From a “molecular” point of view, each electronic state of the organic molecule has a well-defined potential energy hyper-surface (PES), with a minimum corre­sponding to the molecular geometry. There is therefore also a set of characteristic molecular vibrations and the electronic relaxation process is mediated by the molecular vibrations of the corresponding electronic state. One aspect of the energy losses concerns the hot-carrier relaxation of the molecule through excited elec­tronic-states, and the excited state via which the electronic charge is finally trans­ferred to the neighbouring molecule. The electric potential related to an exciton is greater for the higher excited states. In analogy with relaxation to the ground state, relaxation between electronic states is driven by the molecular vibrations in the corresponding electronic states. The carrier relaxation-dynamics are of fundamental importance to understanding the processes underlying carrier transport in both organic and inorganic devices. The advantage of an organic system as a model for research is the rich molecular vibrational-spectrum that can be studied computa­tionally and experimentally. Determining the PES of ground and excited electronic — states would provide insights into the variation in molecular geometry and vibra­tions in these states, with a view to ascertaining whether such changes are signif­icant and measureable and further, whether the structure and vibrations can be chemically tailored to improve the performance of OPVs.

Organic systems, like hexa-peri-hexabenzocoronene with a perylene dye have been successfully used in photodiodes with efficiencies as high as 34 %. However, these systems are too complex for high-level, molecular-modelling studies of vibrations in ground and excited electronic-states. These vibrations play a key role in properties such as: light-harvesting, exciton transport, primary charge-separation, and electron/hole conduction. In this context, the columnar discotic liquid-crystals based on triphenylene derivatives like hexakis(alkyloxy)triphenylenes (HATn), are attractive representative model systems for organic photovoltaic devices. Works on ground-state (valence-band) electron-phonon coupling of HAT6 (HATn with n = 6) have shown that structure and dynamics play a central role in the charge-transfer process of this material, and it was noted that the dynamics of the aromatic cores and the alkyl tails have an important effect on the electronic properties [14-18]. Perhaps surprisingly, it was found that the electronic properties are affected not only by the dynamics of the aromatic cores, but also by the dynamics of the alkyl tails. HAT6 is a convenient model system (Fig. 6.8) because it has the optimum side — chain length giving the broadest mesophase range. Between five and seven carbons in the alkyl tail are required for the columnar phase to form. Further increase or decrease of the tail length reduces the temperature range of the hexagonal columnar mesophase formation. From a practical point of view, it is generally thought that the

BHJ setup leads to optimal performance of DLC solar cells. An interesting design route to achieve this device architecture is to dilute the columnar liquid crystalline phase with a non-discogenic electron acceptor, such as 2,4,7-trinitro-9-fluorenone (TNF) (Fig. 6.8).

A HAT6/TNF diamond-like carbon (D-LC) charge-transfer (CT) system is then realized. The following sections present selected results that highlight the use of neutron powder diffraction (NPD) and quasielastic neutron scattering (QENS) measurements, as well as simulations based on classical molecular-dynamics (MD) and ab initio density functional theory (DFT), to probe structure and dynamics of the HAT6/TNF D-LC-CT system. We also show that this type of study, in which both the nuclear and electronic aspects of structure and dynamics are mixed, extensive work from other techniques must be considered. Consequently, whilst the main thrust is neutron scattering, results from nuclear magnetic resonance, Raman, resonant-Raman, UV-visible, and infrared (IR) measurements are also presented and discussed. We emphasise here that neutron-scattering methods can bring interesting structural and dynamical information for HAT6 systems, largely because of their abundant H-atoms. In contrast to X-rays, neutron scattering from H atoms is strong and by deuterating the HAT6 sample (HAT6D) the related sites and their dynamics are highlighted differently.

The activation of a copper-chromium system in H2 is accompanied by an accumu­lation of H2, which can become active for hydrogenation reactions, in the absence of H2 in the gas phase. This sorption of H2 is due to a specific mechanism of Cu2+ reduction from the CuCr2O4 structure, a reaction which does not release water. The reduction leads to the formation of metallic copper and to protons substituting for copper cations in the vicinity of O2 anions. INS spectra of copper chromite as prepared and reduced at various temperatures are shown in Fig. 2.1. The initial catalyst shows bands at low frequencies, due to the mass of the atoms, these bands having a low intensity because no proton motion is involved. Upon reduction, a large intensity

increase can be measured and OH-groups bending modes are observed in the energy range displayed in Fig. 2.1. Increasing the reduction temperature yields a shift of these bending modes to higher frequencies (from 700-800 to 1,220 cm-1), while the stretching modes shift to lower energy. This indicates that hydrogen bonding strength with neighbouring anions increases with the rise in temperature. The band around 400 cm-1 was assigned to librations of two geminal protons (i. e. HOH-groups) [6].

Due to the complex deformation behaviour of highly-anisotropic Zirconium alloys, these materials have been studied extensively by neutron diffraction, in combination with polycrystalline deformation models, the elasto-plastic and visco-plastic self­consistent models [18], to understand the deformation and texture development.

The difference in tensile and compressive behaviour in Zircaloy-2 was demon­strated by MacEwen et al. [19]. The strains were measured for each lattice plane on a time-of-flight (TOF) instrument (the General Purpose Powder Diffractometer at the intense-pulsed neutron source (IPNS), which is no longer operational). Neutron time-of-flight instruments can measure multiple lattice planes without reorienting the sample, as is required on a constant-wavelength instrument.

The residual stresses, stress tensor, and inter-granular stresses were characterized after 5 % strain by Pang et al. [20]. The measured lattice strains were in good agreement with those predicted by elasto-plastic self-consistent models, which predict deformation modes such as slip and twinning.

Zirconium and its alloys have a hcp structure, and there are too few slip systems for standard plasticity so twinning is an important contributor to plastic deforma­tion. Rangaswamy et al. [21] compared changes in texture and twin volume-frac­tions to predictions from a visco-plastic self-consistent polycrystal model, which described both slip and twinning.

Balogh et al. [22] examined the deformation behaviour of Zr-2.5Nb samples by full-pattern diffraction line-profile analysis (DLPA) to determine the evolution of the density and type of the dislocation-structure induced by irradiation and plas­ticity. Control samples were compared to samples removed from a CANDU nuclear reactor pressure-tube to determine the evolution of microstructure and plasticity characteristics during deformation (27 % cold work during manufacture). The pressure tube was in service for 7 years at *250 °C with a neutron fluence of

1.6 x 1024 m-2 (E > 1 MeV). Results show that fast-neutron irradiation signifi­cantly increases the overall dislocation density, accomplished entirely by an increase in the (a) Burgers vector dislocations.

Since the commercialization of the Li-ion battery by SONY Corporation in 1991, research has focused on identifying better electrode and electrolyte materials. SONY combined LiCoO2 as a positive electrode material with a carbonaceous material as a negative electrode, and LiPF6 in a carbonate solution as the electrolyte. Initially, research focused on replacing the relatively expensive LiCoO2 with other transition-metal oxides where the most important structural groups include spinel and layered transition-metal oxides. Layered Li transition-metal oxides, LiMO2 with M = Mn, Co, and Ni, represent one of the most successful classes of positive — electrode materials. The layered topology offers easily-accessible two-dimensional ion diffusion pathways. In particular, the LiCo1/3Ni1/3Mn1/3O2 composition [18] results in high capacity, safety, and lower material costs than LiCoO2.

An important alternative to LiCoO2 is the spinel LiMn2O4 with the inexpensive and environmentally-benign Mn, which functions with the Mn3+/4+ redox couple. LiMn2O4 operates at 4.1 V versus Li/Li+ offering excellent safety and high power due to its three-dimensional lattice, in principle allowing three-dimensional Li-ion diffusion. Substitution of Mn by M = Co, Cr, Cu, Fe, and Ni, has led to the discovery of the high-voltage spinels LiM0 5Mn15O4 and LiMMnO4 with potentials between 4.5 and 5 V versus Li/Li+ [19], exemplified by spinel LiNi0 5Mn15O4 operating at 4.7 V versus Li/Li+ [20, 21]. However, Mn-based spinels have been plagued by capacity fade, generally considered to be the result of Mn dissolution into the electrolyte and Jahn-Teller distortion of Mn3+. Mn dissolution has been largely inhibited by substitution of dopants in the spinel structure [22].

The introduction of LiFePO4 [14] has initiated research on polyanion-based positive electrodes with the structural formula LiM(XO)4 (M = Fe, Mn, Co, and X = S, P, Si). The strong covalent X-O bonds result in the large polarization of oxygen ions towards the X cation, leading to larger potentials compared to oxides. Phosphates, in particular LiFePO4, have been extensively explored because of their favourable electrode properties, reasonably high potential and capacity, stability against overcharge or discharge, and their composition of abundant, cheap, and non-toxic elements. To date, almost a thousand papers have been devoted to the understanding and improvement of conductivity in the semiconductor LiFePO4. These have encompassed studies of surface-conductive phases [23, 24], modifica­tion of crystallite size [25], and elegant fundamental mechanistic and modelling studies [26-28]. Compared to the phosphates, the silicates (LiMSiO4) exhibit lower electrode-potentials and electronic conductivity [29]. Other promising groups of positive-electrode materials include fluorophosphates with the A2MPO4F (A = Alkali metal) stoichiometry [29].

For many years graphite has been the main negative-electrode material allowing Li-ion intercalation between graphene sheets at *0.3 V versus Li/Li+. The low potential results in a high battery-voltage, partly responsible for the high energy and power density of Li-ion batteries. The main disadvantage of the low voltage of a negative electrode is the restricted stability of the commercial carbonate-based electrolyte solutions at low potentials versus Li/Li+. Depending on the electrolyte — electrode combination a kinetically stable SEI layer can be formed. For graphite this was achieved by the addition of ethylene carbonate [30] which lead to the success of the LiCoO2-C battery. Efforts to develop low-voltage negative electrodes with capacities that exceed that of C have concentrated on reaction types other than insertion reactions. Metals such as Si and Sn form alloys with Li in so-called conversion (or alloying) reactions [31-34]. Although the (gravimetric) capacities of these reactions can reach almost 10 times that of graphite, the volumetric energy — density is not significantly improved. Moreover, the large volume changes inher­ently related to this type of reaction make it a challenge to reach good cycleability. The oxide insertion hosts offer much better stability, highlighted by the Li4Ti5O12 spinel [35] that has almost no volume change upon Li insertion. Other titanium oxides that attract considerable attention as negative electrodes are anatase, brookite, and rutile TiO2, particularly in the nanostructured form [36, 37] and in combination with carbonaceous materials providing good with electronic-conduc­tivity [16, 36]. The disadvantage of titanium oxides is their relatively high negative- electrode potentials, reducing the overall working voltage of the battery and thereby reducing both energy and power density. For example, titanium oxides operate at *1.6 V versus Li/Li+, over four times higher in voltage than graphite. In this context, an interesting oxide exceeding the graphite capacity at similarly low potentials is the layered transition-metal oxide Li1+xV1-xO2 [38].

One of the key strategies that improved electrode performance in general is the nano-sizing of the electrode crystal particles. The nano-sized electrode particles reduce both solid state (Li-ion) diffusion and electron conduction. The latter is achieved by mixing and coating with conducting phases (i. e. with carbonaceous materials). However, numerous recent observations indicate that nano-sizing elec­trode particles also has a large impact on electrode materials properties [39, 40] creating both opportunities and challenges for enhanced Li-ion storage. Changes in properties that are observed upon nano-sizing include smearing of the voltage profile [41—43], changing solubility limits and phase behaviour [41, 44—46], unexpected kinetics [47], and larger capacities [45, 48—51]. The downside of the large surface-area of nanostructured materials is the relative instability of nano­materials, which can promote electrode dissolution and the increased reactivity towards electrolytes at commonly used voltages, e. g. below 1 V versus Li/Li+, which may adversely affect performance. Another potential disadvantage is the lower packing density of electrodes, leading to lower volumetric energy-densities.

Among the materials that benefit from the possibilities of nano-sizing are the relatively stable transition-metal oxides and phosphates operating well within the stability window of the electrolyte.

As the prime mover of carbon through the atmosphere, carbon dioxide (CO2), plays a vital role in enabling the cycle of carbon from the Earth’s crust (where it is found in elemental graphite and diamond, carbonates, and fossil fuels) to our oceans

A. Das • D. M. D’Alessandro (H)

School of Chemistry, The University of Sydney, Sydney, NSW, Australia e-mail: deanna. dalessandro@ sydney. edu. au

A. Das

e-mail: anita. das@sydney. edu. au V. K. Peterson (H)

Australian Nuclear Science and Technology Organisation, Lucas Heights, NSW, Australia e-mail: vanessa. peterson@ansto. gov. au

© Springer International Publishing Switzerland 2015 33

G. J. Kearley and V. K. Peterson (eds.), Neutron Applications in Materials for Energy, Neutron Scattering Applications and Techniques,

DOI 10.1007/978-3-319-06656-1_3

(where it occurs in carbonate minerals formed by the action of coral-reef organisms and aqueous CO2). For hundreds of millions of years, the carbon cycle has main­tained a relatively constant amount of CO2 in the Earth’s atmosphere (approxi­mately 400 ppm by volume). While the contribution from human industry is relatively small, its recent growth has shifted this natural balance. Since the start of the Industrial Revolution around 1760, the concentration of CO2 in the atmosphere has risen dramatically from 280 to 385 ppm today [1, 2]. This significant rise has been attributed to an increasing dependence on the combustion of fossil fuels (coal, petroleum, and natural gas), which account for 86 % of man-made greenhouse-gas emissions, the remainder arising from land use change (primarily deforestation) and chemical processing.

The development of more efficient processes for CO2 capture from major point sources such as power plants and natural-gas wells is considered a key to the reduction of greenhouse-gas emissions implicated in global warming. Numerous national and international governments and industries have established collabora­tive initiatives such as the Intergovernmental Panel on Climate Change [3] (IPCC), the United Nations Framework Convention on Climate Change [4], and the Global Climate Change Initiative [5] to achieve this goal. The capture and sequestration of CO2, the predominant greenhouse gas, is a central strategy in these programmes as it offers the opportunity to meet increasing demands for fossil-fuel energy in the short to medium term, whilst reducing the associated greenhouse-gas emissions in line with global targets. Carbon capture and storage (CCS) will complement other strategies such as improving energy efficiency, switching to less carbon-intensive fuels, and the phasing in of renewable-energy technologies.

Three major technologies are predicted to have the greatest likelihood of reducing man-made emissions to the atmosphere that are implicated in global warming. These processes include postcombustion and precombustion capture from power plants involving CO2/N2/H2O and CO2/H2 separations, respectively, and natural-gas sweetening (CO2/CH4/N2 separation). The separation processes required for each of these capture applications differs with regard to the nature of the gas mixture and the temperatures and pressures involved, imposing constraints on the materials and processes employed [6, 7].

Conventional CO2 capture processes employed in power plants world-wide are typically postcombustion ‘wet scrubbing’ methods, involving the absorption of CO2 by amine-containing solvents such as methanolamines [8]. Power plant flue-gas streams consist primarily of N2, H2O, and CO2 in a 13:2:2 ratio by weight [9]. Prior to the compression and liquefication of the captured CO2 for transportation to storage sites, CCS requires the separation of CO2 from all other flue-gas compo­nents. CO2 is strongly absorbed by the amine to form a carbamate species [10], however, the high heat of formation associated with the creation of the carbamate leads to a considerable energy penalty for regeneration of the solvent. Since the flue streams from coal-fired power plants contain dilute concentrations of CO2 (typically 10-15 %) at relatively low pressures and temperatures (1 atm., 40 °C), it is estimated that CO2 capture and compression will increase the energy requirements of a plant by 25-40 %. Analysis has shown that the thermodynamic minimum energy-penalty for capturing 90 % of the CO2 from the flue gas of a typical coal-fired power plant is approximately 3.5 % (assuming a flue gas containing 12-15 % CO2 at 40 °C) [11]. The transportation and storage of CO2 will necessitate further investment and capital costs. These economic and energy comparisons underscore the immense opportu­nities and incentives that exist for improved CO2 capture processes and materials. Despite improvements in conventional postcombustion chemical-absorption meth­ods, wet-scrubbing methods suffer a number of drawbacks and are therefore not cost — effective for large-scale carbon emissions reduction.

While the retrofitting of existing power plants using postcombustion capture methods presents the closest marketable technology, two major alternatives to postcombustion CO2 capture processes have been proposed, and are currently in the test stages of development [12]. Precombustion processes involve a preliminary fuel-conversion step using a gasification process and subsequent shift-reaction to form a mixture of CO2 and H2 prior to combustion. The high pressure of the product gas-stream facilitates the removal of CO2 from the CO2/H2 mixture at pressures of 30-50 bar and temperatures of 50-75 °C [13]. The significant advantage of precombustion capture is that the higher component concentrations and elevated pressures reduce the energy capture penalty of the process to 10-16 %, roughly half that for postcombustion CO2 capture. A further advantage is that precombustion technology generates an H2-rich fuel, which can be used as a chemical feedstock in a fuel cell for power generation or in the development of an H2 economy. In oxyfuel (or denitrogenation) processes, fuel is combusted in O2 instead of air by the exclusion of N2, thereby producing a concentrated stream of CO2 without the need for separation (in high, sequestration-ready concentrations of 80-98 %). Since the separation of interest in this case is air separation (O2 from mainly N2 at a pressure of around 100 bar and temperature of 50 °C), reducing the cost of O2 generation is key to industrial viability. While the emerging technologies associated with precombustion and oxyfuel processes cannot be readily incorpo­rated (via retrofitting) into existing power plants as can postcombustion CO2 cap­ture processes, the projections from the IPCC indicate that the required extensive capital investments will be compensated by the relatively higher efficiency of the CO2 separation and capture process [3].

Another important application for CO2 capture technologies is the ‘sweetening’ of sour natural-gas wells, where the sweetening refers to the separation of CO2 from CH4. Natural gas reserves (mainly CH4) are typically contaminated with over 40 % CO2 and N2 and the use of such fields is only acceptable if the additional CO2 is separated and sequestered at the source of production. The capture of CO2 from ambient air has also been suggested, however, the low concentration of CO2 in air (0.04 %) presents a significantly higher barrier to capture compared with post­combustion methods, and the expense of moving large volumes of air through an absorbing material presents a further challenge in its implementation [6].

The key factor that underscores significant advancements in CCS is materials to perform the capture process [6, 7]. The challenge for gas-separation materials is that the differences in properties between the gases that have to be separated are rela­tively small. However, differences do exist in the electronic properties of the gases:

CO2 has a large quadrupole moment (13.4 x 10-40 cm2 vs. 4.7 x 10-40 cm2 for N2 and CH4 is non-polar) and CH4 adsorbs preferentially over N2 due to its higher polarizability (17.6 x 10-25 cm3 for N2 and 26.0 x 10-25 cm3 for CH4).

A diverse range of promising methods and materials for CO2 capture applica­tions that could be employed in any one of the abovementioned postcombustion, precombustion, or oxyfuel processes have been proposed as alternatives to con­ventional chemical absorption. These include the use of physical absorbents, membranes, cryogenic distillation, hydrate formation, chemical-looping combus­tion using metal oxides, and adsorption on solids using pressure and/or temperature swing adsorption, where the adsorption and desorption temperature/pressures are different and a “swing” is made between them [14]. The key requirements for these new materials are that they exhibit air and water stability, corrosion resistance, high thermal-stability, high selectivity and adsorption capacity for CO2, as well as adequate robustness and mechanical strength to withstand repeated exposure to high-pressure gas streams. A number of review articles have elaborated the status of a new classes of materials for CO2 capture [6, 7]. In particular, metal-organic framework (MOF) materials are progressing at a rapid pace.

With respect to new materials, the key scientific challenges are the development of a level of molecular control and modern experimental and computational methods. For crystalline materials, adsorption isotherms and breakthrough-curve measurements under conditions that closely resemble working-condition gas mix­tures are essential. In reality, pure gas-adsorption isotherms are often measured, with ideal adsorbed solution theory (IAST) applied in some cases to predict mul­ticomponent adsorption behaviour [15]. A parameter that must be assessed in all cases is the enthalpy of adsorption, since the cost for the regeneration of a capture material is clearly dependent on the energy required to remove the captured CO2.

Characterization of the molecular transport properties of materials is essential to obtain an understanding of transport processes. Important molecular-level infor­mation required includes: how the material structure changes with loading, how adsorbates bind to the material, and how different permeates influence each other’s solubility. In situ techniques are particularly powerful as they allow the interactions between gas molecules and the matrix to be probed and determine the material structure under different loading conditions. The most significant information from such measurements is gained from the ability to correlate absorbate uptake with the absorbent structure and the molecular-level absorbate mobility. A comparison between the molecular-level absorbate mobility and its macroscopic diffusion should provide insights into the mechanism of selective transport through these materials.

In parallel with experimental studies, computational modelling methods are being developed, both as a tool to understand further details of the adorbate — absorbent interaction, and as a tool to predict the performance of materials proposed for a given separation process, with the latter enabling large-scale screening of new materials. Ultimately, a clear understanding of the structure — and dynamics-function relations will direct experimental efforts towards a new generation of materials with improved CO2 capture abilities. Developing force fields for computational work using detailed structures is important for the successful prediction of thermody­namic and transport properties of new materials.

Among various experimental possibilities, structural analysis of compound semi­conductors by diffraction techniques using X-rays or neutrons has become a technique of choice. The reasons are various, and here we place the focus on the use of neutrons as a radiation source.

The highest efficiency thin-film devices consist of a Cu-poor Cu(In, Ga)Se2 absorber layer. The interplay of structural with electronic and optical properties is therefore interesting to study for this quaternary compound and the corresponding ternaries (CuGaSe2 and CuInSe2). In the case of CuGaSe2 or Cu(In, Ga)Se2 the cations Cu+ and Ga3+ have an identical number of electrons (28). This is a problem for the differentiation of these two cations by conventional X-ray diffraction, where the diffraction is at the electron shell of the atoms. Since atomic scattering form factors /for X-rays are proportional to the atomic number Z, the positions of the unit cell atoms of similar atomic number and the fractional occupation of the Wyckoff sites are difficult to distinguish. Hence, a differentiation of these cations in the atomic structure by conventional X-ray diffraction is impossible. In the case of neutron diffraction, the scattering is at the nucleus and the neutron-scattering lengths of copper and gallium are different (bCu = 7.718(4) fm, bGa = 7.288(2) fm [18]). Moreover, in the case of X-rays, destructive interference effects lead to a decrease of the scattering amplitude with angle. In contrast to diffraction at the electron shell, the atomic nuclei cross-sections are very small and the interference effects are also very small. Therefore, neutron-scattering amplitudes do not decrease rapidly with angle, resulting in the advantage of high intensities of Bragg-reflections observed even at high Q-values. This is especially important for the determination of atomic positions and atomic site occupations. For the analysis of point defects in compound semi­conductors and position parameters, for instance of the anion, neutron powder dif­fraction is advantageous.

Small-angle neutron scattering (SANS) is a well-established characterization method for microstructure investigations, spanning length scales from A to micron sizes. Within the SANS approximation, Q simplifies to Q = 2я0/А The SANS Q range is typically from 0.001 to 0.45 A-1. For example, scattering from a simple spherical system in a solvent yields a SANS coherent macroscopic neutron-scattering cross section (scattering intensity in an absolute scale) of:

^dX(Q) = NVpAq2P(Q)S: (Q) (1.2)

where (N/V) is the number density of particles, VP is the particle volume, Др2 the contrast factor, P(Q) is the single particle form factor, and SI(Q) is the inter-particle structure factor. SI(Q) has a peak corresponding to the average particle inter­distance.

Natural-gas sweetening (separation of CO2 from CH4) is an industrially significant separation process as CO2 represents a substantial (up to 70 %) impurity in natural — gas wells [80]. The presence of CO2 reduces the energy content of the natural gas, and its acidity in the presence of water can result in the corrosion of natural-gas lines. Physical solvent-based processes for CO2 removal from natural-gas are abundant, however, the large amount of water recycling needed makes solvent- based processes highly limited in this application due to solvent degradation and loss during operation [81]. Porous solids present a more efficient and environ­mentally friendly way to capture CO2 from natural-gas wells. In this case, sepa­ration largely proceeds based on quadrupole moment, due to the similar properties of the two gases in other respects (kinetic diameter, polarizability, dipole moment). Additionally, the flexible structure of some MOFs upon adsorption-desorption (in contrast with “rigid” adsorbents such as carbons and zeolites), may result in dynamic and stepwise adsorption at different pressures. This is generally known as a “gate opening” phenomenon, and arises mainly from the flexibility of the net­works and their affinity for particular guests [82]. In MIL-53 [Cr(OH)(bdc)], for example, the selective adsorption of CO2 over CH4 is strongly affected by the presence of water which causes dramatic changes in the pore structure [27].

Neutron scattering has also been used extensively to study CH4 confined in porous materials, in particular to study methane confined in MOFs, commensurate

with the increasing work investigating these hosts for application in CO2/CH4 separations.

The metal sites, including open-metal sites, in many MOFs also interact with CH4. NPD studies of the Mg2(dobdc) material show the binding of one CD4 molecule per open-metal site, resulting in the large CH4 storage capacity of 160-174 cm3(at standard temperature and pressure, STP)/cm3, approaching the DOE target of 180 cm3(STP)/cm3 for solid-based CH4 storage at room temperature [83]. Direct determination of CD4 sorption sites in Zn(mIm)2 and Zn4O(bdc)3 were gained using NPD (Fig. 3.8) [84]. The primary CD4 adsorption sites are associated with the organic linkers in Zn(mIm)2 and the metal oxide clusters in Zn4O(bdc)3. In Zn4O(bdc)3 the first binding sites (“cup” sites) were not found to alter the Fm 3m symmetry of the host-guest system. CD4 at these primary sites possesses well — defined orientations, implying relatively-strong binding with the framework. With higher CD4 loading, additional CD4 molecules populate secondary sites and are confined in the framework. The confined CD4 at these secondary sites is orienta­tionally disordered and stabilized by the intermolecular interactions. The CD4 guest is a high symmetry guest whose ordered location (at the primary sites) significantly alters the symmetry of system. The “hex” and “ZnO2” CD4 sites caused a symmetry lowering of the system to I4/mmm as a result of the symmetry incompatibility of the tetrahedral CD4 molecules with the local geometry. At higher CD4 loadings a P4 mm structure was found, where CD4 sites aligned themselves along the c axis and further lowered the symmetry.

Using a similar approach, a comprehensive mechanistic study of CD4 was performed in Cu3(btc)2, Cu2(sbtc) where sbtc = frans-stilbene-3, 3′, 5, 5′-tetracar — boxylate, and Cu2(adip) where adip = 5, 5′-(9, 10-anthracenediyl)di-isophthalate and also known as PCN-14, allowing a comparison of structures that consist of the same dicopper-paddlewheel secondary-building units (the well-known dicopper acetate unit), but contain different organic linkers, leading to cage-like pores with various sizes and geometries (Fig. 3.9). This work revealed that CD4 uptake takes place primarily at two types of strong adsorption site: (1) the open Cu sites which exhibit enhanced coulombic attraction toward CD4, and (2) the van der Waals potential pocket sites in which the total dispersive interactions are enhanced due to the molecule being in contact with multiple “surfaces”. Interestingly, the enhanced van der Waals sites are present exclusively in small cages and at the windows to these cages, whereas large cages with relatively flat pore surfaces bind very little CD4 [85].

The self-diffusion of CH4 was measured directly using QENS in the isostructural Cr(OH)(bdc) and V(O)(bdc) materials [86]. The hydroxyl groups in Cr(OH)(bdc) were expected to hinder CH4 mobility, although this work revealed a global one­dimensional diffusion mechanism of CH4 in both materials, echoing the single-file diffusion regime found for CO2 [49]. An interesting result of this work was that CH4 diffusivities are significantly higher in V(O)(bdc) than in Cr(OH)(bdc) over the whole range of investigated CH4 loadings.

There have been several neutron-scattering studies targeting the separation mechanism of CO2 from CH4. The polymer/selective-flake nanocomposite mem­branes exhibiting selectivity for O2 over N2 (discussed in Sect. 3.3.3) also shows substantial selectivity of CO2 over CH4. Again, SANS results revealed that this occurs within only 10 wt% of the AlPO layers. The Zr6O4(OH)(bdc)6 material, also known as UiO-66(Zr), is a MOF with encouraging properties for CO2/CH4 gas separation, achieved by combining good selectivity with a high working capacity

and the ability for regeneration under relatively-mild conditions [87]. Zr6O4(OH) (bdc)6 is built from Zr6O4(OH)4 octahedra that extend into three-dimensions via bdc ligands, resulting in two types of microporous cages. The dynamics of CO2 and CH4 within Zr6O4(OH)(bdc)6 was measured using QENS and matched with results from MD simulations [87] (Fig. 3.10). Importantly, this work established the

concentration dependence of the diffusivities of CH4 and CO2 (self and transport, respectively) within the material. The flexibility of the framework was found to influence significantly the diffusivity of the two species, and CH4 was found to diffuse faster than CO2 over a broad concentration range, a result that is in contrast to zeolites with narrow windows, for which opposite trends were observed. Further analysis of the MD trajectories for CH4 provided insights into the global micro­scopic diffusion-mechanism, proposed to occur by a combination of intracage motions and jump sequences between the material’s tetrahedral and octahedral cages. The coadsorption of CO2 and CH4 in the material, from both a thermody­namic and a kinetic perspective, was also studied using this approach. It was shown that each type of guest adsorbs preferentially in the two different pores, where CO2 occupies the tetrahedral cages and CH4 the octahedral cages. Further, a very unu­sual dynamic behaviour was also noted in the study of CH4/CO2 mixtures in Zr6O4(OH)(bdc)6 in that the slower CO2 molecule was found to enhance the mobility of the faster CH4, again contrasting with the usual observation for CO2/ CH4 mixtures in narrow-window zeolites, where the molecules diffuse indepen­dently or slow the partner species.

The self-diffusion properties of pure CH4 and its binary mixture with CO2 within the NaY zeolite have also been investigated by the combined QENS/MD approach. This material combines several favourable features including a good selectivity, high working capacity, and potential easy regenerability that make it a good can­didate for the selective adsorption of CO2 over CH4 [88]. The QENS measurements at 200 K led to an unexpected self-diffusivity profile for pure CH4 with the presence of a maximum for a loading of 32 CH4/unit cell, which was previously unobserved for the diffusion of an apolar species in a zeolite with large windows. The QENS measurements report only a slight decrease of the self-diffusivity of CH4 in the presence of CO2 when the CO2 loading increases. MD calculations successfully reproduce this experimental trend and suggest a microscopic diffusion-mechanism in the case of this binary mixture [89].

The fundamental steps are:

(1) Determining the PES of the stacked HAT6 molecules and extracting the corresponding equilibrium parameters in terms of co-facial separation (D) and twist angle (0) between the monomers,

(2) Calculating charge-transfer integrals (CTIs) by modelling a defined structural — disorder.

(3) Evaluating how these quantities vary with the imposed conformational fluctuations.

Three degrees of freedom can be modelled in practice at the ab initio level[4] to mimic structural fluctuations: the co-facial separation D, the twist angle 0, and the lateral slide or offset L. The three-dimensional PES highlights (Fig. 6.9) an energy minimum at 0 * 30° and D * 3.5 A, L being constrained to zero. This agrees well with previous work dedicated to similar systems with smaller aliphatic tails. Fol­lowing the estimation of the equilibrium parameters of stacked HAT6 molecules, the next step is to vary these parameters in order to mimic structural disorder, and to estimate how the HAT6-HAT6 interaction is affected. This is achieved by evaluating CTIs[5] which indicates the importance conformational fluctuations in the columnar phase. This is shown in Fig. 6.9 (middle left and right, and bottom). The dependence of the charge transfer J on the twist angle 0 between two stacked HAT6 molecules is evaluated at a fixed distance D =3.5 A, determined from PES calculations.[6] Due to the D3h point-group symmetry, the angular dependence of J is periodic.

When 0 A n f, there is a reduction in point-group symmetry from D3h to C3 and if 0 = n f then the symmetry is lowered from D3h to C3v. Increasing the twist angle from 0 to 60° results in a decrease of the interaction of individual HAT6 molecules. The dependence of the charge transfer J on the co-facial separation D between two stacked HAT6 molecules, with a fixed twist angle, 0 = 30°, is found to decreases exponentially. J increases rapidly with increase of the co-facial separation D until reaching the long-range interaction limit at higher D, which is the monomeric zero

Є

L — >

overlap limit. Therefore the charge-transport description, in terms of only the neighbouring electronic coupling, is adequate.

The third degree of freedom to be investigated is the lateral slide, or offset, L between the columnar stacked HAT6 molecules. Such a fluctuation can notice­ably perturb charge-transfer processes through symmetry breaking, and hence spatial overlap. The offset L is achieved by sliding one HAT6 molecule with respect to the other along a C2 axis. The complex nodal structure (nodes of the wave — function), which is perpendicular to the plane of the large HAT6 molecule, is well reflected through consecutive and damped oscillations of local maxima and minima, which correspond to a constructive/destructive character of the overlap between the individual HAT6 molecules. A higher offset results in a zero overlap, and hence zero charge-transfer. In the presence of dynamic and/or static structural fluctuations the mobility, and therefore the conductivity, should scale approximately quadrati — cally with the charge-transfer integral as has been reported for similar organic materials [19-21]. Thus, the above data can be used to quantitatively obtain charge — transport properties in triphenylene derivatives. It follows that this level of analysis can be extended to more realistic large-scale structural fluctuations, obtained from NPD measurements and MD simulations in the liquid crystalline phase.

Catalysis by Au nanoparticles has attracted considerable attention since it was discovered that Au particles, with a size less than a few nm, are active. Au nano­particles deposited on oxides can dissociate H2 heterolytically on sites involving one Au atom and a nearby surface oxygen atom. This is in contrast with other metals such as M = Ni, Pd, Pt, etc., where H2 is dissociated homolytically (in a symmetric fashion). The relatively poor activity of Au for hydrogenation reactions was attributed partly to the inability of Au to break the H-H bond and to the instability of Au hydrides (Au-H species), while various M-H species have been clearly evidenced by INS.

The INS spectra obtained after chemisorption of H2 on 3-4 nm Au particles supported on nanoparticulate ceria (5 nm) show apeak at400-600 cm-1 accompanied by a broad band from 750 to 1,200 cm-1. The first peak was assigned to bridging hydroxyl groups, and the second band to librational modes of water present on the catalyst surface and resulting from the reaction of hydrogen with oxygens of ceria [7]. The lack of observation of Au-H species by INS can be explained by the low per­centage of Au on the sample (Au loading: 0.48 wt%). On the other hand, the formation of Au-H could be observed by Fourier-transform infrared (FT-IR) spectroscopy, which seems to indicate a greater sensitivity of IR spectroscopy for this sample.