Texture is the orientation distribution of the single crystals forming the poly­crystalline aggregate. Due to the anisotropy of the single crystals, particularly the hexagonal alpha phase, texture is of great importance in the deformation behaviour of zirconium alloys. Therefore, understanding texture evolution during thermo­mechanical processing steps and service is a necessary precursor to predicting texture. Texture is affected by temperature [18, 26] and strain rate [27]. In situ diffraction is essential in understanding texture evolution.

Figure 4.9 shows an example of texture evolution in Zr-2.5Nb, with both tem­perature changes and deformation. This measurement was made on the HIPPO diffractometer [28, 29] at the Los Alamos Neutron Scattering Center (LANSCE) using a high-temperature deformation furnace [30]. The texture was measured at room temperature and at a number of steps up to 975 °C[ 31]. At this temperature the sample experienced a compressive strain of 20 %. Texture measurements were made at 975 °C, then at the same temperature steps during cooling to room temperature.

At room temperature the bcc beta phase is meta-stable and increases during heating at locations on the rim of the pole figure where the 0001 alpha pole-figure previously showed maxima, showing that alpha grains transform directly to beta grains [33]. At 975 °C only the bcc beta phase exists, at this point the sample had a strain of 20 %. The beta grains transform to alpha during cooling and (by the Burgers orientation relationship) the {0002} planes of the alpha phase become {110} planes of the beta phase, which is reflected in the texture evolution.

Pre and post treatment conditions of the sample are confirmed in Fig. 4.6a/b and e/f, respectively. The Burger orientation-relationship determines the transition of

(2) hcp/a to maxima in (110) bcc/p during a temperature increase to 975 °C. The resulting textures at high temperature are shown in (c). The (222) bcc/p planes align with the applied compression direction (d). To study the evolution of many properties such as texture, it is necessary to measure at intermediate steps, rather than just the start and end of processing.

Using Bragg-edge transmission [34] and neutron imaging in combination makes simultaneous mapping of the strains and texture (Fig. 4.10) [35] of the crystallites within the entire sample possible.

In Li-ion battery electrode research, neutron powder diffraction (NPD) is one of the first techniques used for structural characterization of synthesized materials, often in combination with X-ray diffraction. NPD also plays an essential role in under­standing the insertion mechanisms that may induce phase transitions and/or solid — solution behaviour, both of which may depend strongly on temperature, particle size, doping, and the chemical or electrochemical conditions.

The coherent neutron-scattering cross-sections of Li is relatively greater than that obtained for many transition-elements and greater than the analogous coherent scattering cross-sections of Li for X-ray or electron scattering, making it possible to determine Li-ion positions and occupancies. Thanks to these and other advantages of NPD, the technique has played a key role in the development of the large diversity of Li-ion battery electrode materials that exist to date.

Knowledge of the Li-induced phase transitions in C is primarily based on X-ray diffraction [52]. Although the crystal stages I and II in C are proven, long standing debate exists concerning LiC18 and no evidence of the higher-order lithiated stages exists based on X-ray diffraction studies. Neutron diffraction proved the existence of the LiC18 stage [53] and showed deviations from the common structural picture of lithiated C, including a charge-discharge dependent structural evolution and the appearance of higher-ordered stages [54].

Neutron diffraction has been used to determine the lithiated structures of tita­nium oxides including LixTiO2 anatase [55-57], rutile [58], brookite [59], and Li1+xTi2O4 and Li4+xTi5O12 spinel [55]. Li4Ti5O12 spinel is a state-of-the-art Li-ion battery electrode material [35, 60-63] operating at 1.56 V versus Li and is suitable as an anode using high-voltage cathode materials. Li4Ti5O12 can be lithiated up to the composition Li7Ti5O12 and possibly higher, with end members having a spinel structure adopting the cubic space group Fd-3m. This material is attractive as a Li-ion insertion electrode because the ‘zero strain’ property results in excellent cycle life: upon lithiation from the initial Li4Ti5O12 to the ‘fully-lithiated’ Li7Ti5O12 there is almost no change in lattice parameters (0.2 %) [35, 60-65]. In the defect spinel Li4Ti5O12 all the energetically-favourable tetrahedral 8a sites are occupied by Li. Additionally, 1/6 of the 16d sites are also randomly occupied by Li while the remaining 5/6 of the 16d sites are occupied by Ti atoms, and this can be represented as [Li3]8a[Li1Ti5]16d[O12]32e. Lithiation leads to occupation of all the octahedral 16c sites and emptying of the tetrahedral 8a sites to reach the lithiated Li7Ti5O12 composition, which can be represented as [Li6]16c[Li1Ti5]16d[O12]32e.

In lithiated anatase, neutron diffraction data obtained to relatively high momentum transfers were used to resolve the split Li-ion positions within the material’s distorted oxygen octahedra [56]. In composite anatase TiO2/Li4Ti5O12, neutron diffraction showed that the TiO2 phase was lithiated before Li4Ti5O12, as expected from the lower potential of Li4Ti5O12 versus Li/Li+, enabling tuned Li insertion/extraction based on the choice of voltage range [66]. The increase in curvature of the voltage profile and larger capacities for nano-sized materials appears to be a general obser­vation for the various TiO2 polymorphs such as the tetragonal anatase [42,45,48,58, 67] and rutile [48], orthorhombic brookite [68] and monoclinic TiO2(B) [50, 69]. The sensitivity of neutron diffraction for Li-ions has been decisive in revealing the altered thermodynamics of nano-sized titanium oxides. In anatase the Li solubility increases systematically when particle sizes are reduced leading to a phase-size diagram [45]. In addition, a second phase transition from the known Li-titanate phase towards tetragonal LiTiO2 was discovered, which was later confirmed by neutron diffraction in anatase nano-tubes [70]. Higher Li-ion solubility was also observed in nano­structured rutile using neutron diffraction [71], suggesting similar size effects. A remarkable finding in spinel nano-sized Li4Ti5O12 is that of increased capacity with decreasing particle size, exceeding the maximum composition observed for the micron-sized Li7Ti5O12. Neutron diffraction proved the increase in capacity to be due to simultaneous Li occupation of both 8a and 16c sites, providing an atomic-scale explanation for the larger capacity of the nano-sized materials [43].

A very interesting negative electrode, exceeding the graphitic anode capacity, at similarly low potentials, is the layered transition-metal oxide Li1+xV1-xO2 [72]. Interestingly, LiVO2 does not allow Li insertion whereas Li1+xV1-xO2 with x > 0 leads to a very low intercalation voltage close to 0.1 V with a capacity almost twice that of graphite. Neutron diffraction has shown that in Li1 .07V0. 93O2 part of the octahedral V is replaced by Li, which allows additional Li-ions, responsible for the large capacity, to occupy the neighbouring tetrahedral sites that are energetically unfavourable in LiVO2, and is supported by modelling studies [38].

Gummow et al. used neutron diffraction to show that the structure of the low — temperature cubic phase of LiCoO2 is not ideally layered, and that 6 % of the Co reside in the octahedral (8a) sites of the Li layers [73, 74]. The hexagonal structure of the high-temperature phase of LiCoO2 was also determined from neutron dif­fraction, illustrating that Co and Li planes alternate in the ABCABC oxygen stacking. Aiming at higher capacity cathodes, Li1+xCoO2 has been prepared raising the question of where the additional Li resides [75]. Combined Rietveld refinement using both X-ray and neutron diffraction data excluded both Co in the Li site and the presence of tetrahedral Li and Co [76]. Based on this, it was deduced that excess Li replaces some Co and that the charge is compensated for by O vacancies [77]. In mixed-cation layered transition-metal oxides, such as the so called ‘high capacity’ LiMn1/3Co1/3Ni1/3O2, neutron diffraction continues to be an indispensable tool for determining the cation distributions which have been shown to depend on the synthesis conditions [78—80].

Neutron diffraction has played a pivotal role in understanding the complex insertion and phase transitions in spinel transition-metal oxides. Using NPD, Fong et al. [81] described the crystal structure of LixMn2O4 for x = 1 and 0.2 and Wills et al. [82] determined the crystal structure of LiMn2O4 at low temperatures as well as its magnetic properties. Neutron diffraction revealed the partial charge-ordering in spinel LiMn2O4 at 290 K that further hinders its use as a positive-electrode material in Li-ion batteries [83]. From neutron diffraction data, superstructure reflections were found (related to charge-ordering phenomena) at 230 K which in combination with electron diffraction patterns revealed a 3a x 3a x a super cell of the cubic room-temperature spinel representing the columnar ordering of electrons and holes [83]. Two of the five Mn sites correspond to well-defined Mn4+ and the other three sites are close to Mn3+ as derived from Mn-O bond length analysis. This charge ordering is accompanied by simultaneous orbital ordering due to the Jahn — Teller effect in the Mn3+ ions. Li-excess compounds Li1+xMn2-xO4 were found to provide better cycling performance than the stoichiometric LiMn2O4 as they min­imize the extent of the Jahn-Teller distortion during cycling (i. e. increase the overall oxidation state of Mn during cycling). In addition, Li doping at octahedral 16c sites reduces the exothermicity of the Li insertion/extraction reactions by an amount similar to that associated with the dilution of the Mn3+ ion [84]. Neutron diffraction by Berg et al. [85] showed that Li occupies 16c sites in Li114Mn186O4 which is also accompanied by charge-compensating vacancies at Mn 16d sites. Calculations also showed that the 16d sites should be favourable for Li at low Li contents while at higher contents, the 16c and mixed 16c and 16d site occupation is likely [86]. However, a recent study by Reddy et al. [87, 88] shows that at lower Li doping regimes, x = 0.03 and 0.06, the structural model containing Li at 16c sites still results in a better fit to the neutron diffraction data than models with 16d site Li occupation. In the work by Yonemura et al. [89] samples were synthesized in controlled atmospheres which led to the realization and quantification of O-deficient LiMn2O4 and Li-excess (O-deficient) Lij+xMn2-xO4-y spinels. Using neutron dif­fraction data they determined the quantity of O, mixing of Li and Mn at the 8a and 16c sites, the interatomic bond distances, and the relationship between these crystallographic parameters.

In the high-voltage spinels, neutron diffraction enabled the transition-metal ordering in LiNi0.5Mn15O4 that breaks the cubic Fd-3m to cubic P4332 symmetry to be determined and this was found to be dependent on the cooling rates used in the synthesis [90]. In X-ray diffraction data the small difference in atomic number between Mn and Ni makes it hard to quantify this ordering, whereas it is easily modelled using neutron diffraction [91, 92]. In the low-voltage plateau, using the large difference in the coherent neutron-scattering cross section of Mn and Ni, researchers determined that extensive migration of Ni and Mn was occurring in the spinel structure due to the loss of long range Ni-Mn ordering [92].

In the polyanion-based positive-electrode materials, neutron diffraction data contributes significantly in the characterization and understanding of the electrode properties. The higher potential of the Mn3+/Mn2+ redox couple has initiated the synthesis of LiMnPO4 and LiMnyFe1-yPO4 materials. Neutron diffraction data helped reveal that the reduced activity of the Mn3+/Mn2+ couple is related to the distortion of the MO6 octahedra with M = Mn3+, and this distortion was found to be much larger than the change in the unit cell [93], effectively prohibiting the Mn3+ to Mn2+ transition. Neutron diffraction was also used to characterize the cation distribution in related olivine structures with other transition metals and transition — metal mixtures such as LiCoyFe1-yPO4 [94], LixCoPO4 [95], and V-substituted LiFePO4 [96, 97].

For tavorite, LiFePO4(OH), neutron diffraction showed both the Li and H to be located in two different tunnels running along the a and c-axes, the tunnels being formed by the framework of interconnected PO4 tetrahedra [98]. Another promising class of tavorite-structured cathode materials are the fluorophosphates which exhibit good storage capacity and electrochemical and thermal stability. LiFePO4F exhibits a complex single-phase regime followed by a two-phase plateau at 2.75 V. Neutron diffraction in combination with X-ray diffraction was used to resolve the single phase end-member Li2FePO4F structure showing that Li-ions occupy multiple sites in the tavorite lattice [99]. Additionally, in the pyrophosphate-based positive electrode Li2-xMP2O7 (M = Fe, Co), multiple Li sites were identified using neutron diffraction [100].

In general, the combination of X-ray and neutron diffraction has become the established approach to characterizing electrode materials in great detail. The fol­lowing example concerning the extensively-studied olivine LiFePO4 demonstrates the value of neutron diffraction in revealing the impact of dopants, defects, and particle size on LiFePO4 structure and performance, thereby providing crucial understanding for the design of future electrode materials.

In the last decade LiFePO4 has emerged as one of the most important positive electrodes for high-power applications owing to its non-toxicity and outstanding thermal and electrochemical stability [14]. The first-order phase transition, pre­serving its orthorhombic Pnma symmetry, results in highly-reversible cycling at the 3.4—3.5 V versus Li/Li+ voltage plateau with a theoretical capacity of 170 mAhg 1. The olivine structure is built of [PO4]3 tetrahedra with the divalent M ions occupying corner-shared octahedral ‘‘M2” sites, and the Li occupying the “M1” sites to form chains of edge-sharing octahedra. The magnetic structure of LiFePO4 has been solved using neutron diffraction, with the appearance of extra reflections below the Neel temperature indicating antiferromagnetic behaviour at low tem­peratures for both end-members FePO4 (Fe3+) and LiFePO4 (Fe2+) [101].

In contrast to the well-documented two-phase nature of this system at room temperature, Delacourt et al. [102, 103] gave the first experimental evidence of a solid solution LixFePO4 (0 < x < 1) at 450 °C, and in addition, the existence of two new metastable phases with compositions Li0.75FePO4 and Li0.5FePO4. These metastable phases pass through another metastable phase on cooling to room temperature where approximately 2 out of 3 Li-positions are occupied, again determined using neutron diffraction to be Li*067FePO4 [103]. In Li*067FePO4, the average Li-O bonds are longer than in LiFePO4 due to the shortening of Fe-O bond lengths as shown in Fig. 7.3. It was suggested that this bond-length variation is the origin of the metastability of the intermediate phase, and thus of the two — phase mechanism between LiFePO4 and FePO4. Interestingly, this metastable phase appears to play a vital role in the high charge/discharge rate of the olivine material [104].

The room-temperature miscibility gap in LixFePO4 was determined by Yamada et al. [105] using NPD. These researchers found intermediate Li-poor Lia=0 05FePO4 and Li-rich Li1—p=089 phases, as shown in Fig. 7.4. This explains the compositional range over which the voltage is constant (plateau) and proves the presence of mixed — valence states of iron (Fe2+/Fe3+). These mixed-valence states provide ionic and electronic conductivity, an essential ingredient for the material’s application as a Li-ion battery electrode.

An early report that led to intensive discussions suggested that the poor elec­tronic conductivity of LiFePO4 could be raised by 8 orders of magnitude by su­pervalent-cation doping, which was proposed to stabilize the minority Fe3+ hole carriers in the lattice [106]. It was only after detailed refinement of models against combined neutron and X-ray diffraction data that researchers were able to determine

Fig. 7.5 a LiFePO4 adopting Pnma symmetry with the split Li-ion (medium grey) position in the centre. b NPD data for LiFePO4 and Li0.96Zr0.04FePO4 (target composition) including the difference between the fits and data. The fit residuals are wRp = 1.7 % and Rp = 1.9 %, as well as wRp = 1.7 % and Rp = 1.8 %, respectively. c The same data as in (b) shown for a limited d-spacing range. Reprinted with permission from (M. Wagemaker, B. L. Ellis, D. Luetzenkirchen-Hecht, F. M. Mulder, L. F. Nazar, Chem. Mater. 20, 6313 (2008)) [107]. Copyright (2008) American Chemical Society

the positions and role of the dopants. In this case NPD provides contrast between Li and the dopants at the Li site (M1), and X-ray powder diffraction provides contrast between the Fe and many of the dopants at the Fe site (M2). Moreover, to determine three species (Li orFe, dopants and vacancies) on crystallographic sites (M1 or M2) requires more than X-ray or neutron diffraction alone. The neutron diffraction pattern for one of the doped materials is shown in Fig. 7.5. Although the changes in the diffraction pattern upon doping are extremely small, the accuracy afforded by the data make it possible to conclusively locate the supervalent-cation dopants in LiFePO4. Figure 7.6 shows that supervalent-cation doping of up to * 3 % atomic substitution can be achieved in the LiFePO4 lattice in bulk materials prepared by a solid-state route at 600 °C. The results show that the dopant resides primarily on the M1 (Li) site and that aliovalent-dopant charge is balanced by Li vacancies, with the total charge on the Fe site being +2.000 (± 0.006), within the limit of experimental error [107]. It is thus expected that dopants may have little influence on the elec­tronic conductivity of the material, which is confirmed by calculations [108]. Furthermore, the location of the immobile high-valent dopant within the Li chan­nels is expected to hinder Li-ion diffusion assuming one-dimensional diffusion.

The one-dimensional migration channels through the LiFePO4 olivine structure means that the electrode performance can be severely influenced by defects. In the olivine structure the most favourable defect is predicted to be the Li-Fe anti-site pair, in which a Li-ion (at the M1 site) and a Fe ion (at the M2 site) are interchanged. These defects, with concentrations up to 8 %, were first observed in LiFePO4 syn­thesized at low temperatures, leading to non-thermodynamically favoured materials [109]. Small concentrations of anti-site defects, as high as 1 %, were suggested to remain even up to solid-state synthesis temperatures as high as 600 °C [110]. In addition, combined neutron and X-ray diffraction has indicated that after fast hydrothermal synthesis crystalline-defective LixFeyPO4 coexists with amorphous Li/ Fe-PO4 structures. These techniques also showed that the Fe is included in the structure more rapidly from the amorphous precursor than Li, causing defects in the structure [111]. Anti-site defects are expected to play a decisive role in the Li-ion conductivity and Gibot et al. [112], using combined neutron and X-ray diffraction data, demonstrated that large concentrations (up to 20 %) of these anti-site defects in nanoparticles suppress the first-order phase transition normally observed in LiFePO4 leading to a single-phase room temperature reaction upon (de)lithiation. More detailed insight into the correlation between particle size and Li-ion substoichiom­etry was obtained by the direct synthesis of substoichiometric Li1-yFePO4 nano­particles [113]. Combined neutron and X-ray diffraction data of partially-delithiated substoichiometric olivines revealed segregated defect-free (where Li is extracted) and defect-ridden (where Li remains) regions, as shown in Fig. 7.7. This proved that both the anti-site defects obstruct Li+ diffusion, explaining the detrimental electro­chemistry and that the anti-site defects form clusters.

Further details of the anti-site clustering in LiFePO4 were obtained using a combination of neutron diffraction with high-angle annular dark-field scanning

transmission electron microscopy and ab initio calculations, indicating that they form zig-zag type clusters, completely different from the structurally equivalent LiMnPO4 where the anti-site defects appear to be randomly distributed [114].

Another topic that has been of great interest is the impact of the particle size on the intercalation properties. When insertion electrode materials are downsized to nano­meter dimensions, voltage profiles change considerably reflecting a change in ther­modynamics [37, 39]. First direct evidence of modified electrochemical-structural behaviour in nano-sized insertion electrodes was provided by neutron diffraction on TiO2 anatase, which showed large changes in Li solubility in phases and a strongly — altered phase composition and morphology [45]. Also, the solubility limits during the insertion reaction in LiFePO4 have been under active research, mainly using neutron diffraction as a direct probe [41, 44, 102, 115—120]. This research shows narrow solid-solution domains in micron size particles at room temperature [117] and a solid solution over the entire compositional range above 520 K [102, 121]. Yamada et al. [117] suggested that the extended solid-solution composition-ranges in small parti­cles and a systematic decrease of the miscibility gap was due to strain based on Vegard’s law [41]. Kobayashi et al. [44] isolated solid-solution phases, also sup­porting a size-dependent miscibility gap. Direct evidence of enhanced solubility in the end phases with decreasing primary crystallite-size was provided by a systematic

neutron diffraction study of particle sizes between 22 and 130 nm [46]. The Fourier-density difference maps in Fig. 7.8 illustrate that the Li densities in the Li — poor and Li-rich phases increase and decrease respectively, with decreasing particle size. These observations could be reproduced by calculations based on a diffuse interface model [46, 122]. The diffuse interface introduces an energy penalty for a Li concentration-gradient creating a smoothly-varying Li concentration over an inter­face region with a width of * 10 nm, as shown in Fig. 7.9. The confinement of this interface layer in nano-sized particles moves the observed solubility away from the bulk values. Interestingly, neutron diffraction also proved that the solubility in both phases (LiFePO4 and FePO4) depends on the overall composition, especially in crystallites smaller than 35 nm. Furthermore, this observation could be explained quantitatively by the diffuse-interface model. By varying the overall composition the domain sizes of the coexisting phases change, in this case leading to confinement effects in the minority phase.

The ex situ neutron diffraction studies discussed above have contributed to our current state of understanding of electrode materials. This is in particular based on the sensitivity of neutrons for Li, the charge-carrying element in Li-ion battery elec­trodes. This is vital knowledge not only for the synthesis of new materials, but also for mechanistic understanding of the impact of supervalent doping, defects, composition, and particle size on the intercalation process as illustrated for olivine LiFePO4.

Fig. 7.9 Measured and calculated solubility limits as a function of particle size and overall composition. Left a Symbols Li occupancy for both the Li-poor triphylite а-phase LixaFePO4 and the Li-rich heterosite р-phase LixpFePO4 where xa and xp represent the average solubility limits as a function of particle size, having an overall composition Li05FePO4. Va and Vp represent the corresponding unit-cell volumes. The size of the symbols is approximately the size of the error. Lines Calculated average compositions based on the diffuse interface model. b Calculated concentration profiles based on the diffuse interface model in the а-lattice direction for three different particle sizes at the overall composition Li05FePO4. Right Measured and calculated solubility limits as a function of overall composition. a Symbols Li occupancy derived from neutron diffraction data for both the Li-poor triphylite а-phase and the Li-rich heterosite p-phase representing the average solubility limits as a function of overall composition for different particle sizes. Lines Calculated average compositions based on the diffuse interface model. The size of the symbols is approximately the size of the error. b Calculated concentration profiles based on the diffuse interface model in the а-lattice direction for three different overall compositions all having the particle size 35 nm. Reprinted from (M. Wagemaker, D. P. Singh, W. J.H. Borghols, U. Lafont, L. Haverkate, V. K. Peterson, F. M. Mulder, J. Am. Chem. Soc. 133, 10222 (2011)) [46]

Microporous and mesoporous solid-state materials such as activated carbon, car­bon-based molecular sieves, mesoporous silicas, and zeolites have been demon­strated to have a significant, and in some cases selective, CO2 adsorption capacity. Such materials have advantages for CO2 capture over the amine solvents currently employed in industry as they are endowed with better stabilities and lower energies of regeneration. Zeolites in particular have been widely studied for the purpose of CO2 capture due to their defined and controllable pore size, insensitivity to mois­ture, and high uptake at non-extreme conditions (for example, zeolite 13X has a CO2 uptake of 3.6 mmol. g-1 at 25 °C) [16]. At higher, more industrially-relevant temperatures, these zeolites tend to lose adsorption capacity and also suffer low selectivity for CO2 over other gases (e. g. N2 and H2) as a result of the physisorptive nature of the CO2-adsorbate interaction. To enhance selectivity for CO2, amine — impregnated or amine-modified materials have been explored, which couple the chemisorption approach used in conventional liquid-amine capture with the phys — isorption approach traditionally seen in porous solid materials. This technique has also been employed using a number of porous silicas, such as MCM-41 meso — porous molecular sieves impregnated with polyethylenimine [17] and SBA-15 mesoporous silicas covalently tethered with hyperbranched amines [18]. Despite the increase in CO2 selectivity of such materials achieved using this approach, they often suffer low stabilities over repeated cycles.

For industrial purposes, solid materials with high selectivity and capacity for CO2 uptake, as well as stability to extreme industrial conditions and a low energy for regeneration, are desired. Metal-organic frameworks (MOFs) are a highly promising class of material for this application due to their structural and chemical versatility, arising from different combinations of metal coordination-spheres as well as multidentate bridging ligands with different lengths, shapes, and direc­tionalities of the coordinating groups. While this versatility sometimes comes at the expense of being able to predict structure accurately, the MOF scaffold provides a unique platform upon which to systematically tune the functionalities of known structures to obtain a desirable property [19]. They may be rationally engineered to have a high surface area and porosity, can be post-synthetically modified to allow for increased selectivity for CO2, and can possess excellent stability under indus — trially-relevant conditions [20]. High surface areas and the possibility to possess coordinatively-unsaturated metal sites make MOFs particularly attractive as gas — selective adsorbents. Coordinatively-unsaturated metal centres have been generated in such materials via chelation by post-synthetically modifying bridging ligands or via insertion into open-ligand sites [9]. However, they are most often created through the evacuation of MOFs that have metal-bound solvent molecules. An effective strategy in tuning the selectivity of MOFs for CO2 is the introduction of functional groups into the pores such as amine [21, 22] and sulfone groups (a SO2 group attached to two C atoms) [23], known to specifically interact with CO2 preferentially over other gases of industrial interest.

Topological control is another strategy employed in the design of MOFs for gas separation, however, this approach is frequently serendipitous due to unknown mechanisms of formation of these materials. Despite this, it has been shown that pore size and shape modulation can determine the diffusion dynamics of the molecules to be separated (e. g. metal formates, M(HCO2)2 (M = Mg, Mn, Co, or Ni), have been shown to selectively adsorb CO2 over CH4, suggesting a size — exclusion effect by the small pores) [24]. Generally, attempts to increase pore size through the incorporation of longer ligands results in framework interpenetration. While this is disadvantageous from a gas-storage standpoint, it may be favourable for some guest separations by their kinetic-diameter differences (e. g. CO2 over CH4) [25, 26]. This type of “molecular sieving” approach may also be achieved by taking advantage of the structural flexibility in MOFs. For example, the material Cr (OH)(bdc), where bdc = 1, 4-benzenedicarboxylate and also known as MIL-53(Cr), exhibits a two-step CO2 uptake isotherm compared to a single-step CH4 uptake isotherm, indicative of a specific “gating” effect [27].

To get detailed information about the cation and anion distribution within the materials, the neutron powder diffraction data has to be analysed by the Rietveld refinement method [19, 20]. To obtain reliable results it is important to perform the Rietveld procedure in a physically-reasonable sequence. In the case of the chal — copyrite-type compound semiconductors the chalcopyrite-type crystal structure was used as basis model. The free structural parameters of the fit were the lattice constants, the anion position parameter x, the cation site occupancy factors (SOF) of the 4a and 4b site (SOF4a, SOF4b) and atomic displacement parameters (ADP). The following example sequence of free parameters can be applied for the analysis of Cu(In, Ga)(Se, S)2 compounds:

1. Refined parameters in the first step were the global parameters, zero shift and the scaling factor. Profile parameters like the u, v, w, x and y parameters, defining the full width at half maximum (FWHM), were fixed at the values of a previ­ously-measured standard sample such as Y2O3. Also the background values were fixed in the beginning of the refinement. Other fixed parameters were the structure parameters: lattice constants, ADPs, SOFs, and the atomic position parameters. The SOFs were fixed at values according to the chemical compo­sition as known from chemical analysis. The isotropic ADPs (bSo; A = 4a, 4b; 8d) were kept at 1.0

2. Subsequently, lattice constants were refined, the u, v, w, x and y parameters remained fixed, as well as ADPs, SOFs, and position parameters.

3. The bAso of 4a, 4b, and 8d positions are refined, keeping the SOFs of the respective positions fixed.

4. The cation site occupancy factors (SOF4a, SOF4b) are then refined whilst biso were fixed to their previously-refined values.

5. Step 3 and 4 are repeated until no change of parameters occurs. This procedure has been expanded with the use of anisotropic ADPs, bAA, (A = 4a, 4b; i = 1-3), since the tetragonal system leads to the need for anisotropic atomic displacement.

6. The bA and the SOFs are refined simultaneously until convergence.

7. If necessary, background and profile parameters: u, v, w, x and y as well as asymmetry values are refined.

This refinement strategy results in reliable values for SOFs of the respective species. Using the following method of average neutron-scattering length [22], it is possible to determine the point defect concentration in complex off-stoichiometric compounds over a large sample volume.

The dynamic structure-factor, S(Q, ®), describes scattered neutrons in terms of the wave-vector transfer Q and the neutron energy-transfer h®, where h = h/2n and h is Planck’s constant. The timescales and corresponding energies of the processes that are accessed by neutron scattering from energy materials are illustrated in Fig. 1.2. A particular strength of neutron scattering is that the size and geometry of the volume explored by the dynamic process is also available, and is shown on the horizontal axis of Fig.1.2.

This is particularly useful when measuring local or long-range diffusive pro­cesses, for example in fuel-cell electrolytes (see Chaps. 9 and 10).

Inelastic peaks usually arise from a periodic motion, and the forces controlling this motion are stronger than those that would define a more diffusive motion. Correspondingly, inelastic peaks usually arise at higher energy than the quasielastic broadening, as illustrated schematically in Fig. 1.3.

The analysis of porous materials and their interaction with guest molecules using neutron scattering is experimentally challenging. Even with advances in neutron sources and instrumentation, several hundred milligrams of material are usually required for successful neutron-scattering analysis of these systems. Evacuated materials prepared for guest sorption are air-sensitive, mandating their handling in specialist atmospheres such as a helium-filled glove box, where helium is necessary to avoid the heat-transfer medium freezing where the heat-transfer gas is not removed from activated samples prior to low-temperature (<10 K) measurement.

Obtaining a good neutron-scattering signal from the host or guest being studied can involve isotopic substitution, and often with complex ligands that require deuteration. The requirement of neutrons in this work is demonstrated by the recent synthesis of deuterated forms of complex ligands, such as 4, 4′, 4"-benzene-1, 3, 5- triyl-tribenzoic acid, through a technique developed at a specialist deuteration facility associated with a neutron-scattering centre. Such complex chemical syn­thetic routes are achievements in their own right [90].

The majority of neutron-scattering experiments exploring guest-host interactions in porous adsorbents are in situ in nature. The in situ approach, however, varies in accordance to the experimental need. Most commonly activated materials (porous materials with their pores empty) are analysed at low temperature first, before the introduction of guest molecules to the sample at a temperature where the guest will remain in the gaseous state, and the sample is then cooled slowly to where the guest molecules “lock in” to their equilibrium positions, before the measurement con­tinues. These measurements involve careful control of the temperature of the

sample as well as gas-delivery lines through the use of modified cryofurnaces. Advances in neutron instrumentation, particularly large area-detectors and higher — intensity sources, provide the opportunity to resolve in real-time details for such systems [91].

Recent work on hexabenzocoronene (HBC) derivatives and semi-triangular discotic molecules showed that classical MD simulations are of key importance in deter­mining the influence of structure and dynamics on the conductivity of DLCs [22—25]. MD simulations are capable of predicting the conductivity of the meso — phases and reveal how local conformation, disorder, and dynamics affect efficient charge-transfer along the one-dimensional column pathway. However, to realize the possibility for rational design of compounds with optimal structure-mobility rela­tionships, it is necessary to verify that the MD simulations describe the real liquid- crystalline phase correctly. DLC structures from MD simulations were successfully compared to experiments on lattice constants, density, phase-transition tempera­tures, order parameter, and mutual orientation (twist angle) in specific cases. Fur­thermore, a more thorough and efficient analysis such as Rietveld refinement of crystalline structures can be performed. HAT6 exhibits a rather limited charge — carrier mobility of about 10 4 cm2 V 1 s 1 as measured with pulse-radiolysis time — resolved microwave conductivity [26]. The question is to relate this to the much better mobilities in larger molecules such as HBC.

By comparing classical MD simulations with NPD measurements, and probing dynamics using QENS [27] along with considering the above first-principles DFT calculations (Sect. 6.3.1.1), we can go beyond the identification of the local molecular conformation, structural defects, and thermal motions in HAT6 by investigating and discussing their effect on the charge transport.

In the columnar phase, the neutron diffraction (ND) pattern of a discotic liquid crystal can be conveniently subdivided into three regions, as illustrated in the case of the neutron-diffraction pattern of deuterated HAT6 (Fig. 6.10, left).[7] The red area indicates a region with three sharp peaks originating from the two-dimensional hexagonal lattice, the (100) peak corresponding to an average column-column distance. The blue region originates from the broad distribution in tail-tail distances. The intra-columnar distances correspond to the broad yellow shoulder around 3.65 A. The tail-tail region and shoulder disappear almost completely for the fully — protonated sample. Thus, the shoulder represents intra-columnar distances between whole molecules rather than only the core-core separation.

The crystal structure predicted by the MD simulations[8] reproduces the three regions in the diffraction pattern (Fig. 6.11b). The dynamical behaviour of the liquid-crystalline phase is included in the comparison of the MD simulations with the experimental observations. The dynamic behaviour was included by calcu­lating the anisotropic displacements (Fig. 6.11b). The experimentally-observed

diffraction pattern is intermediate between the simulated patterns of the models. The order in intra-columnar distances and the intensity of the (210) peak is overestimated in both cases. Despite these differences, both the diffraction pattern of TWIST60 and TWIST25 show remarkable agreement with the experimental pattern, while the relaxed structures stem from entirely different starting config­urations (details of the models TWIST60 and TWIST25 are given in the caption of Fig. 6.10). Both models converge to comparable minima in configuration space, close to the actual liquid-crystalline structure, considering the agreement with the experimental diffraction pattern. The competition between the mutual van der Waals interactions of the cores with the steric repulsion of the tails causes a high disorder in the core-core distances, rather than a uniform shift of the core-core distance distribution to higher distances. The tails tend to orient toward the open spaces formed by these core-core defects, since the whole-molecule peak lies at a higher distance than the core-core peak. The inferred values of the twist-angle distribution are found to be around 36°, which is close to the to the DFT-estimated minimum-energy twist angle of 30° [28].

To characterize the dynamical processes on the picosecond timescale, the MD simulations are compared with QENS[9] experiments [27]. In the analyses of the experiments, it was consistently found that in addition to the elastic peak, at least two Lorentzian functions are required to fit the data. The peak widths corresponded to time scales of about 0.2 and 7 ps timescale. The incoherent-scattering function can be extracted from the MD simulation and compared with the measured function (Fig. 6.12, left). The agreement is satisfactory, and the essential observation is that the MD models compare well with the measured temporal and spatial dynamics (Fig. 6.12, right). By examining the simulation trajectories more closely the underlying thermal motions can be characterized. Further, the typical amplitudes of molecular motion on the 0.2 and 7 ps can be estimated.

Figure 6.13 illustrates the characteristic translational and tilt motions for a single molecule on the 0.2 ps timescale. The twist-angle deviations are negligible on this timescale with amplitudes smaller than 0.1°. Apparently, the tilt movement of the molecular core is too fast to be followed by the aliphatic tails. On the 7 ps timescale, the rotational motions are much more prominent: the aliphatic tails start to follow the rotations of the core. A step further towards the understanding of the hot carriers relaxation would be to follow closely the dynamics of both the core and tails at the vibrational level and investigate how the molecular vibrations in the ground and excited electronic-states behave.

Bio-ethanol obtained from biomass has been suggested as a promising renewable source of H2. It is a challenge to find low-cost catalysts (without noble metals) able to break the C-C bond of ethanol at low temperature. A CeNiHZOY catalyst was recently found to convert ethanol at 60 °C only, by steam reforming coupled with partial oxidation [8]. The distribution of products is similar to what is obtained by steam reforming at high temperatures: H2 (about 45 %) and mainly CO2 and CO. The reaction is initiated at 230 °C, but the temperature increases after a short induction period so that a temperature of 60 °C is sufficient to maintain the reaction. This is explained by the occurrence of two exothermic reactions: (i) between hydride species

of the catalyst and O2, and (ii) between ethanol and O2 (partial oxidation). The reaction is sustainable because hydride species are replaced and provided by ethanol.

A pre-treatment at 250 °C under H2 is necessary to obtain the active catalyst, which is an oxyhydride. As in the case of copper chromite, large quantities of hydrogen can be stored in CeNiXOY mixed oxides. H2 is heterolytically dissociated at an anionic vacancy and an O2- species of the catalyst. The insertion of hydride species in the solid was evidenced by INS. The spectra of CeNi1OY and CeNi0.5OY are shown in Fig. 2.2. The INS spectrum of the solid treated in vacuum at 200 °C, which contains OH groups, has been subtracted. The peak at about 460 cm-1 was assigned to hydrides and the band around 870 cm 1 to H adsorbed on metallic Ni particles, because the band intensity decreases when the Ni loading is decreased. While the assignment of the first peak appears to be reliable (after re-oxidation, this peak disappears whereas a band corresponding to OH groups emerges around 630 cm 1), the assignment of the higher frequency band to p3-H species on Ni0 particles is less certain, and the contribution from OH groups cannot be excluded.

Hydride formation in Zircaloys reduces strength and ductility significantly. Hydrides form preferentially in areas of higher stress [36] such as near welds or at crack tips. In loss of coolant accidents (LOCAs), overheated Zircaloy cladding may react with cooling water in a complex manner with different phase-transformation temperatures depending on other species such as oxygen or hydrogen. The solubility of hydrogen in zirconium and Zircaloy differs by almost an order of magnitude between the alpha and beta phases [37], and excess hydrogen may form embrittling hydrides, particularly at

Fig. 4.10 Optical (a) and neutron-radiography images (b, c, d, e) of the welded Zircaloy-4 plates. Figure b was produced using neutrons with wavelengths between 1.4 and 2 A, whilst Figures c, d and e are images of the height of selected Bragg edges. In figure c the start of the HAZ clearly shows the disappearance of the (10-10) Bragg edges, whilst the (11-20) edge in figure d reveals the differences between the outer and inner layers of the plate. Reprinted with permission from (J. R. Santisteban, M. A. Vicente-Alvarez, P. Vizcaino, A. D. Banchik, S. C. Vogel, A. S. Tremsin, J. V. Vallerga, J. B. McPhate, W. Lehmann Kockelmann, J. Nucl. Mater. 425, 218 (2012)) [35]. Copyright (2012) Elsevier

crack tips etc. Despite decades of work, the mechanisms of corrosion of zirconium — based alloys, particularly in reactor conditions, are still an active research field.

As hydrogen is a strong neutron-scatterer, the hydrogen concentration (Fig. 4.11) can be measured by neutron radiography [38] with sensitivities of * 1,000 wt. ppm [39]. Neutron diffraction and radiography can be combined [40, 41] to identify and map hydride phases. The distribution [42] and kinetics [37] of hydrogen in mate­rials can be identified by neutron radiography.

The hydrogen concentration in metals (particularly in zirconium-based alloys) can be also be measured by cold neutron prompt-gamma activation analysis PGAA [44]. However, these techniques are not suitable for imaging and rather provide the
bulk average from the volume illuminated by the incident neutron beam. Cold neutron PGAA is based on measuring prompt gamma rays following the absorption of cold neutrons by hydrogen. This method has to be performed in a close prox­imity to a neutron source. The technique does not change the sample and so is well suited to assessing hydrogen uptake during interrupted corrosion testing.

Neutron scattering holds many opportunities for obtaining unique information concerning the porous-solid adsorbent (host) as well as host-adsorbate (host-guest) system. Measurement of structure and dynamics using neutron scattering, across length and time scales pertinent to these systems (also possible at the same time), has been exploited for the better understanding of guest binding in the host and separation mechanisms, as well as the host’s response to adsorption, all of which are key to progressing the application of such systems in CCS.

Information about both host and host-guest structure, which yields details of the structural response of the host to adsorption, the location of the guest in the host, and guest-host interaction, are important to determining structure-property relations. As neutron diffraction intensity does not reduce with scattering angle, relatively more fine structural detail is gained than using X-ray diffraction, providing important detail concerning the guest-host and guest-guest interactions. The iso- topically-dependent structural information afforded by neutrons allows different contrast between parts of the host framework and/or guest to be gained, providing many advantages for such structural investigations. Examples include distinguish­ing between guests such as N2, O2, and CO2, and obtaining details of both the host’s ligands and metal centres, as well as guests, even within a MOF containing heavy-metal atoms and guests containing light atoms. Additionally, the information obtained can be tuned through isotopic substitution, such as in determining the molecular orientations of CH4 within a host using the isotopically-substituted CD4, where D is deuterium (2H),

The dynamic information obtained through neutron scattering is also isotopically dependent, and spectroscopic neutron techniques allow direct measurement of the local environment and the diffusional transport of the guest within the host. Both structure and dynamics can be measured at the same time, enabling insights into the geometry of the guest motion, in turn allowing the details of the mechanism of diffusion of the guest within the host to be gained.

In situ methods are central in the analysis of MOFs for guest separation and storage applications. Although in situ X-ray single crystal and powder diffraction studies of CO2 in MOFs facilitate the understanding of the functional mechanism of MOFs for CCS applications [28—30], in situ neutron-scattering methods have sig­nificant advantages over X-ray studies of MOF-guest systems, with the penetrating power of neutrons being central to this. Neutrons easily penetrate the often-complex sample environments required for control over temperature of the host at the same time as gas delivery, covering easily the range of temperatures from the relatively cold (about -263 °C) conditions required to “lock in” guests and determine accurate structural details, to the more moderate temperatures required to replicate working post-, pre-, and oxyfuel combustion, as well as natural gas-sweetening conditions (40-75 °C). The relatively high penetrating power of neutrons also allows for the analysis of bulk samples, mg—gram quantities, providing information about the more industrially-relevant “bulk” properties of the material. The bulk-scale analysis also aids in accurately dosing the sample with a known number of guest molecules to determine in detail the nature of their interaction with the host.