Extracting the electrical energy from the simple reaction between hydrogen and oxygen to produce water is an extremely attractive proposition, which is exactly what a fuel cell does:

2H2 + O2—> 2H2O + electricity + heat

A basic fuel cell consists of an anode and a cathode separated by an electrolyte as shown in Fig. 9.1 of the following chapter. At the anode, hydrogen is separated into protons and electrons, and because the electrolyte only conducts protons, electrons are forced through an external circuit, providing the current to do work. At the cathode, oxygen reacts with protons and electrons to produce water and heat. It is possible to capture the heat for cogeneration, leading to overall efficiencies as high as 90%. Fuel cells are fundamentally more efficient than combustion systems, and efficiencies of 40-50% are gained in today’s applications. We note that the ubiquitous internal combustion engine used in transport has an efficiency below 20%. The main challenge facing fuel cell use is probably system cost, with esti­mates for automotive applications being * $50/kW, compared with the * $30/kW that is required to make fuel-cells favourable over internal combustion. Internal combustion engines are highly developed and matching their durability, reliability, weight and infrastructure with fuel cells is a considerable hurdle to the commercial uptake of fuel cells, but one which can be overcome by understanding the funda­mental limitations of fuel cells.

Fuel cells are typically distinguished by the type of electrolyte used in charge transport. The major classes of fuel cells include: alkaline fuel cells (AFC), solid oxide fuel cells (SOFC), phosphoric acid fuel cells (PAFC), molten carbonate fuel cells (MCFC) and proton-exchange (or polyelectrolyte) membrane fuel cells (PEMFCs). In Chap. 9 we concentrate on solid oxide fuel cells (SOFC) which use ceramic materials as the electrolyte, enabling their operation at high temperatures using a variety of fuels. Higher temperatures are required for adequate diffusion rates of protons and other charged species, which are measurable via neutron scattering, but impose materials problems and significant start-up delays. Chapter 10 presents neutron scattering studies of the operation of PEM fuel cells,

where the aqueous environment provides rapid diffusion of the charged species at much lower temperatures, but poses other challenges. These two fuel cell chapters make use of a wide range of neutron techniques because they are concerned with structure and dynamics over a variety of length scales, from atomic through to macroscopic. The reader is referred to Chap. 1 for an outline of these techniques for a more thorough description. Chapter 9, in particular, is an excellent demonstration of how combining information from a variety of neutron techniques of analysis leads to a more complete understanding of structure/dynamics-function relations.

The first INS study on proton-conducting perovskites was done by Karmonik et al.

[53] who investigated vibrational proton-dynamics in SrCe0 .95M0.05O3S (M — Sc, Ho, Nd). The INS spectra of the materials are reprinted in Fig. 9.8a and reveal O-H wag modes at around 115 (Sc) and 105 (Ho) meV. For the Nd-doped equivalent, the O-H wag band overlaps with a band at * 80 meV. It appears that the frequency of the O-H wag mode shifts to higher wavenumbers with decreasing size of the dopant cation (Nd! Ho! Sc), indicating that this band is related to protons in the vicinity of such atoms [53]. This behaviour was later validated by Yildrim et al.

[54] , who performed lattice-dynamics calculations on a л/2 x л/2 x 1 supercell of SrCeO3, replacing one Ce by Sc + H to give a supercell of Sr8Ce7ScHO24, whose composition is close to that of the real material. In particular, the authors calculated the vibrational spectrum for the hydrogen at the undoped (U) and doped (D) site and by comparing the experimental and calculated spectra (Fig. 9.8b) it could be confirmed that the O-H wag mode at * 120 meV indeed is associated with protons close to dopant (Sc) atoms, whereas the O-H wag mode at lower frequency, *80 meV, is associated with protons in the vicinity of host-lattice Ce atoms [54].

More recently, Karlsson et al. [55] addressed the question of how the O-H wag frequency depends on the dopant concentration. Specifically, the authors performed a systematic INS study of the BaInxZr1_xO3_x/2 (x = 0.20, 0.50, and 0.75) system, which exhibits an average cubic Pm3m symmetry independent of the In concen­tration. The INS spectra are shown in Fig. 9.9a. It can be seen that the O-H wag vibrations show up as a strong, broad, band between approximately 600 and 1,300 cm-1, whilst the peak-fit analysis presented in Fig. 9.9b shows that this band can be decomposed into three Gaussian components. Figure 9.9c shows the In concentration dependence of the relative intensities of the three peak fitted Gaus — sians. A significant redistribution of intensity amongst the three Gaussians as the In concentration is varied can be observed (the total integrated-intensity of the O-H wag band increases linearly with increasing In concentration [55]). Most interest­ingly, there is an increased contribution from the two high-frequency components to the overall spectrum, reflected by a band broadening towards higher frequencies, whereas the width and position of each individual Gaussian are found to be essentially independent of the In concentration [55]. The increase in total intensity of the O-H wag band results from the increasing concentration of protons in the sample, whereas the increased contribution of the high-frequency modes is due to an increased fraction of protons in more or less strongly hydrogen-bonding con­figurations [55]. The formation of strong hydrogen-bonds is believed to be the result of dopant atoms and/or oxygen vacancies in the vicinity of the protons, which act as charged defects, pushing the proton towards a neighbouring oxygen and increasing the tendency for hydrogen-bond formation [55]. However, the presence of such strongly hydrogen-bonding configurations may equally well be the result of tilts and/or rotations of oxygen octahedra induced by doping at the acceptor-atom site, which is of purely static origin [56]. Whatever the case, the formation and breaking of hydrogen bonds are crucially important for long-range proton transport, since proton transfer is a hydrogen-bond mediated process. Thus, information about the nature of hydrogen bonds, which can be derived from the O-H stretch and O-H wag frequencies, and how they link to the structural and dynamical details of the material, is highly valuable. Moreover, O-H stretch and O-H wag mode fre­quencies are useful in computer simulations, where they are utilized as prefactors in transition-state models to estimate the rates of proton transfer and — OH reorien­tational motion, respectively [57].

Further information about the behaviour of protons in the perovskite lattice may be derived from the temperature dependence of the INS spectra. In this regard, Karlsson et al. [55] performed a variable-temperature study on BaZr1-xInxO3_x/2 (x = 0.20). The INS spectra measured at T = 30, 100, 200, and 300 K are shown in Fig. 9.9d. As can be seen the spectra measured at the four different temperatures look essentially the same, which suggests that there is only a small change of the Debye-Waller factor as the temperature is raised from 30 to 300 K, i. e. the total root mean-square displacement, UT, increases only slightly within this temperature range [55]. The weak temperature-dependence of UT indicates that there is no particular difference between the proton dynamics in this material at 30 and 300 K.

Since it is unlikely that the protons undergo long-range diffusion at 30 K, it follows that this is also the case at 300 K. This is in agreement with the generally low proton-conductivity for barium zirconates at these temperatures [26, 52].

A unique method to “see” Li-ion concentration profiles is provided by neutron depth profiling (NDP), Fig. 7.32. Previously it has been shown that NDP is capable of determining Li concentration gradients in optical waveguides [200], electrochromic devices [201], and under ex situ conditions in thin film battery electrodes and electrolytes [202]. NDP uses a neutron-capture reaction for 6Li resulting in:

6Li + nthermal!4 He (2.06MeV) +3 H(2.73MeV) (7.1)

The kinetic energy of the products due to AE = Amc2, where E is the energy, m is the rest mass of the particles and c is the speed of light, is distributed over the tritium (3H) and the alpha particle (4He), while the incoming thermal energy of the neutron at *25 meV, is negligible. Due to the small particle flux and the inher — ently-low interaction of neutrons with matter, NDP is a totally non-destructive technique. When such a capture reaction takes place in a Li-ion battery electrode, the particles produced (4He and 3H) lose part of their kinetic energy due to the scattering by the electrode material, referred to as stopping power. The stopping power is directly related to the composition and density of the electrode and hence is a known quantity. Therefore, by measuring the energy of the 4He and 3H ions when they exit from the electrode, the depth of the capture reaction can be reconstructed. Typically, the spatial resolution of NDP for well-defined homoge­neous layers is on the order of tens of nano-meters, Currently, the main restrictions of the NDP technique is the maximum depth that can be probed and the time resolution which, depending on the material investigated and the in situ cell design, are approximately 5-50 microns and 10-20 min, respectively. Ideal solid-state batteries can be designed with high spatial homogeneity for initial experiments, before proceeding to more complicated systems.

NPD has only been applied occasionally to Li-ion battery research, however, in these ex situ studies [202-205] NDP is very powerful in identifying Li-ion transport and aging mechanisms. The possibilities of NDP in Li-ion battery research is demonstrated with the first in situ study on thin film solid-state batteries probing the kinetic processes in these Li-ion batteries.

Oudenhoven et al. brought NDP one step further, demonstrating that Li depth profiles can be measured in situ in an all solid-state micro battery system during (dis)charging [206]. The Li-ion distribution was studied in a thin film solid state battery stack containing a monocrystalline Si substrate with a 200 nm Pt current

collector, 500 nm LiCoO2 positive electrode, 1.5 ^ N-doped Li3PO4 (LiPON) electrolyte, and a 150 nm Cu top current-collector. The basic setup of the experi­ment is shown in Fig. 7.33. By subtracting the NDP spectrum of the as-prepared electrode from the charged and discharged spectra, the changes in Li-distributions can be observed directly, see Fig. 7.33. Upon charging, Li in the positive LiCoO2 electrode is depleted and increased at the negative Cu current collector.

The development of large concentration-gradients in both the LiCoO2 electrode and the LiPON solid electrolyte, Fig. 7.34, reveals that in this system ionic transport in both electrolyte and electrode limit the overall charge-rate. The cathode was enriched with 6Li to highlight the redistribution of 6Li and the natural abundance of 6,7Li in the electrolyte during time-dependent experiments. In this case, the NDP intensity increases by approximately a factor of 13. Apart from being able to observe

Energy (keV)

where 6Li is going, the expected diffusive equilibration of the 6Li concentration within the electrolyte was observed during a 2 h equilibration period. Interestingly, the enriched Li remains in the LiCoO2 electrode, even though the exchange current that establishes the dynamic equilibrium would be expected to redistribute the 6Li equally throughout the LiCoO2 electrode and the LiPON electrolyte. The absence of vacancies at the initial stage probably makes the exchange-current extremely small.

As the battery is charged at 0.5 °C, this results in a large decrease in the Li-ion signal of the LiCoO2 electrode, as shown in Fig. 7.34.

The stronger decrease in Li-ion signal near the interface with the electrolyte suggests an inhomogeneous Li-ion distribution in the electrode. Although this may be the case, a redistribution of the 6Li ions due to exchange with the electrolyte will lead to a lower Li-ion signal in the electrode. That this is indeed part of the explanation is clear from the almost 70 % decrease in Li-ion signal. This decrease is more than would be expected under the mild electrochemical conditions that should lead to Li05-CoO2, and hence at most a factor of two decrease in the Li-ion signal is to be expected. However, the inhomogeneous signal from the electrode indicated that the exchange does not reach the back part of the electrode that is closest to the current collector. The inhomogeneous distribution of the Li-ion signal originating from the electrolyte indicates the presence of an inhomogeneous 6Li and Li-ion distribution. The evolution of this non-equilibrium situation was investigated by relaxing the system after charging during a period of 2 h and taking NDP spectra, shown in Fig. 7.34. After 2 h the 6Li gradient almost vanishes in the electrolyte, whereas it remains in the LiCoO2 electrode. Clearly, Li-ions are much more mobile in the electrolyte compared to in the electrode. The work of Oudenhoven et al. shows for the first time that the evolution of the Li distribution and gradient under dynamic conditions can be studied.

While it is beyond the scope of this chapter to give an extensive review of the field of PEMs, it is necessary that reader be aware of the types of materials that are being developed for fuel-cell applications. Regardless of the particular application (e. g., stationary, portable, or automotive power), these materials must exhibit a certain set of chemical and physical properties that are critical for optimal fuel-cell perfor­mance. Any material being used in a fuel cell must exhibit a list of properties including, but not limited to (1) high proton-conductivity (i. e., a good electrolyte), (2) negligible electrical-conductivity, (3) permeability to ions, but allow only one type of charge, (4) resistance to permeation of uncharged gases, (5) variable membrane-area and thickness, and (6) good mechanical strength. Furthermore, the membrane must be of reasonable cost and durability. Ultimately, it is the polymer chemistry and microstructure that give rise to the macroscopic performance prop­erties that are desired. A range of synthetic approaches have yielded materials that include, but are not limited to, poly(perflourosulfonic acid)s (PFSA), ion-containing polystyrene derivatives, polyarylene ethers, polysulfones, polyimides, and ion — containing block copolymers. There have been extensive studies on each of these classes of materials, but to date neutron techniques have been primarily used to study poly(perflourosulfonic acid)s, namely Nafion® [2]. Therefore, it is necessary to give a brief background on this particular material. Throughout this chapter, where specific breakthroughs have been made, studies involving other PEM materials using neutron techniques will be highlighted. On the whole, however, the overlap of neutron measurements and PEM materials has been dominated by PFSAs.

The most widely studied PFSA, Nafion®, is a product of the E. I. Dupont Chemical Company having the structure given below.

The polar perfluoroether side-chains containing the ionic sulfonate-groups have been shown to organize into aggregates, thus leading to a nanophase-separated mor­phology where the ionic domains, termed clusters [3], are distributed throughout the non-polar polytetrafluoroethylene (PTFE) matrix. In addition, the runs of tetrafl — uroethylene, of sufficient length, are capable of organizing into crystalline domains having unit-cell dimensions virtually identical to that of pure PTFE [4, 5]. The degree of crystallinity in PFSIs is generally less than ca. 10 % as a mass fraction in 1,100 equivalent-weight Nafion® (EW, the grams of dry polymer per equivalent number of SO3- groups) and has been shown to vary with EW. The complex, phase — separated morphology, consisting of crystalline, amorphous, and ionic domains, of PFSIs has been the focus of several investigations [3, 5-16]. Over the last 50 years, a wide variety of studies involving Nafion® have aimed to relate the thermal, mechanical, and fuel-cell performance properties (i. e., transport, ionic conductivity, and dielectric behaviour) to specific morphological features. Many of these studies involve neutron-scattering techniques and will be discussed in detail below.

Maths Karlsson

Abstract This chapter aims to demonstrate the important role that neutron scattering now plays in advancing the current understanding of the basic properties of proton-conducting ceramic separator-materials for future intermediate-temperature fuel cells. In particular, the breadth of contemporary neutron scattering work on proton-conducting perovskite-type oxides, hydrated alkali thio-hydroxogermanates, solid acids, and gallium-based oxides, is highlighted to illustrate the range of information that can be obtained. Crucial materials properties that are examined include crystal structure, proton sites, hydrogen bonding interactions, proton dynamics, proton concentrations, and nanoionics. Furthermore, the prospectives for future neutron studies within this field, particularly in view of the latest developments of neutron methods and the advent of new sources and their combination with other techniques, are discussed.

To obtain the local structure in glassy, nano, disordered and amorphous materials, having unresolved, weak or broad signals, neutron total scattering, or/and inelastic neutron scattering (INS) or quasielastic neutron scattering (QENS) can be powerful tools. Total scattering allows extraction of the local structure in terms of interatomic distances, bond angles and coordination numbers. In this case scattering is detected over a wider Q-range and short-range interactions of a sample are probed and modelled. In addition, rotations and vibrations picked up by INS and QENS are very sensitive to local distortions and allow otherwise difficult to detect relevant species such as protons, OH, water and in rare cases Li to be studied.

The negative-anode carbon is a good example of where neutron total scattering, in conjunction with other neutron-based methods, has been able to quantify important, previously ill-defined, aspects of the material’s function [20, 123, 124], as demonstrated with a range of low-crystallinity C negative electrodes. Addi­tionally, C-based anodes can be analysed in the lithiated and delithiated states and over the course of phase transitions. Typical neutron total-scattering data for graphite is presented as a radial distribution function (or pair distribution function) as illustrated in Fig. 7.10, and peak positions are indicative of interatomic distances. Often, neutron total scattering is combined with INS data to provide supporting information concerning the short-range order in C. For further information on total scattering the reader is directed to a review on the structure and dynamics of ionic liquids [125], and total scattering is likely to become increasingly used as the range of nano-sized active electrode materials increase.

Total scattering and INS are particularly attractive for disordered C where conventional diffraction provides limited information and more generally for Li arrangements in C. Disordered C where a large amount of H is present can exhibit significant Li capacity (one ‘excess’ Li per H) and studies have investigated how Li is taken into these materials [123, 124, 126]. Studies have shown that these materials exhibit randomly-arranged graphene fragments of different sizes with edges terminated by a single H, similar to Si with H at the surface. The spectra also contain a boson peak, an indicator of disorder, and distinct similarities to polycyclic aromatic hydrocarbon (PAH) spectra exist, some of which feature two or three edge-terminating H. Additionally, comparison with PAH spectra allowed the determination of methyl groupings when higher H concentrations are used. The boson peak is at the same position in samples with different concentrations of H and changes in position and intensity with Li insertion. This shows that the Li interacts with the C environment, contrary to the idea of Li accumulation in voids. These findings agree with two models of Li insertion: One where Li resides on both sides of the graphene layers (the so-called ‘house-of-cards’ model) and the other where Li is bonded to the H-terminated C at the edge of the graphene layers (and reside in interstitial sites). INS data also illustrate that the Li-Li interlayer and intralayer

interactions are comparable in strength. Computer modelling showed that there is insignificant energy difference between interstitial Li and those that are bonded to the terminal H. Other models include the formation of covalent Li2 molecules, but no evidence was found in support of these. The key aspect in these studies is that all models satisfy the observed capacity of LiC6. Finally, QENS [124] was used to show that Li jumps between nearest or second-nearest neighbour interstitial sites.

Related work investigated the entropy of intercalation into C [127]. This study shows how the sign of entropy changes from low Li concentrations on initial charge (x < 0.2 in LixC6) to higher concentrations (x > 0.2) indicating that multiple pro­cesses are occurring and that one of these is vibrational in origin. In graphite the entropy remains negative, but reduces in magnitude as lithiation progresses. Similar entropy information from INS data during lithiation of LiCoO2 cathodes has also been reported [128].

Cathode materials pertinent to Li-ion batteries based on olivine LiMPO4 have also been probed with INS, but for magnetic properties (low temperature) rather than Li-ion diffusion or lattice dynamic studies. Studies of LiFePO4 [129], LiNi1-x-FexPO4 [130], and LiMnPO4 [131], show spin-wave dispersions and allow characterisation of magnetic-exchange interactions. Further INS work was moti­vated by the need to understand the electronic conductivity in LiFePO4 and probed the thermodynamics and vibrational entropy of the phase transition in Li0 6FePO4 [121]. The oxidation state of Fe influences its neighbouring O atoms and the polyhedral distortions can characterize the motion of carrier hopping between Fe sites, which results in relaxations or displacements that can in turn be considered as the sum of longitudinal phonons. Similarly, occupation or vacancy of Li can result in distortions of atom positions and are expected to alter the frequency of phonons, in particular longitudinal optical phonons.

The phase evolution of Li0.6FePO4 as a function of temperature, via a two-phase transition to a disordered solid-solution transition at 200 °C [121], can shed light on the reaction mechanism during charge/discharge of this cathode. This is particularly pertinent as the two-phase or solid-solution mechanism of LiFePO4 is a topical issue as discussed above. The difference in two-phase and solid-solution LiFePO4 optical modes above 100 meV (higher energies) was found, with broadening evi­dent for the solid-solution sample. The low-energy region features mostly acoustic lattice modes, translations and librations of PO4 and translations of Fe. By com­parison with infrared (IR) and Raman data, it was found that the PO4 stretching vibrations are damped in the solid-solution sample. The difference in INS data of solid-solution and two-phase samples at higher energy mostly involve optical modes that can arise from motion of Li-ions, charge hopping between Fe-ions, and heterogeneities. The entropy was found to be larger in the solid-solution phase in conjunction with the subtle differences in the dynamics due to different optical modes. The similarity in two-phase and solid-solution phonon density of states (Fig. 7.11) agrees with the ease with which LiFePO4 seems to undergo either transition, and the difficulty in pinning down the experimental evidence related to the reaction-mechanism evolution.

Arguably the most studied materials using INS are manganese oxides and lithiated manganese oxides, predominantly due to the ease of using H as a probe for Li. These compounds are used for both primary and Li-ion batteries and ion-exchange methods have been used to show where Li may reside in these compounds. Although indirect, this information can provide further answers to some of the problems in this field of research. Attempts are also being made to use INS to provide comparisons between H and Li where H is used as a calibrated probe for Li [132].

One approach is to replace structural or surface water present on manganese oxides with protons, which can in turn be exchanged for Li to see how Li might displace water in these compounds. This was undertaken for spinel Li133-x/3CoJ1Mn;L67-2x/3O4 [133] which shows, as is the case in many compounds of this family, that protons are inserted as hydroxyl groups giving a strong incoherent INS signal. The hydroxyl groups are located on the O atoms neighbouring the vacant 16d sites and aligned with the 8a sites in the spinel structure. Conversely, studies on undoped spinels have shown that the Li extraction from the 16d sites allows the insertion of protons. The main features of the INS spectra are strong y(OH) modes, ahighly ordered proton site, a shoulder and smaller features between 300-700 cm 1 showing riding of protons on the oxide lattice and some librational water modes. The hydroxyl groups have characteristic signals around 908 cm-1 and their orientations are also determined using INS [134-137] of spinel-derivative compounds. Interestingly, IR data shows features between 950 and 1300 cm-1 which were considered to arise from protons, but the absence of these features at corresponding frequencies in the INS data indicate a manganese oxide lattice origin. Notable discoveries of this and related studies include the finding that in undoped spinels 40 % of protons cannot be exchanged and form disordered water, the chemical re-insertion of Li in Li-rich spinel Li16Mn16O4 removes most of the hydroxyl groups [137], that generally the reversible Li amount is 50 % in both undoped and doped spinels, and that fewer protons are re-exchanged as

the Co concentration increases. The latter is an interesting way to tune the Li-proton exchange capacity of these materials.

Figure 7.12 shows the INS data from a series of Li-rich Li1.6Mn16O4 spinels formed through various methods. The pure sample (bottom of Fig. 7.12) shows some evidence of protons, OH and water, whilst the acid-treated version, where acid results in H-Li exchange, shows strong characteristic peaks for protons and y(OH) groups. Finally, the acid-treated sample undergoes a chemical Li re-insertion step and results in the loss of the proton and OH signatures. However, the re­inserted material does not replicate the pure sample suggesting some protons remain as structural water and hydroxyl groups [137]. Relative comparisons of the INS intensity can be made between the acid-treated and re-inserted samples, with the 909 cm 1 peak showing a larger drop in intensity compared to the 1,087 cm 1 peak, which is attributed to an H site being easier to depopulate. A comparison of INS data for two Li-rich variants, Li133Mn167O4 and Li16Mn16O4, shows that the proton stability is higher in Li16Mn16O4 than in Li133Mn167O4. This suggests the reason that Li16Mn16O4 has a larger Li-ion exchange capacity than Li133Mn1 .67O4 concerns the stability of the inserted species (or more specifically the stabilized proton sites).

Studies of ^-MnO2 [135] illustrate subtle differences in INS spectra depending on synthesis precursors, noting that precursors and conditions are both important. This work again highlights the need to focus on the protons (often disordered). A related study investigated proton-exchanged spinels that form ^-MnO2 showing that the proton diffusion was dependent on octahedral Mn vacancies [136]. In this study, certain features in the INS spectrum were found to disappear in the highly crys­talline sample, suggesting that motion can be perturbed with crystallinity. Researchers have also looked at the proton and water environments in bare and lithiated MnO2 [132] to demonstrate how lithiation influences the proton and water motions, which can then be used to extract information on lithiation processes. Using neutron total scattering from oxidized and lithiated versions of ^-MnO2 researchers derived models for oxidation and lithiation [138].

Further work on the spinel LiMn2O4 system investigated the cubic to ortho­rhombic phase transition near room temperature, which is associated with Mn3+/ Mn4+ charge ordering [139]. Excess Li was introduced at the 16c site to study why the phase transition is suppressed in this situation. QENS was used here, where data were found to be dominated by magnetic contributions rather than that from Li hopping, with the slight narrowing of the elastic line near room temperature leading to the preliminary conclusion that electrons are localized on the Mn. A dynamic transition in Li-rich compounds seems to coincide with the structural transition in the parent. The magnetic properties of Li096Mn2O4 were explored in a related INS study [140] showing that two short-range magnetic transitions are present and related to spin ordering of Mn3+ and Mn4+.

Quasielastic neutron scattering (QENS) has played a central role in understanding the nature of proton dynamics in proton-conducting perovskites. The usefulness of the technique comes from the fact that it gives access to the relevant time — and length — scales on which the atomic-scale dynamics of protons typically occur. In addition, the very large neutron-scattering cross section of protons provides a good contrast in experiments and enables studies of systems containing only small amounts of protons.

Neutron imaging (radiography) is becoming increasingly important in the study of Li-ion batteries as the spatial and temporal resolution of the detectors continually improve, and more advanced computational methods allow tomographic and/or three-dimensional rendering [207]. Neutron radiography (NR) is used to show macroscopic information concerning the Li distribution within a Li-ion battery, and in some cases while a process is occurring or at different states-of-charge [208]. Additionally, the H distribution in the electrolyte can be probed [209]. Examples of such studies include the Li distribution at the charged state versus the discharged state, during high-temperature battery operation, during fast charge/discharge cycling, and during overcharging [207, 208, 210, 211].

Neutron imaging has also been used to study alkaline [212,213] and Li-air batteries [214]. The future for neutron radiography relies on new instruments with improved spatial resolution, but also temporal resolution to allow time-resolved in situ exper­iments. Another method under considerable investigation is the combination of dif­fraction and imaging, which requires the definition of a gauge volume which is imaged and from which diffraction data can also be collected. This has been demonstrated for physically-larger batteries such as Na metal halide batteries, which usually have larger electrodes [215]. However, to be pertinent for Li-ion battery research, the gauge volume has to be reduced to become comparable to the thickness of electrode layers.

Early work on imaging Li-ion batteries explored the different types of battery construction, e. g. coin, prismatic, and cylindrical cells, and the distribution of Li at various battery states or during charge/discharge [216]. Figure 7.35 (left) shows images of a coin cell (CR1220 from Panasonic) from the charged to the discharged state, where lighter (white) regions at the charged state correspond to the Li anode (Li metal) and electrolyte (arrow). Over the course of discharge the Li-ions move towards the cathode (MnO2) resulting in an even distribution of white regions. The authors comment that if standard components and standardised cells are constructed then a more quantitative description of the Li distribution can be made. They also explored charging rates and other constructions, some of which had further experimental difficulties due to the internal structure of the batteries and the need to account for absorption by various layers. An example of the same electrode chemistry in the cylindrical case (CR1/3-1H) is shown in Fig. 7.35 (right) before and after discharge [158] showing similar Li distributions at the charged and discharged states.

In commercial batteries overcharging can be a potentially-devastating failure mechanism, and imaging studies on commercial graphite//LiNi0.8Co0 .15Al0.05O2 (NCA) batteries show what is deposited on the graphite anode during overcharge [208]. By performing an in situ measurement the deposition of a material on the graphite anode was studied (Fig. 7.36) and later determined to be Li. In addition, the authors were able to characterize ‘where’ the Li deposits during battery processes. Another work explored ‘fresh’ and ‘fatigued’ batteries, where batteries that had been cycled 200 times. The 18650 cylindrical batteries showed no differences at the macroscopic level in the neutron images of between fatigued and fresh batteries [182], even though neutron diffraction data indicated less Li insertion in fatigued graphite.

Another study illustrated that a 14 pm spatial resolution is attainable for battery samples using neutron imaging [210]. This work used a purpose-built graphite — containing cell to quantify Li content during charge/discharge and the residual Li content after each cycle, showing quantification over several cycles (e. g. capacity loss). Figure 7.37 shows the evolution of Li content and its distribution in graphite during the first discharge. The Li distribution was compared during cycling and between cycles. A slight difference in Li content between the separator and current collector was found. Further work investigated LiFePO4llgraphite pouch-cells and revealed Li concentration gradients across electrodes and in their bent regions [211]. Figure 7.38 shows the distribution of Li in the layers of the pouch cell at various states-of-charge. The authors used the Beer-Lambert law to correlate colour gradients, shown in Fig. 7.38, to the Li concentration. One advantage of using a pouch cell is that one image contains many layers, so an increase in electrode thickness can be seen in multiple layers verifying the result (as can Li concentration gradients). Clearly, this information can direct the development of better per­forming electrodes.

More recent work used cold neutrons rather than thermal neutrons, harnessing the stronger interaction of colder neutrons with matter, to visualize Li-ion distributions in Li-I batteries used in pacemakers [217]. This work also used three-dimensional imaging (tomography) and discussed methods to improve the signal-to-noise ratio in the images. The authors collected 50 images at 0.3 s for each angular step (rotation) of 0.91° which were then used to construct the three-dimemsional image. Figure 7.39 is an example of a cross section of a neutron tomography image of the fresh battery (left) and after a certain period of discharge (right). The battery is made of plates of Li and I. An unexpected change in the Li distribution (white) was observed in this study, where the smooth distribution became highly irregular.

The irregularity of the Li distribution after discharge is extracted in the three­dimensional image shown in Fig. 7.39, where the Li formations are clearly seen.

Further experimentation was undertaken on a Li-ion polymer battery using monochromatic imaging with cold neutrons, specifically targeting the anode and the processes that occur within it [218]. The LiC6 compound, but no evidence of the staging phenomenon often observed in LixC6 anodes was seen. This work was the first real-time in situ imaging of a commercial Li-ion battery, with some results shown in Fig. 7.40.

A subset of radiography research using commercial batteries, and in some cases custom-made batteries, is the study of gas evolution [209]. One in situ study showed how excess electrolyte present in batteries is consumed in the first charge-cycle, resulting in the formation of the SEI layer and some volume expansion. Additionally, gases were found to be evolved during the first charge. Figure 7.41 shows the consumption of excess electrolyte in these cells. The authors were also able to approximate the amount of expansion and contraction of the electrodes indirectly.

In situ neutron radiography has been extensively used to study the interface between graphite and a range of gel-based electrolytes [219]. By using this tech­nique, the generation of gas bubbles in the first charge can be visualized and quantified. This information allows the best electrolyte to be proposed, noting that the generation of gas bubbles, particularly on graphite surfaces, leads to perfor­mance degradation. This measurement also provided information on the spatial distribution and kinetic evolution of gas bubbles, as well as the electrolyte dis­placement and volume expansion in graphite. In order to undertake these mea­surements, specialised cells were developed. A neutron image of the cell is shown in Fig. 7.42. For the in situ experiment the exposure time for each image was 20 s and an image was recorded every 2 min. Figure 7.42 shows how channels of gases are formed seemingly-randomly in the cell and their evolution at different times. It was found that LiC6 is formed only where gas emission is absent, illustrating some heterogeneities in the charge distribution and electrode composition. The gel-based electrolytes tested in this study show less gas evolution (3 %) compared to liquid — based electrolytes (60 %) and this was related to the smaller amount of gas evo­lution during the first cycle.

7.7 Perspectives

This chapter has aimed at demonstrating how neutron-scattering methods allow researchers to elucidate crucial structural and kinetic properties of electrodes, electrolytes, and complete batteries. Neutron-scattering techniques play a key role in the development of new materials by relating structure to functional properties. Future battery research and development will in particular profit from the advances in in situ neutron-scattering techniques, probing complete battery systems. This gives the opportunity to relate battery performance to material and electrode structural, morphological, and dynamic properties under non-equilibrium and ageing conditions, which is vital information for the design of future batteries.

10.3.1 Nanoscale Membrane Structure

By and large, the neutron technique most used to study PEM materials has been SANS. It is understood that the nanostructure of PEMs plays a significant role in water uptake and transport, which are materials with properties vital to the performance of a working fuel-cell. However, in order to design a material with desired features, one must have a detailed understanding of the interplay between the nanostructure of the membrane and the overall performance properties. SANS is a versatile tool in elu­cidating the structure of a variety of membrane materials and can also be used to study transport, which will be discussed in the section on water transport.

As previously mentioned, Nafion® has been the most widely-studied PEM material to date and SANS has been employed extensively to study the nanoscale structure of this complex material [8, 10, 11, 14-23]. The earliest structural studies of Nafion® utilizing SANS and small-angle X-ray scattering (SAXS) revealed a broad peak at a Q value between 0.1 and 0.2 A 1, called the ionomer peak, which has been attributed to the correlation between the nanophase-separated ionic domains, termed clusters. The crystalline component contributes to the scattering at multiple length scales including peaks in the Q range 0.6-2.0 A-1, owing to the structure of the amorphous and local crystalline-lattice, in addition to a broad peak entered at lower values of Q (^0.05 A 1), which is related to the inter-crystalline scattering (known as the long period). Moreover, ultra-small-angle scattering reveals an upturn that can be associated with large-scale heterogeneities. The scattering for hydrated Nafion® over a wide range of length-scales can be seen in a review by Gebel and Diat [19, 24].

An example of the scattering using neutrons can be seen in Fig. 10.1, for Nafion® films cast from a dispersion, annealed between 80 and 180 °C, and equilibrated in liquid water. When annealed below the alpha relaxation temperature (Ta) of Nafion® (^ 100 °C), the lack of a peak at low Q values is evidence that there is no apparent long-range crystalline order in the film. Above Ta, however, one observes a peak due to long-range crystalline order and a shift in the crystalline peak to lower Q values with increasing annealing temperature, indicating that higher annealing-temperatures result in larger, more widely separated crystalline domains. An analysis of the scattering curves can be seen in the inset in Fig. 10.1. Clearly, the spacing of the crystallites increases with increasing annealing tem­perature. Moreover, the spacing between the ionic domains decreases with increasing annealing temperature. It is known that the crystalline structure plays an integral part in the mechanical stability and durability of fuel-cell membranes, but these data also reveal the relationship between the crystalline structure and water uptake. The decreased spacing between the ionic domains and the lower incoherent — background with increasing annealing temperature are evidence that these annealed films have a lower water-uptake. In these materials water retention must be bal­anced with annealing temperature in order to achieve desirable proton-conductivity and mechanical integrity. This is just one example of how SANS can be used to probe structure-processing-property relationships.

Over the decades, development and advancement of state-of-the-art scattering techniques were able to reveal the many scattering features, over multiple length — scales, of Nafion® and other perfluorosulfonic acid membranes. As a result, there has been a progression in the complexity and characteristics of the many mor­phological models that have been proposed to explain the observed scattering in effort to gain a deeper fundamental insight into how the structure is related to

material performance. The proposed models generated have included a variety of structural units and range from the earliest spherical cluster-network model to other models including lamellar, sandwich-like, fringed micelle, rod-like, and ribbon-like, as well as a model which includes cylindrical water channels. Each of these models was able to account for the scattering to some degree reasonably well, making it difficult to discern which most accurately describes the morphology. Of course, the models able to capture the many scattering-features over a large Q range, with physically relevant fitting parameters are of the highest value. For the details of the structural models of Nafion® and their respective parameters, the reader is directed to the extensive literature concerning this topic.

In general, scattering techniques provide an excellent way to characterize the global structure of fuel-cell membranes [19, 24]. For polyelectrolytes such as Nafion® there are scattering features that are ubiquitous and considered to arise from favourable structures for fuel-cell membrane application, although that con­ventional wisdom is now being called into question by more recent studies. Typically, polyelectrolyte fuel-cell membranes contain ionic moieties that are able to conduct protons or hydroxide ions in alkaline fuel-cell membranes. These ion — containing, polar-groups phase separate from the more hydrophobic components of the polymer and can form ion-conducting channels which are responsible for ion transport. Quite often these ionic domains give rise to a scattering peak in SANS and, if other hierarchical structures are present, other scattering features are observed. This is especially true when block copolymers are used to provide a structural basis for the membrane. A recent review by Elabd and Hickner [25] has evaluated the state-of-the-art block-copolymer membranes by leveraging the self — assembled nanostructure of block copolymers as a template for creating well — defined transport pathways for use in fuel cells.

In addition to the work on Nafion®, there is a rich body of literature in which SANS has been used to probe the structure of a variety of PEM materials including sulfonated polyimides [26], sulfonated polyetherketones [24, 27], sulfonated tri — fluorostyrenes [28], poly(styrenesulfonic acid)-grafted cross-liked polytetrafluoro — ethylene [29], and a host of other materials [30-34]. Ultimately, one seeks to understand the role that molecular-level structure and chemistry play in the development of material nanostructure and how this nanostructure is correlated with performance properties such as water content and transport, as well as ion con­ductivity. For example, in the work by Iwase et al. SANS (in conjunction with SAXS) was used to investigate the hierarchical structure of graft-type PEMs syn­thesized from cross-linked PTFE [29]. The structure was studied over a large range of length scales (0.6 nm to 1.6 pm) as a function of the degree of grafting, Xg. It was determined that the structure of these materials consisted of conducting layers of polystyrene sulfonic acid (the grafted domains) arranged in lamellar stacks on the surface of the PTFE crystallites. Within the conducting layers, they observed scattering features consistent with correlations between sulfonic acid domains. With less than 15 % grafting the grafted domains were found to reside mainly in the amorphous domains between the PTFE crystalline lamellae. Within this regime, the lamellar spacing increased with increasing grafting content up to a value of Xg of about 5 % and remained constant until 15 %. Above 15 % the grafting domains appeared to phase separate from the hydrophobic matrix and become contiguous, thus forming a highly conductive domain around the crystallites.

While X-ray scattering is certainly more widely accessible for structural char­acterization of membrane materials, neutrons offer the unique benefit of contrast variation, or contrast matching, in the structural determination of systems with complex architectures. Owing to the large differences in scattering-length (SL) between deuterium and hydrogen, one can use isotopic replacement in the polymer, or the solvent, to highlight the scattering from various structural components or phases. One example of this can be found in the work by Gebel et al. [35] in which they used various mixtures of D2O/H2O to swell N(CH3)+-neutralized forms of Nafion® as a way of elucidating the nature of the scattering entities in these hydrated films. By varying the ratio of D2O to H2O and normalizing by the scat­tering of Nafion® in pure H2O they were able to match out the structural component of the scattering due to Nafion® and to observe the counterion condensation at the interface between the hydrophobic components of the polymer and the hydrophilic water domains. This was the first measurement of condensation in a perfluoro- sulfonated ionomer. Using contrast variation to explore neutralized forms, different models could be applied to determine which structure accurately described the scattering curves. While this study was unable to determine the shape of the scattering particles (i. e., spherical or rod-like), it was determined that the features were aggregates of the polymer backbone surrounded by the electrolyte solution, as opposed to the scattering particles being cavities filled with the electrolyte solution.

Recently, a series of studies using in situ SANS, among other techniques, on block-copolymer electrolyte membranes consisting of a polymethylbutylene (PMB) block and polystyrenesulfonate (PSS) block have begun to call into question whether or not ionic aggregates are necessary for effective proton-transport, espe­cially in the presence of structures established by block-copolymer morphology [33, 36, 37]. The composition of the block copolymer was varied in order to tune the size of the domains. Also varied was the degree of sulfonation of the polystyrene block, within a particular composition. This body of work represents an excellent example of the application of neutron scattering to elucidate the structure-property relationships of PEM materials in environments that are application relevant. These in situ measurements were achieved using a specially designed sample chamber at the National Institute of Standards and Technology (NIST) Center for Neutron Research wherein the humidity and temperature of the environment surrounding the sample could be controlled. Moreover, the water reservoir within the sample chamber could be filled with various mixtures of H2O and D2O, allowing contrast — variation experiments to be performed. The scattering was measured over a range of relative humidity and temperature values using D2O. For one particular block copolymer composition, the scattering indicated the presence of a hexagonal phase over the entire range of relative humidity and temperature values studies (Fig. 10.2). At low temperatures (25 °C) and humidity (^25 %), the scattering arises from the block-copolymer morphology. However, at 95 % relative humidity and 40 °C, a shoulder at approximately Q =1.8 nm-1 was observed, which became more pro­nounced and intense upon further heating and humidification. It was acknowledged that this peak was similar to that of the ionomer peak observed in Nafion®, but was referred to as the ‘water peak’ as it was only visible upon hydration. The water domain-spacing was taken to be 2n/Qmax of the peak. For a given block-copolymer system under the same environmental conditions (relative humidity = 95 % and at 60 °C) the position of the water peak was shown to shift to higher values of Q with increasing levels of sulfonation. This result was attributed to a decrease in the average distance between sulfonate groups upon increasing sulfonation. A contrast — variation study was performed to determine the origin of the water peak. The water reservoir was filled with a volumetric mixture of D2O/H2O of 32/68, chosen to

match the scattering-length density (SLD) of the dry PSS block. The water peak was shown to disappear when the samples were humidified with this mixture, indicating that the peak arises due to the presence of a substructure within the PSS superstructure, most likely a heterogeneous distribution of water-rich and water — poor domains as is found in most polystyrene ionomer systems.

One of the most important observations of this study was the absence of the water peak when the size of the hydrophilic domains was below a critical thickness (Fig. 10.3). For PSS-PMB copolymers this critical thickness was on the order of 6 nm to 10 nm for a sulfonation level of about 47 mol percent. It was determined that the water-rich domains were effectively homogenized due to confinement effects. This SANS work has played a critical supporting role in determining the molecular and morphological origins for the enhanced water retention and proton transport observed in the study of these copolymer systems. These results have provided a new perspective in the strategy for developing materials for use in PEM fuel cells.

A particularly promising, yet challenging, clean-energy technology is the solid oxide fuel-cell (SOFC) [1-9]. At the heart of this device is an oxide-ion or proton­conducting oxide electrolyte, which is sandwiched between two porous electrodes (the anode and cathode). The working principle of the SOFC is based on the chemical reaction of hydrogen (at the anode) and oxygen (at the cathode) to pro­duce electricity and water, working temperatures being in the range * 800-1,000 °C. The electrolyte does not conduct electrons, but is permeable to the diffusion of either oxide ions or protons. Schematics of SOFCs based on oxide-ion and proton-conducting electrolytes are shown in Fig. 9.1.

M. Karlsson (H)

Chalmers University of Technology, 412 96 Goteborg, Sweden e-mail: maths. karlsson@chalmers. se

© Springer International Publishing Switzerland 2015 243

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_9

In the 1990s the Siemens Westinghouse SOFC, consisting of an oxide-ion con­ducting yttria-stabilized zirconia (YSZ) electrolyte, a lanthanum strontium manga — nite cathode, and a nickel-YSZ cermet anode, in a unique tubular design, was developed and is now commercially available [1, 10, 11]. Thanks to its high oper­ation temperature, which offers the advantages of fuel flexibility (existing fossil fuels can be used) and high energy-conversion efficiency, the SOFC is particularly attractive for use in combined heat and power applications or efficiently coupled with gas turbines. However, the high operation-temperature also has disadvantages, such as a long startup time, durability issues, and the need for relatively expensive component materials. Therefore, in more recent years, much research has focused on trying to reduce the operation temperature of the SOFC to the so-called intermediate temperature range between approximately 200 and 500 °C. Such a reduction in operating temperature would have a beneficial impact on the total cost and the durability of the fuel cell, as problems associated with thermal cycling and perfor­mance degradation would be reduced and it would be possible to use cheaper materials in interconnects and heat exchangers [12]. Lowering the temperature would also shorten the startup time of the fuel cell, which is of particular importance for mobile applications. Further, in comparison to low temperature (<100 °C) polymer-electrolyte membrane fuel cells (PEMFCs) [13], a temperature of 200 °C is high enough in order to allow for the use of smaller amounts of expensive platinum to catalyze the electrode reactions [12]. In fact, intermediate-temperature fuel-cell technology has the unique potential to be used both for stationary combined heat and power systems and for hybrid and plug-in hybrid vehicles and are indeed expected to produce energy densities per volume and specific energy per weight significantly larger than state-of-the-art Li ion and Ni metal hydride batteries [14].

However, it is discouraging that any decrease of the operation temperature of SOFCs results in a decreased power density of the device, mainly due to a lowering of the ionic conductivity of the electrolyte. Targeted conductivities exceed

7TC

800 500 400 300 200

1000/r (K ’)

Fig. 9.2 Conductivity of state-of-the-art electrolytes over a wide temperature range. Note the gap of highly-conducting materials in the temperature range * 100-500 °C. The figure is redrawn and modified from Ref. [15], copyright Wiley, 2002. The conductivity of Cs2GeS2(OH)2^yH2O is taken from Ref. [16]

0.01 Scm V and in this respect proton-conducting ceramics emerge as the main candidate electrolytes for SOFCs. However, despite intense research, the conduc­tivities of even the best proton-conducting ceramics are still one to two orders of magnitude below the target. This can be appreciated from Fig. 9.2, which displays the conductivities for the state-of-the-art ionic conductors over a wide temperature range, including electrolytes for low temperature PEMFCs and high-temperature SOFCs, i. e. the more mature fuel-cell technologies. A significant advancement in relation to the conductivity of present day proton-conducting ceramics is therefore critical to future breakthroughs in the development of next-generation fuel-cell technology, operating in the intermediate temperature range. Such an enhancement depends on increasing the understanding of the fundamental science of key mate­rials aspects such as crystal structure and proton-conduction mechanisms, in the most promising classes of materials, and the exploration of novel and completely new systems. For this purpose, neutron scattering is a powerful tool that has been applied successfully to studies of both oxide-ion and proton-conducting materials, and its role in this area of research is vast.

The aim of this chapter is not to give an exhaustive account for all the neutron scattering studies in this field, nor is it aimed at providing a review of SOFC technology, as excellent reviews can be found elsewhere [1-9]. Rather, this chapter centres on proton-conducting ceramics as separator materials for intermediate — temperature fuel cells and aims to give a flavour of how a variety of neutron methods can be used to reveal the details of structures and proton dynamics in this class of energy-relevant materials. In particular, the scope of contemporary structural and dynamical studies of perovskite-type oxides, known as the most promising class of proton-conducting ceramics for intermediate temperature fuel-cell applications, will be reviewed. This is followed by a concise review of selected examples of studies of other classes of promising candidate materials, such as hydrated alkali thio-hy — droxogermanates, solid acids, and lanthanum gallates, to illustrate the breadth of information that can be obtained. The chapter finishes with my personal thoughts on prospectives for future work in this timely area of research. The chapter may be viewed as an extended version of my recent publication in Dalton Trans. [17].