the role of different processing approaches on the thermoelectric properties ‐ Seebeck coefficient as well as electrical and thermal conductivity ‐ is demonstrated.The synthesis of phase‐pure ternary solutions of tetradymite‐type materials (Bi Tetradymite‐type ternary (BixSb1‐x)2Te3  nanoparticles were synthesized in an ionic‐liquid approach. The effect of thermal annealing on the chemical composition and morphology of the nanoparticles as well as their transport properties prior and after annealing is reported. Despite an increasing demand for energy, considerable amounts of untapped energy in the form of waste heat derive from almost every electrical and mechanical process, including machine operation, oil refining, steel making and food production. In recent years, waste‐heat harvesting and recovery using thermoelectric (TE) materials have attracted significant attention as a promising technology to generate clean energy and reduce carbon emission.2σT/κ)zT= and bismuth telluride (Bi2Te3) as well as their solid ternary solutions (BixSb1−x)2Te3 have already found their way in many commercialized thermoelectric devices due to their remarkable performance near room temperature. that are typically used in solution‐based methods often lead to Te‐ or Bi‐rich materials due to the incorporation of unwanted metal impurities into the crystal lattice.3[Bi3I12]single source precursor (Et2Sb)2Te and were subsequently subjected to a post‐thermal treatment. This modified route ensures the formation of pure nanoparticles by removing any ionic liquid residues. Aside from the effect of thermal annealing on the chemical composition and nanostructure of the resulting nanoparticles, we herein also report on the effect of the nanoparticle processing on the resulting transport properties. The chemical composition, phase purity, morphology, and particle surface of the obtained nanomaterials were examined by energy‐dispersive X‐ray spectroscopy (EDX), powder X‐ray diffraction (PXRD), scanning electron microscopy (SEM), and transmission electron microscopy (TEM).We herein report on the synthesis of phase‐pure ternary tetradymite‐type materials (Bi22.1single source precursor (Et2Sb)2Te3[Bi3I12] was applied in the wet‐chemical synthesis of binary Bi2Se3 and Bi2Te3 nanoparticles.4mim]3[Bi3I12] compared to many other bismuth precursors eliminated the occurrence of unwanted side reactions such as homolytic bond breakage reactions and formation of Bi double layers. Together, the combined decomposition of these precursors will ensure the production of ternary alloys with the correct stoichiometry and defined concentrations of the metal (Bi/Sb) dopant.Low temperature syntheses of tetradymite type materials often require reactive metal organic precursors with well‐defined decomposition pathways to guarantee the formation of highly stoichiometric products. The thermal decomposition of the Et2Sb)2Te61, 65xSb1‐x)2Te3  nanoparticles from a modified route by reaction of (Et2Sb)2Te with different amounts of [C4mim]3[Bi3I12] in [C4C1Im]I.2Sb)2Te at 150 °C, the reaction was stirred for 12 h, after which the resulting metal chalcogenide powders were repeatedly washed with acetonitrile and isolated by centrifugation. All experimental steps including nanoparticle synthesis and purification were strictly performed under inert gas conditions to avoid surface oxidation of the resulting materials (Scheme xSb1−x)2Te3 nanoparticle series were annealed at 250 °C under a vacuum medium of 10−6−10−7 mbar for 24 h and were then naturally cooled down to room temperature.We herein synthesized BixSb‐x2Te3 (x2.2xSb1‐x)2Te3 nanoparticles  were confirmed as phase‐pure ternary materials 2Te3 . The increasing substitution of the antimony sites by bismuth atoms can be observed from the reflection shifts towards lower reflection angles and from the increase of the lattice parameters. In this case, the gradual exchange of the smaller antimony atoms by larger bismuth atoms would eventually expand the unit cell volume. Most importantly, the steady shift of the highest intensity reflection proves the alloying of a single‐phase material rather than the formation of a separate BixTey component.According to PXRD, the as‐prepared BixSb‐x2Te3 naxSb1−x)2Te3 revealed crystallite sizes of 47, 22, and 20 nm for x=0.25, 0.5, and 0.75 using the Scherrer equation  is consistent with the decreasing size of the crystalline domains in response to the increasing bismuth concentration. Rietveld refinements on the PXRD patterns of BixSb−x2Te3 re2O3, Bi2O3, TeO2), metal Sb or metal Bi, as well as additional BixTey phases were also not observed.However, the exclusive sharpening of the (110) reflection at 41.61° (2θ), that is observed in the sample containing the highest bismuth content, indicates a preferential growth along the ab‐plane perpendicular to the c‐axis. Shouldering of the (015) and (1010) reflections, which would indicate the presence of elemental Te, are absent in any sample of the alloy series. Other impurity phases including crystalline oxidation products (i. e. SbxSb1−x)2Te3 based materials by eliminating crystal defects and modifying the chemical composition, grain size, and carrier concentration.3[Bi3I12]. NMR analysis revealed that the thin amorphous film consists of ionic liquid residues including [C4mim]3[Bi3I12] and [C4C1Im]I . As expected, the color intensity and glass coverage of the ionic liquid residue that was extracted from the (BixSb1‐x)2Te3 samples increased with higher concentrations of bismuth. The heat treatment serves as a final purification step to remove ionic liquid residues that remained on the particle surface after the washing procedure. The presence of such impurities could especially compromise the electrical mobility of these nanomaterials. The XRD peak assignment of the alloy series after annealing verified the phase purity of these materials as ternary solid solutions of (BixSb1−x)2Te3 2Te3 nanoparticles is almost 3.5 times larger (Scherrer's equation) compared to the original samples 2Te3 could be approximated to the desired compositions where x=0.28, 0.51, and 0.65. The EDX results agree with the PXRD measurements, hence confirming the high stoichiometry of the (BixSb1−x)2Te3 nanoparticle alloys.Quantification of the EDX spectra within standard deviations confirmed the stoichiometric compositions of the ternary alloy series which closely correspond to the theoretical Bi,Sb : Te ratio of 40 : 60 Table . The bisEDX elemental mapping analyses revealed a homogeneous distribution of Bi, Sb, and Te within the materials Figure .xSb1−x)2Te3 ternary materials showed the formation of anisotropic nanoparticles with the hexagonal plate‐like morphology 2Te3 nanoparticles ranging from 30 to 200 nm in size and 10 to 20 nm thickness 2Te3 samples. High‐resolution TEM measurements (HR‐TEM) confirmed the crystallinity and clean surfaces of the (BixSb1−x)2Te3 nanoparticles. In particular, the HR‐TEM dark field image of the (Bi0.75Sb0.25)2Te3 ternary phase in [110] orientation displays uniform quintuple layers of 1.02 nm thick, which confirms that there are no additional Bi‐bilayers in the lattice structure before (as‐prepared) and after annealing 2Te3 ternary phase .Due to the elimination of shape and size‐directing capping agents, large size distributions of small spherical particles as well as large intergrown plates can be observed in the resulting 2Te3 2Te3 materials is thus proven to adopt a layered and infinite arrangement of quintuple layers (QLs) that are stacked along the c‐axis . In each QL, one could observe five atomic layers with a Te(1)‐Bi‐Te(2)‐Bi‐Te(1) stacking sequence that is terminated by a Te(1) atomic layer on both sides. More specifically, the bonding between the Bi or Sb and Te atoms inside the quintuple QL are best described as strong covalent and ionic interactions, whereas the interactions among the adjacent QLs consist of weak van der Waals forces.Additionally, the SAED pattern of Bi0.7Sb0.252Te2.3zT materials, but makes an assessment of the influences on the thermal conductivity reduction difficult by the multitude of potential sources like alloy composition itself, nanoparticle structure, or other defects that are typically produced by the ball milling. In part, even the amorphization of the nanoparticles by the high energy ball milling may be expected. Motivated by the excellent purity and crystalline integrity of the here‐synthesized nanoparticles, we assess these influences in our compacted nanoparticle pellets. The synthesis and all subsequent processing steps were performed under inert gas conditions, hence avoiding the effect of oxygen impurities. We further optimize the nanoparticle processing steps and discuss the in‐plane and through‐plane thermal conductivity of the obtained nanocrystalline bulk pellets.Thermoelectric properties of alloyed chalcogenides are often optimized by the approach of high energy ball milling followed by rapid compaction like SPS.2.42Te3 for this study, characterized in‐depth in an earlier work,2Te3 were produced by keeping the processing parameters identical and the complete thermoelectric properties were characterized  of all samples within the given error bars. These results are comparable with published data,The hot‐pressing procedure resulted in reproducible over‐all properties and almost identical 2.5−6−10−7 mbar for 24 h at temperatures below 300 °C. The Seebeck coefficient is the transport coefficient that is most sensitive to any changes in the Fermi level of the semiconductor and therewith the chemical composition of the nanoparticles. Therefore, in Figure The here‐synthesized alloy nanoparticles were in part objected to an annealing procedure. This additional annealing step was intended to verify the cleanliness of the nanoparticles since residual IL could be removed by exposing the nanoparticles to a base pressure of 102.6The consolidation of the powder may affect the transport properties of the pellet. The sintering affects the microstructure of the macroscopic pellets, but also the reconstruction of the individual interfaces at the atomic level.0.75Sb0.25)2Te3 were produced by the optimized hot‐pressing procedure and a spark plasma sintering procedure, respectively. Because of the different technologies used to densify the nanoparticles, parameters like temperature and pressure had to be adapted. Figure −3 (81 %), while the SPS processed sample showed only a density of 5.1 g cm−3 (70 %) after the typical signatures of sample shrinking detected within the SPS machine by tracking the movement of the tools. Note that higher temperatures or longer hold times realized after the observation of sample shrinking typically only result in partial melting.−1 at RT that was higher than the electrical conductivity of 50 S cm−1 of the SPS sample. This was caused by both, an improved charge carrier density as well as an improved charge carrier mobility of this n‐type alloy sample. We conclude that the hot‐pressing procedure generates the conditions for an efficient thermal activation of carriers in the pellet. Further, defects, especially pores, that reduce the carrier mobility are less favored by the hot‐pressing than by the SPS processing. Hence, hot‐pressing is better suitable to allow for a dense arrangement of the here‐used nanoparticles in the compacted pellet, also resulting in improved electrical properties.This experiment was designed in a way to directly compare the influence of hot‐pressing and spark plasma sintering. For this, two pellets of the same nanoparticle batch of the composition 2Te3 and (Bi0.25Sb0.75)2Te3. The change in the thermal conductivity with Bi content is consistent with what is expected from previous literature.2Te3 and Bi2Se3 nanoparticles.Figure 0.25Sb0.75)2Te3 a ball milling and hot‐pressing approach resulted in isotropic samples, which showed thermal conductivities of around 1 W m−1 K−1 at room temperature. While this value corresponds to the here‐presented through‐plane direction, it is significantly reduced compared to the thermal conductivity in in‐plane direction,0.7Sb0.3)2Te3, the in‐plane thermal conductivity was found to be around 1.3 W m−1 K−1 at room temperature for a sample that was also exposed to a combination of ball milling and SPS.xSb1−x)2Te3 nanoparticles followed by SPS processing−1 K−1 at room temperature for a composition of (Bi0.25Sb0.75)2Te3. This comparison shows that the thermal transport properties of the obtained product are not mainly defined by the chemical composition, but rather by the sum of the synthetic details such as the used additives, as well as the processing details including the compaction method. In an earlier study, we already showed that high purity Bi2Te3 and Bi2Se3 compacted nanoparticles exhibited an exceptionally high thermal conductivity close to that of single crystals or even ab‐initio calculations for bulk.Further we find that the thermal conductivity is higher than expected for nanocrystalline bulk of this alloy compositions. Single crystalline reference data of similar compositions is added to Figure zT values of these alloyed nanoparticles range from zT=0.06 to zT=0.15 at 400 K for the composition of (Bi0.25Sb0.75)2Te3 in in‐plane direction, depending on the details of the applied processing.While the here‐synthesized nanoparticles provide a close‐to‐perfect model system to study the influence of chemical composition and nanostructure on the transport properties, they obviously cannot compete with the best thermoelectric materials of the nominally identical composition due to their high thermal conductivity which results from the extraordinary high purity and crystalline integrity of the nanoparticles. Consequently, 3xSb1−x)2Te3 were synthesized by thermolysis of the single source precursor (Et2Sb)2Te with various amounts of the Bi source [C4mim]3[Bi3I12] in a surfactant‐free ionic liquid‐based approach in [C4mim]3I. The phase purity as well as the perfect stoichiometric composition and homogenous element distribution within the resulting nanoparticles was proven by XRD, SEM, EDX and TEM. After compaction of the nanoparticles, the thermoelectric transport properties, Seebeck coefficient, electrical conductivity and thermal conductivity were determined. Our studies clearly proved the crucial role of nanoparticle processing on the resulting transport properties.Highly stoichiometric, phase‐pure ternary solid solutions of the type , 1‐halobutanes  and CH3CN  were commercially available, while [C4C1Im]I, [C4C1Im]3[Bi3I12], and (Et2Sb)2Te were prepared by literature methods.3[Bi3I12] was dissolved in 5 mL of [C4C1Im]I, the red ionic liquid solution was stirred under an Ar atmosphere at 150 °C. 1028 μl (4.1 mmol) of (Et2Sb)2Te were added dropwise yielding a black suspension which was stirred at 150 °C for 12 h. The resulting black precipitate was centrifuged and repeatedly washed with 20 mL of dry acetonitrile (6x). The solid product was dried in vacuum at ambient temperature. An aliquot was transferred to a closed ampoule to be annealed at 250 °C for 24 h under dynamic vacuum (10−7 mbar).The corresponding amount (see Table TGA/DTA studies were performed using a Mettler‐Toledo DSC Star1 system that is operated under inert gas conditions.1H (400 MHz) and 13C{1H} (75.5 MHz) NMR spectra (δ in ppm) were recorded using a Bruker AVNEO 400 MHz spectrometer and were referenced to internal DMSO‐d6 .2Te3 (PDF 15–863) from the ICSD database was used. For each Rietveld refinement, the instrumental correction determined using a standard powder sample LaB6 from NIST  as reference material =4.15689 Å) was considered.PXRD patterns were collected with a Bruker D8 Advance diffractometer with Cu Kα radiation  using a Si single crystal as sample holder to minimize scattering. The powder samples were re‐dispersed in EtOH on the Si surface and investigated in a 2θ range from 10 to 90° with a step size of 0.01° 2θ (counting time 0.6 s). Rietveld refinement was done with the program package TOPAS 4.2 (Bruker) to determine lattice parameters and average crystallite sizes by using the Scherrer equation.The particle morphology and elemental composition of the powdered samples were analyzed by scanning electron microscopy (SEM) using Apreo S Lovac microscopes equipped with a Bruker Quantax 400 units (EDX). TEM studies were carried out on a Jeol JEM 2200 fs microscope equipped with probe‐side Cs‐corrector operated at 200 kV acceleration voltage.D ρ cp where D is the thermal diffusivity, ρ the density, and cp the heat capacity. Bulk literature values were used for cp.Pellets were prepared using a manual hydraulic press (Specac Atlas) and a custom‐made heating mantle equipped with a PID temperature control. A pressure of 100 MPa was applied for 90 minutes while the temperature of the stainless‐steel die was kept at 300 °C. One pellet was sintered by spark plasma sintering procedure using SPS 210‐Gx (AGUS). For sintering, a pressure of 35 MPa was applied for 1 minute at temperature of 240 °C. The density of the pellets was determined geometrically. Thermoelectric properties were measured from room temperature up to 240 °C. The thermal diffusivity was measured along the pressing direction and perpendicular to the pressing direction (Linseis LFA). Thermal conductivity was calculated from the diffusivity data with the relation κ=Supporting Information: PXRDs including the Rietveld refinement of (BixSb1‐x)2Te3  nanoparticles, crystal structure of (BixSb1‐x)2Te3 materials, and 1H‐NMR spectrum of ionic liquid contamination from nanoparticle annealing.The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


The photocatalyst were prepared and characterized by UV–Vis, XRD, TEM, EDS, Nitrogen physical adsorption and XPS technique. UV–Vis analysis revealed addition of F, S and Bi into the lattice of TiO2 led to the expansion of TiO2 response in the visible region and hence the efficient separation of charge carrier. The photocatalytic potential of as prepared catalyst to degrade acrylonitrile wastewater under simulated and natural sunlight irradiation was investigated. The extent of degradation of acrylonitrile wastewater was evaluated by chemical oxygen demand (CODCr). CODCr in wastewater decreased from 88.36 to 7.20 mgL−1 via 14 h irradiation of simulated sunlight and achieved regulation discharge by 6 h under natural sunlight, illuminating our photocatalyst effectiveness for refractory industrial wastewater treatment. From TEM results, we found that SiO2 could disperse the photocatalyst with different component distributions between the surface and the bulk phase that should also be responsible for the light absorption and excellent photocatalytic performance. The XPS analysis confirmed the presence of surface hydroxyl group, oxygen vacancies.Highly active photocatalyst, having certain anti-ionic interfering function, of F, S and Bi doped TiO Various types of organic pollutants are formed during the production of acrylonitrile4 which has definitely induced serious impact on environmental and public health. Owing to its low bioavailability, high toxicity and mingled composition, acrylonitrile production wastewater has been directed as one type of refractory organic wastewater5. Therefore, it is necessary to develop a safe and efficient technology for the treatment of acrylonitrile wastewater. Various methods have been reported for the treatment of acrylonitrile wastewater, among those methods photocatalysis acquired much attention over the past decade. Since, photocatalytic reaction under sunlight irradiation is more energy-advantageous, and a lot of researchers have made vast efforts to realize the industrialization of photocatalytic treatment of industrial wastewater under sunlight8. However, there were few successful reports under sunlight because of the complexity of industrial wastewater11. Thus, photocatalytic treatment of industrial wastewater under sunlight was a great challenge for the researchers. Watanabe12 have reported that photocatalysis would cause a new environment revolution twenty years ago.Acrylonitrile is considered as a significant industrial chemical, originated by the direct oxidation of propylene with ammonia. It is extensively used for the preparation of synthetic rubber and resin, plastic and acrylic fiber13. Across various photocatalysts, titanium dioxide (TiO2) is known as the most determined material because of its chemical stability, high oxidation potential, nontoxicity and physical stability14. However, the use of TiO2 in photocatalysis are limited because of its certain drawbacks like: their large band gap15, which means that in solar energy processes, only UV light can be utilized and their low photocatalytic efficiency because of the fast recombination rate of electron–hole pairs. Therefore, many efforts have been promoted to reduce the bandgap of TiO2 by doping or by band gap engineering19.In our previous study, the high photocatalytic activity of F-doped TiO2 was attributed to the increase in the number and strength of surface acid sites20.It was explained that F-doping led to the creation of surface oxygen vacancies17, or the increase of Ti3+ state21. On the other hand, higher photocatalytic activity of S doped sample was attributed to the increase in the surface Bronsted and Lewis acid sites22. Samantaray indicated that sulphate radical impregnation decreases the crystallite size and stabilized the anatase phase of TiO223. In addition, Bi doped TiO2 exhibited a red shift in the optical adsorption and Bi3+δ+ species played a vital role in minimizing the electron hole recombination16. According to Li et al.24, Bi doping into TiO2 generates a new intermediate energy level below the conduction band edge of TiO2, extending the absorption in the visible region and enhanced their photocatalytic efficiency. On the other hand, SiO2 was used usually as a supporter, and its dispersing effect on nanoparticle size as well as that with oxidativity has not been reported.It is well known that, semiconductor-based photocatalysts have been investigated as an auspicious material for the solar energy conversion in regard to the breakdown of hazardous organic pollutants2 for the decomposition of organic pollutants in acrylonitrile wastewater under solar light irradiation, we modified TiO2 with the combination of silica and F, S, Bi doping (F–S–Bi–TiO2/SiO2).The dispersion of SiO2 produced ultrafine nanoparticles. We have found that silica dispersion changed the aggregation state, constituent distribution in amount and morphology of the nanocatalyst, which was responsible for light absorption and increased photocatalytic activity20. We first degraded the acrylonitrile simulated wastewater and then degraded the acrylonitrile wastewater. This photocatalyst exhibited excellent performances in both the photocatalytic decompositions of organic pollutants under simulated and natural sunlight. So, our approach is an important attempt for the photocatalytic treatment of industrial wastewater.To improve the photocatalytic activity of TiO2 F30 TEM was used to analyze the physical structural characteristics of the photocatalysts. The samples were ultrasonically dispersed in ethanol. The suspension was deposited on a Lacey-carbon film, which was supported on a copper grid. The particle size distributions of catalysts with or without SiO2 dispersant were calculated with NIH software using TEM image treatment. The crystalline phases of the photocatalysts were determined by X-ray diffractometer  at a scanning rate of 6° min−1 in the 2θ angle range of 10°–80° using Cu Kα radiation combined with nickel filter. The accelerating voltage and the applied current were 40 kV and 200 mA, respectively. Crystallite sizes were calculated according to Scherrer equation:mea and bins are the breadths of intrinsic diffraction profile, the test sample diffraction integral profile and instrumental diffraction profile, respectively. UV–visible spectra were measured on UV-2450 UV spectrophotometer . The range of the scanning wavelength was 200–800 nm. The BET specific surface area was measured by BELSOROP-MINI II (Japan) adsorption instrument and pore size distribution was analyzed by Barrett–Joyner–Halenda (BJH) method.The valence states on the surface of catalysts were analyzed by a Thermo ESCALAB 250XI X-ray photoelectron spectrometer (America) using Al Kα  as a radiation source. The irradiation of simulated sunlight and natural sunlight intensity was measured with a FZ-A RADIOMETER irradiance meter (China). From 6 am–8 pm, it was measured at a certain interval as shown in Table −1 to 290 °C (hold for 2 min). The electron energy was set at 70 eV and the ion source temperature was 230 °C. The standard spectra in GC–MS database was used to identify the chemical constituents in wastewater. In the present study, refractory acrylonitrile wastewater was obtained from an acrylonitrile manufacturing plant. The acrylonitrile wastewater was pretreated through adsorption with microporous zeolite, HZSM-5 before photocatalytic degradation, CODCr decreased from 582.4 to 88.36 mgL−1, and after that there were no change in the values. The organic pollutants in the wastewater after adsorption were detected and the results were listed in Table GC–MS  was used for the component’s analysis. The instrument was equipped with a DB-5 capillary column . The injector and MS transmission line temperatures were 250 and 310 °C, respectively. The oven temperature initiated at 40 °C (hold for 5 min), and then increased at 5 °C min20. The quartz glass reactor was sealed after 180 mL of acrylonitrile wastewater (10 mgL−1) and 300 mg of the catalyst was placed in it. The mixture was magnetically stirred in the dark until the adsorption equilibrium attained. Then, the spherical Xenon short arc lamp  was turned on for 12 min. In the process of illumination, the reaction solution of 1 mL was taken out at the interval of 2 min, and after that photocatalyst was filtered out through a filter film of 0.45 μm, acrylonitrile concentration was measured by HP-LC with LC-2030 UV detector and Sunfire TM C18 column. The wavelength of detector was 210 nm and the mobile phase volume ratio of methanol to water was 3:7. Triplicate samples from each batch were taken for the tests.The photocatalytic activities of the photocatalyst used for acrylonitrile simulated wastewater were evaluated in a photocatalytic reaction system−1 of sulfate radical and 22 mgL−1 of chloride ion consistent with Table The ion effects were examined by following procedure: appropriate amount of sodium sulfate and sodium chloride were added to 180 mL of solution to obtain 61.8 mgLCr values were detected with a Fast COD Detection Instrument (LH-5B-3B(V8)). TOC concentrations of the samples were measured via a Shimadzu TOC analyzer . Concentration of inorganic ions, BOD and CODCr in acrylonitrile raw wastewater, the wastewater after adsorption measured according to national standard methods were listed in Table The catalyst contents and reaction device were used as the same as that in simulated wastewater. After adsorption equilibrium, the spherical Xenon lamp was turned on for 14 h. At given intervals of illumination, 4 mL of reaction solution was taken out and was filtered out through a filter film of 0.45 μm . COD2/SiO2 catalyst after 4 min was 63%, which was much higher than TiO2–P25 and F–S–Bi–TiO2.Figure −1 sulfate radical) and sodium chloride (22 mgL−1 chloride ion) were added into acrylonitrile simulated wastewater and the degradation was carried out by the catalyst under the same condition mentioned above. These results revealed that free sulfate radical, chloride ions, or sodium ions inhibited the degradation progress. This was similar with our previous studies which illustrated that, certain inhibition effect by free sulfate radical on the degradation process22.In order to demonstrate the interference of inorganic ions, a certain amounts of sodium sulfate (61.8 mgL2 P-25 (Degussa), F–S–Bi–TiO2/SiO2 and F–S–Bi–TiO2 catalysts. The absorption of F–S–Bi–TiO2/SiO2 catalyst was increased in the UV–visible region at 300–600 nm as compared to TiO2 P-25. However, the UV–visible spectrum of F–S–Bi–TiO2 became flat below 335 nm, showing weakest absorbance. Although, F–S–Bi–TiO2/SiO2 exhibited lower absorption ability towards visible region (400–600 nm) than F–S–Bi–TiO2, showing the higher degradation ratio. The higher degradation ratio was attributed to the better dispersion of SiO2, because in F–S–Bi–TiO2/SiO2 photocatalyst TiO2 was well dispersed, reducing the agglomeration and enhancing the absorption in the UV region22. However, both the photocatalysts have the same weight but F–S–Bi–TiO2 contains more S elements and absorbed more visible light. The Photocatalytic reaction mainly depends upon the UV light.Figure 2/SiO2 and F–S–Bi–TiO2 calcinated at 450 °C were recorded as shown in Fig. 2 as a dispersant agent decreases the average crystal size, and increases the surface area of the photocatalyst.In order to know the crystal structure of the prepared catalyst, XRD patterns of F–S–Bi–TiO2 were shown in Fig. ca. 4 nm of particles were produced on the rough end faces, where borderline disappeared in the interior of the large piece crystal ) and Bi4Ti3O12 ) were exhibited in Fig. 2. The black parts in the figure represented the pores in SiO2. The other general crystal configurations look like bulk type of TiO2 and Bi4Ti3O12 ) crystal with borderline mark aggregated into a large crystal and also possessed identical crystal face  as calculated by XRD  was higher than that of F–S–Bi–TiO2 (45.9 m2/g). And, greater adsorption capacity will lead to greater degradation rate, which was consistent with the experimental results.Moreover, the F–S–Bi–TiO2/SiO2 and F–S–Bi–TiO2 were shown in Fig. 4+ state to Ti3+, and these Ti3+ state was related to oxygen vacancies28, the content of F changes in both types of the catalysts were investigated. In Fig. s spectrum of F–S–Bi–TiO2 shows only one peak centered at binding energy 684.4 eV (represented by red color), indicated only surface fluoride species were present in F–S–Bi–TiO2. However, in case of F–S–Bi–TiO2/SiO2, the F 1s XPS spectra shows two peaks centered at binding energy 688.5 and 686.2 eV, respectively (represented by black color). The peak at binding energy 688.5 eV was attributed to the doped F into the substituted sites of TiO2 lattice and produced mixed oxide structure of O–Ti–F30. The lower binding energy, centered at 686.2 eV was attributed to the surface fluoride species adsorbed on the surface of TiO230. This result shows that silica plays a very important role in stabilizing the dispersion of F ions. Next, the O 1s spectrum of F–S–Bi–TiO2/SiO2 catalyst was shown in Fig. 29, while the peak at binding energy 530.4 eV ascribed to lattice oxygen of TiO216. Next, the S 2p spectra of the catalyst was shown in Fig. p peak at binding energy 169.5 eV was observed, corresponding to S6+ state of SO42− species32. The binding energy at 164.6 eV was attributed to elemental sulfur33, which was overlapped with Bi 4f XPS spectra. Hence, the SO42− species were mainly adsorbed on the catalyst surface, which improved surface acid strength and favored to the adsorption and degradation of the pollutants34. The Bi 4f XPS spectrum was shown in Fig. f7/2, while the peak at binding energy 164.7 eV belonged to Bi 4f5/2. These values of binding energies were higher than the binding energy of Bi3+, indicating that bismuth existed in Bi3+δ state and formed Bi–O–Ti bond16. The intensities of Bi 4f peaks of F–S–Bi–TiO2/SiO2 were decreased as compared to F–S–Bi–TiO2, which indicate that dispersion function of silica was selective, that was only favorable for the substitution of F for the lattice oxygen but not for the formation Bi oxides. The weight percentages of the elements in the F–S–Bi–TiO2/SiO2 and F–S–Bi–TiO2 were shown in Table XPS spectra of F–S–Bi–TiOCr (Table Cr as a function of illumination time under simulated sunlight. It was found that CODCr value decreased from 88.36 to 7.20 mgL−1 and TOC concentration decreased from 39.45 to 2.57 mgL−1 for 14 h irradiation. F–S–Bi–TiO2/SiO2 catalyst was recycled through filtration after reaction and the recovery ratio was more than 95%. The photocatalytic reaction performance of recycled F–S–Bi–TiO2/SiO2 was demonstrated through repeated rounds under the same reaction conditions. The value of CODCr was 30.67 mgL−1 after 14 h of illumination at the second round. After photocatalysis, the value of CODCr was 45.86 mgL−1 at the fourth round.On the basis of above analysis Tables  and 2 weCr Table , the use2/SiO2 catalyst for the degradation of pollutants in industrial acrylonitrile wastewater under natural sunlight irradiation. The value of CODCr decreased to 39 mgL−1 after 6 h irradiation, which has satisfied China national discharge standards.Figure 2 under UV radiation. The degradation percentage of the dye and CODCr were 98.50% and 91.50%, respectively11. Dai et al.25 investigated the adsorption purification of acrylonitrile production wastewater by a microporous zeolite, CS-Z1 and a visible light-driven Ti–β–Bi2O3 photocatalyst. Nanoporous Ti–β–Bi2O3 was prepared via a solvothermal synthesis method in laboratory. Dai et al. was mainly interested in the theory, but not in application. They have not considered the inorganic ions effects. In this paper, we used industrial acrylonitrile wastewater as testing samples. The industrial acrylonitrile wastewater was pretreated by microporous zeolite, HZSM-5. After the treatment, the wastewater contained some inorganic and organic matter, as shown in Tables 2 dispersing sol–gel for the synthesis of highly active ultrafine nanoparticle catalyst with certain anti-ion interfering function and for the first time we successfully degraded industrial acrylonitrile wastewater in only 6 h under natural sunlight. This is the most significant findings for the photocatalysis application in environment.The actual textile dyeing wastewater was effectively oxidized by TiO.10 have demonstrated that photocatalytic purification of phenol containing wastewater by TiO2-P25 using sacrificial agent under ultraviolet part of sunlight. In this study, there was not any sacrificial agent, however, we still realized that the values of CODCr and TOC near zero discharge. These results show the potential of as prepared photocatalyst in near zero discharge and high TOC removal efficiency of wastewater treatment.For most of the pollutants, close to zero discharge is a final goal that people pursue in environment protection. Sixto et al2 was several times faster than without SiO2. The later exhibited some general properties of a large bulk aggregate. On the other hand, there were a majority of ultrafine nanoparticles with acidic sites on the catalyst with SiO220. A large number of bumps increased the quantum size effect and UV absorption at 270–380 nm, which should raise activity. Several nanometers package of irregular crystal films have promotion effect on light absorption. Photon were generated by light which can go through several nanometers of irregular Bi4Ti3O12 out layer film and get into TiO2 (Fig. - ion in the lattice, which could produce holes19 and also contribute the high activity of the catalyst with silica dispersion.In our previous study, we found that the reaction rate of the catalyst with SiOiO2 Fig. . The bum2 with SiO2 dispersing sol–gel particles and successively degraded the organic pollutants in acrylonitrile industrial wastewater to reach national discharge standard under 6 h of natural sunlight irradiation. In the prepared photocatalyst the doping of F, S and Bi causes the enhanced absorbance in the visible region. The results of photocatalytic activity evaluation demonstrated that CODCr value reached to discharge standard still after four recycle uses under the simulated sunlight irradiation, and to near zero discharge for the fresh photocatalyst. These results exhibited effectiveness and potential of our photocatalyst for the treatment of complicated and refractory industrial wastewater. The XPS and EDS analysis implied that S and Bi were doped successfully. The photocatalyst size distribution has been identified by visible nano-aggregates, constructed with 2–4 nm of finer gel particles. It indicated the function of SiO2 dispersing to form ultrafine nanoparticles. The nano-aggregates might form an identical lattice faces or different faces when gel particles crystallized. TEM results revealed that, the bump’s numbers may also be responsible for the increase in the light adsorption and photocatalytic activity. The identical face might originate from silk ribbon film of package. The crystals of several nanometers of out layer films and rough core surfaces of packages also increased the light absorption and enhance their photocatalytic activity.We built up a most simple and low cost approach for the preparation of ultrafine nanocatalyst of F, S and Bi doped TiOSupplementary Information.


The use of lutetium led to an improved SMM performance of DyLu2N@C80, which shows a higher blocking temperature of magnetization (TB=9.5 K), longer relaxation times, and broader hysteresis than DySc2N@C80 (TB=6.9 K). At the same time, Dy2LuN@C80 was found to have a similar blocking temperature of magnetization to Dy2ScN@C80 (TB=8 K), but substantially different interactions between the magnetic moments of the dysprosium ions in the Dy2MN clusters. Surprisingly, although the intramolecular dipolar interactions in Dy2LuN@C80 and Dy2ScN@C80 are of similar strength, the exchange interactions in Dy2LuN@C80 are close to zero. Analysis of the low‐frequency molecular and lattice vibrations showed strong mixing of the lattice modes and endohedral cluster librations in k‐space. This mixing simplifies the spin–lattice relaxation by conserving the momentum during the spin flip and helping to distribute the moment and energy further into the lattice.The substitution of scandium in fullerene single‐molecule magnets (SMMs) DySc What a mix! The diamagnetic metal in DyM2N@C80 and Dy2MN@C80 (M=Lu or Sc) has a considerable influence on the single‐molecule magnetism of clusterfullerenes. DyLu2N@C80 has a higher blocking temperature of magnetization than DySc2N@C80, whereas exchange coupling in Dy2LuN@C80 is much weaker than in Dy2ScN@C80. The different masses of lutetium and scandium also alter the mixing of local cluster vibrations and lattice modes, and hence have a strong impact on the spin–lattice relaxation (see figure). Enclosing lanthanide ions within the fullerene cage is a versatile route to a variety of molecular magnets.3+)3N3−@Cn26−, and in particular the species with Cn2=C80‐hI (hI denotes the symmetry of the carbon cage) and M=Sc, Y, or heavy lanthanides (Gd–Lu), show the largest synthetic yield and are therefore the most studied clusterfullerenes to date.3− in its center and three M3+ ions located at the vertices of the triangle with rather short M−N bonds, offers a convenient platform to create a variety of molecular magnets with divergent properties. First of all, the nitride ion at a distance of only 2.0–2.2 Å from the lanthanideJm is a good quantum number, at least for several lowest‐energy LF states. In particular, the lowest‐energy Kramers doublets for Dy3+ in dysprosium‐scandium NCFs are essentially pure Jm states, with the high‐spin zJ=±15/2 states as the ground Kramers doublet with the magnetic moment oriented along the Dy−N bond.zJ states in terms of Jm composition ensures that the magnetic system is weakly susceptible to external perturbations such as those introduced by dipolar magnetic fields from neighboring molecules or molecular and lattice vibrations.Nitride clusterfullerenes (NCFs) with the composition (MxScx3−N@C80‐hI (x=1–3) as an example.3+ is diamagnetic, this series essentially allows analysis of how two or three dysprosium ions interact magnetically and the effect of these interactions on SMM behavior. DySc2N@C80‐hI was the first endohedral metallofullerene (EMF) proven to be a SMM.2ScN@C80‐hI shows magnetic hysteresis below 8 K without fast QTM relaxation in zero field, which is explained by ferromagnetic exchange and dipolar coupling between the non‐collinear magnetic moments of the two dysprosium ions in the Dy2ScN cluster  series discussed above are determined by the Dy3+ ions, with the diamagnetic Sc3+ ions acting just as placeholders, which are needed to keep the trimetallic cluster composition. However, in this work, we shed more light on the role that the diamagnetic metal can play in the SMM properties of NCFs. For this, we chose lutetium as a diamagnetic lanthanide, the ionic radius (R3+=0.86 Å) of which is noticeably larger than that of scandium (R3+=0.75 Å). Yet, as we reported recently, DyLu2N@C80‐hI and Dy2LuN@C80‐hI can still be synthesized, albeit in lower yields than the dysprosium‐scandium analogues.3+ leads to changes in the internal structure of the trimetallic nitride cluster, and in particular results in shorter Dy−N bonds than in the dysprosium‐scandium NCFs. The shortening of the Dy−N bonds may affect the magnetic anisotropy and the strength of the intramolecular Dy⋅⋅⋅Dy coupling. Furthermore, lutetium is much heavier than scandium, which changes the low‐frequency part of the vibrational spectrum. Because the relaxation of magnetization involves spin–phonon interactions, alterations of molecular vibrations may also result in changes in the SMM properties. Thus, the goal of this work was to study a possible influence of these factors on the SMM properties of the DyM2N@C80‐hI and Dy2MN@C80‐hI NCFs.It is natural to consider that the magnetic properties of the NCFs in the Dy80‐hI fullerene cage  were obtained by arc‐discharge synthesis and separated by recycling HPLC as reported earlier.3N@C80, Dy2LuN@C80, DyLu2N@C80, and Lu3N@C80 substantially complicated the separation of the individual compounds. However, a compositional purity exceeding 90–95 % could be achieved for Dy2LuN@C80, as verified by MS analysis. Due to the very similar retention behavior, the separation of DyLu2N@C80 from Lu3N@C80 was not possible, and the relative content of the two NCFs in the studied sample was around 1:1.35. Because Lu3N@C80 is diamagnetic, its presence in the sample does not lead to strong changes in the magnetic properties of DyLu2N@C80. A possible influence of dilution on the quantum tunneling of magnetization will be specifically considered below.The mixed‐metal dysprosium‐lutetium NCFs with a C2N@C80 and Dy2LuN@C80, measured by SQUID (superconducting quantum interference device) magnetometry, are shown in Figures 2N@C80 exhibits hysteresis up to 9 K at a sweep rate of 3 mT s−1  than for DySc2N@C80 (TB=6.9 K). Note that the magnetic dilution does not affect the TB value,2N@C80 and DySc2N@C80 cannot be caused by the presence of Lu3N@C80 in the former. To conclude, DyLu2N@C80 was found to be a stronger SMM than DySc2N@C80. The substitution of scandium by lutetium in DyM2N@C80 leads to slower QTM relaxation, slower in‐field relaxation, and a broader magnetic hysteresis with remanence.The magnetization curves of DyLu2MN@C80, the influence of the diamagnetic metal on the SMM behavior appears to be weaker than for DyM2N@C80. Similarly to Dy2ScN@C80, Dy2LuN@C80 exhibits open hysteresis without a pronounced QTM step . Thus, the substitution of scandium by lutetium narrows the magnetic hysteresis but does not change the temperature scale of the slow relaxation. Nevertheless, as we show below, the temperature dependence of relaxation times reveals that the mechanisms of the relaxation of magnetization in Dy2LuN@C80 and Dy2ScN@C80 are different.For Dyp Figure  a. Its br Figure  b and thThe opening of magnetic hysteresis in the magnetization curves of molecular magnets indicates that the magnetization attains its equilibrium value slower than the rate of the magnetic field sweep. The slow relaxation of magnetization is a key characteristic of SMMs and needs to be understood in detail. The relaxation of magnetization requires energy exchange between the spin system and a thermal bath, which is mediated by the phonon system. The spin–phonon interaction is crucial for such an energy exchange, and the role of vibrational degrees of freedom becomes paramount. Several mechanisms of phonon‐mediated spin relaxation have been recognized in studies of paramagnetic salts and adapted for SMMs.A direct mechanism implies a single‐phonon process in which phonon frequency matches the energy difference between two opposite spins. Because a magnetic field increases the energy gap between the opposite spins, and the phonon density at near‐zero frequency is low and increases with frequency, the relaxation due to a direct mechanism accelerates in a magnetic field according to Equation τM is the magnetization relaxation time, H is the magnetic field, with two terms describing the relaxation of a Kramers ion in the absence (ca. H4) and in the presence (ca. H2) of hyperfine interactions, and A1 and A2 are fitting parameters. Thus, the relaxation rate scales linearly with temperature, τM,dir-1∼T. However, when at low temperature the phonon density may be very low, the energy exchange between the phonon system and the bath can become the limiting step. This effect, known as a phonon bottleneck, can change the temperature dependence to τM,dir-1∼T2.in which The Raman mechanism implies a spin flip through the absorption and emission of two phonons, with the frequency difference being equal to the energy gap between the opposite spins. As the frequencies can be much higher than the Zeeman energy, and the phonon density increases with frequency, the Raman mechanism is more efficient than the direct mechanism once the temperature is sufficiently high to ensure sufficient phonon population. The original consideration with only acoustic phonons in the Debye model gave the power‐law temperature dependence expressed by Equation n=9 for Kramers ions and n=7 for non‐Kramers ions, and C is the fitting parameter.in which A special case of the Raman mechanism, in which the absorbed phonon energy corresponds to the real excited spin state, is known as the Orbach mechanism. The temperature dependence of the relaxation rate under the Orbach mechanism has an Arrhenius form, expressed by Equation Ueff is the effective barrier (corresponding to the energy of the excited magnetic state) and τ0 is the attempt time.in which Finally, the QTM is a ubiquitous and characteristic relaxation mechanism of SMMs. In QTM, the spin flips to the opposite direction without energy transfer. As such, it should not show temperature dependence. The key condition for QTM is the energy matching of the opposite spin levels, and therefore the application of a magnetic field can quench the QTM when Zeeman splitting becomes large enough.Ueff corresponding to the frequency of the mode. Furthermore, it was shown that the direct mechanism can also demonstrate Arrhenius behavior in some conditions.It has become common practice to analyze the relaxation of magnetization in SMMs as a combination of these processes.2N@C80 and Dy2LuN@C80 in this work were determined at different temperatures by magnetizing the sample to saturation, quickly sweeping the magnetic field to zero or any other required field value, and then following the decay of magnetization with subsequent fitting of the measured decay curves with a stretched exponential . Below, the values obtained here are compared with those of DySc2N@C80 from ref. 2ScN@C80 were re‐measured in this work for better consistency.The relaxation times of magnetization in DyLu2N@C80 are complicated by the relatively long stabilization of the field and fast zero‐field QTM. In addition, the decay curves showed two types of behavior: A fast drop of magnetization for around 90 % of the sample, followed by a much slower relaxation of the remaining magnetization. The fit of the decay curves with two stretched exponents gave values ranging from 54±1 s at 1.8 K to 21±6 s at 5 K for the fast process  gave a Ueff value of 24.2±0.7 K and τ0 of 2.8±0.5 s. The analogous fit for DySc2N@C80 gave Ueff=23.6±1 K and τ0=0.6±0.2 s.Ueff values and differ only in their attempt times. The reason for the low‐temperature Ueff barriers of 24 K in both DyM2N@C80 molecules is not clear. Due to the very strong LF splitting of dysprosium in the NCFs, the energies of the lowest‐energy excited states exceed hundreds of K .τ0 values are many orders of magnitude longer than are usually found for the Orbach mechanism. We tentatively propose that the relaxation of magnetization in SMM EMFs in this temperature range may follow the Raman mechanism with involvement of local vibrations that would also follow Equation (3), but with Ueff corresponding to the vibrational frequency.Short relaxation times cannot be measured very reliably by DC magnetometry, and further relaxation measurements were performed for DyLu2LuN@C80 measured in zero magnetic field has a curved shape in the log (τM) versus T−1 plot  gives C=(1.14±0.28)×10−6 s−1 K−5.45, n1=5.45±0.15, τ0=435±50 s, and Ueff=4.3±0.2 K  with C=(0.51±0.26)×10−6 s−1 K−5.99, n1=5.99±0.33, τ0=56±4 s, and Ueff=8.0±0.1 K ]. To describe the temperature dependence of the relaxation times measured in a field of 0.2 T, we used Equation (4) with the addition of a term describing the direct process, given by Equation Measurements of magnetization relaxation times in Dys Figure  c reveals Figure  b. Such C, n1, τ0, and Ueff were fixed to the values determined for zero‐field relaxation.in which A(H)=(7.59±0.57)×10−5 s−1 K−1.67>and n2=1.67±0.09 (Table n=1) and bottleneck (n=2) direct processes, and indicates that both are likely to take place. If instead of using the fitting procedure, the n2 value is fixed to 1 and A(H) is determined from the field dependence at 2.5 K by using Equation (1)  gives ) Figure  c, then 2LuN@C80, the relaxation times of Dy2ScN@C80 at 2.5 K do not depend on the external magnetic field until it exceeds 0.4 T  are shorter than those in Dy2ScN@C80 (2.105 and 2.108 Å). These geometrical changes may substantially affect the LF acting on the dysprosium ions.The central nitride ion is the main source of the magnetic anisotropy in lanthanide NCFs, and the LF is expected to become stronger with decreasing distance between N2N@C80 and Dy2LuN@C80 molecules and compared the results with those for the dysprosium‐scandium analogues. For Dy2MN@C80 molecules, only one dysprosium ion was treated ab initio at a time, and the other dysprosium was replaced by yttrium. The energies of the Kramers doublets (KDs) in DyM2N@C80 molecules, the pseudo‐spin g‐tensor of the ground state KD, and the transition probability in the first KD are presented in Figure 3+ ions in DyLu2N@C80 and Dy2LuN@C80 exhibit high magnetic anisotropy with overall LF splitting of 1340–1360 cm−1. The quantization axis is aligned along the Dy−N bond, and the ground Kramers doublet is described as an essentially pure state with |Jm|=15/2. The energy of the second Kramers doublet is predicted to be close to 400 cm−1, which ensures that the low‐temperature magnetic properties of both DyLu2N@C80 and Dy2LuN@C80 are determined solely by the ground state of Dy3+, as in the previously studied dysprosium‐scandium NCFs.2N@C80 and DyLu2N@C80 shows that the increase in metal size from scandium to lutetium leads to an increase in LF splitting from 1284 cm−1 (DySc2N) to 1348 cm−1 (DyLu2N). The energy of the second KD (relative to the first KD) also shows an increase from 356 cm−1 in DySc2N@C80 to 391 cm−1 in DyLu2N@C80. As the Dy−N bonds in Dy2LuN@C80 are somewhat shorter than that in DyLu2N@C80, the LF splitting for the dysprosium ion in the former is also slightly higher. Likewise, the LF splitting in Dy2LuN@C80 is somewhat higher than in Dy2ScN@C80, which also correlates with the shorter Dy−N bond lengths (see Tables S8 and S9).To explore the influence of this effect, we performed ab initio CASSCF/RASSI calculations on the LF splitting in DyLuJ,mJ⟩basis are very important  and antiferromagnetic (AF) coupling of the dysprosium ions (J=15/2) can be calculated from Equation To determine the −1 (3.1–4.6 K) for Dy2LuN@C80. This energy difference is close to the Ueff value of 4.3 K determined for the low‐temperature relaxation process with Arrhenius behavior (see above). The assumption that the relaxation proceeds by excitation to the antiferromagnetically coupled state j1,2=0.028±0.001 cm−1. This value also gives a reasonable agreement between the experimental and simulated χT curves .which gives an estimation of 2.1–3.2 cmΔEFM-AFdip, can be calculated by using the well‐known formula for the energy of dipolar interactions between two magnetic moments [Eq. Magnetic Dy⋅⋅⋅Dy interactions have two components, dipolar and exchange. The energy difference between the dipolar interactions in the FM and AF states, ents Eq. ,(8)E1,2dnr→is the normal of the radius vector connecting the two magnetic moments μ1→and μ2→, R12 is the distance between them, and μ0 is the vacuum permeability. For Dy2LuN@C80 with DFT‐optimized coordinates and μ1,2→=10 μB, Equation (8) gives ΔEFM-AFdip=4.8 K. Surprisingly, it appears that the Dy⋅⋅⋅Dy interactions in Dy2LuN@C80 are solely of dipolar nature, with the exchange term vanishing almost completely. For Dy2ScN@C80, Equation (8) gives a very similar value of ΔEFM-AFdip=4.7 K, but as the ΔEFM-AFenergy in Dy2ScN@C80 is twice as large as in Dy2LuN@C80, one can conclude that the exchange term for Dy2ScN@C80 is of a similar size to the dipolar term . For instance, although the relaxation times measured in this work can be well fitted by Equations (1)–(3) and their combinations, the fitted parameters do not provide sufficient insight into the relation between molecular structure and the relaxation of magnetization.Ueff corresponding to half of the vibrational frequency. Furthermore, other research teams analyzed the locality of the spin–phonon interaction and concluded that molecular vibrations spatially localized close to the metal center usually have the strongest contribution to the spin relaxation.2N@C80 and Dy2MN@C80  are compared in Figure Recently, Lunghi et al.−1, whereas the frequencies of metal‐involving modes rarely exceed 230 cm−1, because the metal atoms are much heavier than carbon atoms. Only in the border range of 220–260 cm−1 do the squashing cage modes partially mix with the “breathing” mode of the nitride cluster, in which all three metal atoms move radially in one phase along the M−N bonds. Other metal‐based vibrations occur at frequencies below 200 cm−1. When the nitride cluster is encapsulated inside the fullerene, its external degrees of freedom  are transformed into internal ones . The frustrated rotations  are the lowest‐frequency intramolecular modes predicted to be close to 30–40 cm−1 in dysprosium‐lutetium NCFs and at 40–60 cm−1 in dysprosium‐scandium NCFs. The frustrated translations are mixed with deformations of the cluster (such as in‐plane oscillations of the M−N−M angles). In DyLu2N@C80 and Dy2LuN@C80, these modes are clustered into two groups close to 80 and 160 cm−1 . In dysprosium‐scandium NCFs, such modes are more uniformly spread in the 80–200 cm−1 range. Finally, the nitrogen out‐of‐plane mode also falls in the range of 90–210 cm−1.EMFs have rather peculiar vibrational spectra as their molecules consist of two semi‐independent units, the vibrations of which show almost no overlap in the energy scale, as can be well seen in Figure −1, the spectra of all the NCFs are quasi‐continuous because of densely spaced cage vibrations. The calculations seem to overestimate the cage frequencies by around 10 %. In the cluster frequency range, DyLu2N@C80 and Dy2LuN@C80 exhibit only two Raman peaks, at around 80 and 162 cm−1, close to the predicted frequencies of the mixed translation/deformation modes. The dysprosium‐scandium NCFs exhibit richer spectral patterns, and most of the observed peaks can be reliably assigned to the computed modes, as listed in Table 2N@C80 NCFs). Due to technical limitations, we cannot record the spectra below 50 cm−1, which precludes experimental observation of the cluster librations in dysprosium‐lutetium NCFs. However, good agreement between experiment and theory for the cluster modes above 50 cm−1 ensures that the calculated frequencies are not far from reality. In addition, in some of the earlier Raman studies of the NCFs, peaks at around 30–40 cm−1 were reported for Dy3N@C80, Lu3N@C80, and some other M3N@C80 molecules.The results of the computations agree reasonably well with the experimental Raman spectra. Above 220 cm3N@C80 molecules are expected to mediate energy transfer between the spin and a thermal bath. The rotational motions of the cluster are especially of interest from the point of view of low‐temperature spin relaxation because they not only happen at low frequencies , but also because they may help to conserve the total angular momentum when the spin flips. The relevance of the Einstein–de Haas effect on the single‐molecule level was demonstrated by Wernsdorfer and co‐workers for TbPc2  grafted on a carbon nanotube (CNT).2 molecule around the terbium quantization axis. However, because the molecule was rigidly bonded to the CNT, to fulfill the rotational invariance, the rotational momentum had to be transferred to the momentum of the phonon propagating along the nanotube. As a result, the efficient spin reversal proceeded by the direct mechanism when the external magnetic field created a Zeeman splitting matching the frequency of the longitudinal nanotube phonon. In a dysprosium nitride cluster, the reversal of the spin aligned along the Dy−N bond should induce rotation of the cluster around this bond. But as discussed above, interaction with the fullerene cage restricts the rotational motion of the cluster and turns it into a vibration, albeit retaining its rotational character. In DyLu2N@C80 and DySc2N@C80, the frequencies of the corresponding cluster librations are predicted at 27 and 62 cm−1, respectively  approach.3N@C80 showed a reasonable prediction of the molecular geometry and vibrations.3N@C80 model to determine the Hessian and then computed the dispersion spectra and vibrational eigenvectors for different lanthanide‐containing NCFs by using a proper mass correction in a dynamic matrix.In the absence of experimental information, we performed computational modeling of the lattice phonons of M3N@C80 molecules was simulated with face‐centered cubic (FCC) packing with the optimized unit cell parameter a of 15.5 Å and the distance between the centers of fullerene molecules of 11 Å . The vibrational spectra computed for isolated DySc2N@C80 and DyLu2N@C80 molecules are compared in Figure −1, which is just on the borderline for the acoustic bands of the crystals. The molecular modes away from the frequency range of the acoustic modes show noticeable Davydov splitting but negligible k‐dispersion. But the local modes close to the acoustic bands intertwine with the dispersed lattice bands giving rise to a dense phonon structure starting from zero frequency on. Three clear acoustic modes are perturbed by a set of what appear to be local modes with fluid character across k‐space. To follow the possible mode mixing in k‐space, we chose Γ‐point vibrational eigenvectors as a basis space, in which eigenvectors computed at different k values were projected.The model lattice of Mk. This redistribution is also reflected in the projected DOS in Figure k‐space. However, the 3D model is challenging to grasp due to extensive mixing and a high density of states. For illustrative purposes, we simplified the model to one dimension and considered a linear chain of M3N@C80 molecules with a=10.75 Å .The results of this projection analysis for one pure acoustic mode and for one with cluster libration character are presented in Figure 3N@C80, DySc2N@C80, and DyLu2N@C80. The spectra contain only one prominent acoustic band with a large dispersion of 50 cm−1 and two bands with a smaller dispersion of 10 cm−1 produced by on‐site rotational degrees of freedom. In 1D Sc3N@C80, the lattice and cluster modes are high in frequency, the cluster libration band is flat, and a weak mixing with the acoustic mode can be detected only around the X symmetry point. With the increase in mass in DySc2N@C80, the cluster‐based frequencies decrease faster than those of the lattice phonons, and the degree of mixing increases as manifested in the considerable dispersion of the intramolecular mode propelled by the acoustic band. Furthermore, the local and acoustic mode frequencies are even closer in 1D DyLu2N@C80, and therefore the mode mixing is much more pronounced.Figure k values. Mixing of the cluster rotations with the lattice phonons thus facilitates spin reversal by interaction of the local and lattice modes in k‐space. Based on this conclusion, we tentatively suggest that the Arrhenius behavior with a Ueff of 24 K (17 cm−1) observed in the temperature dependence of magnetization relaxation times of DySc2N@C80 and DyLu2N@C80 . Modeling of the magnetization curves and the spin Hamiltonian solution was accomplished with the PHI program.Powder samples of fullerenes for magnetometry studies were prepared by drop‐casting from toluene or CSλex=514 nm (Ar+ laser) or λex=647 nm (Kr+ laser). The samples for Raman measurements were drop‐cast onto single‐crystal KBr disks.Raman spectra were recorded at 77 K with a T 64000 triple spectrometer (Jobin Yvon) using an excitation wavelength of 3N@C80 molecules were performed at the PBE‐D level with a plane‐wave basis set and the corresponding projector augmented‐wave potentials, treating 4f electrons as a part of the core as implemented in the VASP 5.0 package.−4 eV per atom. The Γ‐point Hessian matrix and then the vibration frequencies (or Γ‐point phonons) were determined by using density functional perturbation theory as implemented in VASP.DFT calculations on isolated MThe ab initio energies and wave functions of LF (ligand field) multiplets for the dysprosium‐lutetium NCF molecules were calculated at the CASSCF/SO‐RASSI level of theory with atomic natural orbital extended relativistic basis set (ANO‐RCC)−1 with 2 k‐point sampling in each periodic direction. In all calculations, the non‐charge‐self‐consistent model was employed with the Slater–Koster parameters developed for similar types of system.Phonon spectra modeling and analysis were performed with in‐house Python scripts based on ASE libraries and with the improved version of some functions.The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


Detailed high‐level quantum chemical calculations support the experimental findings obtained in the highly diluted gas phase using FT‐ICR mass spectrometry. The origin for the existence of these long‐lived excited electronic states and the resulting implications for the varying mechanisms of dihydrogen splitting are addressed.Presented here is that isolated, long‐lived electronic states of ReC As a rare example, the coexistence of the ground and long‐lived, low‐lying electronic states for the bare ReC+ ion is identified. These different electronic states cause distinctly different reactivities towards the thermal activation of dihydrogen. After several control experiments aimed at looking for alternative explanations, the results suggest 1) the coexistence of at least two long‐lived, electronic states of ReC+, and 2) that these two states differ significantly in their reactivity towards dihydrogen splitting. Mechanistic aspects of the H−H bond activation were elucidated by high‐level quantum chemical (QC) calculations.Herein we report an unusual observation in the thermal gas‐phase reaction of diatomic ReC+ was generated by supersonic expansion of helium into a rhenium/carbon plasma generated by laser ablation/ionization of a disk comprised of compressed rhenium/graphite powder  using a Nd:YAG laser, operating at 532 nm inside the external cluster‐source of a Fourier transform ion cyclotron resonance (FT‐ICR) mass spectrometer as described previously .5 times with argon. This procedure ensures thermalization of hot ions and is usually regarded to quench excited electronic states. After thermalization to room temperature in a first and mass‐selection in a second step the ReC+ ions (m/z=199) were reacted with dihydrogen  at a constant pressure and under single collision conditions ], resulting in the generation of ReCH+, signal B in Figure 2 [Eq. (2)]. Here the product ReCD+, signal C, is formed, Figure +/HD couple amounts to KIE(H/D)=0.43 , resulace [Eq. ], resulace [Eq. ], resulace [Eq. ], resul+, in the ion/molecule reactions with dihydrogen are illustrated in Figure +/H2 couple exhibits a marked deviation from the strict linearity of a pseudo‐first‐order reaction kinetics; rather, the intensity decline of the charged educt can be fitted very well by the sum of two exponential decay functions, Equation a1 and a2 represent the relative initial abundances of “different” ions at time t=0 and k1 and k2 their respective rate constants. A two‐parameter function does not fit the experimental data . Clearly, there exist two regimes in that at the beginning the faster reaction prevails while at longer reaction times the slower process dominates , where the fast reaction dominates, a value of 1.5 is obtained. At longer reaction times (12 s), at which the slower reaction matters, this ratio increases to 3.1.Furthermore, the relative abundancies of ReCDk(ReC+/H2) for the fast component has been determined as 1.7×10−10 cm3 molecule−1 s−1, corresponding to a calculated efficiencyφ=11.0 %. For the slower reaction it amounts to only 1.6×10−12 cm3 molecule−1 s−1 (φ=0.10 %). Owing to the uncertainty in the determination of the absolute H2 pressure, an error of ±40 % is associated with these measurements.The rate‐constant + towards dihydrogen poses at least three questions: First, why is it so difficult for the excited state(s) of ReC+ to be converted into the ground state? Second, do the different electronic states react the same way with dihydrogen or are there any noticeable mechanistic differences? Third, which of the electronic states of ReC+ is responsible for the slow and which for the fast process?The unusual reactivity pattern of ReC+/H2 system possesses significant multireference (MR) character. Therefore, we followed a protocol that had been successfully applied previously for the theoretical description of MC+/CH4 couples .To this end, and to obtain mechanistic insight, high‐level quantum chemical calculations were carried out. According to exploratory diagnostic tests, the ReC+ (see Table S1), together with the crossing points (CPs)3Σ−) for the ground state of ReC+. The positive charge is essentially localized on the metal atom (see Table S2). The lowest excited state corresponds to a quintet state, 5Σ−, and is by 77 kJ mol−1 higher in energy. The second lowest excited state of ReC+ is a singlet state, 1Γ, and 100 kJ mol−1 above the ground state. Additional electronic states of even higher energies are listed in Table S1.Figure −1 above the minimum of the triplet surface. Clearly, without substantial energy supply the intersystem crossing (ISC) probability must be extremely low. In addition, the radiative emission of photons from the minimum energy regime of the quintet surface to the ground state triplet is extremely inefficient for the following reasons: 1) the transition from quintet to triplet is spin‐forbidden, 2) the oscillator strength is negligible for this transition owing to the absence of orbital angular momentum for both the 3Σ− and 5Σ− states; a first‐order spin‐orbit correction is not necessary,−1 are required for the singlet minimum to pass over to the quintet surface via CP2, and under the given experimental conditions, this ISC is very likely to happen.Since the potential energy curves of the singlet and the triplet states run almost parallel to each other upon varying the Re−C bond length from 1.6 to 2.2 Å, both states cannot cross in this energy regime. A crossing point, CP1, between the triplet and the quintet states has been located 294 kJ mol1[ReC]+ is necessary to promote this transition . Because of the error bar of the QC method used and the way the ions are generated by laser evaporation/ionization, it can be assumed that electronically excited ReC+ ions initially are present in both their quintet and singlet states, however, in terms of quantity the ReC+ ions in the triplet ground state clearly dominate. Most importantly, while the singlet state could easily cross to the quintet, the latter state cannot be quenched further to the triplet ground state. Obviously, for ReC+ there is a coexistence of a long‐lived excited with the electronic ground state.A frontier orbital analysis indicates that a decoupling of the paired electrons of the σ(Re‐C) orbital in + when exposed to dihydrogen. Simplified two‐dimensional potential energy surfaces (PES's) of the most favorable pathways as well as symmetries and selected structural parameters of key species are shown in Figure Next, we consider computationally the reaction mechanisms of the three electronic states of ReC+ is considered. According to path A (black) in Figure 3[ReC]+ with H2 commences by the barrier‐free and exothermic formation of the encounter complex 3EC. This step corresponds to an electrophilic attack of the positively charged metal center (q=+1.08) of 3[ReC]+ to the incoming ligand H2. Subsequently, an H atom is transferred almost barrier‐free to the carbon atom via the transition‐state 3TS1 (−57 kJ mol−1) to fall down into the rather deep well of 3I1, which is located 266 kJ mol−1 below the reactants, and this species comprises the global minimum on the PES. The energy gained and which, given the absence of a heat bath, is stored in “hot” 3I1, is almost completely reused when the latter decomposes to 4[ReCH]+ by releasing a hydrogen atom. This step is rate‐determining in the almost thermoneutral cleavage of the H−H bond by 3[ReC]+.First the reactivity of the triplet ground state ReC3[ReC]+ reacts with HD, formation of ReCD+ is favored over that of ReCH+. If only the conservation of angular momentum were to play a role and no kinetic isotope effects are operative, an opposite ratio of 2.0 should result in favor of ReCH+. However, based on the calculations, given the differences in zero‐point vibrational energies (ZPVEs), 4[ReCD]+ is by 6 kJ mol−1 more stable than 4[ReCH]+. Thus, the product ratio 4[ReCD]+/4[ReCH]+ should be greater than 1  in Figure 2 with the carbon atom of 5[ReC]+ leads to the formation of the very flexible and loosely bound van der Waals encounter complex 5EC, which then decomposes barrier‐free by the liberation of a hydrogen atom via 5TS1. The computed PES in the 55TS1EC/ area is rather flat, and the obtained stationary points are highly approximate. Obviously, a true transition state (close or similar to 5TS1) should lie somewhat above the adjacent minimum, that is, close or similar to 5EC .The course of the reaction of the quintet state of ReC5TS1, the four atoms assemble in a collinear fashion. In general, such a geometric feature is typical for a classical hydrogen atom transfer.5TS1 involving the carbon and the two hydrogen atoms, and the bond should be featured by a doubly occupied σC‐H‐H‐ and a singly occupied σ*C‐H‐H‐orbital.C‐H‐H‐orbital . Rather, the orbital occupancy evolution diagram  reveals a reaction mechanism which features the electronic structure requirement of a proton‐coupled electron transfer (PECT).C‐H‐H‐orbital and the two singly occupied σRe‐C‐H‐H‐ and σ*Re‐C‐H‐H‐orbitals form a 4c/4e bond. We note that this feature of 5TS1 is rather distinct from the bonding situation discovered in the [M−O]+/CH4 and [M−O]+/H2O couples, in which PCET transition states are characterized structurally by a four‐membered ring and the four electrons originate from two occupied orbitals.In 5[ReC]+/H2, 5[ReC]+, once generated, activates dihydrogen at ambient temperature. When HD acts as a partner for 5[ReC]+ to yield the products 4[ReCH]+ and 4[ReCD]+, the connectivity [Re−C⋅⋅⋅H⋅⋅⋅D] in 5TS1 is lower in energy 1.4 kJ mol−1 than that of [Re−C⋅⋅⋅D⋅⋅⋅H], thus slightly favoring the generation of 4[ReCH]+.+/[ReCH]+ ratio will change with the temporal variation of the 3[ReC]+ versus 5[ReC]+ abundances. As mentioned above, this ratio will increase at longer reaction times where the reactivity is dominated by 3[ReC]+.In any case, as both the intermediate and the transition state on the quintet surface are located below the separated educts 1[ReC]+ with dihydrogen. The association of H2 with 1[ReC]+ to form the encounter complex 1EC is fairly exothermic by 180 kJ mol−1. In 1EC, the calculated Löwdin bond order1EC the H–H σ‐bond is already significantly weakened while the metal center establishes bonds to both hydrogen atoms. Next, 1EC could surmount the transition state 1TS1 to form the intermediate 1I1. 1TS1 geometrically resembles 3TS1, except that the former possesses a much shorter C−H and a longer H−H bond. Finally, the product ion 2[ReCH]+ could be generated by the release of an H• radical in going from 1I1 to 1P1. However, as these products are 13 kJ mol−1 higher in energy than the starting material, this path should not be observed under thermal conditions. Further, although we did not succeed to locate a minimal energy crossing point at the CASSCF level of theory, a conventional two‐state reactivity scenario with efficient transitions between different spin states is not likely to be operative.+ species on the educt side. Also for the singlet/triplet surfaces, shown in Figure 2 activation, that is, the separated reactions of the isolated triplet and quintet state of ReC+.Path C (red) in Figure + can be assigned to the faster and which to the slower reaction components  but also proceeds without encountering any significant, structurally demanding barriers along the reaction coordinate. Clearly, these features are beneficial for the direct H‐atom transfer. Therefore, we conclude that it is the quintet state 5[ReC]+ which is responsible for the faster component shown in Figure Finally, the question is addressed as to which of the two electronic states of ReCs Figure . While t+. The different electronic states exhibit distinctly different reactivities towards the thermal activation of dihydrogen. Obviously, these findings add another facet to the quite complex chemistry and unexpected mechanistic scenarios of small cluster ions.In conclusion, in this combined experimental/computational work we describe the rare, if not unprecedented case for the coexistence of an electronic ground and a long‐lived, isolated excited state in a diatomic ion as exemplified by bare ReCExperimental and Computational Detailsn‐electron valence state perturbation theory (NEVPT2).The ion/molecule reactions were performed within a Spectrospin CMS 47X Fourier transform ion cyclotron resonance (FT‐ICR) mass spectrometer; its operation is described in more detail in the experimental chapter of the Supporting Information.The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


Reactions initiated with Y2O3 hydrating to Y(OH)3, which then dehydrated to YO(OH). Higher temperatures and KOH concentrations led to faster, more complete dehydrations. However, 1 m KOH led to YO(OH) forming concurrently with Y(OH)3 before Y(OH)3 fully dehydrated but yielded a very low phase purity of hexagonal‐YMnO3. Using NaOH mineraliser, no YO(OH) was observed. Dehydration also initiated at a higher temperature in the absence of Mn2O3. The evolution of Rietveld refined scale factors was used to determine kinetic information and approximate activation energies for the reaction. The described hydrothermal synthesis offers a fast, low‐temperature method for producing anisometric h‐YMnO3 particles.The reaction mechanisms, phase development and kinetics of the hydrothermal synthesis of hexagonal‐YMnO Phase development: The hydrothermal synthesis of hexagonal YMnO3 from Y2O3 and Mn2O3 was investigated by in situ X‐ray diffraction. Using this technique, information on mechanism, kinetics and thermodynamics of the reaction has been extracted, as well as the impact that mineraliser type and concentration, and temperature have on these properties. Both methods have the drawback of requiring at least one high temperature crystallisation step above 789 °C, where h‐YMnO3 becomes the most thermodynamically stable phase. The sol‐gel method has been used to prepare nanoparticles with control over the crystal and grain size based on the crystallisation temperatures in the range of 800–1100 °C, giving particle sizes of the order of 100 nm. However, in addition to the high temperature required, another drawback is the amount of time and number of steps required; amorphous precursors require drying from solution at around 150 °C, calcination at ≈500 °C, and crystallization to form h‐YMnO3.3 have highlighted the effect of oxidising and reducing conditions on the outcome of the reaction; oxidising conditions were shown to lead to a higher amount of orthorhombic phase, due to the smaller MnIV B‐site cation causing the perovskite structure to become more geometrically favourable.Solid‐state and sol‐gel methods are well established to prepare highly pure h‐YMnO3 hydrothermally, especially the range of oxidation states that Mn is capable of adopting, since Mn is stable from +2 to +4, with +2 being the most stable state.3+ to disproportionate in solution adds a considerable challenge for the hydrothermal synthesis of YMnO3. Indeed, experimental evidence and theoretical calculations have shown that h‐YMnO3 is capable of accommodating significant oxygen interstitial defects, compensated by the presence of MnIV, up to an oxygen stoichiometry of YMnO3.14.2O5. o‐YMnO3 can also form as it has a Goldschmidt tolerance factor (t) of 0.80 and an octahedral factor (μ) of 0.51 (high spin), which is within the range in which oxide perovskites can form .3 by hydrothermal routes. One method was reported by Stampler et al.3  was synthesised from the oxides with KOH mineraliser . The use of a high‐temperature autoclave was required for the larger rare earth metals due to the higher stability of Ln(OH)3. In an alternative method, aqueous Y(NO3)3, a MnII salt, and KMnO4 were reacted with KOH,3.Hydrothermal synthesis, in contrast to other methods, typically requires only one step, is performed at lower temperatures, and can give phase pure products from reactions lasting only a matter of hours. Still, challenges remain in synthesising h‐YMnO2O3 is converted to Ln(OH)3 which then dehydrates to LnO(OH). However, the reaction progression at different temperatures and what happens with the Mn2O3 during the reaction remain uncertain.The work by Stampler et al.3,x1−NaxNbO3,3,x1−BaxNbO3.In situ X‐ray diffraction (XRD) during hydrothermal synthesis is a technique that has been used since the 1990s, with early papers having been published on zeolites by Norby et al.3 from Y2O3 and Mn2O3 at temperatures in the range of 300 to 350 °C monitored by synchrotron XRD. This study has allowed us to quantitatively measure the phases which appear over the course of the reaction, thereby revealing the reaction progression and kinetics, under different conditions. We have also studied the effect of hydrothermal conditions on Y2O3 to reveal the formation of the different hydroxides under increasing temperature. The results are discussed in relation to those obtained using ex situ techniques from reactions performed in a high‐temperature autoclave.Here we report in situ hydrothermal synthesis of h‐YMnO2O3 and Mn2O3 in 5 m KOH at 300, 320 and 350 °C, and in 5 m NaOH at 320 °C monitored by synchrotron XRD. Integrated XRD patterns and patterns calculated using Rietveld refinement of the final dataset of the reactions are included in Figures m KOH are shown in Figures S1 (a) and (b) in Supporting Information. The reactions show different progressions; the reactions at 300 and 320 °C progresses via an intermediate stage in which the Y2O3 converts mostly to Y(OH)3, and a small amount of YO(OH) which appears several seconds later before both are converted into h‐YMnO3 In addition, o‐YMnO3 and YMn2O5 are present, although some Y(OH)3 remains. At 350 °C, the Y(OH)3 is decomposed more easily and appears only briefly after the heat is applied before completely decomposing into YO(OH). The latter then reacts with Mn2O3 to form h‐YMnO3 and secondary phases, but also remaining in a significant quantity. This reaction progression (Y2O3 to Y(OH)3 to YO(OH)) is in line with previously published work.2O5, in which Mn has an average oxidation state of 3.5, is the major secondary phase at all three temperatures. YMn2O5 forms due to charge disproportionation of Mn, leading to the concurrent formation of Mn3O4, in which Mn has an average oxidation state of 223. YMn2O5 being the most stable phase in the MnOx‐YO1.5 phase diagram below 789 °C explains why this charge disproportionation occurs.2O5 and Mn3O4 with weight fractions of 8.5 and 2.8, respectively (as in the reaction at 320 °C in 5 m KOH) is 3.2. This average oxidation number may be accounted for by a significant error in refining weight fractions with such low intensity, or the presence of soluble MnII or amorphous species not included in the Rietveld refinements. Another challenge with the hydrothermal synthesis of h‐YMnO3 is the appearance of significant amounts of o‐YMnO3 (except in the case of 1 m KOH at 320 °C). The final weight fractions of all observed phases of reactions at 300, 320 and 350 °C are shown in Table S1, where it can be seen that the reaction at 320 °C had a slightly higher proportion of h‐YMnO3. Using 5 m NaOH mineraliser at 320 °C, no YO(OH) was observed  presented in Figure S2 show that the reaction performed in 1 m KOH yields very low purity, while higher KOH concentrations lead to faster reactions. Both 10 m KOH and 1 m KOH reactions led to complete dehydration of Y(OH)3, and the use of Mn3O4 in the reaction resulted in a lower weight fraction of YMn2O5, although Mn3O4 remains after the reaction.Figures 3 was fitted using the Stephens model for anisotropic peak broadening.3 formation are shown in Figure S3. A higher Stephens parameter indicates broader peaks from planes perpendicular to the indicated crystal axis for example, a large (4 0 0) parameter means broad (hk0) peaks.The h‐YMnO2O3 in 5 m KOH was followed under in situ hydrothermal conditions at temperatures in the range 100 to 350 °C. After stabilising at 100 °C, three phases were observed: unreacted Y2O3, Y(OH)3, and a third, unidentified phase. This third phase was not observed at 150 °C, where only Y(OH)3 and a small amount of Y2O3 appeared. In the range from 280 to 330 °C, Y(OH)3 was the only observed phase, and at 340 and 350 °C, YO(OH) appeared alongside Y(OH)3  of the scale factor over time. Therefore, a precursor dissolution limiting mechanism could be possible in this reaction due to the low mineraliser concentration.Using the scale factors extracted from the Rietveld refinement, information on the kinetics of h‐YMnOm KOH, respectively. At 300 °C and 5 m KOH 3, Mn3O4 and YMn2O5 as h‐YMnO3 continues to form. This change in mechanism is the likely cause of the reduction in rate. Additionally, the four datasets at 320 °C and different KOH concentrations (with two reactions at 5 m)  after washing the product with acid to remove Y(OH)3. Using a 10 % excess of Y2O3, the phase purity was improved to 98 %, with minor secondary phases of YMn2O5 and Mn3O4 being observed. The XRD patterns show a large degree of preferred orientation in the (00l) direction, which is clearly indentifed from a simple comparison of the measured diffraction pattern and a simulated pattern of h‐YMnO3 without any orientation . The degree of orientation has been quantified by the March–Dollase model in diffraction pattern refinements (Table S2).Reactions between Y3 formed both in situ and ex situ in the large autoclave showed similar large, flat hexagonal particles of the order of 10 μm across 3 and/or YO(OH), while Mn2O3 disappears as the product appears, as shown by the weight fraction area plots in Figures 3 appears concurrently with the reduction in scale factor of Mn2O3, Y(OH)3, and YO(OH). It is therefore likely that any soluble species would be extremely short lived, and that heterogeneous nucleation occurs on the surfaces of intermediate phases (Y(OH)3, YO(OH) or Mn2O3). It is also likely that a dissolution‐precipitation mechanism occurs rather than a shrinking‐core mechanism due to the fact that the YMnO3 forms in distinct hexagonal shapes which do not resemble the precursors. Stampler et al. suggested from ex situ studies that the reaction requires the breakdown of Ln(OH)3 into the more reactive LnO(OH) before proceeding to form h‐LnMnO3, while Mn2O3 is reactive at least as low as 150 °C. This was inferred from the fact that the required temperature for the reaction to proceed is heavily dependent on the stability of Ln(OH)3. The ionic radius dependence of Ln(OH)3 stability was reported by Klevtsov and Sheina.4− to form h‐YMnO3. Chouaib et al.2O3 has a solubility of ≈2 mm in 5 m KOH solution at 25 °C. This may appear to be corroborated by our results which show the formation of YO(OH) in quantities dependent on reaction temperature before the appearance of h‐YMnO3 3 and Mn2O3, and in the reaction with 5 m NaOH at 320 °C, we do not observe any YO(OH)  much faster; even at the low temperature of 300 °C in 10 m KOH, all of the Y(OH)3 eventually dehydrates ). Related to this, we have also observed that when Y2O3 is heated in 5 m KOH solution, no dehydration of Y(OH)3 occurs until 340 °C. Taken together, this implies that soluble Mn is involved catalysing the dehydration of Y(OH)3, which we observe to take place at 300 °C in the presence of Mn2O3. It is also observed that using 1 m KOH leads to the concurrent formation of YO(OH) and Y(OH)3, before the Y(OH)3 fully dehydrates ), which must result from a different mechanism than the observed fast dehydration at 10 m KOH.The data also show that the KOH concentration is an important factor for the reaction progression and rates. At higher KOH concentrations, the Y(OH)hk0) family and a narrowing of peaks in the (00l) family. This means that there is an increasing crystal size anisotropy over time. Comparing the absolute values of (4 0 0) with (0 0 4) parameters is challenging however as the scales are very different. Interestingly, there appears to be little difference in final values of the (4 0 0) parameters between h‐YMnO3 formed in KOH under different temperature and concentration, however, in NaOH, the (4 0 0) parameter appears to be much larger; 670 for 5 m KOH at 320 °C, and 1460 for 5 m NaOH at 320 °C for the final frames. Whereas for the (0 0 4) parameter, high temperature and mineraliser concentration yielded lower values .The analysis of evolution in Stephens parameters  shows a trend of increasing (4 0 0) parameter and decreasing (0 0 4) parameter for all reactions. This indicates a broadening of peaks in the 3 remained solubilised.By comparison, many other oxide ferroelectrics, such as BaTiO3, most notably the difficulty in synthesising phase pure h‐YMnO3. This is a result of the complexity of the system, with a total of eight crystalline phases having been detected throughout the course of the reaction. We suggest three possible optimization possibilities for reducing the formation of secondary phases during hydrothermal synthesis; 1) improved mixing of the precursor materials to reduce the diffusion distance between reactive species, thereby reducing the time to form secondary phases, and the use of nanoparticle precursors to increase solubility and reduce solid state diffusion distances, 2) adding MnII salt to the reaction and performing the reaction under yttrium rich conditions (then dissolving the yttrium hydroxide secondary phase in dilute acid) to reduce the propensity for YMn2O5 formation or to facilitate its reduction to h‐YMnO3 if it does form, and 3) using a lower pressure to reduce the o‐YMnO3 formation. Performing the reaction under supercritical conditions  may also provide a route to improving the phase purity, as Nørby et al. showed with yttrium aluminium garnet.There are still challenges to be addressed with respect to hydrothermal synthesis of h‐YMnO3 materials as they are able to be synthesised under similar conditions.3 is dependent on the Ln radius, with smaller radii yielding lower stability due to steric effects (Ln is 9‐coordinated in Ln(OH)3 and 7‐coordinated in LnO(OH)) and increased acidity of the smaller Ln ions,The synthesis procedure used here can be extended to other LnMnO3 particles in a range of sizes and shapes which is useful for, for example, catalytic applications, with all reactions reported here taking less than 1 hour. By contrast solid‐state methods for producing powder samples, are only able to produce small (≈100 nm) roughly spherical particles and require much higher temperatures. Thin films are typically required for electronic device application, and the hydrothermal method has been shown to be amenable to deposition on a substrate to make polycrystalline thin films,3.Still, hydrothermal synthesis offers a fast, low‐temperature method for producing h‐YMnO3 from Y2O3 and Mn2O3 as precursors with hydroxide mineraliser, probed through in situ XRD for the first time. The reaction is shown to go via an intermediate period where Y2O3 is rapidly converted to Y(OH)3, which is then dehydrated into YO(OH) before reacting with Mn2O3 to form h‐YMnO3. It is not clear whether dehydration of Y(OH)3 is necessary for the reaction to occur; in the reaction performed in NaOH no YO(OH) is observed which suggests that it is not necessary. From extracted kinetic data as a function of temperature, activation energies for the reaction were estimated to be 240±40 and 270±30 kJ mol−1 for the reaction in 5 and 10 m KOH, respectively, in line with the high‐temperature required for the reaction to proceed due to the high stability of Y(OH)3.In conclusion, we have studied the hydrothermal synthesis of h‐YMnO2O3 , Mn2O3 , KOH , NaOH  were used for the syntheses.Y6 reference. The diffracted beam was detected using a 2D Pilatus 2m detector with acquisition times of 2 or 5 s, and processed to 1D diffraction patterns using software available on the PILATUS@SNBL platform.In situ XRD experiments were conducted at BM01, Swiss Norwegian beamlines (SNBL) at the European Synchrotron Radiation Facility (ESRF). The temperature inside the capillary was calibrated using a hexagonal boron nitride standard prior to experiments. The wavelength of the monochromatic incident beam was 0.77445 Å, refined on a NIST‐660a LaB2O3 and Y2O3 in equal molar ratio into plastic vials, to which 2 or 5 mL KOH or NaOH solution at 1, 5 or 10 m was added, such that the concentration of Mn ions was 0.347 M. A mass ratio of 85.4:14.6 Mn2O3:Mn3O4 was used for the reaction in which Mn2O3 was partly substituted for Mn3O4, giving a 95:5 molar ratio of MnIII:MnII. The reaction of Y2O3 with KOH used 0.347 m of Y ions with 5 m KOH.Samples were prepared for the in situ capillary cellReaction mixtures were loaded into an open cell using a syringe, the cell was then sealed with a Swagelok cap screwed on at the end. For a reaction, the cell was pressurised using a Shimadzu LC‐10ADvp HPLC pump to a pressure of 200 bar. This was then heated using a pre‐set hot‐air blower which was moved into place with a motor. The in situ cell made use of a sapphire capillary which is resistant to strongly basic conditions such as those used in this and our previous studies. Stainless‐steel Swagelok parts and tubing were used to form connections between the HPLC pump and the capillary and to seal the reaction cell.3 was synthesised using a method similar to that described by Stampler et al.2O3 and Mn2O3  were added to the autoclave, to this, 30 mL of 5 m KOH was added, stirred for 30 min, and then heated to 320 °C. It took approximately 40 min to reach the set temperature, the pressure at the set temperature was 90 and 92 bar in the stoichiometric and non‐stoichiometric experiments, respectively. The total reaction time was 6 h (including time to reach the set temperature), after which the autoclave vessel slowly cooled to room temperature over a few hours. The product was collected by vacuum filtration, washed with ≈200 mL of deionised water and dried at 70 °C for 3 h.High‐temperature (>250 °C) ex situ reactions were conducted in a 100 mL Monel  autoclave. h‐YMnO6 standard, from which the peak profile was fitted with a Thompson—Cox–Hastings pseudo‐Voigt (TCHZ) peak type.3,3,3,2O5,3O4,2O3,2O3.All Rietveld refinements were performed using Topas 5.Fits for kinetic data were calculated using the SciPy module in Python.l direction, then in the 1 1 0 direction.Biso values were then refined, and finally atomic positions. In‐house diffractograms were measured using a Bruker D8 A25 DaVinci X‐ray Diffractometer with CuKα radiation, a LynxEye SuperSpeed detector, and a v6 variable incident slit.Refinements of the diffractograms from the ex situ experiments were modelled by adjusting the scale factor and lattice parameters. This was followed by use of the March–Dollase model for preferred orientation in the 00The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.