Rechargeable magnesium batteries were first presented about seven years ago. [1][2][3] Their components included magnesium metal or a Mg alloy anode, Mg x Mo 6 S 8 (0 < x < 2) Chevrel phase cathodes, and electrolyte solutions that contained an ether solvent and a complex electrolyte, a product of the reaction between a MgBu 2 Lewis base and an AlCl 2 Et Lewis acid (Bu = butyl, Et = ethyl). These systems, while demonstrating impressive cycleability, suffered from several drawbacks:i) The micrometric size Mg 0-2 Mo 6 S 8 Chevrel phase cathode suffers from some kinetic limitation and the phenomenon of partial charge trapping (of Mg ions) at low temperatures. [4,5] ii) The electrochemical window of the first generation of electrolyte solutions, THF/Mg(AlCl 2 BuEt) 2 was around 2.2 V, which limited the possible use of cathode materials with a higher redox potential (and higher capacity) than Chevrel phases. iii) For practical use, the synthesis of the components of rechargeable Mg batteries needs simplification. Chevrel phases (CPs), M x Mo 6 T 8 (M = metal, T = S, Se), are of great interest owing to their remarkable electromagnetic, thermoelectric, and catalytic properties [6][7][8][9] . Exceptionally fast cation transport for multi-valent ions (compared to any other inorganic host material) made these materials unique cathodes in Mg batteries. [1][2][3] However, the kinetics of Mg diffusion in the CPs is strongly affected by their composition and temperature. At ambient temperature, the selenide shows excellent Mg mobility in the full intercalation range from 0 to 2 Mg 2+ ions per formula unit, [4] while Mg trapping occurs in the sulfide. During the first magnesiation of Mo 6 S 8 , 2 Mg ions are inserted (i.e., the full theoretical capacity is realized), upon further electrochemical deintercalation of Mg x Mo 6 S 8 , part of the Mg 2+ ions (20-25 %) are trapped and can be removed from the crystal structure, only at elevated temperatures (i.e., only 75-80 % of the theoretical capacity is involved in reversible cycling at low temperatures).[5]Detailed studies [10,11] of the crystal structure of the Mg-containing CPs made it clear that the trapping in the sulfide is caused by a unique ring arrangement of closely located cation sites with low potential energy. The triclinic distortion in the selenide changes the geometry of the cation sites, resulting in the degeneracy of the effect. It can be suggested that the presence of relatively small amounts of Se will be sufficient to improve the kinetics of the Mg 2+ cations in CPs. In fact, in addition to compounds with a single anion, the Chevrel family includes also solid solutions where sulfur and selenium atoms form a common anion framework. [12] Thus, in order to optimize the cathode composition in Mg batteries, it is of great importance to study the influence of the S-Se substitution in the host on the electrochemical behavior. Mg insertion into the binary hosts occurs in two stages: [1][2][3]

Lithium iron arsenide phases with compositions close to LiFeAs exhibit superconductivity at temperatures at least as high as 16 K, demonstrating that superconducting [FeAs](-) anionic layers with the anti-PbO structure type occur in at least three different structure types and with a wide range of As-Fe-As bond angles.

We report Mn K-edge extended x-ray absorption fine structure spectra on La 0.75 Ca 0.25 MnO 3 up to high momentum transfer across the metal-insulator ͑M-I͒ transition. The data show compelling evidence for (i) large or intermediate Jahn-Teller polarons (IJTP), characterized by an anomalous longer Mn-O bond ͑DR 0.09 Å͒ in the metallic phase ͑T , 170 K͒, and (ii) appearance of small JT polarons (SJTP) at T. 170 K, characterized by a longer Mn-O bond ͑DR 0.21 Å͒, which coexist with the IJTP above the M-I transition and has equal probability in the temperature range of colossal magnetoresistance. [S0031-9007(98)06663-0]

to negative, 9 thereby implying that the low-temperature structure plays a key role in defining the electronic properties of these superconductors.The possible mechanism of superconductivity in the REFeAsO 1-x F x and related REFeAsO 1-δ materials is currently unknown. The rapidly developing structural and electronic phenomenology points to considerable similarities with the well-established behaviour of high-T c cuprate superconductors and early theoretical work has suggested that conventional electron-phonon coupling mechanisms are not able to account for the high T c , implying non-BCS origin of the pairing interactions. [10][11][12][13] The parent REFeAsO phases exhibit both a structural and a magnetic phase transition on cooling in a similar fashion to the parent cuprate phase, La 2 CuO 4 . 5,14 Upon doping with fluoride ions, again much like La 2-x Sr x CuO 4 , both the crystallographic and magnetic transitions are suppressed in the superconducting compositions, 6,7 while T c first increases smoothly before passing over a maximum value at an optimal level of doping. Detailed experimental mapping of the structural and electronic phase diagrams as the doping level varies is necessary before we achieve a fundamental understanding of the superconductivity mechanism.Here we probed the temperature evolution of the structural properties of the However, the structural behaviour of the SmFeAsO 1-x F x compositions is very different on cooling. No reflections violating tetragonal extinction rules are evident for the heavily-doped compositions with x = 0.15 and 0.20 ( Fig. 1e and 1f), in which both lattice constants, a and c decrease smoothly with their crystal structure remaining strictly tetragonal down to 20 K ( Fig. 2e and 2f). The rate of contraction, dlna/dT and dlnc/dT at ~5 and ~18 ppm K -1 for the a and c lattice constants, respectively is considerably anisotropic and leads to a gradual decrease of the (c/a) ratio with decreasing temperature. This behaviour is in sharp contrast to the observed thermal structural response of the SmFeAsO 1-x F x (x = 0, 0.05, 0.10, and 0.12) compositions. In these systems, the tetragonal structure is initially robust upon cooling showing a normal contraction of the lattice parameters and interatomic distances. However, as the samples are cooled further, all hkl (h, k ≠ 0) reflections in the diffraction profiles begin first to 4 broaden before splitting at a characteristic temperature, T s (Fig. 1a-1d Supplementary Table S1.The most prominent point arising from the results of the present structural refinements as a function of both temperature and composition is the survival of the orthorhombic crystal symmetry in SmFeAsO 1-x F x well beyond the onset of superconductivity. Crossing the metal-to-superconductor boundary at x ~ 0.07 is not accompanied by the complete suppression of the orthorhombic-to-tetragonal structural phase transition and, as for both x = 0.10 and 0.12 compositions studied here T s > T c , both superconducting phases are orthorhombically distorted (Fig. 3). A...

An exhaustive structural investigation of a Y-doped ceria (Ce1–x Y x O2–x/2) system over different length scales was performed by combining Rietveld and Pair Distribution Function analyses of X-ray and neutron powder diffraction data. For low doping amounts, which are the most interesting for application, the local structure of Y-doped ceria can be envisaged as a set of distorted CeO2- and Y2O3-like droplets. By considering interatomic distances on a larger scale, the above droplets average out into domains resembling the crystallographic structure of Y2O3. The increasing spread and amount of the domains with doping forces them to interact with each other, leading to the formation of antiphase boundaries. Single phase systems are observed at the average ensemble level.

It is widely known that the properties of a zeolite in catalytic or ion-exchange applications depend largely on the crystal structure of the zeolite. When a catalytic process takes place in a porous system with dimensions in the range 3-12 , the reaction pathway is strongly influenced by framework geometry and the steric constraints are fundamental for driving the reaction towards the desired products. [1,2] Even though a few extra-large pore zeolites (with channel delimited by more than twelve tetrahedra) have recently been reported, such as CIT-5, [3] UTD-1, [4] or ECR-34 [5] (characterized by apertures of 18-membered rings), there are several possible applications that involve molecules larger than the pore dimensions of the available zeolites. To overcome this limitation, mesoporous molecular sieves MCM-41, MCM-48, and MCM-50, with pore dimensions larger (about 30-100 ) than those of conventional zeolites have been developed. [6] The alluminosilicates that belong to the mesoporous family M41S have a periodic pore structure (i.e., giving rise to coherent X-ray diffraction), whereas the silica walls are disordered and resemble more the structure of a glass. Unlike conventional zeolites, these materials are not strongly acidic, [7] but they do show promise as supports in other types of catalysts, such as olefin polymerization. [8] Layered zeolitic materials represent another option for treating large molecules. These materials have an advantage in that they combine the good thermal stability of zeolites with active sites of zeolitic nature easily accessible to reactants. In fact, layered zeolitic materials can be pillared or delaminated to produce high-surface-area materials with a majority of their active sites exposed at the crystal surface. [9][10][11] Nevertheless, so far, only a few structures of synthetic layered silicates have been reported, mainly for two reasons: 1) solution of the crystal structure by powder diffraction is a very challenging task and the small crystal size typically do not allow single-crystal X-ray diffraction experiments, piperazine silicate EU 19 being the only excellent exception; [12] 2) as a consequence of the stacking disorder, which occurs between the layers, the powder-diffraction patterns of a layered material often suffer from severe peak broadening that precludes structure solution.Notwithstanding the above difficulties, there has recently been an increased activity in the structure elucidation of layered materials. [13][14][15][16][17][18] Among these materials, those composed by single zeolite sheets like PREFER are particularly interesting, [16] which after calcination leads to the ordered 3D net of the FER-type zeolite. A very similar behavior was reported for the borosilicate named ERB-1, [19] isostructural to MCM-22, the precursor of which is layered in 2D, and the 3D network is formed upon calcination through the condensation of the silanol groups located on the layer surface. Other examples of layered zeolite precursors are EU 19, [12] precursor of the structurally ...

We have applied a combination of spectroscopic and diffraction methods to study the adduct formed between squaric acid and bypridine, which has been postulated to exhibit proton transfer associated with a single-crystal to single-crystal phase transition at ca. 450 K. A combination of X-ray single-crystal and very-high flux powder neutron diffraction data confirmed that a proton does transfer from the acid to the base in the high-temperature form. Powder X-ray diffraction measurements demonstrated that the transition was reversible but that a significant kinetic energy barrier must be overcome to revert to the original structure. Computational modeling is consistent with these results. Modeling also revealed that, while the proton transfer event would be strongly discouraged in the gas phase, it occurs in the solid state due to the increase in charge state of the molecular ions and their arrangement inside the lattice. The color change is attributed to a narrowing of the squaric acid to bipyridine charge-transfer energy gap. Finally, evidence for the possible existence of two further phases at high pressure is also presented.

In this work the first Pair Distribution Function (PDF) study on Ce1‑x Gd x O2‑x/2 (CGO) electrolytes for solid oxide fuel cells is presented, aiming to unveil the complex positional disorder induced by gadolinium doping and oxygen vacancies formation in these materials. The whole range of Gd concentration x Gd (0 ≤ x Gd ≤ 1) of the CGO solid solutions was investigated through high resolution synchrotron radiation powder diffraction. The reciprocal space Rietveld analysis revealed in all the solid solutions the presence of positional disorder, which has been explicitly mapped into the real space. The average structural models, as obtained by the Rietveld method, fit well the experimental PDF data only for a spatial range r > ∼10 Å. The same models applied at lower r values fails to reproduce the experimental curves. A clear improvement of the fit quality in the 1.5 < r < ∼6 Å range was obtained for all the CGO samples applying a biphasic model encompassing both a fluorite CeO2-like and a C-type Gd2O3-like phases. This provides evidence that extended defects at local scale exist in the CGO system. Gd-rich and Ce-rich droplets coexist in the subnanometric range.