Ceramics for electrochemical storage

Ceramics for electrochemical storage

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Yulia Arinichevaa, Michael Wolffa, Sandra Lobea, Christian Dellena, Dina Fattakhova-Rohlfinga,b,c, Olivier Guillona,b,d, Daniel Bo€hme, Florian Zollera,e, Richard Schmuchf, Jie Lib, Martin Winterb,f, Evan Adamczykg, Valerie Pralongg a Forschungszentrum J€ulich GmbH, Institute of Energy and Climate Research, Materials Synthesis and Processing (IEK-1), J€ulich, Germany, bHelmholtz Institute M€unster: Ionics in Energy Storage (IEK-12), M€unster/J€ulich, Germany, cFaculty of Engineering and Center for Nanointegration Duisburg-Essen (CENIDE), University of Duisburg-Essen, Duisburg, Germany, dJ€ulich Aachen Research Alliance: JARA-Energy, J€ulich, Germany, eDepartment of Chemistry and Center for NanoScience (CeNS), Ludwig-Maximilians-Universit€at M€unchen (LMU Munich), Munich, Germany, fUniversity of M€unster, MEET Battery Research Center, M€unster, Germany, gNormandie University, ENSICAEN, UNICAEN, CNRS, CRISMAT, Caen, France

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Introduction: Overview of battery technologies

Electrochemical energy storage refers to a process of a reversible conversion of electrical energy into a chemical form to store or release electrical energy when needed. It plays a key role in the transition toward renewable energy sources (Diouf and Pode, 2015), electric mobility (Notter et al., 2010), and industry 4.0 (Lu et al., 2017) technologies and, therefore, is essential to ensure a sustainable future. Batteries belong to the most promising electrochemical energy storage systems, as they are expected to simultaneously fulfill a large number of criteria in order to meet challenging combinations of consumer demands, such as high-power and high-energy density, long life, low cost, excellent safety, and minimal negative environmental impact (Bieker and Winter, 2016a; Placke et al., 2017). In batteries, electrical energy is stored or generated via various redox reactions at the anode and cathode, depending on the battery chemistries. In this chapter, the history of different battery technologies as well as recent and future trends in its development will be discussed.

1.1 Pioneering work in the development of battery technology The history of battery technology development started almost 300 years ago, when two inventors, Ewald Georg von Kleist and Pieter van Musschenbroek of Leyden, independently invented the “Leyden jar,” the first electrical capacitor, which could store static electricity. This device consisted of a glass jar lined from outside and inside with thin metal foils and partially filled with water, acting as a conductor. A static generator was applied to the electrode, a metal rod, inserted in the jar and Advanced Ceramics for Energy Conversion and Storage. https://doi.org/10.1016/B978-0-08-102726-4.00010-7 © 2020 Elsevier Ltd. All rights reserved.

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Advanced Ceramics for Energy Conversion and Storage

connected to the inner metal foil (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017). The term “battery” was first introduced in 1749 by Benjamin Franklin when he was doing experiments with electricity using a set of linked capacitors, which he compared with a battery of military canons (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017). In 1799, following the Luigi Galvani’s study of bioelectricity, Alessandro Volta invented the first battery (“voltaic pile”), consisting of stacked copper and Zn discs, separated by cloth soaked in salty water. The working principle of the galvanic or voltaic cell, described later by Michael Faraday, is based on the production of electrical energy from spontaneous redox reactions taking place within the cell. When the system is not closed, copper reacts with oxygen from air to build copper oxide cathode. In this case, this primary (nonrechargeable) battery cell can be considered as the first metal-air system (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017). In 1802, William Cruickshank designed the first mass-produced primary electric battery using a voltaic pile: stack of copper and zinc discs were sealed in a wooden box filled with brine (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017). The same year Johann Wilhelm Ritter carried out an experiment with the so-called “Ritter pile.” This cell, consisting of a glass tube filled with a saline solution and closed by corks on both sides, each equipped with gold wires, was charged using the voltaic pile. The electrolysis of water took place during charge and the recombination of oxygen and hydrogen to water occurred during discharge. Thus, this experimental setup can be considered as the forerunner both of the fuel cell technology and of an accumulator, a secondary (rechargeable) battery (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017). These early discoveries paved the way for the development of further primary and secondary electrical batteries. Below, a brief overview of the most common recently relevant and of the next-generation battery technologies will be given.

1.2 Primary batteries In 1866, Georges-Lionel Leclanche patented a primary battery system of technical importance, based on a zinc anode and a graphite cathode, coated by manganese dioxide at the interface with an aqueous electrolyte consisting of ZnCl2 and NH4Cl. It was used to supply railroad telegraphs and electric bells with electricity before the centralized electricity supply was introduced. In 1887, Carl Gassner patented the first dry galvanic element based on zinc container as an anode housing the entire cell with carbon cathode and electrolyte inside. The Gassner’s dry cell had a commercial success and became the prototype for the dry cell industry. In the 1960s, the original zinccarbon technology was improved by substituting the acidic electrolyte by a better conducting alkaline electrolyte, KOH (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017). Today, zinc-carbon alkaline general purpose primary

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batteries are commercially available due to the low cost, environmental friendliness, low internal resistance, and relatively stable discharge plateau (Sayilgan et al., 2009; Yu et al., 2013b). As a result from slow kinetics at the triple-phase boundary solid/ liquid/gas and overpotentials at the air cathode, the maximum energy density is capped around 500 Wh kg
1with an overall energy efficiency of <60% in practice. Therefore, Zn-air primary batteries have found their niche market, primarily in low-power (e.g., hearing aids) and high-capacity applications (e.g., oceanographic navigation instruments).

1.3 Lead-acid battery Lead-acid battery is the first secondary battery technology for practical applications, which has been still technically up to date. Wilhelm Josef Sinsteden reported for the first time in 1854 that lead electrodes immersed in diluted sulfuric acid can store, that is, accumulate, electricity and be used as a coulometer. The first lead-acid accumulator, however, was developed 5 years later by Gaston Raimond Plante (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017; Ruetschi, 1977). The current cell design is represented by a lead dioxide plate as cathode and porous lead plate as anode, separated by plastic or a sponge-type fiberglass mat. The cell can be recharged applying reverse current. Despite of the relatively low practical specific energy of the modern lead-acid batteries (20–40 Wh kg
1), it has high-power capability, low production costs, and high recyclability. Therefore, it became the first commercially viable secondary battery technology, still dominating in the sector of engine starter, lightning and ignition batteries for automotive application as well as backup stationary power applications (May et al., 2018; Kwiecien et al., 2017).

1.4 Nickel-metal-based batteries Nickel-metal-based systems represent another important type of secondary battery technology. In 1899, Waldemar Jungner patented nickel-cadmium and nickel-silver rechargeable batteries with alkaline electrolyte. The invention of nickel-iron alkaline battery was ascribed to Thomas Edison 2 years later. These battery types were forerunners of the nickel/metal hydride (Ni/MH) battery, commercialized in 1989 (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017). The Ni/MH battery solved the toxicity problem of the NiCd battery and had an improved practical specific energy (70–100 Wh kg
1) compared to the NidCd and NidFe batteries (20–50 Wh kg
1). Since then various companies have been working on research and development of Ni/MH batteries to improve its performance and expand applicability targeting higher gravimetric energy, higher delivered power at low temperature, prolonged cycle life at high temperature, and lower material and manufacturing costs. Nowadays, it continues to be an important energy storage technology broadly used for consumer, automotive, and stationary energy storage applications (Chang et al., 2016; Young, 2018; Ouchi et al., 2016; Bernard and Lippert, 2015).

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1.5 Lithium-ion batteries 1.5.1 Primary lithium batteries Gilbert Newton Lewis began pioneering research on lithium battery in 1912. However, it was not until the early 1970s when first primary lithium batteries became commercially available (Scrosati, 2011; Placke et al., 2017; Bieker and Winter, 2016b; Jones, 2017). Lithium batteries, also referred to as lithium metal batteries, are nonrechargeable batteries with lithium metal as an anode with voltages ranging from 1.5 to 3.7 V depending on battery chemistry. Various cathode materials for the primary lithium batteries, such as iodine (I2), manganese dioxide (MnO2), thionyl chloride (SOCl2), sulfur dioxide (SO2), copper oxide (CuO), carbon monofluoride (CFx), silver vanadium oxide (SVO; Ag2V4O11), pyrite (FeS2), copper sulfide (CuS), vanadium pentoxide (V2O5), and silver chromate were developed. The development of primary lithium batteries was driven, that is, by the requirements for a small volume size for consumer electronics application immerged in the market in the 1970s. Primary lithium batteries have been used in implantable electronic medical devices, such as pacemakers, and in aerospace and oceanographic instrumentation (Crabtree, 2015).

1.5.2 Secondary lithium-ion batteries After the success of the primary lithium batteries, attempts to develop rechargeable lithium metal batteries were undertaken (Crabtree, 2015; Blomgren, 2017; Placke et al., 2017; Brandt, 1994). The research focused on the novel intercalation cathode materials leading to discovery of titanium disulfide (TiS2) and tantalum disulfide (TaS2) cathodes. The first secondary lithium battery based on TiS2 and lithium or lithium-aluminum anode was commercialized in 1976–78, followed by four further commercial rechargeable lithium metal batteries with MoS2, V2O5, V3O8, and MnO2 cathodes. However, the instability of the metallic lithium anode in contact with liquid electrolytes caused major safety issues, such as short circuit due to dendrite formation by cycling and subsequent inflammability of the cell, leading to the withdrawal of the secondary lithium metal batteries from the market (Placke et al., 2017; Brandt, 1994). Therefore, the research and development efforts were further concentrated on the replacement of the lithium metal anode with a safer anode material. It led to the development of the first so-called “rocking chair” battery, a lithium-ion cell, based on the reversible transfer of lithium ions between the positive and negative electrodes using intercalation materials of different potentials for the two electrodes. In 1980, a new class of cathode materials, layered transition metal oxides, such as LiCoO2, able to reversibly deintercalate and reintercalate lithium ions at relatively high potentials, was discovered by the team of John G. Goodenough laboratory. At the same time, carbonaceous materials, such as graphitic or amorphous carbons, were developed as suitable negative electrode materials for reversible intercalation-deintercalation of lithium ions. The advantages of the carbonaceous materials are their potential close to that of metallic lithium, leading to a high-energy density, low volume changes upon lithiation as well as high cyclability in combination with a suitable electrolyte. The first commercially successful Li-ion battery (LIB) cell, with high energy (80 Wh kg
1; 200 Wh L
1) and high voltage ( 3.7 V), based on coke anode, LiCoO2

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cathode, and a nonaqueous electrolyte, was introduced in the market by Sony Corporation in 1991 (Crabtree, 2015; Blomgren, 2017; Placke et al., 2017; Brandt, 1994). Since then LIB technology has been undergoing continuous developments leading to superior characteristics in comparison with all the earlier rechargeable battery systems and will continue to improve in terms of cost, energy, safety, and power capability. Rechargeable LIBs have emerged as the dominant energy storage source for consumer electronics, automotive, and stationary storage applications. The most common cathode materials for present Li-ion batteries include LiCoO2 (LCO), LiMn2O4 (LMO), LiNixMnyCo1
x
yO2 (NMC), LiNi0.80Co0.15Al0.05O2 (NCA), and LiFePO4 (LFP) (Mekonnen, 2016; Blomgren, 2017; Whittingham, 2004). The materials show certain advantages and disadvantages and have been used for different applications, which will be discussed in more detail in Section 3. In most of Li-ion batteries, graphitic and hard carbons are used as anodes due to their attractive balance of relatively low cost, abundance, moderate energy density, power density, and cycle lifetime, compared to any other intercalation-type anode materials. Another anode type for LIBs is lithium titanium oxide (Li4Ti5O12/LTO), showing improved low-temperature performance for lower energy, but high-power applications (Choi et al., 2018). Alloying materials, elements which electrochemically alloy and form compound phases with Li at a low potential, such as Si, Sn, Ge, and Ga, have high volumetric and gravimetric capacity, but suffer from mechanical instability due to a large volume change upon lithiation. Silicon-graphite-based composite anodes allow increasing the specific energy, while reducing mechanical instability of pure silicon anodes (Mekonnen, 2016; Blomgren, 2017; Kaskhedikar and Maier, 2009; Nitta et al., 2015). Ceramic anodes will be present in Section 2. Another important cell component is a separator, a thin porous membrane that physically separates the anode and cathode. The separators are often porous polyolefin-based polymer films such as polypropylene (PP), polyethylene (PE), which can be ceramic coated to improve their stability (Blomgren, 2017). The electrolytes are LiPF6-based nonaqueous aprotic organic solvents, such as ethylene carbonate (EC), dimethyl carbonate (DMC), diethyl carbonate (DEC), and/or ethyl methyl carbonate (EMC). The electrolyte may also contain additives, for a better solid electrolyte interphase (SEI) on graphite anode, an increased flame resistance or an intrinsic overcharge protection. Copper and aluminum are used as current collectors on the negative and positive electrode side, respectively (Blomgren, 2017). Exemplarily, the rechargeable LIB cell consisting of graphite anode, LCO cathode, current collectors, porous separator, and liquid electrolyte is schematically represented in Fig. 1 (Ozawa, 1994). The cell is fabricated in the discharged state. The cathode is oxidized according to Eq. (1) during charging. Thereby produced Li+ ions travel through the electrolyte and are inserted into the graphite anode according to Eq. (2). At the same time, the electrons flow to the anode through the external circuit. The equation represents the overall cell reaction (3). The opposite processes take place upon discharge. Cathode : LiCoO2 >Li1
n CoO2 + nLi + + ne


(1)

Anode : C6 + nLi + + ne
>Lin C6

(2)

Cell reaction : LiCoO2 + C6 >LinC6 + Li1
n CoO2

(3)

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Fig. 1 Schematic representation of a rechargeable lithium-ion battery cell based on graphite anode and LiCoO2 cathode.

Advanced Ceramics for Energy Conversion and Storage

Charge e–

e– Discharge Liquid electrolyte

e– e–

Li+

Negative current collector

Anode

Porous separtor

Cathode

Positive current collector

1.6 Alternative rechargeable battery technologies beyond lithium-ion batteries Since the commercial introduction of LIBs in the early 1990s, various alternatives in the field of electrochemical energy storage have been considered (Nagaura and Tazawa, 1990). This is due to the ever-growing demand for high-energy density electrochemical storage devices and the limitation of the energy content of current LIBs. Owing the positive interplay between cost effectiveness and performance, LIB technology serves as a benchmark for future technologies. In this context, so-called beyond lithium-ion batteries (PLIBs) have started to catch the interest of science on a global scale. Besides related battery concepts like sodium-ion (SIB), potassium-ion (KIB), and multivalent-ion (Mg2+, Ca2+, and Al3+), there are also different approaches like metal-sulfur (M-S), metal-air or metal-oxygen (M-Air, M-O2), and all-solid-state battery (ASSB) which are included under the term PLIB.

1.6.1 Molten-sodium batteries Sulfur is a very versatile element and capable to play a role in different kinds of battery types. Among others, liquid sulfur is utilized in the molten-salt Na-S battery, which laid the foundation for e-mobility in the 1960s (Partridge, 1976). It is still being used today, primarily in large-scale, high-temperature energy storage, and conversion applications. The fundamental functionality of the classic high-temperature cell is

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based on a molten S cathode, a molten Na anode, and a beta-alumina solid electrolyte (BASE, also β00 -alumina) with a working temperature in the range between 300°C and 350°C. The electrochemical mechanism of Na-S as for Li-S is involving a cathodic reduction of sulfur to polysulfide Sn 2
and an anodic oxidation of the respective alkali metal. Na-S batteries offer important features such as no self-discharge, low maintenance, and good recyclability. Thermal loss due to the high operating temperatures cannot be completely prevented by thermal isolation. Therefore, only large Na-S cells with a more advantageous surface/volume ratio are energetically efficient. The increased reactivity of liquid sodium compared to solid lithium and the risk of thermal runaway related to deep discharge are the major drawbacks for this technology (Fig. 2). According to the most recent publications, the overall trend goes to low and even room-temperature Na-S cells which started in 2009 by Ceramatec with the development of a Na-Super-Ionic-CONductor (NASICON)-based battery operating at 90°C. For more information about electrolytes, please refer Section 4. Nowadays, lab-scale room-temperature Na-S batteries can be realized with reasonable results (e.g., 580 mAh g
1 after 500 cycles) (Xu et al., 2018b; Qiang et al., 2017). The so-called ZEBRA battery (“Zero Emission Battery Research Activities”) constitutes an alternative high-temperature secondary battery concept and originates from a patent applied by a South African company (“Zeolite Battery Research Africa”) (Dustmann, 2004). The battery consists of a molten sodium anode and a Ni/NiCl2/ NaCl composite cathode which are separated by a beta-alumina solid electrolyte. A preheating time of up to 24 h is needed to get the battery up to the operating temperature of 270°C. While the cell shows a decent efficiency of about 80%, it uses 14% of its own capacity per day to maintain temperature when not in use. Despite the

Fig. 2 Schematic representation of a high-temperature Na-S battery (Dunn et al., 2011).

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Advanced Ceramics for Energy Conversion and Storage

application of ZEBRA batteries in several electric vehicles (EVs) since the turn of the millennium, the commercial success has phased out in the recent years and the focus today is more on battery storage power plants and in the field of military industry.

1.6.2 Lithium-sulfur batteries Among all metal-sulfur batteries, the Li-S battery has probably attracted the most scientific interest in the recent years. While the concept has already been known for >50 years, research and development have experienced a renaissance only one decade ago (Danuta and Juliusz, 1962; Ji et al., 2009). The Li-S cell consists of a sulfur cathode, a lithium metal anode, an organic solvent as electrolyte, and a polymer separator. During discharge, cathodic reduction of sulfur via soluble polysulfide intermediates Li2Sx (x ¼ 2, 4, 6) to Li2S is accompanied with an anodic oxidation of Li metal to Li+ ions. Depending on the polysulfide chain length viscosity, mobility, and solubility of Li2Sx changes during discharge. For many years, the application of Li-S-based systems was limited due to their poor rate performance and low-cycle life. The high gravimetric capacities and significantly higher theoretical energy densities compared to LIB (2500 Wh kg
1 assuming reaction completeness) are important motivation factors to further optimize this promising approach (Peramunage and Licht, 1993). Today, the practical achievable gravimetric energy is often stated to be in the range of 600 Wh kg
1 (Fang and Peng, 2015). Furthermore, Li-S provides a heavy metal free and low-priced battery system which works at room temperature. However, there are several process-related challenges that have to be taken into account. First, Li-S batteries suffer self-discharge related to the reactive polysulfide intermediates Li2Sx (x ¼ 2, 4, 6) and the irreversible formation of an insoluble, isolating Li2S film on the Li anode. Second, the recrystallization of S leads to morphological changes of the cathode composite which reduces the contact between S and C and, therefore, the electric conductivity after a couple of cycles. Another issue is the irreversible reaction of Li with the electrolyte due to dendrite and SEI formation. Excess electrolyte and Li are used to tackle these problems. Together with the large amount of carbon additive needed to improve the electric conductivity, this leads to a high amount of extra mass and is rather disadvantageous in terms of volumetric energy content. In spite of these limitations, there has been a tremendous progress during the past few years in terms of cycle stability (e.g., >4000 cycles at  100% Coulombic efficiency) (Xu et al., 2015b). Although lab-scale systems already look very promising, further improvements have to be done before commercial viability of the Li-S battery.

1.6.3 Metal-air batteries Metal-air or metal-oxygen batteries, respectively, are believed to be potential candidates as next-generation electrochemical energy storage systems employed in EVs due to their high theoretical specific energy density. First, it is necessary to differentiate between aqueous (involving Zn, Mg, Fe, or Al metal anodes) and nonaqueous (involving Li, Na, or K metal anodes) systems. The former concept goes back to

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Zn-air primary batteries in the 19th century followed by commercialization in the early 20th century. The latter was initially proposed in the 1970s for automotive applications but captured worldwide attention just a few years ago, initially for Li-air, and more recently for Na-air and K-air (Heise, 1933; Kinoshita, 1992). Both setups have the open porous cathode architecture in common that permits oxygen supply from the ambient air. Based on their chemical nature to form passivating oxygen or hydroxide layers in aqueous media, metals like Fe, Zn, Al, and Mg are water compatible despite their thermodynamic instability. Herein, the metal is oxidized at the anode during discharge, while O2 from the ambient air is reduced at the catalyst-supported gas-diffusion cathode which possesses architecture similar to gas-diffusion electrodes in hydrogen fuel cells (Fig. 3). For secondary batteries, further challenges arise: poor cycling stability related to catalyst degradation after repeated oxygen evolution and reduction reactions, dendrite formation, and corrosion due to nonuniform zinc dissolution. Thus, optimization of the catalyst material (e.g., Pt, MnOx, and graphene), cathode morphology, anode composition (by alloying with other metals or introducing protective additives) as well as their interactions with each other are all equally important research foci. A more recent approach tries to tackle some of these issues by a bicathode configuration providing charge and discharge functions by separate uni-functional cathodes (Lee et al., 2013a). At the expense of increased cell size, higher price, and lower energy density, an improved cycling performance can be attained. Other aqueous M-air battery systems are either hardly competitive due to the strong discrepancy between theoretical and realistic energy density (Fe-air) or face serious corrosion issues and are limited to mechanically recharging only (Al-air and Mg-air) (Narayanan et al., 2012; Egan et al., 2013; Zhang et al., 2014b).

Fig. 3 Schematic configuration for aqueous and nonaqueous metal-air batteries. Reprinted with permission from Li, Y., Lu, J., 2017. Metal-air batteries: will they be the future electrochemical energy storage device of choice? ACS Energy Lett., 2, 1370–1377. Copyright (2017) American Chemical Society.

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Advanced Ceramics for Energy Conversion and Storage

Nonaqueous M-air batteries follow different charge/discharge mechanisms in dependence of the selected alkaline metal. In accordance with the hard-soft acid base theory, the reactions are as follows: during the first step, O2 molecules get reduced to superoxide O2
anions at the catalyst surface. In dependence of the ion radius M+, the second step involves either the formation of the peroxide M2O2 (M ¼ Li), the superoxide MO2 (M ¼ K), or a mixture of both (M ¼ Na) (Winter and Brodd, 2004). On account of their chemical nature and early development state, nonaqueous M-air batteries face even more challenges compared to aqueous M-air batteries. For the Li-Air battery—as a representative example—the following issues arise: first, unwanted reactions between intermediates singlet oxygen, lithium superoxide, and lithium peroxide with the electrolyte occur. The solvent choice is further limited due to the solubility characteristics of Li2O2 and LiO2. Direct reactions of Li with the electrolyte are hindered to a certain degree because a SEI is formed. Further improvements can be achieved introducing an artificial SEI in the form of protective coatings of lithium-ion conducting ceramics or glasses [e.g., LI Super Ion CONductors (LISICON)] (Wang and Zhou, 2010). Second, the achievable discharge capacity deviates considerably from the theoretical values. This is mostly related to two different effects: the dendrite formation of Li on the anode (which also bears the risk of short circuit) and the deposition of Li2O2 on the porous cathode leading to direct passivation of the electrochemically active surface and pore clogging. Third, recharging requires high overpotentials which can result in electrode and electrolyte decomposition (Grande et al., 2015). Owing to this situation, novel approaches have been followed throughout the last few years. Zhu et al. (2016b) for instance developed a cathode composite material enabling a charge/discharge without any gas evolution. As a summary it may be stated that despite the high theoretical potential of M-air battery technology, its impact has so far fallen short of its potential due to challenges concerning metal anode, air cathode, and electrolyte. Their viability to replace Li-ion batteries in the future remains unclear.

1.6.4 Alternative metal ion batteries While Li is not a rare element, it is discretely located on Earth and a Li shortage is predicted by the year 2050 (Vaalma et al., 2018). With Na as one of the most abundant elements, SIBs clearly have advantages in terms of availability and cost efficiency. However, this only represents a small percentage of the cell production costs and this alone is not sufficient to replace LIBs in the future. The particular appeal in this technology is based on other aspects: Na, in contrast to Li, does not alloy with Al and allow the use of cost-effective Al as an anode current collector instead of expensive Cu. A special feature of SIBs is their deep discharge protection up to an end-point voltage of 0 V, whereas LIBs suffer irreversible capacity loss (ICL) (Slater et al., 2013) (Fig. 4). At first glance, SIBs work similar to LIBs with an ion transfer of Na+ instead of Li+ between two electrodes. To a certain degree, the monovalency of both carrier ions allows a knowledge transfer from the already well-studied LIBs to SIBs. However, due to the larger ion radius of Na distinct differences in intercalation behavior and

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Fig. 4 Current research progress in sodium-ion batteries. Reproduced from Hwang, J.Y., Myung, S.T., Sun, Y.K., 2017. Sodium-ion batteries: present and future. Chem. Soc. Rev., 46, 3529–3614 with permission from the Royal Society of Chemistry. Available from: .

crystal structure preferences are observed. As a consequence thereof, Na-based materials usually show an inferior performance compared their Li-analog. While graphite does not intercalate Na ions reversibly, other carbon-based (e.g., graphene, petroleum cokes) and titanate-based anodes (e.g., Na2Ti3O7) have shown an improved intercalation behavior (Zhecheva et al., 2002; Mei et al., 2016). Analogs to LIBs, several metal oxides, sulfides, and phosphides offer conversion mechanisms, allowing their use as anode material for SIBs (Hwang et al., 2017a). Alternatively, the sodiation/

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desodiation mechanism consists of an alloying reaction using group IVA or VA metals and metalloids. Despite their high theoretical specific capacities, the latter two anode systems suffer rapid capacity fading due to high volume changes. One approach to tackle this issue is the usage of metal/carbon fiber composites (Zhu et al., 2013). For SIB cathodes, a wide variety of materials is available, including polyanions, layered metal oxides, and organic compounds (see more details in Section 2). In contrast to the anode, all cathode materials follow a Li-analog intercalation reaction but involving continuous structural changes due to the large ion Na-ion radius. In accordance with the early studies of Delmas et al. (1980), layered transition metal oxides Na1
xMO2 can be classified into two different structural types in dependence of the sodiation state (Delmas et al., 1980). Until today, a lot of different compositions have been evaluated, for example, NaxNi1/3Co1/3Mn1/3O2 which has great potential for high-voltage applications (Xu et al., 2017b). Besides fluorosulfates, oxychlorides, and NASICON, fluorophosphates like A2FePO4F (A ¼ Na, Li) are probably the most prominent compound group among the polyanion-based cathode materials (Ellis et al., 2007). Besides SIBs, a variety of similar battery concepts based on alternative mono- or multivalent metals have been discussed in the recent years. Introduced in 2004, the KIB comes with the advantages of low cell fabrication costs and long cycling life (Eftekhari, 2004). Herein, the highly stable Fe4[Fe(CN)6]3 (Prussian blue) and KBF4 have been employed as cathode material and electrolyte, respectively (Komaba et al., 2015; Eftekhari et al., 2017). Moreover, secondary cells based on Mg2+ (MIB) and Ca2+ (CIB) have been considered as possible LIB replacements but are still very much in their infancy (Singh et al., 2013). In comparison to Li as an anode material, Mg and Ca possess lower gravimetric but higher volumetric energy densities and do not exhibit dendrite formation at low current densities (Ponrouch et al., 2015). The comparably low ion mobility, low working voltage, and improvable cycling stability can be considered as major drawbacks (Wang et al., 2018a). In contrast to these more recent approaches, the aluminum-ion battery (AIB) has been introduced >30 years ago and still is an ongoing topic pursued by several research groups (Zafar et al., 2017). Despite the advantages (high theoretical capacity, low production cost, and low flammability), the low shelf life and the capacity dependence on temperature, number of charge cycles, and rate of charge have hampered their commercial utilization. Overall, alternative metal ion batteries are either in an early development state or rather suitable for long-life applications with low to medium currents. Especially, large-scale production of SIB cells (filled with liquid electrolyte) has already started. They can offer a cost-effective alternative for stationary energy storage systems, where weight and energy density are of minor importance.

1.6.5 Lithium(-ion) and sodium(-ion) all-solid-state batteries Although commercialized for some niche markets, ASSB technology, a secondary lithium(-ion) or sodium(-ion) battery technology based both on solid electrodes and solid electrolytes, is still under development. It allows in principle overcoming safety concerns of the conventional LIBs due to the lack of leakage of liquids and flammable organic components and also offers the potential for a significant improvement of

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energy density (Manthiram et al., 2017; Schnell et al., 2018; Sun et al., 2017; Kim et al., 2015b; Kato et al., 2016). Rechargeable Na(-ion) ASSBs have been considered as a potential competitive alternative to Li(-ion) ASSBs, especially for large-scale applications, due to the ubiquitous abundance of sodium reserves, lower cost, and comparative performance to that of Li(-ion) ASSBs (Numata et al., 2000; Karim et al., 2013; Palacı´n et al., 2000). However, there are still several technological challenges to be solved before the Li(-ion) and Na(-ion) ASSB technologies will be ready for commercialization (Yu et al., 2017; Kerman et al., 2017). In general, Li(-ion) and Na(-ion) ASSBs have electrochemical storage mechanisms similar to those of corresponding liquid electrolyte cells described above. A typical design of an ASSB cell with lithium or sodium anode is schematically presented in Fig. 5.

Solid electrolytes One of the key components enabling rechargeable ASSB technology is a solid electrolyte. Solid electrolytes as described in detail in Section 4 should satisfy such technological requirements as high ionic conductivity in combination with negligible electronic conductivity, wide voltage window, chemical compatibility with cathode Charge e– – e–

Anode current collector

Li- or Na-metal anode

Solid electrolyte

e–

Li+/Na+

Mixed cathode

e–

Cathode current collector

Discharge

e– +

Fig. 5 Schematic representation of a typical all-solid-state battery with a lithium or sodium metal anode.

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and anode materials, as well as relatively simple fabrication on a large scale with low cost (Manthiram et al., 2017). Generally, solid lithium- or sodium-ion conductors have been divided into three classes, which can complement each other to satisfy these requirements: (1) inorganic glassy or ceramic compounds; (2) organic polymers, and (3) composite or hybrid electrolytes consisting of a combination of the first two classes of materials (Manthiram et al., 2017; Hou et al., 2018a; Zheng et al., 2018). Ionic transport in solid inorganic electrolytes is determined by the concentrations of mobile ions and vacancies, relative sizes of connected conduction pathways in crystal structures with Schottky and Frenkel point defects as well as by ion diffusion properties at the grain boundaries (Hou et al., 2018a,b; Zheng et al., 2018). Promising solid inorganic Li-ion electrolytes comprise amorphous lithium phosphorous oxynitride (LiPON) and with room-temperature conductivities up to several mS cm
1 lithiumsulfide-based glass ceramics, NASICON-type phosphate [Li1+ xAlxTi2
x(PO4)3 (LATP)] and garnet-type oxide [Li7La3Zr2O12 (LLZO)] ceramic electrolytes. Typical solid inorganic electrolytes for Na(-ion) ASSBs, also possessing relatively high room-temperature ionic conductivities over 1 mS cm
1, include Na-β00 -alumina, Na superionic conductors [NASICON, i.e., Na3.1Zr1.95Mg0.05Si2PO12 (Song et al., 2016)], sulfides (i.e., Na3PS4, Na10.8Sn1.9PS11.8), and complex hydrides (e.g., sodium borohydride) (Yu et al., 2018b; Hou et al., 2018a,b). Solid polymer electrolytes have generally significantly lower ionic conductivity than ceramic electrolytes but show mechanical flexibility, light weight, convenience of fabrication process, and accommodation of volume changes of electrodes during charge/discharge. In solid polymer electrolytes, Li- or Na salts are solvated by polymer chains in, for example, polyethylene oxide (PEO)- or polysiloxane-based polymers and Li or Na ions move through the connected polymer chains. The ionic conductivity of a solid polymer electrolyte is related to the number of mobile ions and the segmental motions of the polymer chains. The ionic transport in the connected polymer chains can be blocked by crystalline chain segments, which form below the glass transition temperature (Tg). Tg can be lowered, for example, by adding nanosized fillers. However, a relatively low ionic conductivity at room temperature still represents the major drawback of polymer electrolytes (Zheng et al., 2018; Hou et al., 2018a,b). Composite or hybrid electrolytes, combining the advantages of (glass-)ceramic and polymer ionic conductors, provide improved ionic conductivity with high flexibility for reducing the interfacial resistances between solid electrolytes and electrodes (Zheng et al., 2018; Hou et al., 2018a,b).

Anode materials for Li(-ion) and Na(-ion) ASSBs As anodes for all-solid-state LIBs, carbon and silicon-based materials as well as inorganic compounds, such as SnOxNy, LiTi2(PO4)3 (LTP), can be used (Yersak et al., 2013). Solid electrolytes may also enable the use of lithium metal, which can lead to an increase in volumetric energy density up to 70% in comparison to LIBs with conventional anode materials (e.g., graphite). However, lithium dendrites can

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penetrate under given conditions dense solid electrolyte, including LLZO, which is intrinsically stable against metallic lithium (Porz, 2017). To improve the performance of ASSBs the contact resistance at the electrolyte—Li-metal anode interface should be decreased, for example, by electrolyte surface treatment and coating, to prevent the formation of lithium dendrites (Tsai et al., 2016; Porz, 2017; Basappa et al., 2017). Adapting LIB anode materials to sodium-ion ASSBs is challenging primarily due to the larger size of the sodium ions impeding Na-intercalation and influencing crystal structure preferences. Recent advances have been made on the candidate negative electrodes for Na-ion ASSBs, including carbon materials (i.e., graphite and carbon quantum dots), elementary substances (i.e., Sb and P), phosphorene-graphene hybrid material (Sun et al., 2015), and transition metal oxides (i.e., Fe2O3 and CuO), sulfides and selenides (i.e., Cu2Se, MoSe2, FeSe2, MnSe2) (Qian and Lau, 2018; Cao et al., 2019; Sun et al., 2015). Sodium metal anode can also be used with solid electrolytes to significantly improve the energy density of Na-ASSBs. However, besides possible Na-dendrite formation, low melting point of Na 97.7°C should be taken into consideration to minimize safety-related concerns (Wan et al., 2018; Sun et al., 2015; Gao et al., 2018).

Cathode materials for Li(-ion) and Na(-ion) ASSBs Cathode materials used in lithium(-ion) and sodium(-ion) ASSBs are similar to those in the traditional LIBs and SIBs, respectively [i.e., lithium transition metal oxides, such as LiCoO2 (LCO), LiNixCoyAl1
x
yO2 (NCA)], V2O5 (Yao et al., 2018), for Li(-ion) ASSBs and NaCrO2, Na2MnFe(CN)6, NaTi2(PO4)3, Na-based vanadate phosphates for Na(-ion) ASSBs (Zhao et al., 2018; Lu et al., 2018a). The use of highvoltage cathodes can further boost the energy density of ASSBs. Suitable high-voltage cathode materials include, for example, LiNi0.5Mn1.5O4 (LNMO) (Bini et al., 2018) and Na3V2(PO4)3 (Kovrugin et al., 2018) for Li(-ion) and Na(-ion) ASSBs, respectively. In order to improve the ionic transport within the ASSB-cathodes, cathode materials should be mixed with solid electrolyte (Lu et al., 2018a). The fabrication of mixed cathodes still requires optimization in terms of diffusion pathways as well as due to reactivity of cathode materials in contact with ceramic electrolytes at high temperatures (Park et al., 2016; Kerman et al., 2017; Lu et al., 2018a; Miara et al., 2016).

2

Ceramic anodes for Li and Na-ion batteries

2.1 Introduction Anodes (negative electrodes) of rechargeable batteries should ideally possess a possible low insertion/extraction potential and a high reversible gravimetric and volumetric capacity to provide maximum energy density. Metallic lithium is the best possible anode for LIBs as it is the most electropositive element (
3.04 V vs SHE) with the highest theoretical gravimetric capacity (3862 mA g
1) due to its light weight (6.94 g mol
1) and low density (0.534 g cm
3). However, the high reactivity of

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metallic lithium and the tendency to mossy growth and dendrite formation upon cycling in conventional aprotic solvent are still serious challenges preventing its application as an anode in secondary lithium batteries. A similar situation applies for the use of metallic sodium in secondary SIBs, as sodium is chemically even more reactive than lithium and has a relatively low melting point (97.7°C) (Chen et al., 2017; Aravindan et al., 2015a). As a practical solution of this problem, graphitic carbons were introduced as anode materials, and currently, they dominate the market of the commercially available rechargeable lithium cells. Graphitic carbon is an insertiontype anode material acting as a stable host structure for lithium intercalation able to withstand several hundreds of charge/discharge cycles without structure degradation. Due to a slightly more positive lithium insertion potential (0.1 V vs Li/Li+), graphitic carbons do not suffer from lithium plating and lithium dendrite issues, rendering them safe anode material at moderate charging rates. Shortcomings of graphitic anodes, however, are a much lower insertion capacity as compared to lithium metal (372 mAh g
1) and a limited rate capability (power density). In spite of the progress achieved in the performance of state-of-the-art rechargeable lithium cell, there is a strong need for anode materials with high-energy and power density, low-cycle life, and safe operation. Ceramic materials are in the focus of intensive research activity as possible anode materials due to a huge variety of their compositions and hence, possible solid-state chemistries available for ion/electron storage. Ceramic materials can be produced by scalable methods with an unlimited variety of different morphologies and microstructures, making them attractive for commercialization and large-scale production. From the point of view of electrochemical energy storage mechanism, ceramic anodes can be broadly divided in three classes: insertion-type, conversion-type, and conversion and alloying-type materials (Fig. 6). Insertion-type materials can incorporate additional ions in their structure without significant structural changes.

Fig. 6 Structural changes during charge-discharge of insertion, alloying, and conversion anodes materials. Bensalah, N., Dawood, H., 2016. Review on synthesis, characterizations, and electrochemical properties of cathode materials for lithium ion batteries. J. Mater. Sci. Eng., 5 (4).

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The benefits of such materials are low volume changes and low structure reorganization energies ensuring reversible insertion/extraction and long cycle life. The penalty for the structure stability is, however, the unavoidably lower insertion capacity, which is determined by the available ion vacancies in the host structure. So far, a variety of transition metal-based insertion-type anodes for alkali metal ions were introduced but among them ceramic compositions including titania-based materials are projected to be able to replace graphitic anodes as they provide the required cell safety, cost effectiveness, compatibility with cathodes, and improved electrochemical activity for high-rate applications (Aravindan et al., 2015a). The electrochemical potential of these non-carbonaceous anodes is, however, comparably high (>1 V vs Li/Li+), lowering the applicable potential window of the cell. In addition, the practical reversible capacity is less than that of graphite and most ceramic materials possess a rather limited electronic conductivity. Conversion-type anode materials (CTAMs) undergo a complete transformation from oxidized (transition metal ions) to fully reduced (metal state) during charge and discharge cycles. In some cases, the metal phase can be reduced further with formation of lithium (or sodium) alloys (conversion and alloying mechanism). Because of the multiple electron reactions, conversion and combined conversionalloying-type anodes can provide much higher storage capacities as compared to the insertion-type anodes. However, a significant drawback of such anodes is a huge structure reorganization accompanied by significant (up to several 100%) volume changes, which typically results in a very fast degradation during charge/discharge and rapid capacity loss. It should be noted that practically all the ceramic materials feature in their bulk (macroscopic) state low electronic and ionic conductivities, and often also lower reversibility and higher degradation as compared to graphite anodes. These shortcomings can be, however, successfully mitigated via nanoscaling of ceramic materials. Particularly, a very small size of crystalline domains of only a few nanometers drastically improves the performance of ceramic anodes due to the change in electrochemical storage mechanism as it will be shown below. Furthermore, coating of ceramic powders with carbonaceous materials, optimization of nanostructures, and combination of nanostructures with conducting carbonaceous compounds to form nanocomposites are extremely efficient strategies to enhance storage capacity, improve the rate performance, sustain the volume variation, and thus improve the cycling stability of ceramic anode materials. A dominating majority of ceramic anode materials are prepared as nanostructured materials or nanocomposites, with the number of available compositions and morphologies increasing daily. As it is practically impossible to hold step with this dynamically developing field, this chapter intends at introducing only the most established and promising classes of ceramic anode materials for rechargeable lithium batteries, with a special focus on their structure and the basic electrochemical properties. Moreover, for the interested reader there are several reviews dedicated to the anode materials for SIBs (Dahbi et al., 2014; Hwang et al., 2017a; Irisarri et al., 2015; Kim et al., 2014; Lao et al., 2017; Yabuuchi et al., 2014) that are at an earlier stage of development compared to LIBs.

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In the following, we summarize the main advances regarding the three main types of ceramic anodes for SIBs and LIBs.

2.2 Insertion-type ceramic anodes 2.2.1 Titanate-related battery anodes—Structure and key properties Titanium dioxide (TiO2) anodes TiO2-based materials are among the most studied and attractive ceramic anodes for the Li-ion (Madian et al., 2018) and increasingly also for Na-ion batteries (Ni et al., 2018; Li et al., 2017b; Zhou et al., 2016a; Chen et al., 2016; Wang et al., 2016a; Wu et al., 2016a). TiO2 exists in eight polymorphs, namely rutile, anatase, brookite, TiO2-B, TiO2-R, TiO2-H, TiO2-II, and TiO2-III (Cromer and Herrington, 1955). Among these phases, anatase, rutile and bronze are the most investigated as battery anodes due to the abundance of TiO2, availability, and suitable electrochemical performance ( Jiang and Zhang, 2013). For all TiO2 phases, irrespective of their crystalline structure, storage of 1 mol Li+ per mole is theoretically possible due to the reversible transformation of Ti4+ to Ti3+ state, resulting in a theoretical insertion capacity of 335 mAh g
1. In practice, however, typically only 0.5 mol Li can be reversibly inserted/extracted per mole TiO2 resulting in a moderate insertion capacity of 170–180 mAh g
1 (Madian et al., 2018). Insertion of Li in all titania polymorphs takes place at potentials of around 1.5–1.8 V vs Li/Li+. The relatively high insertion potential renders titania a safe anode material due to a greatly minimized risk of lithium dendrite growth, but on the other hand it limits the voltage (and thereby the energy density) of the respective full cells. TiO2 anatase Anatase is one of the most investigated titania phases for lithium storage (Kavan, 2012). Bulk anatase is a metastable phase (Shannon and Pask, 1965), however, it is the thermodynamically most stable phase on the nanoscale. Anatase crystallizes in a tetragonal symmetry with the body-centered I41/amd space group in which distorted edge-sharing [TiO6] octahedra are stacked in one-dimensional (1D) zigzag chains (Fig. 7A) (Dachille et al., 1968; Yang et al., 2009; Su et al., 2012). Upon insertion, Li ions can diffuse along the empty [TiO6] zigzag channels connecting octahedral interstitial sites. The diffusion coefficient for Li ions in the anatase phase is relatively low, being in the range of 2–6  10
13 cm2 s
1, as reported for single crystals by Kavan et al. (1996) to 4.7  10
12 cm2 s
1 determined by 7Li MAS nuclear magnetic resonance (NMR) (Wagemaker et al., 2001). Li insertion into the anatase phase induces the phase transformation from the tetragonal to the orthorhombic Li0.5TiO2 phase (space group Pnm21) because of a loss of symmetry in the y-direction, resulting in an increase of the unit cell volume by 3.7% (Aravindan et al., 2015a). In addition to nanostructuring, carbon coating of the titania surface and fabrication of carbon-titania nanocomposites are other efficient and intensively explored ways to overcome the poor lithium-ion diffusion and low electrical conductivity in titania materials. In typical synthesis methods, incorporation of carbon into TiO2 is

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Fig. 7 (A) The structure of anatase (Liu et al., 2013a) and (B) potential-capacity profiles of anatase anodes with different particle sizes: 6 nm (A6), 15 nm (A15), and 30 nm (A30) ( Jiang et al., 2007b). Modified from Liu, Z., Andreev, Y. G., Armstrong, A. R., Brutti, S., Ren, Y., Bruce, P. G., 2013a. Nanostructured TiO2 (B): the effect of size and shape on anode properties for Li-ion batteries. Prog. Nat. Sci.: Mater. Int., 23(3), pp 235–244, Reproduced with permission from Jiang, C., Wei, M., Qi, Z., Kudo, T., Honma, I., Zhou, H., 2007b. Particle size dependence of the lithium storage capability and high rate performance of nanocrystalline anatase TiO2 electrode. J. Power Sources, 166(1), pp 239–243. Copyright 2007 Elsevier.

conducted by hydrothermal heating of metal salts in presence of carbon sources such as glucose (Chen et al., 2016; Aravindan et al., 2017). Li insertion/extraction in the anatase phase takes place at a potential around 1.75 V vs Li/Li+. The major part of Li ions is stored in a two-phase reaction manifested by a plateau at 1.75 V vs Li/Li+, with a total insertion of 0.5 mol Li and a gravimetric capacity of 175 mAh g
1 (Fig. 7B). Further Li-ion storage into the TiO2 framework is restrained because of the strong repulsive forces among the Li ions (Madian et al., 2018). At the initial steps of the Li-insertion process, a solid solution is formed, observed as a monotonous curve at potentials below 1.75 V vs Li/Li+, while for a fully lithiated material some amount of Li can be stored interfacially at a potential above 1.75 V vs Li/Li+ (Shin et al., 2011). The mechanism of lithium insertion, the amount of Li ions stored via different mechanisms, and also the rate of the Li-insertion process greatly depend on the size of crystalline domains and the microstructure of the anatase phase (Fig. 7B) (Wagemaker et al., 2007b). Generally, the slower two-phase reaction mechanism dominates the storage in bulk macroscopic materials, while the faster processes of solid solution and interfacial lithium storage play an increasing role in the nanostructured materials (Wilhelm et al., 2004). The strong impact of microstructure on the lithium insertion process provides broad possibilities to optimize the performance of anatase anodes by controlling their morphology, which boosted research

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on nanostructured titania materials (Hu et al., 2006). Systematic investigation of the impact of particle size on the electrochemical performance of anatase was performed by Wagemaker et al. (2007a). The authors have demonstrated that, in contrast to the bulk phase, nanoparticles below 7 nm in size can deliver up to 1 mol of Li ions per formula unit (f.u.) due to a greatly increased solubility on the nanoscale. The very small particles size, however, often negatively affects the cycling stability. Therefore, an optimum range exists for different materials to achieve the highest possible capacity, fast insertion rate, and cycling stability. For anatase, the optimum particle size is considered to be between 8 and 25 nm (Kang et al., 2011), thereby the anodes demonstrate a rather high cycling stability. Thus, full-Li cells with respective anodes were reported to retain 90% of their initial capacity after 700 cycles (Aravindan et al., 2013b). The electrochemical properties of anatase TiO2 vs Na/Na+ offer an appealing working voltage of 0.8 V, and researchers have already demonstrated high-rate capability, suggesting potential use in commercial SIBs (Hwang et al., 2015; Li et al., 2017b; Tahir et al., 2016). TiO2 rutile In contrast to anatase, the most thermodynamically stable polymorph, rutile, attracts much less attention as an anode material. The Li-insertion potential of 1.85 V vs Li/Li+ is the most positive among all titania polymorphs, which makes it less suitable for the application as an anode (Kavan et al., 1999). Furthermore, bulk

Voltage (V vs Li+/Li)

3.5

3.0

2.5

R15 R30 R300

2.0

1.5

1.0 0

50

100

150

200

250

300

350

400

Specific capacity (mAh/g)

(A)

(B)

Fig. 8 (A) Structure of rutile (Liu et al., 2013a) and (B) potential-capacity profiles of rutile electrodes with particle sizes of 15 nm (R15), 30 nm (R30), and 300 nm (R300) (the initial cycles at 0.05 A g
1). Jiang, C., Honma, I., Kudo, T., Zhou, H., 2007a. Nanocrystalline rutile TiO2 electrode for high-capacity and high-rate lithium storage. Electrochem. Solid-State Lett. 10 (5), A127–A129. Modified from: Liu, Z., Andreev, Y. G., Armstrong, A. R., Brutti, S., Ren, Y., Bruce, P. G., 2013a. Nanostructured TiO2 (B): the effect of size and shape on anode properties for Li-ion batteries. Prog. Nat. Sci.: Mater. Int., 23(3), pp. 235–244.

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rutile features a lower Li-insertion capacity given by its crystalline structure. Rutile crystallizes in a tetragonal symmetry with the space group P42/mnm, with [TiO6] octahedra sharing edges along the c-direction and corners located along the ab-planes (Fig. 8A) (Yan et al., 2015b). This property explains the different availability of tetrahedral and octahedral sites and a strong anisotropy for Li diffusion. The Li-ion diffusion into rutile is thermodynamically favorable only along the c-axis channels between the tetrahedral sites, with a diffusion coefficient of 10
6 cm2 s
1. In contrast to that, the octahedral interstitial sites are practically unavailable for Li insertion (Koudriachova et al., 2003), with diffusion along the ab-planes being prohibitively slow (Koudriachova et al., 2003; Wagemaker et al., 2001). In addition, the strong repulsive Li-Li interactions in the c-channels, together with trapped Li-ion pairs in the ab-planes, block the c-channels and restrict insertion to well below its theoretical limit (Yang et al., 2009). As a result, in the bulk crystalline form, it is able to accommodate up to 0.1 mol Li only (Aravindan et al., 2015a). The storage capacity of rutile, however, strongly benefits from nanostructuring (Borghols et al., 2008; Hu et al., 2006; McNulty et al., 2017). Thus, it was reported that nanosized rutile shows a stable capacity of 346 mAh g
1 and a superior rate capability (Fig. 8B) (Hong et al., 2016). The absence of phase boundaries in nanoparticles is believed to be the main reason for the enhanced Li-ion solubility. TiO2-B (bronze) Among other polymorphs, TiO2-B (bronze) has attracted particular attention as a promising anode for lithium batteries due to a combination of favorable properties (Pan et al., 2018). TiO2-B features the lowest insertion potential of 1.55 V vs Li/Li+ (Fig. 9C) among all TiO2 materials, which is beneficial for its use as an anode (Pan et al., 2018). TiO2-B crystallizes in a monoclinic structure with C2/m symmetry. It is composed of corrugated sheets of edge and corner-sharing [TiO6] octahedra (Vittadini et al., 2009), which are joined together to form a three-dimensional (3D)

Fig. 9 (A) The structure of TiO2-B, (B) variation of voltage with state-of-charge for discharge then charge of bulk TiO2-B, TiO2-B nanowires, nanotubes, and nanoparticles on the second cycle, and (C) corresponding differential capacity plots at 50 mA g
1 (Liu et al., 2013a). Modified from Liu, Z., Andreev, Y. G., Armstrong, A. R., Brutti, S., Ren, Y., Bruce, P. G., 2013a. Nanostructured TiO2 (B): the effect of size and shape on anode properties for Li-ion batteries. Prog. Nat. Sci.: Mater. Int., 23(3), pp 235–244.

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framework containing 1D channels (Fig. 9A) (Dylla et al., 2013). In the unique open structure of TiO2-B, practically all octahedral sites are accessible, resulting in an insertion capacity close to the theoretical value of 335 mAh g
1 (Fig. 9B) (Pan et al., 2018). Moreover, the structure is capable of buffering structure changes during lithium-ion insertion without lattice deformation, promoting a high reversibility of the insertion/extraction process and a long cycle life (Dylla et al., 2013). Different to other TiO2 polymorphs whose lithium insertion rate is diffusion limited, lithium insertion into the bronze phase has a pseudocapacitive behavior (Zukalova et al., 2005) resulting in a high-power capability. Although practically infinite unidirectional channels provide fast pathways for Li diffusion, long transport distances are responsible for poor rate capability of bulk TiO2-B material (Dylla et al., 2013). Nanostructuring was demonstrated to be a very efficient means to overcome this limitation and to improve the charging rate of TiO2-B (Zhang et al., 2016c). The numerous reported nano-morphologies include, among others, nanoparticles (Ren et al., 2012), nanosheets (Wu et al., 2015), nanowires (Armstrong et al., 2005), nanorods (Aravindan et al., 2013a), and microspheres (Liu et al., 2011) which show a greatly improved rate performance (Fig. 9C). A serious drawback of nanosized TiO2-B is, however, an ICL in the first charging cycle, which was explained by irreversible reactions on the surface (Vittadini et al., 2009). Coating with carbonaceous species, like for other titania polymorphs, is a promising way to minimize these contributions. Another efficient strategy to boost the rate performance of TiO2-B-based anodes and to increase their cycling stability is hybridization with different conductive carbon materials such as reduced graphene oxide (rGO) for the fabrication of nanocomposites (Etacheri et al., 2014; Yan et al., 2015a). Lithium titanate Li4Ti5O12/LiTi2O4 (LTO) Among all Ti-based ceramic anodes, Li4Ti5O12 (LTO) is the most prominent member of the Li solid-solution family Li3+ xTi6
xO12 (0  x  1) (Deschanvres et al., 1971) that has been developed to a marketable product for electrochemical energy storage (including batteries for EVs) (Vezzini, 2014) since its discovery as a Li-insertion host (Murphy et al., 1983) and first electrochemical measurements in 1989 by Colbow et al. (Colbow et al., 1989; Zhao et al., 2015a; Yi et al., 2015). The unique and advantageous properties of LTO as a LIB anode material are mainly caused by its spinel (AB2X4) crystal structure (Fd3m) that can accommodate three Li+ ions per f.u. with a negligible change of 0.07% (Ronci et al., 2002) of the lattice parameters also known as “zero-strain” Li+ insertion (Yi et al., 2015; Zhao et al., 2015a). The spinel structure of Li4Ti5O12 [alternative writing: Li(Li1/3Ti2/3)O4] is built from a cubic-close-packing (ccp) of oxygen atoms (32e sites) with tetrahedrally (8a, 8b, and 48c) as well as octahedrally (16c and 16d) interstitial sites partially occupied by Li and Ti atoms which can be expressed in space notation as [Li3]8a[Ti5Li]16d[O12]32e (Fig. 10) (Zhao et al., 2015a; Yi et al., 2015). Upon cell charging the LTO anode can reversibly accommodate up to three Li+ ions per f.u. in differing tetrahedral (8a) and/or octahedral (16c) interstitial sites within the packed oxygen array. Following Eq. (4) and Fig. 11 (left) describe the insertion of Li+ ions (beyond the present Li+ ions of the Li4Ti5O12) into empty octahedral sites

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Fig. 10 Schematic representation of spinel-type Li4Ti5O12 crystal structure with octahedral and tetrahedral coordination of Ti (16d site) and Li (8a and 16d sites) by O atoms (A) (Schmidt et al., 2015) and ball and stick representation (B) (Vijayakumar et al., 2011) of the unit cell with Li occupation sites (16d, 16c, and 8a for fully lithiated structure). O atoms represented with big gray spheres and Ti atoms represented by small light gray spheres. Remaining small dark gray spheres represent Li occupancies. Panel (A) reproduced with permission from Schmidt, W., Bottke, P., Sternad, M., Gollob, P., Hennige, V., Wilkening, M., 2015. Small change—great effect: steep increase of Li ion dynamics in Li4Ti5O12 at the early stages of chemical Li insertion. Chem. Mater., 27(5), pp 1740–1750. Copyright 2015 American Chemical Society; (B) reproduced with permission from Vijayakumar, M., Kerisit, S., Rosso, K. M., Burton, S. D., Sears, J. A., Yang, Z., Graff, G. L., Liu, J., Hu, J., 2011. Lithium diffusion in Li4Ti5O12 at high temperatures. J. Power Sources, 196(4), pp. 2211–2220. Copyright 2011 Elsevier.

(16c) that cause a further population of octahedral (16c) sites (and removal from tetrahedral 8a sites) by Li+ ions initially present in the spinel structure by electrostatic repulsion. Fully lithiated Li7Ti5O12, also noted as [Li6]16c[Ti5Li]16d[O12]32e, thereby constitutes a rock-salt structure with coinciding lattice symmetry (Fd3m) (Zhao et al., 2015a; Yi et al., 2015).   ½Li8a Li1=3 Ti5=3 16d ½O4 32e + x e
+ x Li +   Ð ½Li1 + x 8a,16c Li1=3 Ti5=3 16d ½O4 32e + x e
  Ð ½Li2 16c Li1=3 Ti5=3 16d ½O4 32e

(4)

Galvanostatic charge/discharge curves reflecting the electrochemical potential of the LTO phase according to its lithiation state are governed by three regimes (Fig. 11B). During charging, initial Li+ ions (α) are inserted into an isostructural Li4+ aTi5O12 phase resulting in a continuous drop of the electrochemical potential stabilizing with the formation of a parallel existing rock-salt Li7
γ Ti5O12 structure. Ongoing Li+ insertion into the Li4+ aTi5O12
Li7
γTi5O12 two-phase system results in a flat operating potential of 1.55 V vs Li/Li+ until the remaining spinel phase is fully converted to the rock-salt phase, which again results in a single-phase system with a final insertion potential of 1.0 V vs Li/Li+ (Zhao et al., 2015a; Yi et al., 2015).

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Fig. 11 Schematic illustration of Li+ intercalation and deintercalation into spinel-type Li4Ti5O12 structure (A) (Hsieh et al., 2011) and galvanostatic charge/discharge curve with indicated one and two-phase potential regions (B) (Zhao et al., 2015a). Panel (A) reproduced with permission from Hsieh, C.-T., Chen, I.L., Jiang, Y.-R., Lin, J.-Y., 2011. Synthesis of spinel lithium titanate anodes incorporated with rutile titania nanocrystallites by spray drying followed by calcination. Solid State Ion., 201(1), pp 60–67. Copyright 2011 Elsevier; (B) reproduced with permission from Zhao, B., Ran, R., Liu, M., Shao, Z., 2015a. A comprehensive review of Li4Ti5O12-based electrodes for lithium-ion batteries: the latest advancements and future perspectives. Mater. Sci. Eng. R. Rep., 98(1–71). Copyright 2015 Elsevier.

The electrochemical potential of 1.55 V vs Li/Li+ of LTO (Ohzuku et al., 1995) limits the overall cell voltage compared to carbonaceous anodes but significantly increases the LIB safety and stability. The relative high potential prevents the SEI formation, the reduction of an organic electrolyte and dendritic lithium growth. Beyond that, LTO is nontoxic, possesses a high thermodynamic stability (compared to carbonaceous anodes) and has a well-characterized stable crystal structure upon lithium insertion making it a “zero-strain” material (Ronci et al., 2002; Zhao et al., 2015a). In addition, LTO anodes have a negligible amount of ICL, are not prone to solvent co-intercalation or electrolyte decomposition and are in principle capable of delivering high-power density (Aravindan et al., 2015a; Zhao et al., 2015a; Yi et al., 2015). Bulk phase spinel Li4Ti5O12 shows a rather poor electrochemical performance due to its low electronic (<10
13 S cm
1) (Chen et al., 2001) and ionic conductivity (1.6  10
11 cm2 s
1) (Takami et al., 2011), which requires nanostructuring and the fabrication of carbon composite anodes to shorten the Li+ ion diffusion length and enhance the conductivity within the electrode (Aravindan et al., 2015a; Zhao et al., 2015a; Yi et al., 2015). Due to a simple and inexpensive synthesis of nanostructured LTO, primarily by sol-gel-based synthesis routes, research advanced rapidly in the recent years so that composite anodes with a high reversibility (Coulombic efficiency >95%) even at high current rates of up to >1000C could be obtained (Aravindan et al., 2015a). Nano- and microstructures accessible primarily by sol-gel-based synthesis range from zero-dimensional (0D) (nanoparticles) over 1D in the form of nanofibers, -wires,

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-tubes, or -rods to two-dimensional (2D) nanoplate or -sheet structures. By use of advanced templating techniques and/or suitable synthesis conditions, a variety of porous, hierarchical, or array 3D structures could be fabricated (Zhao et al., 2015a; Yi et al., 2015; Zhu et al., 2012a). As an example, a mesoporous thin film built from nanocrystalline spinel LTO particles could reach the theoretical 175 mA g
1 up to a rate of 25C (144 s charge/discharge duration) and maintain up to 115 mA g
1 at a rate of 800C (4.5 s charge/ discharge duration) with a capacity retention of 89% after 1000 cycles at a rate of 100C (Fig. 12) (Feckl et al., 2012). For the preparation of anodes and full cells with thicker electrodes enabling highrate and high gravimetric and volumetric capacities, composites with carbonaceous materials [e.g., (reduced) graphene oxide (GO, rGO), carbon nanotubes] and carbon coatings are prepared (Zhao et al., 2015a; Yi et al., 2015). Zhu et al. (2011) prepared spherical LTO microparticles with a porous morphology resulting from close packing of nanoparticle building blocks. The conductivity of the material was significantly increased to 10 S cm
1 by applying a 6 nm homogeneous carbon coating. Due to a relatively high electrode/electrolyte interface of the porous structure increasing and facilitating the Li+ ion flux, a reversible capacity of 126.4 mAh g
1 could be demonstrated at a rate of 20C (corresponding to 3 min charge or discharge time). In addition, the high stability of the composite anode was shown by a capacity retention of 95% after 1000 cycles at a rate of 1C (1 h charge/discharge time) (Zhu et al., 2011). The electrochemical performance of Li4Ti5O12/C electrodes is dependent on various parameters such as the carbon coating thickness and uniformity, the choice of precursor, the preparation technique as well as the synthesis conditions. Thin, <10 nm carbon coatings can be applied by using sugar (glucose), polyvinylpyrrolidone (PVP), polyacrylonitrile (PAN), citric acid (CA), polyacrylate (PAA), maleic acid (MA), or LTO acetate precursors with an adjacent thermal decomposition and graphitization under nitrogen atmosphere (Yi et al., 2015; Zhao et al.,

Fig. 12 Transmission electron micrograph of spinel Li4Ti5O12 nanocrystals forming mesoporous thin film electrode (A). Gravimetric capacity of mesoporous thin film LTO nanocrystal electrodes at different C-rates (B) and capacity retention at a rate of 100C (C) (Feckl et al., 2012). Modified from: Feckl, J.M., Fominykh, K., D€oblinger, M., Fattakhova-Rohlfing, D., Bein, T., 2012. Nanoscale porous framework of lithium titanate for ultrafast lithium insertion, Angew. Chem. Int. Ed. Engl., 124(30), pp 7577–7581. Copyright 2012 John Wiley and Sons.

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2015a). Beyond carbon composites, also metal (Ag, Au, or Cu) composites were prepared to increase the overall conductivity within the electrode and to enable a highcapacity retention of 75% [131 mAh g
1 for LTO/Ag composite (Liu et al., 2012b)] at a rate of 30C (2 min charge/discharge time) (Yi et al., 2015; Zhao et al., 2015a). The electrochemical performance of active LTO phases is greatly affected by the way of their fabrication. Possible preparation techniques include solid-state reactions at high temperatures of 600–1000°C, microwave-assisted synthesis, molten-salt synthesis, hydrothermal, and combustion synthesis (Yi et al., 2015; Aravindan et al., 2015a). Among them, combustion and sol-gel-based LTO anodes show a superior performance as compared to the anodes containing solid-state derived active material. Furthermore, electrodes with hydro- or solvothermally derived LTO product also show an excellent cyclability and high-rate capability (Yi et al., 2015). Full-cell assemblies with high-voltage cathodes like LiCoO2 (LCO) or LiMn2O4 provide cell voltages of 2.5–3.0 V, whereby LTO shows a better compatibility with LiCoO2 than with Mn-based layered oxide cathodes (Zhao et al., 2015a; Aravindan et al., 2015a). For a LTO/LCO full-cell configuration, an ultralong cyclability of 117,000 cycles could be demonstrated ( Jansen et al., 1999; Aravindan et al., 2015a). A cell composed of a LiMn2O4 cathode with a nanostructured LTO anode has exhibited an extended stability for 30,000 cycles at a rate of 10C (6 min charge and discharge time) with a capacity retention of 95% (Takami et al., 2013). A full cell with LiFePO4 cathode in a 18,650 cell configuration retained its full capacity for 20,000 cycles at a charging rate of 10C (6 min charge/discharge time) and still retained 95% after 30,000 cycles at a charging rate of 15C with discharge rates of 5C, respectively (Zaghib et al., 2011). In total, the high structural stability upon Li+ insertion, cyclability, and rate capability render Li4Ti5O12-based anodes suitable for the application in battery electric vehicles (BEVs) or stationary power supply with high-power and safety demands. LTO has been also investigated as an anode material for SIBs (Sun et al., 2013). A capacity of 155 mAh g
1 at 0.91 V was reported and presented the best cyclability among all reported oxide-based anode materials. Two sodium ions were inserted per f.u. of the zigzag layered Na2Ti3O7 phase through a biphasic process at a very low potential of 0.3 V vs Na/Na+—the lowest voltage ever reported for a transition metal oxide in Na-ion batteries (Senguttuvan et al., 2011). The fully reduced phase Na4Ti3O7 formed in this process was reported to crystallize in an ordered rock-salt-type structure (Rousse et al., 2013). TiNb2O7 (TNO) Besides the already commercialized LTO, TiNb2O7 (TNO) is in the focus of research for an application as novel high rate, high-capacity LIB anode material based on its discovery as a Li+ ion insertion host by Cava et al. (1983) and the work of Goodenough and Kim (2011) proposing its application as an anode material. TiNb2O7 crystallizes in a layered monoclinic structure (C2/m space group), where all Ti4+ and Nb5+ metal ions are octahedrally coordinated by oxygen. Edge and corner shared octahedral [MO6] units form a crystallographic shear structure that can reversible accommodate Li+ ions in its interstitial sites (Fig. 13A (a–d)) oriented along the crystallographic [010] direction (Lu et al., 2011). Due to the similar ionic radii of Ti4+

Li4 Li O Ti Nb

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Fig. 13 Schematic representation of calculated primitive cell (A) of (a) LiTiNb2O7, (b) Li2TiNb2O7, (c) Li3TiNb2O7, and (d) Li4TiNb2O7 with Li-ion insertion sites (Lu et al., 2011). Cyclic voltammetry of solid-state synthesized TiNb2O7 particle (B) (Ise et al., 2018). Galvanostatic charge/discharge curves of mesoporous TiNb2O7 templated the block copolymer at different current rates (C) (Guo et al., 2014a). (A) Modified from Lu, X., Jian, Z., Fang, Z., Gu, L., Hu, Y.-S., Chen, W., Wang, Z., Chen, L., 2011. Atomic-scale investigation on lithium storage mechanism in TiNb2O7. Energy Environ. Sci., 4(8), pp 2638–2644. Copyright 2011 Royal Society of Chemistry (Great Britain); (B) Reproduced with permission from Ise, K., Morimoto, S., Harada, Y., Takami, N., 2018. Large lithium storage in highly crystalline TiNb2O7 nanoparticles synthesized by a hydrothermal method as anodes for lithium-ion batteries. Solid State Ion., 320(7–15). Copyright 2018 Elsevier; (C) Reproduced with permission from Guo, B., Yu, X., Sun, X.-G., Chi, M., Qiao, Z.-A., Liu, J., Hu, Y.-S., Yang, X.-Q., Goodenough, J.B., Dai, S. 2014a. A long-life lithium-ion battery with a highly porous TiNb2O7 anode for large-scale electrical energy storage. Energy Environ. Sci., 7(7), pp 2220–2226

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and Nb5+ ions, an antisite disorder can be found for this structure (Guo et al., 2014a; Aravindan et al., 2015a). In contrast to LTO, the unit cell volume of TiNb2O7 is enlarged by 7.22% upon Li+ insertion, which can be a reason for experimentally observed capacity fading (Aravindan et al., 2015a; Guo et al., 2014a). The maximum theoretical capacity of the material depends on the number of transferred electrons and thus on the applied potential window for cycling. For a fiveelectron transfer involving the Ti3+/Ti4+ and Nb3+/Nb5+ redox couples, a theoretical capacity of 387.6 mAh g
1 is obtained. Cycling in a smaller potential window involving only the Nb3+/Nb5+ redox couple and thus a four-electron transfer process TiNb2O7 shows a maximum capacity of 310 mAh g
1 (Lu et al., 2011). The experimentally observed maximum reversible capacity is about 280 mAh g
1 (Fig. 13, right) (Tang et al., 2013) when operated in a potential window from 1.0 to 2.5 V. This corresponds to an insertion of 3.6 Li+ ions per f.u. of TiNb2O7 and is associated with a full reduction of Nb5+ and Ti4+ to Nb4+ and Ti3+, respectively, and a partial contribution by the Nb4+ to Nb3+ redox couple (Tang et al., 2013; Lu et al., 2011). The clearly visible feature between 1.5 and 1.75 V vs Li/Li+ in the cyclic voltammogram of NTO (Fig. 13, middle) originates thereby from the Nb4+/Nb5+ and Ti3+/Ti4+ redox couples, resulting in an average insertion voltage of about 1.64 V vs Li/Li+ at a rate of 0.1C (Lu et al., 2011) (10 h charge/discharge time). The broader Nb3+/Nb4+ redox couple around 1.25 V vs Li/Li+ thereby only partially contributes to the total capacity when cycled down to 1.0 V vs Li/Li+ (Lu et al., 2011). Band structure calculations of TiNb2O7 suggest that the material is an indirect semiconductor with a bandgap of 2.17 eV (Lu et al., 2011), whose conductivity significantly increases upon Li+ insertion due to a transformation to a metal-like band structure (Lu et al., 2011; Tang et al., 2013). To alter the band structure and to increase the conductivity of delithiated TNO, Nb(IV) doping into the Ti sites can be used as was demonstrated for a Ti0.9Nb0.1Nb2O7 stoichiometry (Han et al., 2011). Besides doping, carbon coating or nanostructuring of the TiNb2O7 active phase is necessary to enable fast electron and Li+ ion migration at high charge/discharge rates similar to other low-conducting ceramic anode materials previously discussed in this chapter (Aravindan et al., 2015a). In addition to the increasing power density, nanostructuring also mitigates capacity fading of TNO upon extended cycling attributed to its unit cell volume expansion (Aravindan et al., 2015a), which is the prime issue of using TNO as LIB anode material. In the recent years, different synthesis approaches for the fabrication of TiNb2O7 as high-rate LIB anode material were used, including solid-state reactions of TiO2 and Nb2O5 (Lu et al., 2011), sol-gel routes with polymer templates to generate mesoporous structures (Guo et al., 2014a; Jo et al., 2014) (Fig. 14), and electrospinning of TNO nanofibers (Tang et al., 2013; Aravindan et al., 2015b). Nanostructured mesoporous sol-gel-based TiNb2O7 (Fig. 14) anode material shows a high reversible capacity of 281 mAh g
1 at a low rate of 0.1C greatly surpassing that of LTO. At a rate of 20C (3 min charge/discharge duration) still over 180 mAh g
1 could be realized. The cyclability of nanostructure TiNb2O7 is promising with a capacity retention of 84% after 1000 cycles at 5C (12 min charge/discharge duration) and Coulombic efficiencies close to 100% (Guo et al., 2014a).

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Fig. 14 Schematic diagram of the synthesis and formation of block-copolymer-templated mesoporous TiNb2O7 (Guo et al., 2014a). Reproduced with permission from: Guo, B., Yu, X., Sun, X.-G., Chi, M., Qiao, Z.-A., Liu, J., Hu, Y.-S., Yang, X.-Q., Goodenough, J.B., Dai, S. 2014a. A long-life lithium-ion battery with a highly porous TiNb2O7 anode for large-scale electrical energy storage. Energy Environ. Sci., 7 (7), pp 2220–2226.

Research on nanofiber-based TiNb2O7/LiNi0.5Mn1.5O4 full cells showed an operating voltage of 3 V with a reversible capacity of 100 mAh g
1 at high C-rates and a capacity retention of 90% over 700 cycles (Aravindan et al., 2015b). Due to the superior reversible gravimetric capacity together with its high-rate capability, TiNb2O7 is a potential LIB anode material for future commercialization (BEVs and stationary energy storage batteries) and replacement of LTO, if capacity fading upon extended cycling can be affectively addressed and minimized (Aravindan et al., 2015a; Lu et al., 2011).

Other transition metal oxides: Li3Nd3W2O12—A garnet-type ceramic anode Garnet framework materials composed of Li3A3B2O12 structures gained lots of interest since Weppners’ group reported Li5La3M2O12 (M ¼ Nb, Ta) garnet materials as potential solid-state lithium-ion conductors exhibiting ionic conductivities of >10
4 S cm
1 at ambient conditions (Luo et al., 2018; Aravindan et al., 2015a; Satish et al., 2014). Li3Nd3W2O12 is a member of the garnet framework structure family and was investigated by Goodenough and his coworkers as a possible insertiontype anode material. The crystal structure of Li3Nd3W2O12 can be described as a Li3A3B2O12, where Li occupies square antiprismatic, octahedral, and tetrahedral sites in a 3:2:3 ratio (Fig. 15) (Satish et al., 2014; Cussen and Yip, 2007). The tetrahedral Li-sites are bridged by empty octahedral sharing opposite faces with two tetrahedral sites. Every face of a Li-site is bridged to neighboring Li-sites by the means of octahedral sites creating a 3D interstitial space, which can theoretically host up to 9 mol of Li. However, there is a practical limit of 7 mol of Li per f.u. The W4+/6+ redox couple [Fig. 16A] in Li3Nd3W2O12 can be used to host reversibly Li ions resulting in a theoretical capacity of  106 mAh g
1. This material benefits from a low operating voltage of  0.3 V vs Li/Li+ with high-power capability (Fig. 16). Nevertheless, an irreversible huge capacity loss in the first cycle and the high

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sensitivity toward H2O and CO2 with the formation of insulating carbonates are the main drawbacks of this material (Satish et al., 2014; Xie et al., 2012). In order to address the last point, Satish et al. were able to protect the lithium garnet network and keep it electrochemical active by a uniform carbon coating (Satish et al., 2014). Moreover, Luo et al. show that Li3Nd3W2O12 synthesized via a sol-gel route consists of smaller primary particles and shows a superior rate performance and cycling stability as compared to Li3Nd3W2O12 prepared by a conventional solid-state method (Luo et al., 2018).

Fig. 15 Schematic representation of Li3A3B2O12 garnet framework (A) and octahedral and tetrahedral Li+ ion occupancies (B) (Cussen and Yip, 2007). Reproduced with permission from: Cussen, E.J., Yip, T.W.S., 2007. A neutron diffraction study of the d0 and d10 lithium garnets Li3Nd3W2O12 and Li5La3Sb2O12. J. Solid State Chem., 180(6), pp 1832–1839. Copyright 2007 Elsevier.

Fig. 16 CV curve of Li3Nd3W2O12 with indicated Wolfram redox couples (A) and galvanostatic charge/discharge curve for first two cycles (B) (Luo et al., 2018). Modified from: Luo, M., Yu, H., Cheng, X., Ye, W., Zhu, H., Liu, T., Peng, N., Shui, M., Shu, J., 2018. Sol–gel synthesis and in situ X-ray diffraction study of Li3Nd3W2O12 as a lithium container. ACS Appl. Mater. Interfaces, 10(15), pp 12716–12721. Copyright 2018 American Chemical Society.

Satish et al. and Luo et al. also reported the electrochemical performance on a fullcell level with Li3Nd3W2O12 as anode and LiMn2O4 and LiFePO4 as a cathode, respectively (Luo et al., 2018; Satish et al., 2014). Satish et al. reported that the

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Li3Nd3W2O12/LiMn2O4 full cells deliver an initial discharge capacity of around 115 mAh g
1 at 1C (100 mA g
1) and 82 mAh g
1 after 100 cycles (capacity retention of 71%) within a working window of 2.95–3.9 V (Satish et al., 2014). The Li3Nd3W2O12/LiFePO4 full cell reported by Luo et al. exhibits an initial discharge capacity of around 110 mAh g
1 within a working window of 2–3.4 V, which only slightly decreases during the following 20 cycles to reach 100 mAh g
1 (Luo et al., 2018).

Silicon (SiON, SiCN, SiSnO, N and related)-based ceramic anodes Polymer-derived silicon oxycarbide (SiOC) and silicon carbonitride (SiCN) have emerged as potential anode materials in the middle of the 1990s. SiOC and SiCN ceramics are typically prepared by pyrolysis of organic polymers containing Si, H, C and N, or/and O in an inert atmosphere at 1000–1600°C, or alternatively via a sol-gel approach (Feng et al., 2017b; Rohrer et al., 2017). The microstructure of SiOC consists of an amorphous polymer-like Si-O-C network interpenetrated by a disordered free carbon phase. Silicon atoms in the Si-O-C network are tetrahedrally bonded to oxygen and carbon coincidently forming SiO4
xCx (x ¼ 1–4) building units, which include also SiO2 and carbonrich regions (David et al., 2016; Sang et al., 2018; Rohrer et al., 2017; Mera et al., 2013). The amorphous free carbon phase is composed of isolated carbon nanodomains (lower free carbon content) or a carbon percolation network (higher amount of free carbon) (Fig. 17) (Rohrer et al., 2017). The major lithium storage sides aren’t fully ensured, yet, whether they are the Si-O-C units or the free carbon phase (Graczyk-Zajac et al., 2015). Nevertheless, it is expected that a reversible capacity of up to 1350 mAh g
1 for SiOC ceramics should be possible (Sang et al., 2018).

Fig. 17 Schematic synthesis routes of polymer-derived SiOC and SiCN ceramics and sketches of the resulting microstructures (Mera et al., 2013). Reproduced with permission from: Mera, G., Navrotsky, A., Sen, S., Kleebe, H.-J., Riedel, R., 2013. Polymer-derived SiCN and SiOC ceramics—structure and energetics at the nanoscale. J. Mater. Chem. A, 1(12), pp 3826–3836. Copyright 2013 Royal Society of Chemistry (Great Britain).

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SiCN can be obtained with a broad variety of different microstructures tunable by the choice of polymers, typically polysilylcarbodiimides or polysilazanes. Polysilylcarbodiimides lead to tetrahedrally coordinated silicon SiN4
xCx (x ¼ 1–4) building units and an amorphous carbon phase analogous to SiOC. Polysilazanes initiate nanocomposites of Si3N4, SiC, and a free carbon phase (Fig. 17) (Reinold et al., 2013; Mera et al., 2013). Nevertheless, the intrinsically poor electrical/ionic conductivity of SiOC and SiCN ceramics results in a rather poor electrochemical performance (Sang et al., 2018; David et al., 2016; Feng et al., 2017b) (Fig. 18). SiOC and SiCN ceramics with high free carbon content exhibit a better electrical conductivity and usually a better electrochemical performance (Sang et al., 2018; Feng et al., 2017b; Reinold et al., 2015). Introducing conductive carbonaceous compounds such as carbon nanotubes, graphite, or graphene is another strategy to further improve the conductivity as was demonstrated by Sang et al. (2018) by comparing SiOC and a 3D-graphene-SiOC composite. Using SiCN or SiOC ceramics together with Si or Sn nanoparticles is a promising approach to increase the capacity, as was demonstrated by Rohrer et al. (2017) and Kaspar et al. (2014).

Fig. 18 Galvanostatic charge/discharge curves (A) and cycling performance (B) of SiCN synthesized from polysilylcarbodiimide at 800°C [HN1-800(black)] and at 1300°C [HN11300(gray)], respectively (Reinold et al., 2015). Reproduced with permission from: Reinold, L. M., Yamada, Y., Graczyk-Zajac, M., Munakata, H., Kanamura, K., Riedel, R., 2015. The influence of the pyrolysis temperature on the electrochemical behavior of carbon-rich SiCN polymer-derived ceramics as anode materials in lithium-ion batteries. J. Power Sources, 282(409–415). Copyright 2015 Elsevier.

Na-Super-Ionic-CONductor (NASICON)-type battery anodes NASICON-structure-related anodes—Structure and key properties The abbreviation NASICON standing for Na-Super-Ionic-CONductor describes a class of ceramic compounds with a stable 3D framework built from two types of transition metal-oxygen (MO6 and M0 O6) octahedra that share all corners with sulfate,

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581

phosphate, silicate, or arsenate tetrahedra. The resulting interstitial space in the prototype structure of the AxMM0 (XO4)3 composition can accommodate up to five alkali metal ions depending on the oxidation state of the transition metals M, M0 (Fe, V, Ti, Zr, Sc, Mn, Nb, In) and the polyanion forming center atom X (S, P, Si, As) ( Jian et al., 2017). Due to the interconnected voids in all directions of the structure, an outstanding sodium (also a high lithium and a reasonable K+, Mg+, Ca2+) conductivity is reached which led to an investigation of NASICON structures as solid electrolytes from the discovery in 1976 onwards ( Jian et al., 2017). Details can be found in Section 4. Besides the use as solid ion conductors, the capability of sodium and lithium-ion insertion for energy storage applications was also systematically investigated for this class of materials. The electrochemical potential for (de)sodation or (de)lithiation of a NASICON compound is in the first place affected by the comprised transition metals (M, M0 ) and their oxidation state but is in addition altered by the electron-withdrawing properties of the polyanions (XO4), termed as induction effect in that context. Furthermore, there exist four structural polymorphs (Fig. 19) for a given composition, namely an orthorhombic (Pbna), monoclinic (P21/c), triclinic (C1), and a corundum-like structure that mainly differ in the alignment of M-M0 dimers ([MO6]2[XO4]3) along the crystallographic c-axis and which affects the electrochemical potential of inserted ions ( Jian et al., 2017). The electrochemical potentials for (de)lithiation of selected NASICON structures with differing transition metal redox couple can in principle be ordered as following: V4+/V5+ (4.6 V vs Li/Li+) > V3+/V4+ (3.8 V vs Li/Li+) > Fe2+/Fe3+ (2.8 V vs Li/ Li+) > Ti3+/Ti4+ (2.5 V vs Li/Li+) > Nb4+/Nb5+ (2.2 V vs Li/Li+) > Nb3+/Nb4+ (1.8 V vs Li/Li+) > V2+/V3+ (1.7 V vs Li/Li+) > Ti2+/Ti3+ (0.4 V vs Na/Na+) ( Jian et al., 2017; Senguttuvan et al., 2013). The relatively high operating potentials of most NASICON compounds mainly led to the development and application as cathode materials in solid-state batteries (Huggins, 2008). NASICON structures applicable as anodes are mainly limited to zirconium, vanadium, and titanium-containing compounds with electrochemical potentials ranging from 0.4 V vs Na/Na+ (roughly equals 0.7 V vs Li/Li+) for Na3Ti2(PO4)3 (NTP) Ti2+/Ti3+ (0.7 V vs Li/Li+) up to 2.1 V vs Na/ Na+ NaTi2(PO4)3 (NTP) based on the Ti3+/Ti4+ conversion (Senguttuvan et al., 2013). In a full-cell solid-state battery configuration with suitable solid electrolytes (e.g., NASICON or garnet type) and high-voltage cathodes (e.g., NASICON, layered transition metal oxides or spinel structures) cell voltages of 1.7 V and more are obtained (Yu et al., 2018; Senguttuvan et al., 2013). Although the reversible capacity of NASICON-type anodes is below that of graphite, the ICL due to the initial formation of a metal ion intercalated anode structure is omitted. A further advantage of the NASICON-type anodes conditioned by its relatively high electrode potential is an avoided SEI formation (Chen et al., 2017).

Zr-based NASICON anodes (NZP) NaZr2(PO4)3 (NZP) was among the first compounds structurally characterized in 1968 and later known to be a representative composition of the NASICON structure family (Hong, 1976). The structure is built up from edge-sharing [PO4] tetrahedra with [ZrO6] octahedra that form interstitial voids

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Fig. 19 Schematic representation of NASICON polymorphs: (A) orthorhombic (Pbna), (B) monoclinic (P21/c), (C) triclinic (C1), and (D) corundum like (Anantharamulu et al., 2011). Reproduced with permission from: Anantharamulu, N., Koteswara Rao, K., Rambabu, G., Vijaya Kumar, B., Radha, V., Vithal, M., 2011. A wide-ranging review on Nasicon type materials. J. Mater. Sci. 46(9), pp 2821–2837. Copyright 2011 Springer Nature.

in shape of tunnels for an efficient sodium-ion migration and storage (Fig. 20) (Chen et al., 2017). In the range of x ¼ 1–2.5 Na+ per f.u. NaxZr2(PO4)3 no change in the crystal structure is observed, but to reach the full theoretical capacity with three Na+ ions a change in the structure can be observed with an accompanied volume change of 10% from NaZr2(PO4)3 to Na3Zr2(PO4)3. In a half-cell configuration, an initial capacity of 150 mAh g
1 could be obtained which is close to the theoretical value of 153 mAh g
1 for solid-state synthesized material. The reversibility and stability

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A Na1

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Fig. 20 Schematic representation of NaZr2(PO4)3 (NZP) structure with Na+ ion migration pathways (Hueso et al., 2013). Reproduced with permission from: Hueso, K.B., Armand, M., Rojo, T., 2013. High temperature sodium batteries: status, challenges and future trends. Energy Environ. Sci., 6(3), pp 734–749. Copyright 2013 Royal Society of Chemistry (Great Britain).

of the NZP structure shown by galvanostatic measurements and also by high Coulombic efficiencies over 100 cycles render this composition promising for the application as SIB anodes (Chen et al., 2017; Wang et al., 2014a). Ti-based NASICON anodes (NTP/LTP) Titanium-based NASICON A3Ti2(PO4)3 (A ¼ Na, Li) shares the common 3D framework structure of [metal-O6] octahedra and [PO4] tetrahedra connected by shared oxygen corner atoms. Sodium-containing Na3Ti2(PO4)3 NTP has been shown to be ordered as a triclinic variant of the NASICON prototype structure shown in Fig. 21 (Senguttuvan et al., 2013). Remarkable for the Ti-containing NASICON structure (LTP) that can be used as an anode is its high room-temperature ionic conductivity of 10
6 S cm
1(Aono et al., 1990b), which renders the material a suitable solid electrolyte for SIBs and LIBs. However, comparable to other ceramic NASICON materials, pure NTP and LTP suffer from a low electronic conductivity which requires carbon coating, particle size reduction and transition metal doping to achieve higher rate performances ( Jian et al., 2017; Chen et al., 2017). NTP as well as Li3Ti2(PO4)3, LTP are widely investigated electro-active materials for SIB and LIB applications in terms of their use as anode, cathode, and electrolyte material. However, the insertion potential related to the Ti3+/Ti4+ redox couple with 2.5 V vs Li/Li+ or 2.2 V vs Na/Na+ in LTP or NTP, respectively, are considered to be too high to be used as anodes and too low to be employed as cathodes in nonaqueous LIBs or SIBs ( Jian et al., 2017). To apply Ti-based NASICON as anode in

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Fig. 21 Schematic representation of triclinic Na3Ti2(PO4)3 (NTP) structure (A) and galvanostatic cycling curve (B) at potential region of Ti3+/Ti2+ plateau (Senguttuvan et al., 2013). Reproduced with permission from: Senguttuvan, P., Rousse, G., Arroyo y de Dompablo, M. E., Vezin, H., Tarascon, J. M., Palacı´n, M.R., 2013. Low-potential sodium insertion in a NASICON-type structure through the Ti(III)/Ti(II) redox couple. J. Am. Chem. Soc., 135(10), pp 3897–3903. Copyright 2013 American Chemical Society.

nonaqueous environment, insertion at 0.4 V vs Na/Na+ assigned to the Ti2+/Ti3+ redox couple (Fig. 21B) is used, which provides a reversible capacity of 60 mAh g
1. The reduction of Na3Ti2(PO4)3 to Na4Ti2(PO4)3 thereby causes a structural rearrangement from triclinic to rhombohedral unit cell symmetry (Senguttuvan et al., 2013). Beyond solid-state or nonaqueous battery application, Ti-based NTP (Pang et al., 2014) and LTP (Weng et al., 2017) NASICON structures are investigated as anodes for aqueous batteries in combination with high-voltage layered oxide cathodes (LiMn2O4, LiCoO2, etc.). For this application, the anode is operated at a potential of 2.6 V vs Li/Li+ corresponding to the Ti3+/Ti4+ redox couple located within the stability window of H2O ( Jian et al., 2017; Chen et al., 2017). High-rate NTP/LTP anode material may be synthesized by various approaches; an example is a hydrothermal synthesis of small nanocrystals that are embedded in a carbon matrix to form the anode. In combination with a layered oxide cathode and liquid electrolyte a capacity of 40 mAh g
1 with respect to the combined mass of cathode and anode is reached at high charge-discharge rates. Furthermore, for aqueous full cells with LTP-C composites, a high cycling stability and capacity retention of up to 90% after 1000 cycles at high rates of 6C (equals 10 min for full charge/discharge cycle) could be obtained (Wu et al., 2014; Jian et al., 2017).

2.2.2 Conversion-type ceramic anode materials A lot of attention has been given to CTAMs for LIBs and SIBs, including mainly transition metal compounds (MaXb, M ¼ Mn, Fe, Co, Ni, Cu and X ¼ O, S, Se, F, N, P, etc.). The MaXb materials undergo reversible electrochemical redox reactions with

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585

Li+/Na+ leading to the formation of transition metal particles M and a binary lithium/ sodium compound LinX/NanX. After delithiation, the initial state is (theoretically) formed again (Eq. 5): Ma Xb + ðb  nÞLi + + ðb  nÞe
Ð aM + bLin X Ma Xb + ðb  nÞNa + + ðb  nÞe
Ð aM + bNan X

(5)

The reaction potential determined by the ionicity of the MdX bond is normally in the range of 0.5–1.0 V vs Li/Li+, with increasing ionicity the potential shifts to higher values. The low and safe lithiation potential together with a high theoretical specific capacity are the main advantages of the CTAMs (Lu et al., 2018b; Tang et al., 2015; Zhang et al., 2018a). Nevertheless, the conversion-type compounds suffer from poor reaction kinetics, low electronic and ionic conductivity, huge volume expansion (>200%), large potential hysteresis, and continuous electrolyte decomposition (Lu et al., 2018a,b; Tang et al., 2015; Zhang et al., 2018a). CTAMs used in SIBs reveal additionally a low initial Coulombic efficiency and a poor cycling stability (Zhang et al., 2018a). Nanoengineering is a very promising strategy to address these drawbacks. Downsizing is of great interest because the reaction of the formed nanosized transition metal M with the binary lithium/sodium compound is more favorable due to an increased electrode-electrolyte interphase and a decreased diffusion length enabling faster reaction kinetics. In addition, it has been demonstrated by numerous research groups that nanostructured materials can at least partially accommodate volume changes during cycling. The conductivity of the CTAMs can be further improved by introducing carbonaceous support materials, which inhibits the electrolyte from decomposition (Zhang et al., 2018a; Lu et al., 2018b). Even more efficient is using the CTAMs together with an alloying-type material as composites, which enable to efficiently reduce the voltage hysteresis (Lu et al., 2018b). Iron oxides are well-studied CTAMs in both LIBs and SIBs. Besides the high specific capacities, additional advantages of this class of materials are their low cost, high abundance, environmental friendliness, and the high corrosion resistance. Especially, the hematite α-Fe2O3 and the magnetite Fe3O4 phase are thoroughly investigated. The hematite phase α-Fe2O3 has a theoretical specific capacity of 1008 mAh g
1, involving a six-electron redox reaction (Zhang et al., 2018a; Aravindan et al., 2015a; Zhao et al., 2015b). Fe2 O3 + 6Li + + 6e
Ð 2Fe + 3Li2 O Fe2 O3 + 6Na + + 6e
Ð 2Fe + 3Na2 O

(6)

Nevertheless, the low electrical conductivity and the large volume changes during electrochemical cycling are serious drawbacks of hematite. Nanoengineering, as mentioned before, and the addition of carbon support materials are successful strategies to alleviate this problem, as, for example, proved by Zhao et al. (2015b), Zhang et al. (2018a), Aravindan et al. (2015a), and Kong et al. (2016).

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Zhao et al. (2015b) reported a yolk-shell Fe2O3/C composite anchored on multiwalled carbon nanotubes (MWNT). They could demonstrate that this composite exhibits a promising cycling stability together with high specific capacities for both lithium and sodium storage (Fig. 22). That excellent performance was attributed to the good electrical conductivity provided by the MWNT and the carbon coating and the 3.0 0.2

Current (mA)

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Fig. 22 Hematite phase α-Fe2O3/MWNT/C anode in LIBs: CV curves at a scan rate of 0.5 mV s
1 (A), discharge/charge curves at 0.2 A g
1 (B), cycling performance at 0.2 and 2 A g
1 (C) and in SIBs: CV curves at a scan rate of 0.5 mV s
1 (D), cycling performance at 0.16 A g
1 (E) (Zhao et al., 2015b). Reproduced with permission from: Zhao, Y., Feng, Z., Xu, Z.J., 2015b. Yolk–shell Fe2O3 ⊙ C composites anchored on MWNTs with enhanced lithium and sodium storage. Nanoscale, 7 (21), pp. 9520–9525. Copyright 2015 Royal Society of Chemistry (Great Britain).

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void space which has a buffering effect on the volume changes of Fe2O3 (Zhao et al., 2015b). Magnetite Fe3O4 is also a promising CTAM. It has a better electrical conductivity compared to α-Fe2O3 due to an electron exchange between the Fe2+ and Fe3+ centers. The theoretical specific capacity is a little bit lowered with 926 mAh g
1, involving an eight-electron process (Zhang et al., 2018a; Aravindan et al., 2015a): Fe3 O4 + 8Li + + 8e
Ð 3Fe + 4Li2 O Fe3 O4 + 8Na + + 8e
Ð 3Fe + 4Na2 O

(7)

The reduction of Fe3O4 to Fe0 takes place at a potential around 0.7 V vs Li/Li+ and the back reactions from Fe0 to Fe2+ to Fe3+ occur at around 1.7 and 1.8 V vs Li/Li+. Unfortunately, similar to hematite and the other CTAMs, magnetite also suffers from relatively high volume changes during electrochemical cycling that rapidly lead to the deterioration of the electrode morphology. Among others, Fu et al. (2016) successfully demonstrated that nanosizing and nanostructuring are beneficial for the electrochemical performance of Fe3O4. Combining Fe3O4 nanoparticles with a conductive support like graphene in the form of a composite material is an even more promising approach to stabilize and fully utilize the material (Fu et al., 2016). Spinel oxide NiCo2O4 was first reported as a conversion material for SIBs (Alca´ntara et al., 2002). A reversible conversion reaction was then reported: NiCo2 O4 + 8Na ! Ni + 2Co + 4Na2 O Following this work, many transition metal oxides have been reported as anode materials for SIBs showing a conversion reaction, such as iron oxides Fe2O3 and Fe3O4 (Komaba et al., 2010), cobalt oxide Co3O4 (Rahman et al., 2014), nickel oxide NiO ( Jiang et al., 2014), tin oxides SnO and SnO2 (Lu et al., 2015a), and copper oxide CuO (Lu et al., 2015b). Similar to the conversion-type anodes for the LIBs discussed above, nanosizing and carbon coating were shown to tremendously improve the performance of conversion-type anodes also in SIBs to reach reversible capacities up to 400 mAh g
1. Similar to the above-mentioned CTAMs, Sn, Sb, and Zn-based compounds (MaXb, M ¼ Sn, Sb, Zn and X ¼ O, S, P, etc.) undergo conversion reactions with Li+ or Na+ up to respective metal state. In addition to the conversion process, the formed metals can accommodate further amounts of Li+/Na+ upon subsequent polarization to form LixM or NaxM alloys. The theoretical specific capacity of such conversion and alloying materials is, therefore, higher than that of “pure” conversion-type materials, which is one of their biggest advantages. The alloying/dealloying reaction suffers from huge volume changes upon lithiation and delithiation leading to internal stress and the so-called pulverization, which is probably the most challenging problem of conversion-alloying materials (Fig. 23). Those problems are even more severe in the case of SIBs. The volume expansions are more pronounced compared to the Li+ counterparts due to the larger ˚ ; Li+: 0.59 A ˚ ). Resulting volume changes also cause an ionic radius of Na+ (1.02 A

588

Advanced Ceramics for Energy Conversion and Storage

Fig. 23 Schematic representation of volume changes accompanying conversion and alloying processes in LIB and SIB anode materials.

unstable SEI. Both, pulverization of the electrode and an unstable SEI thereby result in a limited cycling stability and performance (Aravindan et al., 2015a; Zhang et al., 2018a). Tailoring the bulk materials down to the nanosized range is again proven as a promising strategy. In addition, the use of soft carbonaceous support material can be beneficial to buffer volume changes. Among others, tin oxides are considered to be promising conversion-alloying anode materials for SIBs and LIBs due to their environmental friendliness, the low costs, and their high specific capacities. The reaction of SnOx with Li+/Na+ can be divided into two parts. At first, SnOx undergoes a conversion reaction (Aravindan et al., 2015a; Zhang et al., 2018a): SnOx + 2xLi + + 2e
Ð Sn + xLi2 O SnOx + 2xNa + + 2e
Ð Sn + xNa2 O

(8)

Followed by the alloying step: Sn + xLi + + xe
Ð Lix Sn ð0  x  4:4Þ Sn + xNa + + xe
Ð Nax Sn ð0  x  3:75Þ

(9)

According to these equations, a theoretical specific capacity of 1138 and 1494 mAh g
1 for SnO and SnO2 are obtained upon the reaction with Li+, respectively. In the case of Na+, values of 1022 (SnO) and 1378 mAh g
1 (SnO2) are theoretically achievable. Tin oxides suffer from huge volume changes which result in an inferior electrochemical performance as pointed out above. It should be also mentioned that the conversion step is often discussed to be irreversible, especially in the case of bulk material. It is, however, reported that the conversion step becomes reversible or at least partially reversible in the case of nanosized tin oxides. Hence, nanoengineering and also the use of carbonaceous support materials is a promising way to improve the performance. Nanomaterials can provide voids that accommodate the internal structural strain during cycling. Together with the buffering effect of the carbonaceous support materials and the increased overall electrical conductivity, this leads to an

589

Specific capacity / mAhg –1

Ceramics for electrochemical storage

Current / mA

0.5 0.0

1st cycle 5th cycle 20th cycle

–0.5 –1.0 –1.5 0.0

(A)

0.5

1.0 1.5 2.0 2.5 + Potential / V vs Li/Li

ATO70/rGO30

1600 1C

1200

(B)

5C

1C 10C

800

20C 40C

400

60C

SnO2, 70/rGO30

0 0

3.0

2C

20

40 60 Cycle number

80

Fig. 24 Cyclic voltammetry of Sb:SnO2 nanoparticles on rGO (A) and cycling performance (B) of conversion and alloying-type (Sb:)SnO2 nanoparticles on rGO (Zoller et al., 2018). Reproduced with permission from: Zoller, F., Peters, K., Zehetmaier, P.M., Zeller, P., D€ oblinger, M., Bein, T., Sofer, Z.K., Fattakhova-Rohlfing, D., 2018. Making ultrafast highcapacity anodes for lithium-ion batteries via antimony doping of nanosized tin oxide/graphene composites. Adv. Funct. Mater., 28(23), pp 1706529. Copyright 2018 John Wiley and Sons.

improved electrochemical performance. Introducing dopants (such as antimony) into tin oxides is an additional strategy to ameliorate the conductivity and thus the cycling behavior (Chen Jun and Lou Xiong, 2013). Combining these strategies is also very promising as it was demonstrated recently by the Fattakhova-Rohlfing group using nanosized Sb-doped SnO2 and reduced graphene as support material exhibiting an excellent electrochemical performance (Fig. 24) (Zoller et al., 2018). Since the MdS bonds in metal sulfide are weaker than the corresponding MdO bonds in metal oxides, sulfides can be kinetically favorable for conversion reactions with Na+ ions. Therefore, many transition metal sulfides such as MS or MS2 with M ¼ Co, Mo, Fe, Sn, Cu, Mn, Zn, Ni, or Ti, have been studied (Hwang et al., 2017a; Zhou et al., 2016b). Depending on the transition metal elements, the Na+ ion storage mechanism of metal sulfide materials could be classified as the conversion reaction and/or combined insertion and the alloying reaction. The introduction of carbon additives such as graphene or carbon nanotubes into active materials is indispensable to have advantages over conversion materials such as accommodation of large volume expansion/shrinkage resulting in effective stress relief. The alloying materials have been also studied as anodes for SIBs for the same reasons as the conversion materials, that is to say the possibility to react with a large number of sodium reversibly at relatively low operating voltage (<1 V) (Lao et al., 2017). However, the main drawback of these materials (metals, metalloids, or polyatomic nonmetal compounds) is a large volume change during the alloying-dealloying reaction. Because of constrains imposed by the battery packaging, such volume change leads to mechanical stress of the active particles. Due to the relatively large abundance in the Earth crust, silicon has been widely studied but Morito et al. (Morito et al., 2009; Xu et al., 2016), who report that contrary

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Advanced Ceramics for Energy Conversion and Storage

to the 4.4 lithium-ions uptake, only one sodium per Si could be uptaken with a very low diffusion. Nevertheless, Xu et al. (2016) demonstrate an excellent reversibility of Si nanoparticles composed of both amorphous and crystallize Si with a capacity of 279 mAh g
1 at 10 mA g
1. Later, the sodiation/desodiation behavior of microsized and nanosized crystalline Si was investigated and it was shown that the amorphization of Si in the course of the first sodiation leads to the break of the majority of the SidSi bonds, and crystalline Si is transformed into an amorphous Na-Si alloy. Finally, after the desodiation process, the amorphous structure is maintained. Similarly to silicon, germanium could form bonds with only one sodium and cannot store it in its crystalline structure due to the large ionic size of sodium. Then, numerous studies have been performed in order to design nanostructured materials as nanowires or thin films achieving almost the theoretical capacity of 350 mAh g
1 (Hwang et al., 2017a; Dai et al., 2014). Another element of choice from the group 14 of the periodic table remains tin. Indeed, the metal could in theory form Na15Sn4 alloy (847 mAh g
1) through several intermediates. Ellis et al. (2012) and Ong et al. (2011) have experimentally demonstrated that Sn undergoes a reversible electrochemical redox reaction to reversibly form Sn-Na intermetallic phases. The microstructural evolution and phase transformation study proves, however, the volume expansion from 56% for the amorphous NaSn2 to 336% and 420% to Na9Sn4 and Na15Sn4, respectively (Wang et al., 2012, 2015b). Therefore, the major part of the published work deals with the serious volume changes during alloying-dealloying reactions. In order to buffer the volume changes, carbon coatings using sophisticated core-shell architectures were mostly explored (Luo et al., 2016a; Xie et al., 2015), leading to the delivery of a high specific capacity of 443 mAh g
1 and reversible sodium storage properties with negligible capacity fading after 100 cycles. Similarly, the group 15 elements of the periodic table (Sb, P, Bi, and As) offer the possibility of a large specific capacity as anode materials with the same drawback of volume expansion during the charge. Antimony delivers a theoretical capacity of 660 mAh g
1 according to the formation of Na3Sb (fully sodiation state) (Darwiche et al., 2012). The same strategy of carbon coating was then investigated to limit the large volume expansion (390%). Nanostructuring of Sb to 10–20 nm significantly improved the kinetics. For example the preparation of uniform nanofiber structures with Sb nanoparticles embedded homogeneously in the carbon nanofibers leads to a large reversible capacity of 631 mAh g
1 at C/15, a greatly improved rate capability of 337 mAh g
1 at a rate of 5C and an excellent cycling stability for over 400 cycles (Wu et al., 2014a). With a theoretical capacity of 2596 mAh g
1 associated with the formation of Na3P, phosphorous is another element of choice as anode for SIBs. Among three known allotropes of phosphorous, the white one cannot be used due to its instability. The crystalline black phosphorous transforms at 550°C to the amorphous red form. These two forms of phosphorous were tested as electrode materials, showing a huge volume change (490%) during the electrochemical sodiation/desodiation process. An amorphous red phosphorous/carbon composite anode has been reported (Kim et al., 2014) showing a reversible capacity of 1890 mAh g
1 and good rate capability delivering 1540 mAh g
1 at a high current density of 2.86 A g
1. Recently, a variety of

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591

amorphous phosphorous with nano-architectures and 2D or 3D carbon matrices exhibiting high conductivities were applied to achieve high capacities and stable cycling (Ramireddy et al., 2015; Song et al., 2015; Zhang et al., 2016b). Bismuth reacts with sodium as well to form Na3Bi with an intermediate NaBi alloy (Ellis et al., 2014) leading to a capacity of 385 mAh g
1. Binary intermetallic alloys have also been studied which implies the use of an electrochemically inactive transition element (Ni, Cu, Zn, and Mo) together with an active element (Sn, Sb, and Bi). The beneficial role of the inactive element is to buffer volume changes in the course of the alloying-dealloying process. As proposed by Liu et al. (2014) with highly porous Ni3Sn2 microcages composed of tiny nanoparticles, the mechanical strain of Sn during charge/discharge processes is effectively suppressed by the hollow core structure and the presence of a Ni matrix in the hollow microcages. Moreover, homogeneously encapsulated Ni converted from the sodiation of Ni3Sn2 is beneficial for the required electron transport. As a result, a high reversible capacity of 348 mAh g
1 and a stable cycle retention of 91% after 300 cycles at 1C were demonstrated.

2.2.3 Protective ceramic coating of alkali metal anodes There is an increasing need for high-energy density storage to boost the performance of BEVs or large energy storage systems. Solid-state LIBs and SIBs employing (mostly) ceramic electrolytes are intrinsically safer than their conventional moisture-sensitive counterparts employing organic solvents. To further increase the safety of LIBs and SIBs and prevent possible thermal runaways, the use of carbonaceous anodes (graphite, etc.) should therefore be avoided. The use of ceramic anodes, as discussed in the previous part of this chapter, provides a possible solution. However, due to the relative high potential vs Li/Li+ (e.g., most NASICON or LTO anodes >1 V vs Li/Li+) or Na/Na+ (e.g., most NASICON or LTO anodes >1 V vs Na/Na+), the overall capacity of the cell employing ceramic anodes is rather limited. A breakthrough in the energy density of LIBs and SIBs can be achieved only when light-weight lithium or sodium-containing anodes with a very high specific capacity are used as anodes, and light-weight multiion reaction enabling elements such as sulfur and oxygen are used as cathodes. The main reason for the capacity limitation of today’s LIBs is the cathode and its single-ion intercalation mechanism (valid for most layered transition metal oxides like LiCoO2) exhibiting a theoretical capacity at a maximum of only 250 mAh g
1. In contrast, the cathodes based on multiion reactions like S and O2 have a really significantly higher theoretical capacity of 1672 mAh g
1. Metallic lithium itself is an ideal candidate as anode for S and O2 cathodes due to its highest theoretical capacity of 3860 mAh g
1 and the most negative electrochemical potential of
3.04 V vs SHE. Furthermore, the theoretical energy densities of Li-O2 (3505 Wh kg
1) and Li-S (2567 Wh kg
1) systems are remarkably higher than that of present LIBs (387 Wh kg
1; commercialized graphite/LiCoO2). Using metallic lithium directly as anode has been long seen as not achievable due to the following reasons: first, and probably the most important issue is the unrestrained

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formation of dendrites which can lead to short circuits, fire, or explosion in the presence of flammable electrolytes. Moreover, irreversible reactions of the metallic lithium with the electrolyte, or the cathode materials result in a loss of active material and a fast increase of the cell impedance which means a rapid capacity fading. Another problem is caused by the volume changes within the metal anode upon cycling. This can cause lithium corrosion, pulverization of the anode, and the formation of a large amount of inaccessible (electrochemically inactive) lithium. There are different strategies to address those problems, for example, tailoring the anode structure, using new and/or optimized electrolytes or building a protective layer [or artificial SEI] (Guo et al., 2017b).

100

Discharge capacity (mAh/gs)

Cycle # 1400 1200 1000 800 600 400 200

14 nm AI2O3

0

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Cycle # Fig. 25 Discharge capacities and Coulombic efficiencies of bare (A) and ALD-protected (B) Li metal anodes of Li-S cells (Kozen et al., 2015a). Reproduced with permission from: Kozen, A. C., Lin, C.-F., Pearse, A. J., Schroeder, M. A., Han, X., Hu, L., Lee, S.-B., Rubloff, G.W., Noked, M., 2015a. Next-generation lithium metal anode engineering via atomic layer deposition. ACS Nano, 9(6), pp 5884–5892.

Ceramics for electrochemical storage

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Protective coatings have to fulfill certain requirements. They should be stable against lithium and the electrolyte, both mechanically strong and flexible, ionically conductive, and electrically insulating and have a high transference number. Protective layers based on polymers or ceramics are probably the most promising material classes (Guo et al., 2017b; Lin et al., 2017). LiF, Al2O3, Li3N, SiO2, Li3PO4, or LiPON have been, for example, explored as possible inorganic surface protective layers. Atomic layer deposition (ALD) is a very promising method for preparing those layers because it enables a high uniformity with a very low layer thickness down to the subnanometer range at the same time (Lin et al., 2017; Guo et al., 2017b; Li et al., 2018b; Kozen et al., 2015a). Kozen et al. (2015a) prepared ALD-protected Li-metal anodes with a 14 nm layer of Al2O3 and used them in Li-S batteries. They could demonstrate the superior electrochemical performance of the ALD-protected anode compared to the bare Li-foil exhibiting after 100 cycles a specific capacity of around 1080 mAh g
1 compared to 600 mAh g
1, respectively (Fig. 25) (Kozen et al., 2015a).

3

Ceramic cathodes for Li and Na-ion batteries

3.1 Lithium-based ceramic cathodes for energy storage 3.1.1 Introduction Active materials are the foundation of a battery, as they contain the chemical energy, which is to be converted into electrical energy (Placke et al., 2017). Ceramic cathode materials (Whittingham, 2004; Andre et al., 2015) are, therefore, a critical cell component for LIBs. By definition, a cathode is being reduced upon the discharge of a battery, as it accepts electrons from the external circuit. In rechargeable batteries such as LIBs, this process is reversed upon charge: the cathode is then being oxidized and releases electrons into the external circuit. A vast number of inorganic ceramic materials with different crystal structures exist, which can reversibly insert/release lithium ions. There are, however, some material parameters which define the practical feasibility of a cathode material. These include the average redox potential V of the underlying electrochemical process, the number of electrons x (and Li+) transferred per f.u., the molar mass M, and the crystallographic density ρ of the candidate material. The number of reversibly cyclable electrons x is directly correlated with the attainable cathode capacity (Kasnatscheew et al., 2017c). While V and x should be maximized, M should be minimized in order to achieve a high specific energy. It should be noted that practical redox potential and capacities are often significantly lower than their theoretical values, due to overpotentials or instability of the host structure beyond a certain lithium content (Kasnatscheew et al., 2016b). A high crystallographic density ρ is critical for attaining batteries with high volumetric energy density. There should be minimal volume changes during operation to avoid mechanical stress and to allow a long cycle life. Also, the cathode should operate within the stability window of the used electrolyte (Kasnatscheew et al., 2017d). Table 1 lists critical material properties for selected cathode materials for lithium metal and LIBs.

594

Table 1 Physical and practical electrochemical properties of selected cathode materials for lithium batteries at the material level (NCA ¼ LiNi0.80Co0.15Al0.05O2; LMR-NMC ¼ Li1.15Ni0.15Co0.15Mn0.55O2). Cathode material

Electrons transferred x in practice

Crystallographic density [g cm23]

Discharge potential vs Li/Li+ [V]

Specific capacity [mAh g–1]

Volumetric capacity [mAh cm21]

Specific energy [Wh kg–1]

Energy density [Wh L–1]

86.9 160.1 112.0 112.8 97.9 96.1 87.8 157.8 180.8

1 1 1 3 0.6 0.7 0.85 1 1

5.08 4.62 3.22 3.87 5.05 4.60 4.25 3.58 4.31

3.0 1.7 2.1 2.1 3.9 3.7 3.4 3.5 4.0

300 170 240 690 160 200 300 160 130

1524 785 768 2670 758 920 1275 572 560

900 290 500 1449 625 740 1020 560 520

4550 1330 1620 5608 3150 3400 4300 2000 2200

182.7

1

4.40

4.7

140

658

660

2900

The Li-free layered cathodes (MnO2, MoS2, TiS2, FeF3) (Zhou et al., 2010) require a lithium metal anode. The calculated energy densities at the material level neglect volume changes upon (de)lithiation.

Advanced Ceramics for Energy Conversion and Storage

MnO2 MoS2 TiS2 FeF3 LiCoO2 (LCO) NCA LMR-NMC LiFePO4 (LFP) LiMn2O4 (LMO) LiNi0.5Mn1.5O4 (LNMO)

Molar mass [g]

Ceramics for electrochemical storage

595

In addition, practical cathode materials should fulfill additional requirements, which include good electronic and ionic conductivity for a sufficient rate capability and avoid polarization during the charge/discharge process. Nonsufficient electronic conductivity can be compensated by addition of carbon additives to the electrode composite (Zhou et al., 2010; Qi et al., 2015). Large cathode particles (>10 μm) are preferred, as they are beneficial for achieving a high electrode packing density and thus high volumetric energy density. Larger particles result in a lower specific surface area, which in turn reduce the extent of side reactions (e.g., metal dissolution and electrolyte decomposition) with the electrolyte. Ideally, the constituting elements of the cathode should also be nontoxic and low cost. This, however, cannot always be realized, as the currently widely used layered oxide cathodes contain substantial amounts of nickel and cobalt. The first rechargeable lithium batteries in the early 1970s relied on layered transition metal chalcogenides like TiS2, MoS2, or MnO2 as cathode materials and combined with a lithium metal anode, which acted as the main lithium source. Among the chalcogenides, oxides are not only lower in mass, but also have a higher electronegativity, which results in higher voltages due to the inductive effect of the anion. During discharge, x lithium ions are intercalated into the interlayer space of neighboring MX2 slabs (M ¼ metal, X ¼ anion), forming LixMX2. Despite the high reliability and reversibility of the intercalation processes at the cathode, secondary Li/MX2 batteries could never be commercialized at a large scale, due to the insufficient safety of the lithium metal anode, that is, dendrite formation and thermal runaway (Winter et al., 2018). By the use of graphite as a complementary intercalation material at the negative electrode combined with a lithiated layered oxide cathode (i.e., LiCoO2 or LCO), this issue was overcome. The introduction of the LCO/C-based LIB in 1991 marked a major breakthrough and lead to the large commercialization of LIBs (Winter et al., 2018). An additional advantage of this cell chemistry lies in the fact that cells are now assembled in their discharged state, which facilitates production significantly. As a result, lithium ions are since then provided by the positive electrode in all commercial LIBs. Today, three major material classes are used in commercial LIBs, namely (i) layered oxides (Noh et al., 2013), (ii) spinel oxides (Manthiram et al., 2014; Kim et al., 2012b), and (iii) polyanionic cathodes (particularly olivine-type phosphates) (Andre et al., 2015; Ellis et al., 2010a,b). The crystal structures of the corresponding parent compounds LiCoO2, LiMn2O4, and LiFePO4 are depicted in Fig. 26 While layered oxides currently dominate the market of high-energy LIBs, spinel- and olivine-type cathodes are deployed in various high-power applications. Olivine-type materials are also preferred in heavy-duty motive applications, such as buses and trucks. Furthermore, (iv) conversion-type (Cabana et al., 2010) and (v) anionic redox-type materials are considered as possible alternative high-capacity cathode chemistries. In the following chapters, the structure, electrochemistry, optimization strategies, and future challenges of these materials classes will be presented.

3.1.2 Layered oxides Layered transition metal oxides with the generic formula LiMO2 (M ¼ Ni, Mn, Co) are the most widely used cathode materials for LIBs, owing to several favorable properties, including their good cycle life, rate capability, and high volumetric energy density (see also Table 1).

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Fig. 26 Crystal structure (with unit cells highlighted) of (A) layered oxide lithium cobalt oxide ˚ , c ¼ 14.05 A ˚ , α ¼ β ¼ 90°, γ ¼ 120°], LiCoO2 [space group R3m (166), a ¼ b ¼ 2.81 A (B) spinel-type lithium manganese oxide LiMn2O4 [space group Fd3m (227), ˚ , α ¼ β ¼ γ ¼ 90°], and (C) olivine-type lithium iron phosphate LiFePO4 a ¼ b ¼ c ¼ 8.17 A ˚ , b ¼ 6.05 A ˚ , c ¼ 4.73 A ˚ , α ¼ β ¼ γ ¼ 90°]. [space group Pnma (62), a ¼ 10.39 A

As previously mentioned, lithium cobalt oxide (LiCoO2, LCO) is widely considered as the archetype cathode material for rechargeable LIBs. LCO crystallizes in a rhombohedral, layered structure with alternating layers of CoO2 and Li in an ABCABC oxygen stacking sequence along the c-axis (see Fig. 26A). Co and Li occupy the sites 3a and 3b are octahedrally coordinated within a cubic closed packed oxygen array. This structure is often being referred to as the O3-structre, since Li is octahedrally coordinated and the unit cells stacks three layers of lithium along c. The 2D structure facilitates fast in-plane lithium diffusion, which is also expressed in the high lithium diffusion coefficient of DLi ¼ 10
11–10
7 cm2 s
1 (Dokko et al., 2001). The relatively high room-temperature electronic conductivity of 10
4 S cm
1 is beneficial for its use in composite electrodes with low amounts of conductive carbon (Levasseur et al., 2002).

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During the charge process, lithium cations are extracted from the interslab region, while Co3+ ions are being oxidized to Co4+, according to the equation below: III

Charge

+
III IV ƒ! Li C oO2 ƒƒƒ ƒƒƒ ƒ Li1
x Co1
x Cox O2 + xLi + xe

(10)

Discharge

The deintercalated lithium ions migrate through the electrolyte, cross the porous separator toward the negative electrode, while the ejected electrons move through the external circuit. The redox processes of LCO take place at a potential of 3.9 V vs Li/Li+ (see Fig. 27) and are highly reversible with a low potential hysteresis. While the theoretical capacity of LCO amounts to 274 mAh g
1, corresponding to its complete delithiation, its practical capacity is limited due to structural instability of the O3-structure at high states of charge, when the amount of extracted Li exceeds x  0.6 in Li1
xCoO2. In turn, its practical capacity only reaches up to 150–160 mA g
1, as the delithiation process is controlled through an upper cutoff potential of 4.2 V vs Li/Li+, while coated LCO can be charged slightly higher (Kloepsch et al., 2013). Nevertheless, LCO provides a high specific energy of 625 Wh kg
1, also owing to its high redox potential. Its high crystallographic density of 5.05 g cm
3 is unmatched by competing cathode materials for LIBs and is the key to its high volumetric energy density of 3150 Wh L
1. LCO can be synthesized through a facile solid-state process (at 900 °C) at a large scale. The abundance of the key metal cobalt in the earth crust is low (30 ppm) and geographically highly concentrated,

Fig. 27 Comparison of the discharge voltage profile of layered oxide cathodes LiCO2 (LCO), LiNi1/3Mn1/3Co1/3O2 (NMC111), LiNi0.80Co0.15Al0.05O2 (NCA), and Li1.15Mn0.55Ni0.15Co0.15O2 (Li/Mn-rich NMC) vs a metallic lithium anode.

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especially in the Democratic Republic of Congo (Olivetti et al., 2017). The working conditions within a number of cobalt mines in central Africa often raise ethical concerns. Despite its high content of cobalt, a toxic and costly metal, LCO is still widely used in small LIBs for portable consumer electronics (<10 Ah). The high raw material cost as well as safety issues (i.e., reactive oxygen release upon overcharge), however, disqualify LCO for larger-scale applications (e.g., EVs). The isostructural LiNiO2 (LNO), where Ni replaces Co, is not easily accessible, since Ni2+ ions with a 3d8 (t62ge2g) electron configuration are not easily oxidized to Ni3+ state in an octahedral crystal field ( Julien and Mauger, 2013). Synthesis attempts thus often result in the Li-deficient product Li1
xNixO2 (Bianchini et al., 2018, 2019). The tendency of Li+ and Ni2+ to exchange positions due to their similar ionic radii ˚ and Ni2+: 0.69 A ˚ ) is commonly being referred to as “cation mixing” (Li+: 0.76 A (3a/3b), is known to occur in LNO (Myung et al., 2017). As a result, Ni ions can be found within the Li layer, which not only decreases discharge capacity by lowering the amount of accessible lithium, but also impedes rate capability by blocking diffusion pathways for Li+. The tendency of LNO for exothermal oxygen release during operation renders it an unsafe cathode material, which has hindered its practical application despite high practical capacity of 200 mAh g
1. A mixed occupancy of the transition metal site by partial substitution of Ni by Co in LiNi1
xCoxO2 (with x ¼ 0–0.3, e.g., LiNi0.8Co0.2O2) was later found to suppress the unwanted Li/Ni cation mixing and thus support structural ordering (Cho et al., 2000). This strategy leads to an overall improved electrochemical performance, including a higher reversible capacity, faster lithium diffusion, and higher electronic conductivity. The manganese-based layered oxide LiMnO2, whose chemical composition is attractive due to the lack of Ni and Co, is not directly accessible through synthesis, as the formation of the spinel structure is thermodynamically favored. While it can be obtained through Na+/Li+ ion exchange from NaMnO2 (Capitaine et al., 1996), it still transforms to the spinel structure (LiMn2O4) upon charge/discharge cycling. Thus, the complete substitution of Co in LCO by either Ni or Mn results in structurally instable isotypes, which are practically not feasible. Fig. 27 shows an overview of the characteristic discharge voltage profiles of various layered cathodes discussed in this chapter. Layered oxides with mixed occupancy of two [e.g., LiNi0.8Co0.2O2 (Cho et al., 2000) and LiNi0.5Mn0.5O2 (Yabuuchi et al., 2011; Ohzuku and Makimura, 2001a)] and three transition metals [e.g., LiNi1/3Mn1/3Co1/3O2 (Ohzuku and Makimura, 2001b)] have been studied intensively. Solid solutions of LiNixMnyCozO2 (Liu et al., 1999) are prevalent in most of today’s high-energy LIBs, due to their high-energy density, good rate capability, and long cycle life (Noh et al., 2013). These cathodes are typically abbreviated as NMC, followed by three digits, which indicate the transition metal composition, for example, NMC111 for x ¼ y ¼ z ¼ 1/3 or NMC811 for x ¼ 0.8, y ¼ z ¼ 0.1. These NMC cathodes are typically synthesized via a coprecipitation route, where the mixed metal hydroxides or carbonates are precipitated in the desired ratio within an alkaline, aqueous medium (Liang et al., 2014). In the next step, the precipitate is mixed with a Li-containing salt (Li2CO3 or LiOH) and calcined at 800–900°C to form the desired layered structure.

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In NMC, each transition metal has its individual impact on the physical and electrochemical properties of the cathode: Ni: Nickel increases the attainable capacity and enhances the electronic conductivity. Mn: Manganese occurs as electrochemically inactive Mn4+ and enhances the thermal stability and thus safety characteristics of the cathode. Co: The addition of cobalt is known to improve electronic conductivity and structural ordering of the layered structure (suppression of Li/Ni-mixing).

Fig. 28 illustrates the impact of the transition metal composition of NMC, especially the Ni-content x, on their resulting electrochemical properties such as the discharge capacity and cycle life as well as the thermal stability of delithiated (i.e., charged) material. Moving from NMC111 toward Ni contents up to 0.85, both Mn and Co are gradually being substituted. In doing so, the discharge capacity increases from 163 to 206 mAh g
1 within the same voltage window. For the correct interpretation of discharge capacities of NMC and NCA, cf. Kasnatscheew et al. (2016a, 2017b). Since the average discharge voltage remains relatively unchanged, the additional capacity of Ni-rich composition results in higher-energy densities. Furthermore, the electronic conductivity of 1.7  10
8 S cm
1 as well as the lithium diffusion coefficient of 10
8–10
9 cm2 S
1 for NMC811 is roughly three orders of magnitude higher than for NMC111. On the contrary, other critical properties such as cycling stability

Capacity retention Thermal stability

320

Preferred performance

Li[Ni1/3Co1/3Mn1/3]O2

100 95

300 Thermal stability / ºC

85

260

80

Li[Ni0.6Co0.2Mn0.2]O2

75

Li[Ni0.7Co0.15Mn0.15]O2

240

Li[Ni0.8Co0.1Mn0.1]O2

70

220

65

Li[Ni0.85Co0.075Mn0.075]O2

160

170

180

190

200

Capacity retention / %

90

Li[Ni0.5Co0.2Mn0.3]O2

280

210

Discharge capacity / mAh g–1

Fig. 28 Illustration of the impact of the Ni-content x in LiNixMnyCozO2-type cathodes (with x ¼ 1/3, 0.5, 0.6, 0.7, 0.8, 0.85) on the attainable discharge capacity (bottom axis; 3.0–4.3 V), thermal stability of delithiated material (blue dots, left axis) and capacity retention (red dots, right axis) (Noh et al., 2013).

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Advanced Ceramics for Energy Conversion and Storage

and thermal stability at the charged state are negatively affected by extraction of large amounts of Li+ from Ni-rich compositions (Kasnatscheew et al., 2017a). The lower thermal stability, which is also reflected in the higher tendency toward oxygen release ( Jung et al., 2017), reduces the safety characteristics of the resulting battery. Ni-rich cathode compositions are also more reactive, which shows in their higher sensitivity toward exposure to air and moisture, as surface impurities such as Li2CO3 or LiOH form more easily upon storage and processing (Noh et al., 2013; Jung et al., 2018). Furthermore, Ni-rich NMC also exhibits a higher reactivity toward the electrolyte, likely due to the presence of Ni4+ at high states of charge. Recent studies have shown that a structural layered to spinel reconstruction may take place at the surface of NMC particles under high-voltage cycling conditions. This surface layer is likely to contribute to the impedance buildup at the cathode and thus to the capacity fading of the battery (Lin et al., 2014a,b). As Ni contents of x > 0.9 are being investigated for NMC, in order to achieve even higher-energy densities, the stabilization of the cathode/electrolyte interphase (CEI) (Qian et al., 2016; Wagner et al., 2016c, 2017), for example, through “doping” (which is actually a substitution) ceramic or organic coating becomes increasingly important (Bianchini et al., 2018, 2019; Kim et al., 2018). Novel material design concepts, which involve core-shell (Sun et al., 2009; Jun et al., 2017) or concentration gradient (Lim et al., 2015; Pajot et al., 2018) NMC particles, also result in promising electrochemical results. These NMC particles typically involve a high-capacity, Ni-rich bulk, which toward the surface contain less Ni and more Mn, which render them more stable toward the electrolyte [reactions involving the electrolyte include transition metal dissolution (Evertz et al., 2016) or electrolyte oxidation (Wagner et al., 2016a)]. This strategy combines the advantages of Ni- and Mn-rich NMC compositions and have thus resulted in materials with high capacity and extended long-term stability. Currently, NMC532 and NMC622 are the state-of-the-art cathode for batteries in most advanced EVs, while Ni-rich NMC (Ni  80%), due to its higher-energy density and lower raw material cost (low Co), becomes increasingly attractive (Schmuch et al., 2018; Beltrop et al., 2018; Tiax, 2012). The Ni-rich cathode LiNi0.80Co0.15Al0.05O2 (NCA) exhibits very similar electrochemical properties as NMC811 (both 200 mAh g
1) and is a well-established cathode for high-energy LIBs. NCA contains a small fraction of light-weight and low-cost Al3+ ions, which are electrochemically inactive, but stabilize the layered structure and impede cation mixing (similar to Co3+). While NCA is superior to NMC811 in terms of cycling stability and power, it brings slight disadvantages with regard to cost and thermal stability (Myung et al., 2017). Li/Mn-rich “layered-layered” oxides (LMR-NMC) (Lu et al., 2001; Thackeray et al., 2007) which are nanocomposites of classical LiMO2-type layered oxides and the structurally related layered Li2MnO3, represent a potential alternative class of cathode materials (Rana et al., 2014a,b). The addition of the Li2MnO3 phase results in a Li/Mn-rich composition, which enables high specific capacities at low raw material cost (Schmuch et al., 2018). Typically, a molar ratio of x ¼ 0.33–0.5 is chosen for xLi2MnO3 (1
x) LiMO2, which corresponds to compositions with a Li excess of y ¼ 0.15–0.2 in the solid-solution notation Li1+ yM1
yO2. Specific capacities of up

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to 300 mAh g
1 have been reported for LMR-NMC, which are much higher than for stoichiometric NMC (Qiu et al., 2016). Fig. 27 shows the voltage profiles of prominent layered oxides in comparison. As can be seen, LMR-NMC operates in a wider voltage window of 2.0–4.6 V, compared to 3.0–4.3 V for NMC or NCA. The average discharge potential as well as the crystallographic density of LMR-NMC are slightly lower than for NMC (see Table 1). Nevertheless, very high-energy contents of 1 kWh kg
1 and 4.3 kWh L
1 are attainable for LMR-NMC at the material level. Major challenges, however, lie in the material’s rapid capacity and voltage fade upon electrochemical cycling (Croy et al., 2012), which has to date impeded its commercial use. Additional disadvantages are a low initial Coulombic efficiency and energy efficiency (Meister et al., 2016), pronounced voltage hysteresis and low rate capability, which are all linked to the partial anionic redox contribution of the oxygen lattice (O2
/On
) (Assat and Tarascon, 2018). While the origin of the excess capacity stemming from LMR-NMCs remained poorly understood for long, the discovery and understanding of the underlying anionic redox mechanism (Koga et al., 2013a,b) opened a new and wide research field for novel cathode materials involving a reversible anionic redox, which will be presented in Section 3.1.6 (Saubane`re et al., 2016; Xie et al., 2017; Li and Xia, 2017). Fig. 29A illustrates the synthesis processes of transition metal oxide cathodes (Ahmed et al., 2017; Schmuch et al., 2018). The oxides with mixed transition metal

Fig. 29 (A) Schematic illustration of synthesis process of layered LiMO2-type layered oxide cathodes (with M ¼ Ni, Co, Mn, or Al) (Schmuch et al., 2018). (B) and (C) SEM micrographs of commercial NMC-532 cathode at 2.5 k-fold and 500-fold magnification.

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Advanced Ceramics for Energy Conversion and Storage

composition are typically prepared over a co-precipitation route. An aqueous solution of the transition metal salts (e.g., MSO4: M ¼ Ni, Mn, Co), a base (e.g., NaOH or Na2CO3) and a chelating agent (NH4OH) are pumped into a continuously stirred tank reactor (CSTR), in which the transition metal precursor precipitates, which is then filtered and washed. After drying and sieving, the precursor is mixed with a stoichiometric amount of a lithium salt. Lithium carbonate is lower in cost, however, high-nickel NMC cathodes have more favorable physical and electrochemical properties, when calcined with lithium hydroxide (Yakovleva, 2017). The final cathode materials are formed through calcination in air or pure oxygen atmosphere. The as-prepared cathode powder is then de-agglomerated and may optionally be surface treated in order to stabilize the electrode/electrolyte interface and, thus, electrochemical performance (Myung et al., 2017; Chen et al., 2010). Fig. 29B and C shows exemplarily electron microscopy images of NMC532 (LiNi0.5Mn0.3Co0.2O2) cathode material at different magnifications. The typically spherical secondary particles have an average particle size of 10–12 μm (D50) and consist of primary particles in the 100–200 nm range. The ceramic cathode powder has a specific surface area of <0.5 m2 g
1 and a tap density of 2.5 g cm
3.

3.1.3 Spinel oxides Spinel materials, which often refer to lithium manganese oxide and its derivatives, have been extensively investigated as cathode materials for LIBs since the first report on LiMn2O4 (LMO) by Thackeray et al. (1983). Due to the abundance and environmental friendliness of Mn, LMO is considered as low-cost cathode candidate and was commercialized by NEC in 1996 (Numata et al., 2000). It crystallizes in a cubic structure with the space group Fd3m, in which lithium and manganese cations occupy tetrahedral and octahedral sites, respectively, in the cubic closed-packed array of oxygen anions. As shown in Fig. 26B, in the framework of LMO, the edge-sharing MnO6 octahedra form a 3D Mn2O4 sublattice in which the interstitial space is interconnected and can function as 3D pathways for Li-ions diffusion ( Julien and Mauger, 2013; Karim et al., 2013). Being attributed to the reversible (de)intercalation of one lithium ion with redox couple Mn4+/Mn3+, LMO shows two plateaus at 4.1 and 3.9 V (vs Li/Li+) in the potential profile. It has a theoretical specific capacity of 148 mAh g
1 and can provide 120 mAh g
1 in practice (Nitta et al., 2015). The potential profile of Li1
xMn2O4 is divided into two consecutive plateaus, whereas the lower plateau at 3.9 V vs Li/Li+ can be assigned to a two-phase reaction (x ¼ 0–0.4) with the coexistence of two cubic phases and the flat, upper plateau at 4.1 V vs Li/Li+ (x > 0.4) to a single-phase process (Ohzuku, 1990; Kosilov et al., 2017; Xia, 1996). These features are made responsible for the superior rate capability and high redox potential of LMO compared to other cathode materials and favor its use in high-power applications. The main disadvantage of LMO lies in its poor cycling stability. Several mechanisms have been proposed to address capacity fade, including: (i) lattice distortion (Jahn-Teller effect of Mn3+) on the particle surface (Sun and Jeon, 1999), (ii) generation of oxygen defects/vacancies (Wang et al., 2002), (iii) manganese dissolution into the electrolyte ( Jang, 1996), (iv) formation of new phases and loss of

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crystallinity (Aurbach et al., 1999; Huang, 1999; Palacı´n et al., 2000) as well as (v) structural instability arising from the coexistence of two cubic phases during cycling process (Lee et al., 2000; Xia and Yoshio, 1997). It is generally considered that Mn dissolution is the most critical factor, since the dissolved Mn2+ ions are reprecipitated on the surface of anode, thus impairing reversibility of charging-discharging (Kim et al., 2012b). Mn dissolution is reported to follow two mechanisms, one involves acid corrosion by HF through the reaction Li1
xMn2O4 + HF $ λ-Mn2O4 + LiF + Mn2++H2O, and the other one regards the formation of soluble Mn2+ through the reaction 2Mn3+(solid)$Mn4+(solid) + Mn2+ (solution) (Lee et al., 2013a,b). Besides, it is also greatly dependent on the lattice orientation of the surface interfacing with electrolyte (Liu et al., 2016). Consequently, the suppression of Mn dissolution is an effective strategy to overcome the poor cycling stability of LMO. In an effort to improve the cycling stability of LMO, Mn has been partially substituted by different other transition metals in LiMxMn2
xO4 (M ¼ Ti, V, Cr, Fe, Co, Ni, Cu, Zn, etc.) (Obrovac et al., 1998). The resulting products often lead to changes in the potential profiles due to the different distribution of oxidation states in each material. For example, in addition to the process of Mn3+ $ Mn4+ at 4 V, the Co- and Fe-substituted materials LiCoxMn2
xO4 and LiFexMn2
xO4 exhibit a working potential at 5 V, where Co and Fe are electrochemically active with redox couples of Co3+/Co4+ and Fe3+/Fe4+ (Kawai et al., 1998a,b; Shigemura et al., 2001), respectively. However, such a high voltage also leads to a serious decomposition of conventional battery electrolytes, which will finally result in poor cycling stability. In contrast to the Co- and Fe-containing analogs, the materials substituted by Ni and Cr undergo a different electrochemical process (Mikhailova et al., 2013). The pristine structures only contain Mn4+ that the (de)lithiation should be accompanied by the activation of Ni2+/Ni4+ and Cr3+/ Cr4+ redox couples. Since Mn4+ is electrochemically inactive in the operation potential window, Mn dissolution originating from the presence of Mn3+ can be avoided and the substituted materials show an improved cycling stability (Nikolowski et al., 2008). LiNi0.5Mn1.5O4 and LiCrMnO4 are ideal formulas and both compounds have been investigated intensively. Unfortunately, the Cr substitution material suffers from the serious migration of the transition metals from the octahedral sites to the tetrahedral sites (Sigala et al., 1997). This unfavorable occupation impedes the diffusion pathways of Li ions and contributes to capacity degradation. Thus, among all the LMO derivatives, LiNi0.5Mn1.5O4 (often referred to as “highvoltage spinel”) is widely regarded as the most promising one. It owns a practical specific capacity of 135 mAh g
1, and high working potential of 4.7 V vs Li/Li+ (Manthiram et al., 2014), which show an acceptable compatibility with state-ofthe-art electrolytes. Ni substitution leads to a shortening of the M(Mn, Ni)–O bond length, which strengthens the structural stability of the spinel (Kawai et al., 1998a, b). Thus, LiNi0.5Mn1.5O4 usually exhibits a better cycling stability than pristine LMO. Two different crystal structures of LiNi0.5Mn1.5O4, depending on the ordering of nickel and manganese ions in the lattice, have been reported. One is face-centered spinel (space group Fd3m), which is similar as LMO phase except that the Ni and Mn ions, instead of only Mn ions, are randomly distributed in the 16d sites. The other one is primitive simple cubic crystal (space group P4332), where Mn and Ni ions are

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Advanced Ceramics for Energy Conversion and Storage

located in the 12d sites and 4a sites, respectively. The O ions occupy the 24e and 8c positions, while the Li ions are distributed in the 8c sites (Manthiram et al., 2014). In general, the synthesis route has a strong effect on the resulting structure, low calcination temperatures, and a postannealing process lead to ordered phase (Liu et al., 2016; Kunduraci and Amatucci, 2006). Ideally, the oxidation state of Mn in ordered phase is fixed to +4 and remains electrochemically inactive. Nevertheless, the existence of Mn3+ can be detected in disordered phase due to the oxygen loss from the LiNi0.5Mn1.5O4 framework, and the rock-salt impurity phase Li1
xNixO or NixO further appears. The typical voltage profile of ordered and disordered LiNi0.5Mn1.5O4 can be viewed as Fig. 30. The long and flat plateau at 4.7 V is signed to the reactions of Ni2+/Ni4+ redox couple in both phases. The small and sloping plateau near 4 V represents the Mn3+/Mn4+ couple (Liu et al., 2018a,b,c). Since Mn3+ only exists in the disordered phase, the length of this plateau directly correlates with the amount of Mn3+ and furthermore the level of disordering. The electronic conductivity of disordered phase is 2.5 orders of magnitude higher than that of the ordered phase (Kunduraci et al., 2006), which may explain why the rate capability of disordered phase is better than of the ordered phase, although it shows a low specific discharge capacity due to the existence of a phase impurity. Although many advantages have been reported on the LiNi0.5Mn1.5O4 spinel material, its insufficient cycling performance, particularly at elevated temperatures, still hinders practical application. Various factors are responsible for this, for example, the degree of Ni/Mn ordering, electrolyte instability within the high-voltage region, unfavorable particle morphology, and surface properties. Accordingly, different strategies have been proposed, cation doping (substitution) and surface coating are commonly applied. Several cations, for example, Fe (Liu and Manthiram, 2009), Cr (Katiyar et al., 2009), Al (Zhong et al., 2011), Ti (H€oweling et al., 2015), Ru (Kiziltas-Yavuz et al., 2014), etc., have been employed as doping elements for the

Fig. 30 Charge/discharge voltage profile of disordered and ordered LiNi0.5Mn1.5O4 (LNMO).

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high-voltage spinel, and the doped materials show enhanced cycling stability or rate capability. For example, Ru doping can suppress electrode polarization at high current rates and stabilize the crystal structure by preventing the pileup of Li+ on the particle surface. Thus, LiNi0.4Ru0.05Mn1.5O4 can deliver 121 mAh g
1 initial capacity at 10C and maintain 82.6% after 500 cycles (Wang et al., 2011). The role of the surface coating is either to form a protective layer against corrosive HF or to improve the electronic conductivity. Metal oxides and phosphates [such as ZnO (Arrebola et al., 2010), Al2O3 (Kim et al., 2015c), Li3PO4 (Chong et al., 2014), AlPO4 (Liu and Manthiram, 2009), etc.] and conductive carbon (Yang et al., 2011) have proved to be good coating materials for LiNi0.5Mn1.5O4 spinel. Besides, many efforts have also been put on finding compatible electrolytes. Both electrolyte additives, for example, borate- or phosphorus-based derivatives, aromatic/heterocyclic compounds, and alternative solvents or lithium salts, for example, sulfones, organic nitriles, and fluorinated carbonates, show their ability to improve the electrolyte stability at high potentials (>4.6 V vs Li/Li+). An interesting approach to suppress Mn dissolution is the morphology controlling, especially surface orientation. The (110) orientation has proved to be most prone toward Mn2+ dissolution (Hirayama et al., 2007). In contrast, the {100} surfaces are more stable against dissolution and have positive effects on the rate performance. Thus, tailoring the LiNi0.5Mn1.5O4 particles with more {100} surface and less {110} could improve the long-term cycling performance significantly. Finally, because of the high operation potential of LiNi0.5Mn1.5O4, anion intercalation into conductive carbon can be nonnegligible side reaction during charge (Qi et al., 2014). Overlithated (Li-rich) spinel materials, Li1+ xMn2O4 (Hung Vu et al., 2017; Bianchini et al., 2014a,b) and Li1+ xNi0.5Mn1.5O4 (Mancini et al., 2017), attract attention, due to the very high theoretical specific capacity (347 mAh g
1 for Li2.5Ni0.5Mn1.5O4) and energy (1100 Wh kg
1 for Li2.5Ni0.5Mn1.5O4) based on the completed utilized of the redox couples Ni2+/Ni4+ and Mn3+/Mn4+. The capacity can be increased further by decreasing the lower cutoff potential to employ the Mn2+/Mn3+ reaction, but this is associated with rapid capacity fading (Mancini et al., 2016). The poor cycling performance, in general, limits the feasibility of overlithated spinel materials. The phase transformation from cubic to tetragonal structure is the main issue. Further improvement on electrochemical performance and deeper mechanistic studies are required prior to practical application.

3.1.4 Polyanionic cathode materials Polyanionic compounds (Andre et al., 2015; Ellis et al., 2010a,b), with 3D frameworks based on the tetrahedral polyanion units (XO4)n
and their derivatives (XmO3m+1)n
(X ¼ P, S, As, Mo, or W) have been intensively investigated in the past years. The large (XO4)n
polyanions and the strong X-O covalency stabilize the structure and elevate the redox potential of transition metals (Mn+/M(n
1)+ vs Li/Li+) (Manthiram and Goodenough, 1989) through a M-O-X inductive effect (Manthiram and Goodenough, 1989) and thus enable long-term cycling stability and excellent safety properties of the materials (Gong and Yang, 2011; Masquelier and Croguennec, 2013).

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Phosphates: Among the polyanionic cathodes, the phosphate family and its most well-known member LiFePO4 (LFP) are widely regarded as the most representative. LFP is characterized by good thermal stability and high-power capability. It adopts an olivine structure in the space group Pmnb (see Fig. 26C) and was first reported as a lithium insertion material in 1997 (Padhi, 1997). The theoretical and experimental capacities are 170 and 165 mAh g
1, respectively, and its redox behavior is characterized by a flat working potential at 3.5 V vs Li/Li+, which involves the Fe2+/Fe3+ redox couple. In contrast to layered and spinel oxide cathodes, the Li+ diffusion pathway occurs almost exclusively along one dimension along the crystallographic b-axis. The major weaknesses of LFP lie in its low ionic and electronic conductivities. Reduction of the primary particle size, surface modification with conductive agents and cation doping are found to be very effective ways to improve the electrochemical performance of LFP. Many materials, including carbonaceous materials (Ravet et al., 1999), metals (Park et al., 2004; Croce et al., 2002), metal oxides (Muruganantham et al., 2018; Sivakumar et al., 2015), conductive polymer (Park et al., 2007), etc., have been successfully utilized as coating layer. Carbon is by far the most popular choice and the carbon type, the coating amount and the preparation method, etc., have a strong impact on the electrochemical performance of the final product. Since sp2-type carbon exhibits better electronic properties than sp3-type carbon, graphene exhibits a superior behavior compared to the other coating materials (Lung-Hao Hu et al., 2013; Wang et al., 2018a,b,c,d). A breakthrough concerning the commercialization of LFP took place in 2002, when an increase of the electronic conductivity by eight orders of magnitude through very low amount (<1%) cation doping, for example, Zr, Mg, Nb, and Ti (Chung et al., 2002). However, the use of nanosized particles leads to low tap density that result in a low volumetric energy density. Also, nanoparticles exhibit a high sensitivity toward moisture/air and require much care during storage and transportation (Martin et al., 2011). Other well-studied phosphate compounds with olivine structure include LiCoPO4 and LiMnPO4 (LCP and LMP), both of them offer higher working potential than LFP, that is, 4.8 and 4.1 V (vs Li/Li+), attributed to the redox couple of Co2+/Co3+ and Mn2+/Mn3+, respectively, but suffer from the even lower electronic conductivities (Amine et al., 2000). In addition, LCP also brings the risk of electrolyte decomposition and LMP cannot avoid the capacity fade originating from the Mn3+ Jahn-Teller distortion and large volume change during cycling. Recently, much research effort has been devoted to binary and ternary materials, for example, Li(Mn,Fe)PO4, Li(Co,Fe)PO4, and Li(Mn,Fe,Co)PO4, which should ideally combine the advantages of each system. Especially the Li(Mn,Fe)PO4 system shows a very promising performance. The general strategies to enhance the electrochemical performance of LFP are also followed in binary and ternary materials. For example, a composite of LiMn0.8Fe0.2PO4/carbon nanospheres/graphene nanoribbons shows large initial discharge capacity (168.8 mAh g
1), high average working potential (3.9 V vs Li/Li+), good rate capability (83 mAh g
1 at 50C), and excellent longterm cycling stability (90% capacity retention at 2C) (Hou et al., 2018b). Silicates: Polyanionic lithium-ion orthosilicates Li2MSiO4 (M ¼ Fe, Mn, Co, Ni) are a very appealing material class, since they could theoretically allow the utilization of two lithium ions per f.u. through the redox couple of M2+/M4+, thus doubling the

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theoretical capacity that can be obtained from the corresponding phosphates. The reversible (de)lithiation performance with silicates was first reported for Li2FeSiO4 in 2005 (Nyten et al., 2005), which is still considered as the most promising material after comprehensive investigation within the family. Li2MnSiO4 shows fast capacity fading caused by the existence of instable Mn3+ and the phase transformation upon cycling (He and Manthiram, 2014), while Li2CoSiO4 and Li2NiSiO4 arise electrolyte stability issues, since the extraction of the second Li ion would take place at potentials above >5.0 V vs Li/Li+ (Arroyo-de Dompablo et al., 2006). In contrary, Li2FeSiO4 exhibits two redox plateaus at 2.8 V by Fe2+/Fe3+ and 4.0–4.8 V by Fe3+/Fe4+, with a very high theoretical capacity of 331 mAh g
1. Several polymorphs have been confirmed for Li2FeSiO4, such as monoclinic P21/n, orthorhombic Pmn21, and orthorhombic Pmnb (Chen et al., 2013a,b). The differences lie in the connectivity of the LiO4, FeO4, and SiO4 tetrahedra. The formation of certain polymorph is highly dependent on the synthesis temperature. The lithium diffusion pathways in silicate cathodes have not been well studied due to their vast structural variety and the difficulty in experimental probing, but they all suffer from slow kinetics of lithium diffusion and poor cycling stability. The phase changes during the initial cycles are the main reason for the decay (Armstrong et al., 2011; Masese et al., 2014). The design of a rational test protocol may retain a stable structure in the charged state in order to obtain a better cycling performance. Besides, other approaches to address slow kinetics of lithium diffusion, for example, surface modification (Zhang et al., 2013), cation doping (Li et al., 2018a; Kumar et al., 2017) and construction of nanosized particles (Wu et al., 2013b), or porous structure (Zhang et al., 2015; Chen et al., 2013a,b) have also been successfully employed. NASICON and anti-NASICON compounds: The investigations on the (de)intercalation abilities of alkali cations (Li or Na) into NASICON and anti-NASICON structural compounds started in the 1980s (for a detailed information on NASICON, refer to Part XX on ceramic electrolytes). The materials used as cathodes in the LIBs have a general formula LixMM0 (XO4)3. The structure is built on a framework of MO6 and M0 O6 octahedra sharing their corners with XO4 tetrahedra and vice versa (Masquelier and Croguennec, 2013), and are commonly described with the R3 rhombohedral (NASICON) or the C2/c monoclinic (anti-NASICON) phase. Li3V2(PO4)3 is the most studied material and the electrochemical performance is highly influenced by its structure. The rhombohedral phase undergoes a reaction based on the V3+/V4+ redox couple between the composition Li3V2(PO4)3 and Li1V2(PO4)3, and monoclinic phase has all three lithium ions active and thus shows better electrochemical properties (Rui et al., 2014). Although the rate performance of NASICON materials is better compared to the other polyanionic cathode materials due to the 3D network of interconnected conduction pathways, the relatively low electronic conductivity still requires carbon surface coating or other modifications (Tan et al., 2018). Many other polyanionic frameworks were also investigated as the LIBs cathode materials. Fluorophosphates (LixMPO4F, M ¼ V, Fe, Mn, etc.) combine the strong electronegativity of the fluorine with the inductive effect of polyanion units (XO4)n
and show an enhanced structure stability with higher working potential vs the phosphate at the same composition ( Julien et al., 2015). A similar concept was

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applied in the design of fluorosulfates (LixMSO4F, M ¼ Fe, Co, Ni, Mn, Zn, etc.), besides the strong electronegativity of both sulfur and fluorine (Bensalah and Dawood, 2016), the substitution of (PO4)3
by (SO4)2
can result in an increase of the redox potential by 0.8 V (Gong and Yang, 2011). Borates (LiMBO3, M ¼ Mn, Fe, Co, etc.) are another attractive class of materials. They show the opportunity for high specific capacity because the (BO3)3
unit is the lightest polyanion group (Allen et al., 2002). Though great progress has been made in enhancing the electrochemical performance of polyanionic cathodes, most materials still have technical issues such as a low gravimetric capacity, capacity fade, and complicated synthesis route that have to date hindered their commercialization and need to be addressed in the future work.

3.1.5 Conversion-type cathode materials Besides the previously discussed classical cathode materials for LIBs, which are based on an intercalation or insertion mechanism for lithium uptake, also materials exist which react reversibly with Li+ via a conversion reaction (Andre et al., 2015). A generic reaction formula for a conversion-type active material is described by the following equation (Andre et al., 2015; Cabana et al., 2010): Ma Xb + ðb nÞ Li $ a M + b Lin X

(11)

l

with a metal M (mostly a transition metal) and the corresponding anion X at the oxidation state n. The scheme of a characteristic conversion reaction is shown in Fig. 31 for the generic conversion material MO. Upon lithium uptake, the active material MaXb reacts to nanosized domains of a metal M within a (often electronically insulating) lithium salt matrix LinX. While 3.0 M O

MO

Potential (V vs Li/Li*)

2.5 O

2.0

Li MO 1–5nm

1.5

1.0

Rev. delithiation

1st cycle irr. capacity

M 1–5nm

Electrolyte decomposition

MnO emf

Rev. lithiation

Li2O + M

0.5 1st lithiation

0.0 0

20

40

60

80

100

Capacity (%)

Fig. 31 Reaction scheme for the (de)lithiation of a characteristic conversion-type active material MO (also applicable for X ¼ S, N, P, F) (Courtel et al., 2012).

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classical cathodes involve a host structure, which is typically retained with small changes of the lattice parameters upon uptake/release of lithium, the structure of conversion-type materials is “converted” to two spatially separated phases. The underlying conversion mechanism often brings slower reaction kinetics as well as a pronounced voltage hysteresis, which results in a low energy efficiency of this material class (Meister et al., 2016). The repeated structural reorganization upon electrochemical cycling, which often involves significant volume changes, results in an insufficient cycle life. In addition, many conversion-type materials have a low initial Coulombic efficiency (see Fig. 31) as well as low electronic and ionic conductivities. To compensate the latter, large amounts of conductive additive may be required within the composite electrode. In contrast to current LIB cathode materials, many conversion-type materials are inherently Li free and thus require either a lithium metal anode or a prelithiation process for their operation. Furthermore, the average discharge potentials of conversion-type cathodes reported to date are significantly lower than for state-of-the-art cathodes in LIBs (see Table 1). The appeal of conversion-type active materials lies in their extraordinarily high specific capacities, which in many cases exceed intercalation materials by far and are often reasoned by a multiple electron transfer in conjunction with a low molecular weight. The low cost and in some cases low toxicity as well as good thermal stability are further advantages. While there are many CTAMs, which include metal oxides, sulfides, nitrides, phosphides, and fluorides (X ¼ O, S, N, P, F) (Cabana et al., 2010; Courtel et al., 2012; Jia et al., 2016), there are only few candidates for positive active materials (Cabana et al., 2010). The most promising materials include fluorides such as FeF3 (Plitz et al., 2005; Wu et al., 2009), phosphates such as BiPO4, chlorides such as CuCl2 (Dey, 1989), and sulfides such as CuS. The electron-withdrawing character of fluorides and chlorides give rise to high theoretical redox potentials. The ionic bond of metal halogenides typically gives rise to a high thermal stability (Zhou et al., 2010). Among the fluorides (Amatucci and Pereira, 2007) (MFx with M ¼ Ni, Cu, Co, Mn), especially FeF3 is attractive, due to its high theoretical specific capacity of 712 mAh g
1 at a potential of 2.74 V vs Li/Li+ (Courtel et al., 2012). The lithiation of FeF3 upon discharge occurs in two consecutive steps. First, lithium is inserted into interstitial sites of FeF3 in a voltage range of 4.5–2.0 V, giving rise to LiFeF3 and capacity of 237 mAh g
1. In the second step, the actual conversion of LiFeF3 to nanosized metallic Fe and LiF takes place at voltage of 1.75 V, providing an additional capacity of 475 mAh g
1. While practical capacities of up to 690 mAh g
1 have been reported, the experiment was carried out with 50 wt% of active material in the electrode composition, at a low current rate of C/50 and an elevated temperature of 60°C. Despite the low current, a pronounced potential hysteresis of at least 800 mV between charge and discharge process remains (Martha et al., 2014). Nevertheless, the achievable energy contents of FeF3 at the material level of roughly 1450 Wh kg
1 and 5600 Wh L
1 are outstanding compared to most other cathode materials, when assuming a practical discharge voltage (Andre et al., 2015) of 2.1 V (see Table 1). Although relatively stable cycling (230 mAh g
1 for 150 cycles) and rate capability (>200 mAh g
1 at 10C, 60°C) (Martha et al., 2014) have been reported, further work on the improvement of FeF3 [e.g., doping

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(Liu et al., 2012a), coating (Zhang et al., 2012), or composite formation (Prakash et al., 2011)] needs to be carried out. Same applies for FeF2, which shows similar characteristics with lower theoretical and practical capacities (Qtheo: 571 mAh g
1; Qprac: 230 mAh g
1) (Gu et al., 2015). CoF2 is another promising representative of the fluoride family, which provides a slightly higher practical discharge voltage of 2.6 V with a marginally lower theoretical and near-theoretical practical capacity (550 mAh g
1 at C/50) (Wang et al., 2015c). In summary, conversion-type cathode materials hold out highly attractive theoretical specific capacities and energy densities at the material level. As their electronic conductivity is often low, their energy density is in practice reduced through the addition of conductive additives at the electrode level. Despite numerous efforts to overcome the many technical challenges of this material class, such as their voltage hysteresis, low rate performance, and short cycle life, further research needs to be conducted in the future for conversion-type cathodes to be used commercially.

3.1.6 Anionic redox-type cathode materials The classic cathode materials, as being discussed in the previous part in this chapter, such as layered LiMO2, spinel-type Li2MnO4, and polyanionic LFP, are mostly operated through reversible insertion/extraction of lithium ion out/in the host structure involving cationic redox of transition metal ions. In the past few years, compounds with the activation of anionic redox (mostly oxygen), mainly layered materials, have attracted much research interest. The redox reaction of oxygen has been first proposed during the study of electrochemical behavior of LiCoO2 at high potential (Tarascon et al., 1999), but the identification of the electrochemical activity of Li2MnO3 caused much attention (Rana et al., 2014b). Since the further oxidation of Mn4+ in octahedral coordination is not possible within the electrochemical stability window of electrolyte, the obtained capacity was believed to be associated with oxygen redox. The Tarascon group has contributed an excellent work to the understanding of the working mechanism of oxygen redox (Assat and Tarascon, 2018): part of the O 2p orbitals is weakly bound, when pointing toward Li in the structure due to the relatively small overlap with the Li 2s orbital. They behave like O nonbonding states and are located above the stabilized (M–O) bonding band in the band structure. Meanwhile, the partially filled (M–O)* antibonding band splits into empty upper- and filled lowerHubbard bands (UHB and LHB, respectively). The relative position of LHB and O 2p nonbonding band directly determines the occurrence of oxygen redox reaction and its reversibility. As shown in Fig. 32, (i) when the LHB lies higher than the O 2p nonbonding band, the electrons are exchanged only from the filled LHB and the material undergoes one band redox with cationic ions. (ii) When the LHB is lower than the O 2p nonbonding band, electrons are removed from the O 2p band and gaseous oxygen may be released irreversibly, the process is still on band reaction but with anionic redox. (iii) The situation in between, for example, the LHB overlap with O 2p nonbonding band, electrons can be exchanged from both LHB and O 2p band, the distortion of MO6 octahedra is consequently viewed and the OdO distances are shortened to enable the emergence of stabilizing M–(O2)n
interactions, as the results, the anionic redox is reversible and two band redox reaction is realized for extra capacity.

Ceramics for electrochemical storage

Li1/3M2/3O2 slab in Li-rich layered Li2MO3

MO2 slab in layered LiMO2

O EF

(M–O)*

M

(M–O)*

EF O 2p non bonding

Li

(M–O)

(M–O)

(A)

(B) Case 1, U << D UHB

U

Case 2, U/2 ≈ D e–

UHB

U D

(M–O)

U

e–– e

LHB Δ

Case 3, U >> D

LHB

UHB e–

MO6 distortion, O–O shortening

D LHB

(M–O)

O2 (gas) release

(M–O) Cationic redox

(C) (one-band redox)

Reversible anionic redox

(D) (two-band redox), extra capacity

(E)

Irreversible anionic redox (one-band redox)

Fig. 32 Crystal structure and the corresponding band structure of LiMO2 (A) and Li/Mn-rich Li2MnO3 (B). Different cases of the Li2MnO3 band structure (C–E) when taking Mott-Hubbard splitting into account. Modified from: Assat, G., Tarascon, J.-M., 2018. Fundamental understanding and practical challenges of anionic redox activity in Li-ion batteries. Nat. Energy, 3 (5), 373–386. Available from: . 611

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Unfortunately, the Li2MO3 (M refers to 3d transition metals) as well as the Li/Mnrich NMC materials have proved to react according to (ii): the oxygen loss can be always detected from the initial charge process of these materials that cause fast capacity fading (Luo et al., 2016b; Hy et al., 2014; Yu et al., 2009). Considering the position of LHB strongly depends on the properties of transition metals, that is, it moves down from left to right in the periodic table (Tin+ ! Nin+) because of orbital contraction, but up from 3d to 5d metals due to orbital expansion, the fabrication of compounds (Li-rich type in the formula LixMO3) by employing 4d and 5d metals, for example, Mo, Ru, Ir, etc. or group 4 and group 5 metals, for example, Ti, V, Nb, etc. could be a reasonable approach. The first result was published on Li2Ru1–ySnyO3 (Sathiya et al., 2013) and was followed by Li1.211Mo0.467Cr0.3O2 (Lee et al., 2014), Li3NbO4-based and Li1.3NbxM0.7
xO2 (Yabuuchi et al., 2015), Li2In1
xSnxO3 (McCalla et al., 2015), Li1.3Nb0.3V0.4O2 (Yabuuchi et al., 2016), Li4=3 Mo6 + 2=9 Mo3 + 4=9 O2 (Hoshino et al., 2017), Li1.15Ni0.45Ti0.3Mo0.1O1.85F0.15 (Lee et al., 2017), Li2MnO2F (House et al., 2018), Li2Mn2/3Nb1/3O2F (Lee et al., 2018), etc. Much higher specific capacities compared to conventional cathodes can be expected from these materials, for example, both Mo- and Nb-contained materials are found out to be able to deliver capacities as high as 320 mAh g
1 at low specific current of 10 mA g
1. It is worth to note that the reversible oxygen redox is not only found in layered structures with 4d or 5d metals, but also in other phases with 3d metals. Freire et al. (2016) have reported the electrochemical performance of Li4Mn2O5 material, it crystalizes in a rock-salt-type structure and exhibits highest capacity of 355 mAh g
1 (initial cycle with redox couples Mn3+/Mn4+/Mn5+ and O2
/O
) so far on Li-Mn-O family (Freire et al., 2016). In general, the materials with oxygen redox face the shortages of poor cycling stability and low ionic conductivity that call for the further efforts from both theoretical and experimental points of view.

3.1.7 Outlook Unlike the anode, where graphite, silicon, and lithium metal have been the materials for practical consideration (Placke et al., 2017), various cathode materials have been commercialized. Key performance and cost indicators are significantly different for the various cathode materials (Fig. 33). Depending on the application requirements and forthcoming improvements, we will still see parallel development and commercialization of multiple cathodes in the future. Also, cathode materials will still play a key role for the performance, and in particular for the costs, of LIBs and lithium metal batteries.

3.2 Cathodes for Na-ion batteries 3.2.1 Introduction A cathode material must be composed of an active material having a high potential (high redox couple potential) and a high specific capacity (sodium-rich material and of low molar mass) in order to obtain an important specific energy. Cyclability

Ceramics for electrochemical storage 613

Fig. 33 Comparison of key performance indicators of major cathode chemistries in spider diagrams (LMP ¼ LiMnPO4).

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and lifetime are also important parameters and these are governed by the structural stability of the compound. Finally, the power delivered by the material will be due to kinetic phenomena such as ionic diffusivity. In addition to all these important electrochemical characteristics, the chosen material should also be obtained at low cost, be nontoxic, and be chemically stable. The choice of the elements of the periodic table is, therefore, particularly limited. Four major families of materials are available for SIB cathode applications; transition metal oxides (Clement et al., 2015; Han et al., 2014b; Hwang et al., 2017a; Kubota et al., 2014), polyanionic materials (Barpanda et al., 2018; Fang et al., 2017a,b; Ni et al., 2017), metal hexacyanometalates (Ma et al., 2017b), and organic compounds (Xu et al., 2018a,b,c). Among them, polyanion-type and oxides compounds, both ceramic materials are perceived as the most promising for future Na-ion batteries thanks to their structural stability, safety, and appropriate operating potential. As mentioned previously, this distinction also exists for LIB technology with LiFePO4 (LFP) (Padhi et al., 1997) on one side and, on the other hand, LiCoO2-derivatives (Mizushima et al., 1980) such as Li[NixMnyCoz]O2 (NMC) or Li [NixCoyAlz]O2 (NCA). The efforts trying to generate suitable cathode electrode materials for Na-ion batteries have been mostly focused on the duplication of the reported lithium-based electrode materials adapted to the sodium system. This is the reason why rapidly the same families of materials have been reported these last 5 years in literature (Fig. 34). Hereby a detailed description of these various families of materials is given. Energy density (Wh/kg) 300 400 500 600

200

5.0

Na[Ni1/2Ti1/2]O2 Na0.71CoO2 Na2FePO4F Na3V2(PO4)2F3 Na V (PO )

4.5

3

Volts vs Na+/Na

4.0

2

Na2/3[Mn2/3Ni1/3]O2 Na[Sb1/3Ni2/3]O2 NaFePO4

4 3

Na[Mn1/3Ni1/3Co1/3]O2

NaFeO2

Na[Fe0.5Co0.5]O2

Na0.44MnO2

3.5

Na2/3[Mg0.28Mn0.72]O2 β−NaMnO2

Na2FeP2O7

3.0

Na2/3[Fe1/2Mn1/2]O2 Na4Fe3(PO4)2P2O7

2.5

Na0.6MnO2

Na[Fe1/2Mn1/2]O2

2.0

α−NaMnO2

NaCrO2

NaFe(SO4)2(OH)6

V2O5

Na[Ni1/4Fe1/2Mn1/4]O2

NaxVO2

Na2Mn3O7

NaTi2(PO4)3

1.0

500 400

Na1.5+yVO3

1.5

600

300 NaTiO2

200

0.5 0.0 0

25

50

75

100

125

150

175

200

225

250

Capacity (mAh/g) Fig. 34 Comparison of some cathode materials for Na-ion batteries; in green the oxides and in red the polyanionic-based materials.

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615

3.2.2 Transition elements oxides The transition metal oxides are mostly represented by the formula NaxMO2 (0 < x < 1), in which M is composed of one or several transition metals. These materials can be classified into two categories: 3D structure for x < 0.5 and 2D structure for x > 0.5. The presence of tunnels and layers facilitates the insertion and extraction of sodium into these oxides. However, it is important to note that some other examples of tunnel or lamellar compounds cannot be assimilated as NaxMO2. The mono-metal oxides present some advantages, such as the good ionic diffusivity of cobalt oxides, the high theoretical specific capacity of manganese oxides, and the high redox potential of iron- or nickel-based materials. However, their respective overall electrochemical activities are relatively limited. The substitution of electrochemically active or inactive elements has proved to be an interesting approach to address concerns about low capacity retention, low operating voltage, and/or structural instability (Xiang et al., 2015). Incorporation of active elements makes it possible to adjust the redox reaction to a target voltage in order to obtain a smoother charge/ discharge profile and a higher operating voltage. The use of inactive elements allows stabilizing the oxide layers, especially after desodiation, which results in better longlife cycling. However, this substitution of inactive elements reduces the amount of sodium inserted/extracted and thus reduces the specific capacity of the material. For industrial applications, three of the most interesting elements for use in cathode materials for SIBs are iron, nickel, and manganese. Fe has the advantages of being very abundant, not harmful to the environment, and has a high operating voltage. Ni has high electrochemical performance because of the Ni4+/Ni2+ redox reaction which has a high operating voltage. Mn has the particularity of giving strong specific capacities but a lower average potential. Coupling these elements with Co, Ti, Li, or Mg can lead to specific capacities and higher operating voltages while maintaining a good structural stability. Some examples of 2D and 3D metal transition oxides Nax[MM0 M00 ]yOz (with M, M0 ,and M00 ¼ Ni, Fe, Mn, Ti, …) are presented here.

3D structure Na4Mn4Ti5O18-type structure In Na4Mn4Ti5O18, the titanium and manganese ions are located in two different sites: all the Ti4+ ions and half of the Mn3+ ions are in octahedral sites (MO6), while the other Mn3+ ions are in square-based pyramidal sites (MO5) (Mumme, 1968). The organization of these polyhedra leads to the existence of two types of tunnels: the first S-shaped with two Na sites (Na1 and Na2) and the second, smaller, with only one Na site (Na3). The trivalent manganese ions in pyramidal sites cannot be oxidized in the tetravalent state, thus, 20% of sodium ions cannot be extracted from this structure. Guo et al. (2014a,b) were the first to study the material Na0.61Ti0.48Mn0.52O2 as cathode material for Na-ion batteries. Despite the high amount of Na (<0.5), the material shows a tunnel structure (Guo et al., 2016). The authors obtained a reversible capacity of 86 mAh g
1 at 0.2C between 1.5 and 4.0 V, indicating good potential for this tunnel-type material. By carbon coating and nanostructuration, the performance was improved to 122 mAh g
1 after 150 cycles at the same rate ( Jiang

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et al., 2015). The substitution of a part of the Mn by Fe increased the average potential to 3.56 V, for a capacity of 90 mAh g
1 (Xu et al., 2015a,b). Moreover, in aqueous medium, Na4Mn4Ti5O18 forms a nonstoichiometric phase NaxMn4Ti5O18 (2.8 < x < 3.2), leading to a capacity of 36 mAh g
1 in 1 M Na2SO4 electrolyte ( Jayakumar et al., 2015). Na4Mn9O18 (Na0.44MnO2 in NaxMnO2 conventional formula) is isostructural with Na4Mn4Ti5O18 and shows the same tunnel structure. Sauvage et al. (2007) tested Na0.44MnO2 in Na-ion cells between 2 and 3.8 V for a composition varying from x ¼ 0.18 to 0.64 (NaxMnO2) and obtained a capacity of 140 mAh g
1 at a slow rate of C/200. The capacity decreased quite rapidly, and only half of the initial capacity retained after 50 cycles. They assumed that the Na1 and Na2 sites were easily accessible, whereas the ions in the Na3 sites were generally not extracted. This material also showed a good stability in the presence of the aqueous electrolyte Na2SO4 (Whitacre et al., 2010) and in the NaTi2(PO4)/Na0.44MnO2 system the cell was able to operate for >1000 cycles at 100C. By forming nanowires, Cao et al. (2011) improved the lifetime of the material with a capacity retention of 77% after 100 cycles at C/2 rate. Indeed, the nanostructuration generally makes it possible to improve the performance of the electrode materials. Nanoparticles, nanowires, or nanosheets allow higher intercalation/deintercalation rates by shortening diffusion paths. Lifetime is also greater because of the smaller structural changes that apply to smaller volumes. A multiangular rod-shaped Na0.44MnO2 also reached retention of 99.6% over 2000 cycles at 1000 mA g
1 (8.3C) (Liu et al., 2017b). Na0.44MnO2 was also mixed with the cryptomelane-type KMn8O16 to form the composite Na0.5K0.1MnO2 by a coprecipitation method. This material showed a highly reversible capacity of 82 mAh g
1 after 300 cycles at 1C, in the potential window 4.3–1.5 V (Wu et al., 2016b). Other 3D materials In manganese oxides presenting diffusion/intercalation in the form of tunnels, the study of MnO2 was quite obvious. This material, which has at least nine polymorphs, has many different structures in tunnels and layers. Due to the relatively large size of their tunnels, the α- and β-MnO2 phases have already been studied for the insertion of sodium (Su et al., 2013a). As shown in Fig. 35, the α-MnO2 phase is made of double chains of edge-sharing MnO6 octahedra, interconnected by the corners to form a structure in tunnels 2  2 and 1  1. With an adequate morphology of nanowires and thanks to the existence of these large 2  2 tunnels, the insertion of Na+ has already been carried out in α-MnO2. Su et al. (2013a) obtained high first discharge and charge capacities of 278 and 407 mAh g
1, respectively. This capacity decreases rapidly after this first cycle and after 100 cycles, the discharge capacity is 75 mAh g
1. β-MnO2 is only made of 1  1 tunnels formed by chains of edge-sharing octahedra and this material showed a higher retention capacity than α-MnO2 (Su et al., 2013b). The discharge capacity (298 mAh g
1 for the first discharge) is maintained at 145 mAh g
1 at the hundredth cycle. Out of the Na-Mn-O system, it is also interesting to present the case of the nanotube-type compound Na2V3O7. One sodium ion can be reversibly extracted from this material in the potential range 4.5–2.0 V at C/20 rate (Adamczyk et al., 2018).

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617

Na0.44MnO2 a Na1 b

Na2

Na3

Mn4+ Mn3+ Na+ O2–

(A) b a

(B)

α-MnO2

β-MnO2

Fig. 35 (A) Structural view of Na0.44MnO2 along the c-axis with the three sodium sites labeled Na1, Na2, and Na3. (B) Structures of the two MnO2 polymorphs along the c-axis, showing the tunnels 2  2 et 1  1 for α-MnO2 and the tunnels 1  1 for β-MnO2.

At C/5, the retention capacity reaches 84% over 60 cycles despite the amorphization of the material. In the vanadium oxides, one has to cite the case of the amorphous material Na1.5+ xVO3 that delivers a reversible capacity of 150 mAh g
1 at an average potential of 1.8 V (Venkatesh et al., 2014).

2D structure Among the many candidates for cathode materials, the layered oxides NaxMO2 (with M ¼ Ti, V, Cr, Mn, Fe, Co, Ni, or a mix of two or more elements) offer many advantages because of their 2D structures, their great theoretical capacities, and their ease of synthesis. The layered materials are interesting intercalation hosts because they have weak interlayer interactions and a gap between these layers, which allows easy ionic diffusion. Among them, AMO2-type lamellar compounds consist of MO6 octahedra planes joined by the edges and the alkaline ions are between these MO6 sheets.

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Delmas et al. (1981) have developed a nomenclature to describe alkaline ion stacking arrangements between layers. Each structure is designated by Xn, where X is the coordination of the alkali metal and n is the number of layers of MO6 octahedra that makeup the elementary cell (taking values 2 or 3). The alkali metal may be in octahedral coordination (O), trigonal prismatic (P), or tetrahedral (T). It is important to note that lithium, because of its small size, cannot occupy a prismatic site in the AMO2 compounds. Sodium, which is larger, can be found in the two types of configurations O or P. Finally, an apostrophe indicates a distortion of an ideal polyhedron. Following this notation, the most widespread lamellar sodium oxides can be classified into two categories: O3 and P2 types. The O3 (generally obtained for 0.7  x  1) are, therefore, composed only of alternating layers of NaO6 and MO6 octahedra joined by the edges and are thus considered as ordered “rock-salt” structure materials. For these materials, the phase transition O3 ! O0 3 ! P0 3 ! P3 occurs easily since it requires only the gliding of MO2 planes. Therefore, O3 phases are generally considered to be less stable because they can undergo sliding of MO6 layers during electrochemical processes. The P2-type phases (x  0.7), on the other hand, tend to maintain their structure during electrochemical extraction and insertion because a phase transition from a P2 to a P3 or O3 requires breaking and reforming of MO bonds and, therefore, requires a higher temperature. The diffusion of Na+ ions being better in the prismatic sites, the high-speed performances are better in the case of the materials P2 (Qi et al., 2017). Indeed, in the O3 compounds, the direct jump from one octahedral site to another adjacent site requires a very high activation energy and the Na ions thus migrate through the interstitial tetrahedral sites shared by two octahedral sites. On the contrary, in the P2-type, a diffusion path is open between neighboring prismatic sites and a lower diffusion barrier is expected. However, limitations exist for P2 materials with higher potential. For example, in the case of the material P2-Na2/3[Ni1/3Mn2/3]O2 (Lu and Dahn, 2001), a transition P2 ! O2 appears when Na < 1/3 and it limits the cyclability of the material. The capacity retention is greatly improved when this phase transition is excluded. In addition, O3-type phases have higher first charge capacities because of a larger amount of sodium (Deng et al., 2018; Han et al., 2014b). Studies are, therefore, carried out to limit the harmful structural transformations of the O3 phases or to increase the capacities of the P2 phases. These studies involve the use of substituents in O3 (Deng et al., 2018; Mu et al., 2015; Oh et al., 2014; Zheng and Obrovac, 2017) phases or the synthesis of “hybrid” P2/O3 (Bianchini et al., 2018, 2019; Qi et al., 2017) materials. Several examples of mono and multimetallic materials are presented below with a particular focus on manganese-based materials. NaxCoO2 The NaxCoO2 material was one of the first oxides to be restudied for electrochemical sodium insertion. Because of the results obtained with its lithiated analog LiCoO2 (Mizushima et al., 1980), NaCoO2 could appear as a promising candidate. However, the electrochemical behaviors of LiCoO2 and NaCoO2 are very different. As shown in the 1980s by Delmas et al. (1981), the initial phase O3-NaCoO2 undergoes a succession of phase transition O3-O0 3-P0 3 as the amount of sodium decreases in the compound. Delmas group (Berthelot et al., 2011) was able to isolate

Ceramics for electrochemical storage

619

up to nine phases from P2-NaxCoO2 for 0.50  x  1.00 at room temperature. Phase P2-Na0.74CoO2 provides a capacity of 107 mAh g
1 in the potential window 2.0–3.8 V but suffers from a capacity decay of about 0.1% for each cycle up to 40 cycles and an increase in its polarization (Ding et al., 2013). However, there are some encouraging points for NaxCoO2. Indeed, the ionic diffusion is greater in NaxCoO2 (0.5–1.5  10
10 cm2 s
1) than in LixCoO2 (<1  10
11 cm2 s
1) (Shibata et al., 2015) thanks to clusters of interconnected vacancies in the material (Willis et al., 2018). Using a particular morphology of microspheres, Fang Yongjin et al. (2017) reached a capacity of 125 mAh g
1 with a capacity retention of 86% after 300 cycles for the P2-Na0.7CoO2 phase. Inspired from Li technology, a study was conducted on the Na-NMC phase, of formula Na[Ni1/3Mn1/3Co1/3]O2. This compound allows the reversible extraction of 0.5 Na for a capacity of 120 mAh/g in the potential window 2.0–3.75 V (Sathiya et al., 2012). The phase transitions follow the sequence O3-O1-P3-P1 with an increase of the c parameter. Carlier et al. (2011) showed the good reversibility of the Na+ insertion/extraction process in the Co-rich P2 phase Na2/3[Co2/3Mn1/3]O2 phase. This electrochemical process takes place mainly by solid solution, except for x ¼ 0.5, where the formation of an ordered phase P2-Na0.5[Co2/3Mn1/3]O2 is foreseen. By changing the stoichiometric ratios of Mn and Co precursors, several P2-Na2/3[MnyCo1
y]O2 (with y ¼ 0, 1/6, 1/3, 1/2, 2/3, 5/6, 1) compounds were compared from an electrochemical point of view (Wang et al., 2013a). Substitution of Co by Mn improves the specific capacity of the compound, but the greater the amount of cobalt and the higher the redox potential. Moreover, the stability in cycling is better with a higher amount of Co. NaxNiO2 The phase O0 3-NaNiO2 provides a first charge capacity of 199 mAh g
1 (0.85 Na extracted) and a first discharge capacity of 147 mAh g
1 (0.62 Na reinserted) in the voltage range 2.2–4.5 V (Vassilaras et al., 2013). However, the cyclability is relatively low due to the formation of an inactive phase above 3.75 V, Na0.19NiO2 (Han et al., 2014a; Wang et al., 2017c). By limiting the window to 2.0–3.75 V, 0.63 (147 mAh g
1) and 0.52 Na (123 mAh g
1) are extracted and inserted, respectively, with a better capacity retention (94% of the initial capacity) after 20 cycles. Thanks to the activity of the Ni2+/Ni4+ and the O3-P3 phase transition, Na [Ni0.5Mn0.5]O2 also provides a large reversible capacity (185 mAh g
1) (Komaba et al., 2009, 2012). In charging up to 4.5 V, almost all sodium is extracted without irreversible structural change. Better capacity retention is achieved when the potential window is limited to 3.8 V. The substitution of Ni by Ti in Na[Ni0.5Ti0.5]O2 allowed a very good stability in cycling with a retention of capacity of 93.2% after 100 cycles at 0.2C (Yu et al., 2013a). In the pure titanium phase, O3-NaTiO2, only 0.3 Na can be reversibly extracted (75 mAh g
1) in the voltage window 0.8–1.6 V (Maazaz et al., 1983). The material undergoes an irreversible phase transition if it is oxidized at higher voltage because of titanium ions migration from their TiO2 layers to the free interlayer space during sodium extraction. Because of the low operation voltage of the Ti4+/Ti3+ redox couple, the titanium substitution will be essentially used to stabilize the structure during sodium extraction.

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Advanced Ceramics for Energy Conversion and Storage

NaxFeO2 The phase α-NaFeO2, which is a typical example of O3-type lamellar structure, is easily obtained by solid-state synthesis (Takeda et al., 1980) and has shown its electrochemical activity from the earliest studies in the 1990s (Takeda et al., 1994). By limiting the potential to 3.4 V, 0.3 Na are reversibly extracted and reinserted from NaFeO2 due to the Fe4+/Fe3+ redox couple (Yabuuchi et al., 2012b; Yabuuchi and Komaba, 2014; Zhao et al., 2013). The electrochemical process is then marked by a plateau at 3.3 V with a very low polarization. At higher potential, NaFeO2 suffers from an irreversible transition due to the migration of Fe3+ ion into tetrahedral sites close to the free octahedral sodium sites (Yabuuchi et al., 2012b; Yabuuchi and Komaba, 2014). The reinsertion of the sodium ions is, therefore, disturbed by iron in tetrahedral sites and this leads to a significant degradation of the electrochemical properties. To increase the capacity retention, Yoshida et al. (2013) synthesized the O3-Na[Fe0.5Co0.5]O2 phase. The decrease in irreversible capacity is attributed to the suppression of iron migration due to the occupation of transition metal sites by cobalt. More sodium can be extracted and reinserted and the specific capacities are more important. In addition, the presence of Co enabled an improvement in electrical conductivity, so that the electrode was active up to 30C regimes for a capacity of 102 mAh g
1. These performances for very high rates are obtained thanks to the O3-P3 phase transition. A second compound that does not contain cobalt, O3-Na [Fe0.4Ni0.3Mn0.3]O2, also shows the O3-P3 phase transition with low polarization (Yabuuchi et al., 2013). However, the capacities at higher rate are lower than the compound Na[Fe0.5Co0.5]O2 and this can be explained by the high conductivity of the cobalt oxides. The use of the Fe4+/Fe3+ redox couple in the material P2-Na2/3[Fe1/2Mn1/2]O2 gives a capacity of 190 mAh g
1 at an average voltage of 2.75 V (Yabuuchi et al., 2012a). The study of P2-Na0.62[Fe1/2Mn1/2]O2 (Koga et al., 2013a,b) showed that the best specific capacities were obtained between 1.5 and 4.0 V and that oxidation up to 4.3 V lead to the appearance of a new structure which caused irreversible structural changes, harmful to the cycling of the electrode material (Mortemard de Boisse et al., 2014). NaxVO2 and V2O5 The charge/discharge curves of O3-NaVO2 and P2-Na0.7VO2 compounds are similar and present numerous plateaus for 0.5 < x < 1 (Didier et al., 2011; Guignard et al., 2013; Hamani et al., 2011) with first charge capacity of 120 and 100 mAh g
1, respectively (Fig. 36). However, the P2 phase has a lower polarization than the O3 phase in the potential window 2.4–1.2 V and the derived curves are drastically different. This difference in polarization can be explained by ion diffusion mechanisms that do not follow the same paths. In the O3 phase, the Na+ ions diffuse through the tetrahedral sites, which require a lot of energy for the large Na+ ions whereas in the case of the P2 phase, the ions seem to diffuse in a less energetic way. Therefore, the pair V4+/V3+ is not the only active redox couple in the system Na-V-O and other vanadium oxides have been studied in Na-ion batteries by using the couple V5+/V4+. From the 1980s, West et al. reported the insertion of sodium into the lamellar phase α-V2O5 (West et al., 1988). By nanostructuration or by controlling the morphology, this material provides a capacity close to the theoretical capacity expected for the formation of Na2V2O5 (up to 250 mAh g
1) (Li et al., 2015a; Mahadi et al., 2016;

Ceramics for electrochemical storage

621

2.6

3.5

Volts vs Na+/Na

Volts vs Na+/Na

2.4 2.2 2.0

O3-NaVO2

1.8 1.6

P2-Na0.7VO2

1.4

3.0 2.5 2.0 1.5 1.0

1.2

NaxV2O5

0.5 0.5

(A)

0.6

0.7

0.8

x in NaxVO2

0.9

1.0

0.0

(B)

0.5

1.0

1.5

2.0

2.5

3.0

x in NaxV2O5

Fig. 36 Potential composition curve at C/20 for (A) NaxVO2 and (B) NaxV2O5.

Safrany Renard et al., 2017; Tepavcevic et al., 2012; Wang et al., 2015a). Moreover, starting with ball-milled material, three sodium can be electrochemically and chemically inserted into α-V2O5 leading to the formation of the disordered rock-salt structure material Na3V2O5 (Adamczyk et al., 2017). Manganese-based materials Due to low-cost and nontoxic character of manganese, the Na-Mn-O system represents a major challenge for the development of Na-ion batteries for large-scale applications. However, although these materials have high reversible capacities, they undergo low cyclic stability because of the structural collapse and the important Jahn-Teller effect in the trivalent Mn. Stabilization of these materials and improvement of cycling retention are essential for its practical applications and, this can be provided by substitutions. Despite the large number of reported phases in the phase diagram Na2O-MnO-MnO2 (Fig. 37), only few materials have been studied as electrode material for Na-ion batteries. O3-NaxMnO2 (x > 0.7) NaMnO2 exists in two polymorphic forms. The lowtemperature monoclinic phase O3-type α-NaMnO2 and the high-temperature orthorhombic phase β-NaMnO2 (Fig. 38). On first charge, 0.85 Na (210 mAh g
1) are extracted from α-NaMnO2 and 0.8 are reinserted (197 mAh g
1) in the potential window 2–3.8 V at C/30 (Ma et al., 2011). The charge and discharge reactions of this compound do not take the same reaction paths. Indeed, the charge and discharge curves show eight and five plateaus, respectively. The long plateau at 2.63 V is associated with the biphasic reaction (first-order transition) between Na0.93MnO2 and Na0.7MnO2. The addition of Li in the transition metal layers “delays” the O3-P3 transition in charge for the Mn-rich O3-Na[Mn0.5Ni0.25Fe0.25]O2. Since the LidO bond is stronger than the MndO or NidO bonds, the lithium-substituted phase will be more stable. Then, the specific capacity is improved because the phase transition appears for a larger amount of sodium extracted (Oh et al., 2014). The β-NaMnO2 phase was studied as a cathode in 2014 by Billaud et al. (2014a). On first discharge, 0.82 Na are reinserted into this phase, leading to a capacity of 190 mAh g
1. Moreover, this capacity is maintained at 130 mAh g
1 after 100 cycles.

622

Advanced Ceramics for Energy Conversion and Storage

Na2O 0 0.

1.0

0. 1

0.9

0. 2

0.8

Na14Mn2O9

no ctio fra

5

n tio

0.

Na10Mn4O9

0.5

of

Na2MnO2

6

O Na 2

Na4Mn2O5

0.

At.

c fra At.

4 0.

fM

nO

3 0.

Na6MnO4 0.7 Na5MnO4 Na4MnO4 0.6

0.4

7 0. Na2Mn2O3

NaMnO

Na2Mn3O7

2

8 0.

Na26Mn39O55

Na1-x MnO2

0.3

NaMn2O4

0.2

0. 9

0.1

1.

1.0

0.9

0.0 0.8

0.7

0.6

0.5

0.4

0.3

0.2

0.1

MnO

0.0

0

NaMn7O12

MnO2

At. fraction of MnO2

Fig. 37 Ternary diagram Na2O-MnO-MnO2 showing the reported phases in the system Na-Mn-O.

Fig. 38 Structural view along the b-axis of the monoclinic phase (A) α-NaMnO2 and (B) β-NaMnO2. The primary cell is represented in black and the sodium polyhedra in yellow.

Ceramics for electrochemical storage

623

At the structural level, β-NaMnO2 has a lamellar structure different from the conventional O3- or P2-NaxMO2-type structures. Instead of plane layers of MnO6 octahedra, β-NaMnO2 is composed of zigzag layers between which the Na+ ions are in octahedral sites. Ex situ solid NMR study of the 23Na allowed to observe that stacking defects increase during the extraction of sodium, this being accompanied by a loss of crystallinity. However, during the reinsertion of the sodium ions, the structure is recovered, although presenting an increased disorder. Given these important structural transformations, the good reversibility and the capacities obtained are quite remarkable. P2-NaxMnO2 (0.5 < x  0.7) Used as a cathode material, Na0.6MnO2 gives a first charge capacity of 150 mAh g
1 when oxidized from 2.0 to 3.8 V (Caballero et al., 2002). Electrochemical performance was then improved by Kumakura et al. (2016), reaching a capacity of 198 mAh g
1 for the phase P2-Na2/3MnO2. By controlling the synthesis conditions, the authors could also stabilize the distorted P0 2Na2/3MnO2 and obtained a capacity of 216 mAh g
1 with a superior cycle stability. The distortion is due to Jahn-Teller effect of the Mn3+. The loss of significant capacity in this material could be limited by the use of nanostructured materials. Indeed, P2-Na0.7MnO2 shows a capacity of 163 mAh g
1 in the potential range 2.0–4.5 V with a retention capacity of about 67% after 50 cycles (Su et al., 2013c). This nanostructuration and the presence of dominant crystalline planes allow facilitated ionic transport for the insertion and extraction of Na+. Ex situ X-ray diffraction (XRD) identified that there was no phase transition in the fully charged state. When all the Na has been extracted from the material, the “MnO2” obtained crystallizes in the same type of orthorhombic cell with the same space group Cmca. Only the cell parameters are modified, with a contraction of the intersheet spacing. Li et al. (2017a) obtained a capacity of 135 mAh g
1 at a rate of C/10 in the potential range 2.0–4.3 V with the material P2-Na0.53MnO2, synthesized by a coprecipitation reaction. As early as 1999, Paulsen and Dahn (1999) have studied the structural properties and stability of P2-Na2/3MnO2 phases and have shown that the substitution of Mn by Li, Ni, or Co reduces the Jahn-Teller distortions and extends the stability range of the P2 phase. This study was continued on the material P2-Na2/3[Mn2/3Ni1/3]O2 and almost all the Na ions were reversibly extracted from this material, which corresponds to a charge capacity of 161 mAh g
1 (173 mAh g
1 for the theoretical capacity) (Lu and Dahn, 2001). On the basis of the Ni4+/Ni2+ redox couple, a plateau appears at a high voltage of 4.2 V for x < 1/3 and two phases seem to coexist: P2-Na1/3[Mn2/3Ni1/3]O2 and [Ni1/3Mn2/3]O2. A large volume change in the 0 < x < 1/3 region, associated with the O2 phase transition, significantly reduces the cyclability of the electrode (Lee et al., 2013a,b). P2-Na0.8Ni0.4Mn0.6O2 presents good electrochemical properties, especially at higher temperature (55°C) with discharge capacities of 92.0 and 85.3 mAh g
1, respectively, at C/10 and C in the potential range 2.0–4.0 V (Liu et al., 2017a). The electrochemical activities of two phases of the P2-Na0.67[Mn0.65Fe0.35
xNix]O2 system (with x ¼ 0 and 0.15) were compared by Yuan et al. (2014). They obtain similar initial capacities of 205 mAh g
1 for both phases. The unsubstituted phase, however, has low capacity retention.

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Advanced Ceramics for Energy Conversion and Storage

The substitution with Ni allowed maintaining 71% of the initial capacity after 50 cycles and this would be due to the attenuation of the Jahn-Teller distortion of the trivalent Mn. The substitution of a part of the Mn by nonelectrochemically active Ti increases the stability by suppressing the phase transition P2-O2 (Yoshida et al., 2014). Smoothing of the electrochemical curve can also be observed. When substituting by Mg, another nonactive element, the orthorhombic distortion of the ideal hexagonal P2 structure is reduced by increasing the amount of Mn4+ within the compound. Slow cooling during synthesis also reduces distortion by promoting the appearance of vacancies Mn and thus raises the Mn4+ concentration. By mastering the synthesis parameters and with a substitution of only 5% of Mg, an undistorted P2 phase is obtained (Billaud et al., 2014b). The addition of Mg results in a decrease in the specific capacities of the first cycles, but after 50 cycles the capacity is increased by approximately 30% (110 mAh g
1 for the undoped phase against 140 mAh g
1 for the doped one) (Clement et al., 2016). By co-doping with Ni and Mg, the lamellar structure Na0.67[Mn0.80Ni0.1Mg0.1]O2 is stabilized and provides a capacity of 110 mAh/g at a high rate of 2C. In the same way as in the material doped solely with Mg2+, the addition of Ni2+ will promote the formation of Mn4+, however, in the case of this co-doping, the nickel is electrochemically active and, therefore, limits the loss of capacity (Li et al., 2016c). Wang et al. (2016b) were also able to completely remove the phase transition P2-O2 with the material Na0.67[Mn0.67Ni0.23Mg0.10]O2 for a capacity retention of 92% after 100 cycles (100 mAh g
1 at C/10). Working in a higher potential window (1.5–4.4 V), Yabuuchi et al. (2014b) have obtained an abnormally high reversible capacity >200 mAh g
1 in the highly doped compound P2-Na2/3[Mn0.72Mg0.28]O2. This capacity is greater than the theoretical capacity based on the Mn4+/Mn3+ couple and certainly comes from an anionic contribution. Na2Mn3O7 The lamellar phase Na2Mn3O7 (Na4/7[☐1/7Mn6/7]O2 in conventional notation NaxMO2) cannot be classified by Delmas notation. Indeed, it consists of Mn-vacancy-[Mn3O7]2
layers built up with edge-sharing MnO6 octahedra, separated by NaO6 and NaO5 polyhedra. Two Na per f.u. can be reversible inserted in this material, leading to a capacity of 160 mAh g
1 through a plateau at 2.1 V with a low polarization of 100 mV (Adamczyk and Pralong, 2017). Then, a reduced phase Na4Mn3O7 is formed by a biphasic electrochemical process. Interestingly, an additional reversible solid-solution process, corresponding to the extraction of 1.5 Na+, is observed on oxidation at 4.1 V (Fig. 39) due to the oxygen redox activity, consistent with density functional theory (DFT) calculations (Zhang et al., 2017). Based on the theoretical explanation given by Ceder et al. (Seo et al., 2016), this oxygen redox activity is explained by the presence of ☐–O-Na axes due to the Mn vacancies in the [Mn3O7]2
layers (de Boisse et al., 2018). Therefore, by cycling this material in the potential range 4.5–1.5 V, the reversible specific capacity reaches 200 mAh g
1 (Wang et al., 2018b). The copper-doped Na2.3Cu1.1Mn2O7
δ shows a solid-solution behavior in the range 4.5–2.0 V for a capacity of 100 mAh g
1 at C/10 and a capacity retention of 95.8% for 1000 cycles at a fast rate of 20C ( Jiang et al., 2018). Thereby, such as the phase P2-Na2/3[Mn0.72Mg0.28]O2, Na2Mn3O7 shows a participation of oxygen ions in redox activity. This particular case of the high potential oxidation in transition metal oxides is detailed in the following section.

Ceramics for electrochemical storage

5.0

–150

625

–100

+

Volts vs Na /Na

4.5

Capacity (mAh/g) –50 0 50

100

150

O2-/On-

4.0 3.5 1.5 Na extracted

3.0

2 Na inserted

2.5 2.0

Mn4+/Mn3+

1.5 0.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

4.0

x in Na x Mn3O7

Fig. 39 Potential-composition curve of NaxMn3O7 showing both the Mn4+/Mn3+ and O2
/On
redox activity.

Oxygen anionic activity in Na-ion batteries Mn4+-rich compounds As presented so far, the transition metal oxides studied as intercalation electrodes are based on redox reactions that take place in the transition metals. However, if the redox reactions also occur on the oxygen anions, it is in principle possible that the capacity of the accumulators is significantly improved. The use of the oxygen redox in the oxides is, therefore, an important strategy to further increase the reversible capacity. This could also reduce the cost of the positive electrode materials because they are not limited by the absence of oxidizable transition metals as redox centers. Thus, negatively charged oxide ions can potentially give electrons in place of the transition metal(s) and this concept has already been realized in LIBs (Delmas, 2016; Grimaud et al., 2016; McCalla et al., 2015; Sathiya et al., 2013). During the extraction of alkaline ions, the transition metals are oxidized to their maximum oxidation state and an extra capacity, characterized by a plateau of potential, is observed at higher voltage. This plateau indicates oxidation of O2– ions with the possible release of gaseous oxygen. In the case of lamellar transition metal materials, this loss of oxygen can be explained by two mechanisms (Koga et al., 2013a,b): -

-

Oxygen is lost on the surface and there is migration of oxygen from the heart to the surface of the material. The resulting oxygen vacancies are distributed throughout the material. In this case, the Na (or Li) vacancies, formed during the charge, remain in their initial alkaline sheets and are likely to be reoccupied during the next discharge. Loss of O2 occurs at the surface with, this time, the migration of transition metal ions from the surface to the core of the material. The transition metal ions present on the surface should, indeed, be unstable in MO5 configuration and thus migrate through the vacant alkaline sites. This mechanism would induce a decrease in the cell volume, but also a decrease in the number of sites that can be occupied during the next discharge. This second hypothesis is the most likely to occur because of the irreversible reorganization of the cationic network and the densification of the structure.

626

Advanced Ceramics for Energy Conversion and Storage

The sodium counterpart Na2MnO3 (Na[Na1/3Mn2/3]O2) isostructural with monoclinic Li2MnO3 has never been stabilized. Indeed, the large size difference between Na+ and Mn4+ ions blocks the formation of [Na1/3Mn2/3]O2 layers. Therefore, Komaba et al. (Yabuuchi et al., 2014a) wanted to synthesize the Na[Li1/3Mn2/3]O2 phase in which layers of sodium octahedra would be surrounded by layers of lithium and manganese octahedra. The organization of this mixed layer would then give the “honeycomb” network of Li2MnO3. Instead of Na[Li1/3Mn2/3]O2, this group synthesized Na5/6[Li1/4Mn3/4]O2. This material shows high concentration of inactive Mn4+ with average valence of 3.89 +. However, Na5/6[Li1/4Mn3/4]O2 leads to a high reversible capacity of 200 mAh g
1. As for Li2MnO3, a well-defined plateau is observed at a potential of 4.2 V vs Na+/Na, corresponding to the potential of 4.5 V vs Li+/Li. The disappearance of the plateau at the second charge is attributed to a partial release of gaseous oxygen. This same plateau at 4.2 V has also been reported in the case of the oxidation of the phase P3-Na0.6[Li0.2Mn0.8]O2 (Du et al., 2016). Then, Komaba and Yabuuchi’s group were among the first to treat the oxidation of oxygen in SIB. They carried out the same kind of study on the material presented previously: P2-Na2/3[Mn0.72Mg0.28]O2 (Yabuuchi et al., 2014b). In this case, the plateau at 4.2 V is visible on the cycles following the first charge but the capacity decreases progressively from 210 to about 150 mAh g
1 at the 30th cycle. Bruce et al. (Maitra et al., 2018) have resumed the study of the same compound in order to compare it with alkaline-rich systems such as Li2MnO3. An operando mass spectrometry analysis allowed analyzing the gas evolution during charges and discharges. This showed that there was no loss of gaseous O2 in P2-Na2/3[Mn0.72Mg0.28]O2. The presence of Mg2+ within the transition metal layers is sufficient to observe oxygen redox activity and it also suppresses the oxygen loss during the 4.2 V anionic activity. Spectroscopic techniques such as SXAS, XANES, or RIXS confirmed that Mn4+ is not active on the oxidation plateau at 4.2 V and that the capacity is only due to oxygen. Although it shows no loss of oxygen, this material has a relatively weak cyclability and this would be due to irreversible structural changes during the P2-O2 phase transition. In the same 4+
1 way, P2-Na0.78[Ni2+ at 0.23Mn0.69]O2 offers a reversible capacity of 138 mAh g an average potential of 3.25 V (Ma et al., 2017a). The capacity is due to the redox couple Ni4+/Ni2+ and an irreversible reaction on first charge gives an excess of 60 mAh g
1 thanks to anionic activity. The former leads to the appearance of oxy4+ gen vacancies on the surface of the compound. P2-Na0.78[Ni2+ 0.23Mn0.69]O2 is marked by a good cyclability because of the suppression of the P2-O2 transition. Other Na2MO3-type model compounds Li2MnO3 is not the only Li-rich layered material to have been studied for its electrochemical activity. In fact, other “model compounds” allow to better understand the oxygen redox phenomena, such as the demonstration of the formation of (O2)n
dimers in Li2Ru1
yMyO3 phases (M ¼ Ru, Sn, Ti) (Sathiya et al., 2013, 2015). These peroxo/superoxo groups, which represent up to almost a quarter of the oxygen atoms, are sufficiently stable not to be extracted from the material as gaseous O2. Thus, these model compounds show very good stability in cycling, unlike the xLi2Mn4+ O3 + (1
x)LiMO2 composites.

Ceramics for electrochemical storage

627

In 2013, Tamaru et al. studied the electrochemical properties of a Na2MO3-type material: Na2RuO3. This phase crystallizes in a lamellar structure of α-NaFeO2-type consisting of Na and Na1/3Ru2/3 layers and shows metal-type conduction that allows rapid diffusion of Na+ ions. A capacity of 147 mAh g
1, higher than the theoretical capacity for one sodium (137 mAh g
1 corresponding to the Ru5+/Ru4+ couple), seems to come partially from the activity of oxygen. Rozier et al. (2015) substituted Ru4+ by nonelectrochemically active Sn4+ in the Na2Ru1
xSnxO3 solid solution. By X-ray photoelectron spectroscopy (XPS) analysis, they were able to show that the “extra-capacity” in Na2Ru0.75Sn0.25O3 resulted from a combination of cationic
(Ru4+ ! Ru5+) and anionic O2
! O2 n
redox processes. Unlike the lithium analog Li2Ru1
xSnxO3 (Sathiya et al., 2013), an irreversible loss of O2 occurs in the case of Na2Ru1
xSnxO3 for the Sn-richest compounds. The differences in polarity and size between Na+ and Li+ would be responsible for this loss of O2. These first two studies were carried out on samples with so-called disordered Na1/3(Ru1
xSnx)2/3 layers, in which the NaO6 and MO6 octahedra are randomly distributed. Then, de Boisse et al. (Yamada et al., 2016) compared disordered Na2RuO3 with the ordered honeycomb phase. With the latter, the redox activity of oxygen could be activated, and the specific capacity was improved from 130 mAh g
1 (1 Na) to 180 mAh g
1, that is, 1.3 Na extracted. The stabilization of the intermediate phase O1-NaRuO3 is at the origin of this activation. This process requires oxygen rearrangement and gives distorted ˚ ) increase the energy level of RuO6 octahedra whose shortest OdO distances (2.580 A the σ* antibonding orbital of the OdO bond at the Fermi level. So, the redox reaction of oxygen is activated. It has also been estimated by DFT calculations that ruthenium contributes to 34.7% of the charge compensation for 1 < x < 2 and 19.6% for 0.5 < x < 1 in NaxRuO3 (Assadi et al., 2018). The calculated potential curve corresponds to the experimental curve obtained by De Boisse et al. (Yamada et al., 2016) with a first 2.4 V plateau for 1 < x < 2 and a second at 3.6 V for 0.5 < x < 1. The anionic oxidation in Na2RuO3 can, therefore, be activated by the substitution or ordering of Na1/3Ru2/3 layers. The yttrium-doped Na2ZrO3 lamellar material gives a reversible capacity of 120 mAh g
1 at C/20 with a small capacity decrease of about 0.023% per cycle (Song et al., 2017). At a higher current density of C/2, the material has good cyclability up to 1500 cycles for a capacity of about 60 mAh g
1. The anionic deficiencies resulting from yttrium doping seem to activate the oxygen redox process and thus allow to reach a first charge capacity of nearly 400 mAh g
1 between 1.5 and 4.5 V. By substituting a portion of Ti for Cr to create oxygen vacancies in Na2TiO3, Song et al. (Xu et al., 2017a) synthesized the compound Na2Ti0.94Cr0.06O2.97. The latter has initial discharge capacities of approximately 340 and 130 mAh g
1 at current densities of 18.9 and 378 mA g
1 between 1.5 and 4.2 V. After 1000 cycles at 378 mA g
1, the capacity of 100 mAh g
1 corresponds to capacity retention of 74%. One of the main disadvantages of this material is its low conductivity which could certainly be improved by carbon coating. The study of Na2IrO3 showed that it exhibits an electrochemical curve very similar to that of ordered Na2RuO3 (Perez et al., 2016). Na2IrO3 has the advantages of being

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protected against cationic migrations and of not suffering the O2 release that can occur in alkali-rich materials. It is a good model compound because the integrity of the transition metal layers is preserved during cycling and this leads to a very good reversibility of the oxygen redox process. Therefore, two main limitations of the use of alkali-rich oxides have been demonstrated. First, the material Na2IrO3 undergoes phase transitions O3 ! O10 ! O1, which limits the extraction of sodium until the formation of the O1-Na0.5IrO3 phase. Moreover, the second limitation is related to the high reactivity of this type of materials with the electrolyte. The choice of a good electrolyte is essential. The study of transition metal oxides as Na-ion battery cathode materials is a strong subject of scientific interest for the last 5 years. Mastery of synthesis and physicochemical and electrochemical characterizations are essential engines for obtaining high-performance materials. In addition, strategies such as substitution and modification of the morphology of the compounds can, sometimes, significantly increase the properties of the electrode materials. It is then possible to combine the advantages of certain elements such as the high conductivity of Co, the high redox potential of the Ni4+/Ni2+, and the high specific capacity of Mn. Another method for gaining capacity is the contribution of oxygen ions to the redox reaction. This oxidation may, however, generate partial loss of oxygen which would ultimately lead to irreversible structural changes and a sharp decrease in specific capacity through cycling. Mechanisms related to anionic redox activity are still controversial and understanding of these phenomena requires the use of sophisticated characterization techniques and model compounds. In order to improve the stability of cycling to meet the requirements of industrialization, these complex electrochemical processes will have to be better mastered. Thus, new materials with high-energy density could emerge.

3.2.3 Polyanionic frameworks Polyanionic frameworks, with polyanions like (PO4)3
, (P2O7)4
, (SO4)2
, or a mix of them, have also been studied as positive electrode for Na-ion batteries (Fig. 40). These polyanion groups allow increasing the redox potential of the transition metal thanks to inductive effect, compared to oxide compounds. In fact, when another atom X (with a stronger electronegativity than M) is introduced, it increases the ionicity in M–O bonding. This reduction of the covalency reduces the splitting between bonding and antibonding orbitals. In this way, the energy difference between the latter and the vacuum state is augmented and the redox voltage also increases (Barpanda et al., 2018). Moreover, the strong X–O covalent bonds provide more safety to the battery. Various polyanionic frameworks have been explored as electrode materials and several reviews are dedicated to this class of materials (Barpanda et al., 2018; Guo et al., 2017a,b; Hwang et al., 2017a; Ni et al., 2017; Wang et al., 2018c). In Fig. 7, an overview of the different structures is shown.

The case of vanadium-based phosphates NASICON-type structure Due to structural similarity with Na1+ xZr2P3
xSixO12 (Goodenough et al., 1976), the vanadium-based phosphates NaxV2(PO4)3 are named as Na+ SuperIonic CONductor (NASICON). The structure is made up of isolated VO6

Ceramics for electrochemical storage 629

Fig. 40 Overview of the various structural type of polyanionic frameworks reported in the literature (Guo et al., 2017a,b).

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Advanced Ceramics for Energy Conversion and Storage

octahedra and PO4 tetrahedra interconnected by corners, forming a 3D framework with open transport channels allowing high diffusion rates. Two distinct Na sites (Na1 and Na2) coexist in the crystal structure of Na3V2(PO4)3 (Chotard et al., 2015; Delmas et al., 1978) with one Na ion in Na1 sites and two Na ions in Na2 sites. The occupancies of the Na1 and Na2 sites are 1.0 and 0.67, respectively. Thus, one more sodium could be inserted in the Na2 sites, leading to the f.u. Na4V2(PO4)3. Nevertheless, because of the stable oxidation state of V3+ in Na3V2(PO4)3 [compare to the V2.5+ in Na4V2(PO4)3], the synthesis of the later is more favorable. The first electrochemical characterizations of Na3V2(PO4)3 in Na-ion cells were performed by Masquelier et al. (2003). They showed the presence of two voltage plateaus at 3.4 and 1.6 V vs Na+/Na involving the V4+/V3+ and V3+/V2+ redox couples, respectively. This work was followed by Plashnitsa et al. (2010) using a full symmetric cell NVP// 1 M, NaClO4 in PC//NVP but the cycling stability was not sufficient. The capacity retention can be greatly improved by carbon coatings. Chen’s group ( Jian et al., 2012) first reported a one-step solid-state reaction derived from sugar, NaH2PO4, and V2O3 and obtained a capacity of 93 mAh g
1 between 2.7 and 3.8 V, focusing on the high-voltage plateau. The capacity was, then, improved to 107 mAh g
1 by using NaFSI/PC as electrolyte ( Jian et al., 2013) and Saravanan et al. (2013) demonstrated that 50% of the initial capacity can be maintained after 30,000 cycles at 40C with a porous NVP/C composite. The 3.4 V plateau was attributed to a biphasic process thanks to ex situ/in situ XRD and DFT calculations (Lim et al., 2012; Saravanan et al., 2013). Starting from Na3V2(PO4)3, two Na/f.u. can be reversibly extracted, leading to a theoretical capacity of 117 mAh g
1 for the formation of NaV2(PO4)3. Solidstate NMR and annular bright-field scanning transmission electron microscopy (ABF-STEM) analysis indicated that the two extracted sodium ions were initially in Na2 sites ( Jian et al., 2014). Looking at the ionic mobility, the Na ions located in Na1 seem immobile and appear to maintain the structural integrity. The electrochemical properties of this material were then improved by controlling the synthesis parameters (Li et al., 2015a; Ren et al., 2016) or by optimizing the carbon coatings (Li et al., 2014; Mao et al., 2015). Therefore, the NVP@C@rGO compound, consisting of NVP nanocrystals coated with amorphous carbon and wrapped by reduced grapheme oxide nanosheets, showed an outstanding cycling stability with 64% retention after 10,000 cycles at 100C (Rui et al., 2015). The Mg doping (Li et al., 2015b) and the partial substitution of the vanadium by another transition metal is also a way to increase the rate kinetics and the energy density, respectively. By giving access to the V5+/V4+ redox couple, Na3Al0.5V1.5(PO4)3 showed a first charge capacity of 120 mAh g
1 following two plateaus at 3.37 and 3.95 V (Lale`re et al., 2015). The activation of the V5+/V4+ couple has also been showed in Na4VMn(PO4)3 (Chen et al., 2019). NaxVO(PO4)y Due to the high potential of V5+/V4+, the vanadyl phosphate family (VOPO4) has been recognized as promising cathode materials for LIB (Dupre et al., 2004) and was also studied as cathode material for Na-ion batteries. There are seven polymorphic structures of VOPO4: tetragonal (αI, αII, δ, ω), orthorhombic (β, γ), and monoclinic (ε). All these polymorphs are composed of VO6 octahedra and

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PO4 tetrahedra. The VO6 share common corners to form infinite chains and the PO4 act as bridges between the octahedra. In fact, the VOPO4 forms differ mainly in the orientation of the vanadyl bond (Dupre et al., 2004; Ling et al., 2014). Considering the insertion of 1 sodium in VOPO4, the theoretical capacity is of 165 mAh g
1 with an expected voltage of 3.4–3.5 V vs Na+/Na (energy density of around 570 Wh kg
1). The sodium insertion was performed in a layered VOPO4, prepared by chemical delithiation of αI-LiVOPO4. This material, marked by a high operating potential (3.4 V), gave a reversible capacity of 110 mAh g
1 at 0.05C and the incorporation of rGO improved it to 150 mAh g
1 (0.9 Na+) (He et al., 2016). VOPO4 nanosheets, obtained by ultrasonication of VOPO4 2H2O chunks, also displayed reversible capacities of 136 and 70 mAh g
1 at C/10 and 5C, respectively (Zhu et al., 2016a). Owing to this high redox potential, Li et al. (2016a) have developed a full cell made of VOPO4 as cathode and Na2Ti3O7 as anode. This cell led to an operating voltage of 2.9 V for a capacity of 114 mAh g
1 at C/10. This battery also provided a good cycling stability of 92.4% after 100 cycles at 1C. A fully sodiated vanadyl phosphate phase NaVOPO4 (Lii et al., 1991) was first studied by Goodenough et al. (Song et al., 2013) as cathode material for SIB. The authors needed to reduce the particle size by ball milling in order to increase the specific capacity and to decrease the polarization. Thus, 0.6 Na+/f.u. were reversibly reinserted in the ball-milled material giving a capacity of 90 mAh/g (30 mAh/g without ball milling) in the voltage range 2.0–4.4 V. Due to the initial formula NaVOPO4, the theoretical capacity is limited to 145 mAh g
1 in this case. The specific capacity was extended to 101 mAh g
1 at 363 K with an ionic liquid electrolyte (Chen et al., 2015). As for VOPO4, NaVOPO4 can be found in several different forms and these two studies were done on the monoclinic α-NaVOPO4 which has a 3D tunnel structure. By chemical lithium extraction from β-LiVOPO4 and microwave-assisted solvothermal sodium insertion, He et al. synthesized β-NaVOPO4, another 3D-type polymorph. This material is marked by a biphasic process at 3.5 V and a discharge capacity of around 60 mAh g
1 after 50 cycles. The capacity and the cyclability was improved with a layered form of NaVOPO4, giving 67% capacity retention over 1000 cycles at C/2 (Fang et al., 2018). The sodium-rich oxy-phosphate Na4VO(PO4)2 also showed electrochemical activity due to the versatility of the vanadium redox V5+/ V4+/V3+. Almost one sodium was extracted from this material through a biphasic process and one sodium can be inserted leading to the formation of the reduced phase Na5VO(PO4)2 (Deriouche et al., 2017). l

Pyrophosphates Among phosphates family, the pyrophosphate group is another interesting subject of study as a result of their high operating voltage. Pyrophosphates are formed by heating phosphates and they derive their name from this origin (the prefix pyro in Greek means fire). Indeed, when exposed to high temperature, phosphates undergo oxygen evolution, leading to the formation of the oxygen deficient [P2O7]4
2
4
following: 2 PO3
4 ! P2O7 + O . Electrochemical behaviors of three different vanadium pyrophosphates were reported, NaVP2O7, Na2VOP2O7, and Na7V3(P2O7)4. Despite a theoretical capacity of 108 mAh g
1, NaVP2O7 (monoclinic P21/c) showed only 38 mAh g
1 at C/20 involving V4+/V3+ activity at 3.4 V (Kee et al., 2015).

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The limited electrochemical activity was due to very low phase transition kinetics between NaVP2O7 and Na1
xVP2O7. Sodium ions exist as structural stabilizers and desodiation should break the lattice. The potential can be increased by activating the V5+/V4+ redox of tetragonal Na2(VO)P2O7. This later has a structure of infinite layers of [VP2O8]∞ made of corner-sharing VO5 square pyramids and PO4 tetrahedra showing pentagonal tunnels. Sodium atoms are located along the tunnels and between the layers. The large tunnels serve as ion diffusion pathways and then Na2(VO)P2O7 showed a reversible capacity of 80 mAh g
1 (theo. cap. ¼ 93 mAh g
1) with a V5+/V4+ redox potential centered at 3.8 V (Barpanda et al., 2014a). Kim et al. (2016) were the first to synthesize Na7V3(P2O7)4 and to exhibit its possible use as high potential cathode for SIB. Indeed, the redox potential of V4+/V3+ in Na7V3(P2O7)4 is higher than in NaVP2O7 and even higher than the V5+/V4+ of Na2VOP2O7. Because of the opencrystal framework and the low volume changes during desodiation/sodiation, the cyclability of Na7V3(P2O7)4 is stable with 75% retention after 600 cycles at 1C. In all, 3 Na/f.u. can be reversibly extracted giving a capacity of 80 mAh g
1 at C/10. At 8C, 84% of the theoretical capacity is maintained. Carbon-coated fine particles of Na7V3(P2O7)4 also showed that 92% of the initial capacity remained after 100 cycles at 10C (Deng et al., 2016) and a 3D hybrid foam of Na7V3(P2O7)4/C allowed to improve the electronic conductivity and the capacity rates of the material (Li et al., 2016b). Mixed polyphosphates Mixed phosphates are made when isolated phosphates (PO4)3
are present along with diphosphates (P2O7)4
(Sanz et al., 2001). With vanadium as transition metal, the electrochemical activity of Na7V4(P2O7)4PO4 was first reported by Lim et al. (2014). The structure consists of a central PO4 tetrahedron sharing corners with four VO6 octahedra. The four P2O7 bi-tetrahedra act as bridges between VO6 octahedra by sharing corners. The resulting (VP2O7)4PO4 units form a 3D framework with ionic channels for Na. Na ions adopt bipyramidal, square planar, or tetrahedral coordination. This material exhibited a first charge capacity of 82 mAh g
1 at C/40 involving a well-defined plateau centered at 3.88 V (Lim et al., 2014). Galvanostatic intermittent titration technique (GITT) characterization allowed to identify two oxidation and reduction plateaus, with the formation of the intermediate phase Na5V4(P2O7)4PO4 (Deng and Zhang, 2014). Considering V4+/V3+ redox activity, 4 Na can be reversibly extracted from Na7V4(P2O7)4PO4 (capacity ¼ 93 mAh g
1). As for Na7V3(P2O7)4, carbon coating and hybrid 3D framework increased the specific capacity and capacity retention of the electrode material (Li et al., 2016b; Zhang et al., 2014a) and led to good electrochemical properties even in aqueous sodium batteries (Deng et al., 2015). Masquelier’s group also demonstrated that Al substitution into Na7 V4
x 3 + Alx ðP2 O7 Þ4 PO4 increased the capacity because of the activation of the V5+/V4+ couple (M. Kovrugin et al., 2017). Due to their combination of safety and good electrochemical performances, mixed phosphates are an example of possible future commercial Na-ion batteries (Barpanda et al., 2018).

Others transition metal phosphates NaMPO4 (M 5 Fe, Mn) Inspired by the commercial success of LiFePO4 (Yamada et al., 2001), especially for EVs application (Zhang, 2011), the material NaFePO4 has been explored as cathode material for SIB (Barpanda et al., 2018), with the highest

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theoretical capacity of phosphate cathode materials (154 mAh g
1). Unlike the lithium counterpart, NaFePO4 can be obtained in two olivine polymorphs: the metastable triphylite phase (t-NaFePO4), isostructurale with LiFePO4 and the thermodynamically stable maricite phase (m-NaFePO4). Over 480°C, the triphylite undergoes an irreversible transition to form the m-NaFePO4 (Avdeev et al., 2013). Both structures are made of slightly distorted FeO6 octahedra and PO4 tetrahedra but the main difference is the arrangement of octahedra. Indeed, m-NaFePO4 and t-NaFePO4 show edge- and corner-sharing octahedra, respectively. The latter has Na+ transport channels, whereas the former has sodium ions trapped in unconnected cavities, resulting in much lower ionic conductivity. Thus, the maricite polymorph was considered almost electrochemically inactive (Sun et al., 2012). However, the metastable t-NaFePO4 cannot be directly synthesized by solid-state synthesis and must be obtained by chimie douce, such as chemical and electrochemical insertion of Na in delithiated LiFePO4 (Moreau et al., 2010; Zhu et al., 2012a,b). By this way, t-NaFePO4 showed an interesting first specific capacity of 125 mAh g
1, a capacity retention of 90% after 50 cycles at C/20 rate, in the voltage limits 2.2–4.3 V (Oh et al., 2012). Started from carbon-coated-LiFePO4 (C-LiFePO4), the same capacity retention is maintained for 100 cycles at C/10 and these results are comparable with that of C-LiFePO4/ Li cells (Zhu et al., 2012a,b). However, owing to higher charge transfer resistance, the rate capability of C-NaFePO4 was not as good as C-LiFePO4. Interestingly, the electrochemical process of t-NaFePO4 showed a two-step voltage profile with two plateaus due to Fe3+/Fe2+ redox couple. Unlike LiFePO4, the extraction of sodium from t-NaFePO4 lead to the formation of the intermediate phase Na0.7FePO4 and finally to FePO4. The formation of Na0.7FePO4, with the same crystal structure as the triphylite, makes it possible to buffer the stresses due to the large volume changes between FePO4 and NaFePO4 (Casas-Cabanas et al., 2012). During discharge, the three phases FePO4, Na0.7FePO4, and NaFePO4 appear simultaneously. Thanks to nanostructuration, the electrochemical behavior of the m-NaFePO4 was reinvestigated and revealed a capacity of 142 mAh g
1 (92% of its theoretical capacity) with a very good cyclability over 200 cycles at C/20. Experimental and computational studies allowed highlighting the formation of amorphous FePO4 (a-FePO4) during the first desodiation (Kim et al., 2015d). The capacity is increased to 154 mAh g
1 when m-NaFePO4 is in the form of hollow nanospheres. This peculiar morphology allowed an ultrastable capacity over 300 cycles at C/10 between 1.5 and 3.75 V (Li et al., 2015a). Interestingly, crystalized m-NaFePO4 also delivered 100 mAh g
1 in the 120th cycle at C/2 using ionic liquid electrolyte at 363 K without showing any structural alteration (Hwang et al., 2017b). NaMnPO4 also adopts the two polymorphs maricite and triphylite and in the same way as NaFePO4, the maricite structure is the thermodynamically stable form and it is electrochemically inactive (Ong et al., 2011). NaMnPO4 and Na(Fe0.5Mn0.5)PO4 triphylite phases have been stabilized by topotactic reaction using solid-state synthesis in molten salts but both compounds offered poor activities (Lee et al., 2011b). The triphylite-NaMnPO4 was also synthesized via ion-exchange reaction from KMnPO4H2O and delivered a reversible capacity of 80–85 mAh g
1 in lithium cell (Boyadzhieva et al., 2015). During first charge, almost all the sodium was extracted from NaMnPO4 following a voltage plateau at 4.25 V vs Li+/Li.

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Advanced Ceramics for Energy Conversion and Storage

Pyrophosphates Na2FeP2O7 was the first pyrophosphate studied as Na-insertion material by Honma et al. (2012) and Barpanda et al. (2012). This material was synthesized via different ways such as conventional one-step solid-state synthesis (Barpanda et al., 2012) and by glass-ceramics method (Honma et al., 2012). Na2FeP2O7 has a triclinic 3D framework (space group P-1) made of corner-sharing FeO6 octahedral dimers (Fe2O11) interconnected in zigzag by diphosphate groups (P2O7). This framework creates large tunnels allowing Na+ diffusion. Operating at around 3.0 V (Fe3+/Fe2+), this iron pyrophosphate gave a first discharge capacity of 88 mAh g
1 at C/10 (Honma et al., 2012). First-principle calculations and quasiequilibrium measurements determined two kinds of reactions in the voltage window 2.0–4.5 V; a single-phase reaction around 2.5 V and a biphasic process between 3.0 and 3.25 V (Kim et al., 2013b). Both as-prepared Na2FeP2O7 and oxidized-phase NaFeP2O7 showed good thermal stability up to 560°C without decomposition, oxygen evolution, or structural changes (Barpanda et al., 2013a). The Na-deficient phase Na1.66Fe1.17P2O7 [or Na3.32Fe2.34(P2O7)2] exhibited a better rate capability with good capacity retention up to 10C (Ha et al., 2013). The iron excess allows the extraction of 1.17 Na/f.u. instead of only 1 Na/f.u. for Na2FeP2O7 and then raises the theoretical capacity from 96 to 113 mAh g
1. Indeed, the potential of the redox couple Fe4+/ Fe3+ (5 V) is too high to be used with conventional liquid electrolyte. The same synthesis was followed to prepare Na3.32[Fe0.5Mn0.5]2.34(P2O7)2 and Na3.32Mn2.34(P2O7)2 but these phases delivered very low capacities probably due to the poor electrochemical activity of Mn in this system. According to the authors, nanosizing or doping with Mg or Ca is expected to improve the capacities (Ha et al., 2013). However, atomistic simulations techniques predicted that Na2MnP2O7 should provide favorable 3D Na+ diffusion pathways with activation energy close to Na2FeP2O7 (Clark et al., 2014). Then, the manganese pyrophosphate Na2MnP2O7 was expected to be electrochemically active as cathode material. Park et al. reported the activity of Na2MnP2O7 showing that this material could deliver a reversible capacity of 90 mAh g
1 (0.9 Na) at 3.6 V, in accordance with DFT calculated value of 3.65 V (Barpanda et al., 2013b). The electrochemical curve can be decomposed in one single-phase region at 3.32 V and three consecutive two-phase reactions at 3.66, 3.98, and 4.15 V. Nevertheless, due to the lower electronic conductivity of Mn (compare to Fe), Na2MnP2O7 is suffering from limited rate capability. By substituting Mn with Fe, the idea was to increase the high-rate capability and to improve the thermal stability. Na2Fe0.5Mn0.5P2O7 showed a typical single-phase reaction between 2.0 and 4.5 V (Shakoor et al., 2016) with participation of both Mn3+/Mn2+ and Fe3+/Fe2+ redox couples and a better rate capability than the nonsubstituted Mn-pyrophosphate. Mixed phosphates Na4Co3(PO4)2P2O7 was among the first example of mixed phosphate compounds (Sanz et al., 2001). These phases crystallize in the orthorhombic structure with a Pn21a space group forming a 3D framework made of MO6 octahedra and PO4 tetrahedra. PO4 share edges and corners with neighboring MO6, leading to a pseudo-layered structure parallel to the bc plane and (P2O7) units bridge these (M3P2O13)∞ blocks along the a-axis making large tunnels where the sodium ions

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635

are located. Nose et al. (2013a) explored the electrochemical extraction/insertion of this material and showed a multistep redox process in the high-voltage region 4.1–4.7 V. In all, 2.2 Na+ can be reversibly extracted (95 mAh g
1) at C/5 and the capacity is maintained at 80 mAh g
1 at the high rate of 25C with small polarization. This high diffusivity is due to multiple 3D diffusion pathways in the structure. Partial substitution with Mn and Ni in Na4Co2.4Mn0.3Ni0.3(PO4)2P2O7 smoothed the electrochemical curve and gave a capacity of 103 mAh g
1 at 5C. XAS analysis showed that all the transition metals participate simultaneously in the oxidation process (Nose et al., 2013b). From first principles calculations and experimental confirmation, Kim et al. (2012a) showed the activity of Na4Fe3(PO4)2P2O7. This material gave 105 mAh g
1 at the average potential of 3.2 V due to the oxidation of 3 Fe2+/f.u (theo. cap. ¼ 129 mAh g
1). The rate capability of this iron-mixed phosphate was lower than the cobalt counterpart but avoid the use of this expensive and toxic transition metal. The desodiation/sodiation in Na4Fe3(PO4)2P2O7 occurs via a single-phase process with a small volume change of only 4% thanks to the (P2O7) dimers which are able to accommodate the structural changes (Kim et al., 2013c). Na4Fe3(PO4)2P2O7 also showed interesting specific capacity of 84 mAh/g in aqueous cell between 3.4 and 2.5 V vs Na+/Na at 1C (Ferna´ndez-Ropero et al., 2018). Simulations suggested that Na+ diffusion is done through a 3D network of migration pathways with low activation energy of around 0.2 eV and would explain the high-rate capability obtained within these materials. Moreover, Ni doping in Na4Fe3(PO4)2P2O7 is expected to significantly increase the operating voltage [3.7 V for Na4Fe2Ni(PO4)2P2O7] (Wood et al., 2015).

Fluorophosphates

The high redox potential of PO4-based anions can be increased by the use of F
, because the strong inductive effect of fluorine is added to the effect of phosphate. This phenomenon led to the study of some fluorophosphates compounds such as Na2FePO4F, NaVPO4F, Na2MnPO4F, or Na3V2(PO4)3F3 (Fig. 41). In 2003, Barker et al. first explored the tetragonal NaVPO4F (space group I4/mmm) in full-cell vs hard carbon. The structure of this material consists of corner-sharing [V2O8F3] bioctahedra and [PO4] tetrahedral. Sodium ions are located in open channels made by this 3D framework. The full cell showed a two-step voltage profile when cycled between 2.0 and 4.3 V and delivered a capacity of 82 mAh g
1. >50% of the initial capacity is maintained after 30 cycles. The average potential of 3.7 V corresponds to the V4+/V3+ redox activity and based on 1 Na/f.u., the theoretical capacity is 143 mAh g
1. Based on these preliminary results, the study was continued on the low-temperature polymorph, the monoclinic (C2/c) NaVPO4F. The phase transition from monoclinic to tetragonal structure occurs at 750°C (Zhao et al., 2010). Cr-doping in the monoclinic form gave a similar capacity of 80 mAh g
1 (between 3.0 and 4.5 V) and enhanced the capacity retention (91.4% after 20 cycles) (Zhuo et al., 2006). Carbon coating (Lu et al., 2014) or formation of NaVPO4F/graphene (Ruan et al., 2015) composite were also strategies to increase electrochemical

636 Advanced Ceramics for Energy Conversion and Storage

Fig. 41 Structural view of Na3V2(PO4)2F3, Na2FePO4F, NaVPO4F, and Na2MnPO4F.

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performances. NaVPO4F/C composite nanofibers, obtained via an electrospinning process, provided 103 mAh g
1 at 1C with very good rate capability (56 mAh g
1 at 50C) ( Jin et al., 2017). Long-term cycling at 1C was also achieved with NaVPO4F synthesized from V2O5 precursor with 82% of the initial capacity retained after 2500 cycles. Moreover, this compound showed an important capacity of 133 mAh g
1 at C/10 (Law and Balaya, 2018). In order to clarify the composition of the tetragonal NaVPO4F reported by Barker, Sauvage et al. (2006) reinvestigated this material and were able to synthesize the single-phase (VIV) Na3[VO]2[PO4]2F (or Na1.5VOPO4F0.5) instead of NaVPO4F, thanks to 10%wt excess of NaF reactant. Structural data on the monoclinic NaVPO4F correspond to the NASICON-type Na3V2(PO4)3 and, moreover, VF3 sublimation can occur during the synthesis. It is then complicated to differentiate NaVPO4F and Na3V2(PO4)3 and, depending on the oxidation state of the vanadium, several composition Na3V2O2x(PO4)2F3
2x are obtained (Serras et al., 2012). In Na1.5VOPO4F0.5, [V2O8F3] bioctahedra of NaVPO4F are replaced by VO5F octahedra connected by F vertices, leading to [V2O10F] units. In all, 0.56 Na ions/f.u. (87 mAh g
1) can be extracted from this phase through a stepwise process with two plateaus at 3.6 and 4.0 V (Sauvage et al., 2006). A Na1.5VOPO4F0.5/rGO “sandwich structure” also showed 91.4% capacity retention after 200 cycles at C/10 (Xu et al., 2013). By tuning fluorine stoichiometry, Na1.5VOPO3.8F0.7 delivered 134 mAh g
1 at the average voltage of 3.8 V at C/10. Because of the initial oxidation state of the vanadium (3.8 +), > 1 e
/f.u. can be reversibly extracted (Park et al., 2013). Theoretical calculations and experimental analysis made on a series of Na3(VO1
x)2(PO4)2F1+2x (0  x  1) reported that the combination of V4+/V3+ and V5+/V4+ redox couples and the F/O distribution has a direct impact on Na deinsertion/insertion mechanisms and voltage profiles. The good cycling stability of these materials was due to low volume changes (2%) during sodium extraction/insertion. The electrochemical activity of the V3+ phase Na3V2(PO4)2F3 as cathode for SIB has been performed by Shakoor et al. (2012). The voltage profile is marked by two plateaus at 3.7 and 4.2 V, for an average potential of 3.95 V (V4+/V3+). Na3V2(PO4)2F3 delivered 120 and 94 mAh g
1 at C/20 and 4C, respectively. From Chihara’s work (Chihara et al., 2013), the best cyclability was obtained when the material was cycled between 4.3 and 2.3 V. The expansion of the voltage window to lower potential activated the V3+/V2+ redox reaction and affected the lifetime. In full cell against NaTi2(PO4)3 as anode, Na3V2(PO4)2F3 reached a stable capacity of 110 mAh g
1 over 50 cycles. Crystal structure of this fluorophosphates was determined in 1999 as a tetragonal P42/mmm (Le Meins et al., 1999). However, Bianchini et al. (2014a,b) restudied the structure by means of high angular resolution synchrotron radiation diffraction and revealed an orthorhombic distortion leading to the Amam space group. Moreover, the discrepancies in the cell parameters in the literature originated from a F/O mixing in the structure and the control of this composition has a great impact on electrochemical properties (Broux et al., 2016). Ponrouch et al. (2013) also highlighted the presence of four redox processes at 3.67, 3.70, 4.17, and 4.18 V in reduction instead of the two plateaus previously reported.

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The low-voltage domain (3.7 V) reveals three biphasic processes with the formation of the intermediate phases Na2.4V2(PO4)2F3 and Na2.2V2(PO4)2F3 and Na2V2(PO4)2F3. The high-voltage region (4.2 V) is made of a biphasic process, leading to Na1.8V2(PO4)2F3, and a solid-solution process for NaxV2(PO4)2F3 (1.8  x  1.3). The end-member NaV2(PO4)2F3 (space group Cmc21) is obtained via a last two-phase mechanism (Bianchini et al., 2015). Two Na can be reversibly extracted from Na3V2(PO4)2F3 for a theoretical capacity of 128 mAh g
1. Material optimization via carbon coating (Liu et al., 2015), particular morphology (Qi et al., 2016) or nanostructured composite (Zhang et al., 2016a) have been used in order to increase rate capability and specific capacity. In comparison to Na3V2(PO4)2F3, Na3Ti2(PO4)2F3 (60 mAh g
1) and Na3Fe2(PO4)2F3 (44 mAh g
1) provided poor performances (Chihara et al., 2013). However, another iron fluorophosphate, the Na2FePO4F phase, showed promising electrochemical properties as cathode for SIB. This compound crystallizes in the orthorhombic space group Pbcn and is made of Fe2O7F2 bioctahedra interconnected by PO4 tetrahedra forming [FePO4F]∞ layers. The two sodium cations are located in the interlayer space and can diffuse along these 2D pathways (Sharma et al., 2017). Around 0.8 Na were electrochemically extracted from this material and the phase NaFePO4F was stabilized by chemical oxidation with NOBF4 in acetonitrile (Ellis et al., 2007). Na2FePO4F, prepared via low-temperature ionothermal reaction, showed a reversible capacity of 110 mAh g
1 over 10 cycles (Recham et al., 2009) through two well-defined voltage plateaus at 3.06 and 2.91 V (Kawabe et al., 2011), corresponding to Fe3+/Fe2+ activity (Ellis et al., 2010a,b; Lee et al., 2011a). Ex situ 23 Na solid-state NMR revealed a two-phase mechanism associated with the extraction of one Na from Na2FePO4F, leading to the fully charged phase NaFePO4F (Smiley and Goward, 2016). Carbon-coated hollow spheres (Langrock et al., 2013) and ball-milled Na2FePO4F (Law et al., 2015) retained 80% of their initial capacity after 750 and 200 cycles, respectively. The Mn-fluorophosphate phase Na2MnPO4F does not reveal a layered structure but a 3D-tunnel framework (Ellis et al., 2010a,b; Recham et al., 2009). The substitution of Fe by only 0.3 Mn/f.u. in Na2[Fe1
xMni]PO4F led to the 3D structure of Na2MnPO4F (Wu et al., 2011). Specific capacities and cyclabilities of Mn-substituted and Na2MnPO4F were poorer than the iron phase but the operating voltage was improved with participation of Mn3+/Mn2+ redox couple at 3.53 V (Kawabe et al., 2012). First-principles calculations predicted the extraction of 1 Na/f.u. through two voltage plateaus at 3.71 and 3.76 V with the formation of the stable intermediate phase Na1.5MnPO4F and the oxidized phase NaMnPO4F (Zheng et al., 2013). Experimentally, a plateau was observed at around 3.6 V at C/20 (Wu et al., 2018). Proper material optimization led to a discharge capacity of 120.7 mAh g
1 (i.e., 96% of the theoretical capacity based on 1 Na/f.u.) (Zhong et al., 2015). However, the cycling stability and rate capability of this material is still insufficient for practical application (Lin et al., 2014a,b; Wu et al., 2018). As another example of transition metal fluorophosphate, we can cite Na4NiP2O7F2 which presents a very high potential of 5.0–5.2 V on the basis of the Ni4+/Ni2+ redox couple (Kundu et al., 2015).

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Sulfates Regarding the larger electronegativity of sulfur vs phosphorous, SO4-based compounds offer greater redox potential than PO4-based (Recham et al., 2010). Due to structural similarity with NASICON compounds, rhombohedral form of Fe2(SO4)3 has been demonstrated as sodium intercalation host and around 1 Na/f.u. was electrochemically inserted in this material (Mason et al., 2014) at an average potential of 3.2 V (Fe3+/Fe2+). The average potential of the redox couple Fe3+/Fe2+ can be slightly increased with hydrated and dehydrated sodium bisulfates. Indeed, the bloedite-type Na2Fe (SO4)2 4H2O (Reynaud et al., 2014), the kr€ohnkite-type Na2Fe(SO4)2 2H2O (Barpanda et al., 2014b), and the Na2Fe(SO4)2 (Reynaud et al., 2014) showed potentials of 3.30, 3.25, and 3.40 V, respectively. However, the reversible capacities of these compounds are lower than 100 mAh g
1 and are of limited interest for practical applications. Starting with Fe3+ phase, a similar capacity was obtained with eldfellite NaFe(SO4)2 thanks to the insertion of one Na/f.u. at 3.0 V (Singh et al., 2015) through single-phase reaction. As another iron sulfate, Yamada’s group (Barpanda et al., 2014c) developed a new class of alluaudite-type material of formula Na2Fe2(SO4)3 (Fig. 9). This phase deviates from the NASICON structure of the AxM2(XO4)3-type materials and does not show Fe2(SO4)3 units (Fig. 42). Instead, it is made of edge-sharing FeO6 octahedra forming Fe2O10 dimer units. These isolated dimers are bridged together by sulfate tetrahedra. The 3D framework leads to three different Na sites and the partially occupied Na2 and Na3 sites are located in cavities leading to 1D diffusion along the c-axis (Nishimura et al., 2016). Na2Fe2(SO4)3 showed a very high potential of 3.8 V and an initial reversible capacity of 102 mAh g
1, corresponding to 85% of the theoretical capacity. The high-rate capability of this compound was also proved with around 71 mAh g
1 at a fast rate of 10C (Barpanda et al., 2014c). Subsequent l

l

Fig. 42 Structural view of the alluaudite Na2+2xFe2
x(SO4)3 along the c-axis and NASICONtype structure Na3V2(PO4)3 along the a-axis.

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investigations indicated that stoichiometric Na2Fe2(SO4)3 was not stable and, instead, an off-stoichiometric Na2+2xFe2
x(SO4)3 (x  0.25–0.3) was synthesized (Oyama et al., 2015) with partial Na+ substitution in Fe sites. Mn substitution was done in Na2.5[Fe1
yMny]1.75(SO4)3 (y ¼ 0, 0.25, 0.5, 0.75, and 1.0) by using a low-temperature solid-state reaction in order to increase the voltage of Fe3+/Fe2+. However, due to Mn2+, partial substitution of Fe by Mn reduced the capacity. Authors attributed this inactivity to an excessively high potential of the redox couple Mn3+/Mn2+ in alluaudite phases. From ab initio DFT calculations, it was estimated that Mn-alluaudite Na2.44Mn1.79(SO4)3 should be electrochemically active around 4.4 V (Dwibedi et al., 2015) thanks to a 2D-like ionic diffusion (Araujo et al., 2016). As for phosphates, the use of fluorine increases the operating potential in sulfate materials. Thus, synthesis of sodium fluorosulfates NaMSO4F was investigated by Barpanda et al. and it consists of low-temperature reactions between NaF and MSO4 H2O (M ¼ Fe, Co, Ni, Mn, Mg, Cu, Zn) precursors (Barpanda et al., 2010; Reynaud et al., 2012). NaMSO4F are made of chains of corner-sharing MO4F2 octahedra along the c-axis. Octahedra are linked by fluorine and each oxygen atoms are attached to sulfate tetrahedra, forming M-O-S-O-M chains. This 3D framework leads to cavities in which sodium atoms are located (Tripathi et al., 2011). Among these phases, only NaFeSO4F has been found to be electrochemically active. But only 6% of the theoretical capacity (137 mAh g
1) was obtained through a plateau centered at 3.6 V (Barpanda et al., 2010) mainly due to low Na+ conductivity and large volume change during desodiation. Chemical oxidation of KFeSO4F allowed the synthesis of an orthorhombic FeSO4F in which 0.85 Na+ were reversibly inserted/extracted via a two-step process at 3.3 V (Recham et al., 2012). In order to prepare fluorine-free materials, the substitution of F
species can be done by OH
, leading to hydroxysulfate phases. Four phases has been reported in the system Na-Fe-SO4-OH and, among them, the natrojarosite NaFe3(SO4)2(OH)6 shows the greatest theoretical capacity (166 mAh g
1) since three sodium should be intercalated. This phase was synthesized via a simple precipitation method at 90°C from Na2SO4 and Fe2(SO4)3 5H2O in 0.01 M H2SO4 solution and it crystalizes in the trigonal R-3m space group. Its structure consists of layers of corner-sharing FeO4(OH)2 octahedral chains connected with SO4 tetrahedra. By discharging this material up to 1.5 V, 2 Na (120 mAh g
1) were reversibly inserted at an average potential of 2.72 V (Gnanavel et al., 2015). Interestingly, the reduced phase is amorphous but the desodiation led back to the natrojarosite crystalline structure. l

l

4

Ceramics as separators and solid electrolytes

Ceramics play a major role in the improvement of nowadays conventional LIBs with liquid electrolytes as well as in the realization of future ASSBs. In case of LIBs with liquid electrolytes, the separator is commonly an ion-permeable polymer sheet. By incorporating ceramic particles in these polymer separators or by applying a ceramic coating on the polymer sheets, certain properties of these separators can already be significantly improved, for example, the thermal stability or the wettability. In case of ceramic solid-state batteries, the liquid electrolyte and the separator is replaced by a ceramic ion conductor that provides a pathway for Li ions while

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electronically isolating the electrodes from one another at the same time. The pursuit for solid ion conductors that can compete with or even outperform their liquid analogs used in current batteries with respect to, for example, ionic conductivity or electrochemical stability, has steadily pushed the materials development in the recent years. As a result, a variety of promising Li- and Na-ion conducting materials might replace the currently used liquid electrolytes and enable solid-state batteries with improved battery performance. In order to realize feasible solid-state batteries, the ionic conductivity ideally should be above 10
4 S cm
1 or even higher, while the electron conductivity is strictly limited (<10
12 S cm
1) (Zhang et al., 2018a,b,c,d). Since some of the present solid Li- and Na-ion conductors are already fulfilling this requirement, the interface resistance between solid electrolyte and electrode materials is becoming more and more the performance-limiting factor ( Janek and Zeier, 2016). In this context, the material compatibility of electrolyte and electrodes is getting of major importance (Zhang et al., 2018a,b,c,d). Especially, the compatibility with metallic lithium (or sodium) on the anode side is a crucial factor to gain a real improvement in energy density for future solid-state batteries. Taking these key factors into account, Sections 4.2 and 4.3 will give an overview about the state-of-the-art oxygen- and sulfur-based ceramic solid electrolytes.

4.1 Ceramic composite separators The primary function of a separator in a conventional LIB is the physical separation of the two electrodes, while also providing a pathway for the transfer of Li ions through its porous structure, which is filled with liquid electrolyte. Accordingly, some of the most important properties a good separator should display are a high porosity, a good wettability with the liquid electrolyte, a high thermal as well as chemical stability, and good mechanical properties (Cheng et al., 2016). Nowadays, different single or multilayered polymer membranes made of, for example, polyolefin, PP, or polyethylene are usually used as separators in commercial Li-ion batteries (Huang, 2011). Although, these different kinds of separators are already widely used in conventional Li-ion batteries, the low thermal stability of these materials presents a serious issue and also a limiting factor for the safety of the battery. One approach to increase the battery safety during thermal runaway is the preparation of thermal shutdown separators that contain a combination of polymers with different melting points. In this way, during a thermal runaway of the cell, the lower melting polymer fills up the pores of the membrane to block the Li-ion transport while the higher melting polymer still provides a physical barrier between the two electrodes (Arora and Zhang, 2004; Zhang, 2007; Liu et al., 2018b). Another approach to further increase the safety of the separators and hence of the whole battery is the design of ceramic composite separators, where the properties of the polymer separators are improved by adding ceramic particles. There are different ways to realize these ceramic composite separators. One way is by coating a polymer sheet with ceramic particles and another way is the infiltration of a polymer matrix with ceramic particles (Fig. 43). Besides, showing an improvement in the thermal stability also the wettability of the composite separators with the liquid electrolyte can be enhanced by the ceramic particles. A general overview about the different

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Advanced Ceramics for Energy Conversion and Storage

Fig. 43 Sketches of ceramic composite separators realized by (A) coating an ion-permeable polymer membrane with ceramic particles and by (B) infiltrating a polymer membrane with ceramic particles.

combinations of ceramic materials and polymers that were used to prepare coated and infiltrated composite separator was recently published by Cheng et al. (2016). In case of the coating strategy, different materials like Al2O3, SiO2, or TiO2 have already been tested (Cheng et al., 2016). By using chemical vapor deposition (CVD) (Kim and Park, 2012), for example, coated a polyolefin-based separator with SiO2. Their experiments proved that the absorption of the liquid electrolyte is significantly improved by the SiO2 coating in comparison to the untreated PE separator. Furthermore, they showed that the thermal stability of the coated separators is drastically increased. The composite separator can maintain its dimensional integrity at a temperature of 150°C which is already beyond the melting point of PE (130°C) (Kim and Park, 2012). Beyond the fundamental research, there are also already commercially available ceramic-coated separators produced by, for example, Degussa Evonik or LG Chem (Nestler et al., 2014). A disadvantage of the coating strategy is the fact that the ceramic coating can delaminate from the polymer sheet during battery operation. By infiltrating the ceramic particles into a polymer matrix, the delamination problem can be circumvent. In literature, the infiltration of a variety of ceramic materials like γ-LiAlO2, TiO2, MgO, Al2O3, and SiO2 into different polymer matrices has already been tested (Cheng et al., 2016). Prosini et al. (2002), for example, already reported in 2002 about the usage of γ-LiAlO2, Al2O3, and MgO as fillers in a PVdF-HFP polymer matrix. In their tests, only the composite separator infiltrated with MgO particles showed a good cell performance while the Al2O3 and γ-LiAlO2 infiltrated ones were not stable with the anode, respectively, cathode. Overall, the addition of ceramic particles can drastically improve certain properties of the separators and improve the battery performance. The next step on the realization of an inherently safe battery could be the replacement of separator and flammable liquid electrolyte by a solid electrolyte that combines ion transport and separation of the electrodes in one component.

4.2 Oxygen-based solid electrolytes In this section, different oxygen-based solid electrolytes are discussed in the view of properties, which define their performance in future solid-state batteries, for example, ionic conductivity and electrochemical stability. While oxygen-based ceramic electrolytes generally show a better electrochemical stability than their sulfur-based counterparts, they are usually inferior in terms of ionic conductivity. Especially, the grain

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boundary resistance and the density, which can be influenced by processing, have a major influence on the total ionic conductivity of oxygen-based ceramic electrolytes (Zhang et al., 2018b; Manthiram et al., 2017). For the following considerations, some of the most relevant oxygen-based solid electrolytes are further separated into Na- and Li-ion conductors with respective subclasses, which combine different materials according to their structure.

4.2.1 Li-ion conductors A general overview about the temperature-dependent conductivities of some of the most promising oxygen-based Li-ion conductors is shown in Fig. 44. In the following sections, a more detailed discussion of the different material systems and their possible application in future solid-state lithium batteries is given. Temperature [°C] 300

100

200

100

0

Li-ion conductivity [S cm–1]

10–1 10–2 10–3 10–4 10–5 10–6 10–7 1.5

2.0

2.5

3.0

3.5

4.0

4.5

5.0

1000/T [K–1] LiPON Li3PO4 LiPON Li1.6Al0.6Ti1.4(PO4)3 Li1.5Al0.5Ge2–x(PO4)3 Li6.65Ga0.15La3Zr1.90Sc0.10O12 Li7La3Zr2O12 Li3x La2/3–xTiO3 (bulk) Li3/8Sr7/16Ta3/4Zr1/4O3 (total) Li4Al0.33Si0.33P0.33O4 Li14Zn(GeO4)4 Li-b-alumina LiPF6 in EC/DMC

Fig. 44 Comparison of ionic conductivities of different Li-ion conducting ceramics with liquid electrolyte (Stallworth et al., 1999; Yu et al., 1997; Bates et al., 1992; Fu, 1997a,b; Buannic et al., 2017; Murugan et al., 2007; Inaguma et al., 1993; Inada et al., 2014; Farrington et al., 1981; Alpen et al., 1978; Deng et al., 2017).

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Advanced Ceramics for Energy Conversion and Storage

Garnets The first ion conductor with garnet structure, Li5La3Ta2O12, was discovered in 2003 by Thangadurai et al. (2003). This material showed a room-temperature ionic conductivity of around 10
6 S cm
1, which is comparable to LiPON, and an outstanding chemical stability with molten lithium metal. Later, Murugan et al. (2007) described Li7La3Zr2O12 with an improved Li-ion conductivity of around 2  10
4 S cm
1 at room temperature and also a high chemical stability against lithium. Starting from that point hundreds of papers were published dealing with the optimization of chemical composition, processing conditions, and battery development. Li-ion conducting garnets were obtained in two different crystal systems, the tetragonal and the cubic crystal system. In general, cubic modifications show higher conductivities than the tetragonal structures. Li7La3Zr2O12 has a tetragonal structure at low temperatures, whereas the cubic high-temperature phase is not quenchable. The cubic structure can be stabilized by different substitutions on the Li- [e.g., Al (Hubaud et al., 2013) and Ga (El Shinawi and Janek, 2013)] or Zr-site [e.g., Ta (Logeat et al., 2012) and Nb (Ohta et al., 2011)]. Since there are further cubic structures [e.g., I-43d for Li7
3xGaxLa3Zr2O12, x > 0.07 (Wagner et al., 2016b)] known for Li-ion conducting garnets, in this section the structural properties of the tetragonal space group I41/acd and the cubic space group Ia-3d are explained. The latter one is the most detected space group for Li-ion conducting garnets. The cubic structure consists of lanthanum surrounded by eight O-atoms forming a dodecahedron (Wyckoff position 24c), an octahedral coordinated zirconium (16a) and two lithium sites: the 24d Li-site shows a tetrahedral environment, whereas the 96h site has octahedral coordination (Logeat et al., 2012). In Fig. 45, the cubic structure with one of each Li-O polyhedra is visualized. The occupation of the 96 h site is directly related to the ionic conductivity of the garnet, whereas the 24d site has much less influence. If the amount of lithium is too low, the conductivity is insufficient due Fig. 45 Structure of cubic Li7La3Zr2O12, generated from ICSD 422259 (Awaka et al., 2010).

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645

to too small amount of charge carriers. If the amount is too high, the charge carriers have no vacancy to hop on, so that ionic conductivity will be low, too. In comparison, when Li7La3Zr2O12 crystallizes in the tetrahedral space group I41/ acd, the Li-atoms show an ordering so that all octahedral sites and one-third of tetrahedral sites are occupied, leading to a lower symmetry. Since the Li-ion conduction is caused by hopping on the octahedral coordinated sites, a lower ionic conductivity is a consequence of the ordering (Awaka et al., 2009). Miara et al. used a DFT method to calculate the stability and site preference of different substitutions. Only few elements are able to substitute on the Li-site, for example, Al3+, Ga3+, Zn2+, and Fe3+, whereas a lot of elements will occupy the Zr-site, for example, Sc3+, Sn4+, W4+, Nb5+, and Ta5+. Most of the rare earth elements as well as the heavier alkali and earth alkali elements and also Y3+ prefer the La-site (Miara et al., 2015). These findings were supported by the experimental work of numerous researchers. A good overview about the experimentally investigated substitutions can be found in two review papers (Thangadurai et al., 2014; Ramakumar et al., 2017). High conductivities at room temperature were achieved by substitution of Zr by Ta or Nb. The composition Li6.75La3Zr1.75Nb0.25O12 showed a total ionic conductivity of 8  10
4 S cm
1 at 25°C and activation energy of 0.31 eV (Ohta et al., 2011). Li6.4La3Zr1.4Ta0.6O12 synthesized by solid-state reaction and sintered at 1230°C showed a high total conductivity of 1  10
3 S cm
1 at room temperature and an activation energy of 0.35 eV (Li et al., 2012). Hot-pressed Li6.6La3Zr1.6Ta0.4O12 was reported to have an even higher conductivity of 1.18  10
3 S cm
1 at 30°C, while the activation energy was 0.4 eV (Tsai et al., 2016). Li6.65Ga0.15La3Zr1.90Sc0.10O12 has the highest ionic conductivity in a garnet-structured electrolyte reported so far, namely 1.8  10
3 S cm
1 at 27°C. The high ionic conductivity was obtained by applying a dual-substitution strategy. First, the cubic structure was stabilized by introduction of Ga3+ on the tetrahedral 24d-site. Since this substitution reduces the total amount of lithium in the structure, Sc was substituted on the Zr site at the same time so that the amount of lithium remains at a large value of 6.65 per f.u. (Buannic et al., 2017). In order to understand the contributions of grains and grain boundaries in the material, Tenhaeff et al. investigated the temperature dependence of the Li-ion transport in hot-pressed Li6.28Al0.24La3Zr2O12. In the temperature range between
100°C and +60°C, the activation energy for the intragranular Li-ion hopping is 0.36 eV, whereas the activation energy for the Li-transport across the grain boundary is 0.44 eV. The activation energy for the total conduction process is reported to be 0.4 eV, which is in the range of most garnet materials (Tenhaeff et al., 2014). Garnet electrolytes are stable in air, that is, no decomposition will occur. However, the exchange of hydrogen from moisture with lithium was reported, leading to a stepwise formation of a Li2CO3 surface film due to CO2 via LiOH (Uhlenbruck et al., 2018). The Li-H exchange is only occurring on the 96 h-site, whereas the 24d site is unaffected (Galven et al., 2012). The electrochemical window of Li7La3Zr2O12 was calculated to be between 0.05 and 2.91 V vs Li +/Li. At higher voltages decomposition in La2O3 and La2Zr2O7 as well as oxygen release is predicted, however, experiments detected wider

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Advanced Ceramics for Energy Conversion and Storage

electrochemically windows than thermodynamically calculated, most likely due to slow kinetics of decomposition reactions (Zhu et al., 2015). Due to their high chemical and electrochemical stability against Li metal, garnets are expected to be implemented as solid electrolyte in all-solid-state lithium batteries. In the beginning of solid-state battery research, it was assumed that the formation of Li-dendrites is not possible due to the high mechanical strength of ceramic materials. First, galvanostatic lithium cycling experiments showed the opposite, as dendrites were observed already at low current densities (Ishiguro et al., 2014). Several explanations were suggested for this unwanted behavior. As an example, the effect of microstructure was investigated. Cheng et al. showed that a fine-grained surface improves the Li-cycling behavior (Cheng et al., 2015). Tsai et al. (2016) introduced an interlayer between lithium metal and garnet to improve the contact at the interface, suggesting the formation of point contacts as driving force for dendrite formation. As a solution, gold as interlayer that lead to reduced interfacial resistance and so an improved cycling behavior with lithium. However, at higher current densities lithium dendrites were still observed (Tsai, 2016). Following work introduced alternative interlayer materials, like thin films of ZnO (Wang et al., 2017a), Al2O3 (Han et al., 2016), and Ge (Luo et al., 2017). These publications also showed that the adhesion of lithium increases by adding an interlayer. Sharafi et al. (2017) proposed later that the main reason for bad lithium adhesion is an inaccurately cleaned surface. By applying a special wet-chemical polishing process, they were able to eliminate residual surface contaminations, for example, Li2CO3. By removing this unwanted species and by applying heat during the processing of lithium, the adhesion between garnet and lithium metal was increased, removing the need for additional interlayer. Galvanostatic Li-cycling confirmed the success of this process (Sharafi et al., 2017). Experimental investigations, as well as calculations, show that the thermal stability of garnet electrolytes with state-of-the-art electrode materials is limited. Hightemperature stability is necessary for processing of low-resistance solid-solid interfaces in all-solid-state ceramic-based batteries. LiCoO2 shows the highest thermal stability with garnet electrolytes up to 1085°C (Uhlenbruck et al., 2016), whereas other state-of-the-art active material like LiFePO4, spinels like LiMnMO4 (M ¼ Ni, Co, Fe), and NMC materials are decomposed during heating above ca. 500°C in contact with garnet (Thangadurai and Weppner, 2005; Miara et al., 2016). Significant differences in stability are observed for different garnet compositions, for example, Li5La3Ta2O12 shows higher-temperature stability with LiCoMnO4 than Li7La3Zr2O12 (Lobe et al., 2018). Several batteries with garnet electrolytes were already demonstrated in literature. One crucial point of cell development is the design of a cathode showing high capacity as well as well-defined interface with the solid electrolyte. To avoid a reaction during sintering with the cathode material while maintaining a good contact between cathode and electrolyte, in the beginning of cell development the cathode was applied by thin film processes, like sputter deposition (Kotobuki et al., 2011b) and wet-chemical techniques (Kotobuki and Kanamura, 2013). However, these processes result in only low cathode loadings. To increase the cathode loading, composite cathodes consisting of active material and additional components, like a Li-ion conductor, an electronic

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conductor, or a sintering aid, were developed. As a major method, screen printing was established as a method for application of a composite cathode with thicknesses in the double-digit micrometer range. A common sintering additive and Li-ion conductor is Li3BO3. Its low melting temperature allows the bonding of active material to the garnet electrolyte at comparable low temperature. Cells with a composite cathode consisting of Li3BO3 and LiCoO2 show initial discharge capacities of 58% (Ohta et al., 2014) and 74% (Ohta et al., 2013) of the theoretical value. For a mixture of Li3BO3 and LiMn2O4, a discharge capacity of 102.6 mAh g
1 was reached (Feng et al., 2017a). As further additives, indium tin oxide (ITO) is used as an electron-conducting additive. A composite of ITO, LiCoO2, and Li3BO3 shows a cathode utilization of around 70% at 150°C, after cycling crack formation is observed close to and in the cathode (Liu et al., 2017c). A surface-modified NCM-based (ASSB) which is also working with Li3BO3 as Li-ion conductor and ITO as electronic additive shows a discharge capacity of around 123 mAh g
1 in the first cycle. However, the cell shows a fast capacity fading (Liu et al., 2018a, b,c). To increase the overall cell capacity, a cell with a 25 μm thick LiCoO2/ Li6.6La3Zr1.6Ta0.4O12 mixed cathode was developed at Forschungszentrum J€ulich. This cell shows a capacity of 0.84 mAh cm
1 at a current density of 0.1 mA cm
2, which is the highest reported capacity for inorganic, garnet-based cells so far (Finsterbusch et al., 2018; Tsai et al., 2019). ASSBs with garnet electrolyte reported up to now have a bulk, thick electrolyte, which is detrimental for the cell’s energy density. These electrolytes are mainly processed by conventional pressing and sintering. Separators in ASSBs should be as thin as possible (down to a few microns as in solid oxide fuel cells, see corresponding chapter) in order to increase the volumetric and gravimetric energy density to a maximum, while still ensuring a long-term stable separation of the electrodes. Therefore, several deposition processes for garnets were developed. Up to now, no working cell with thin film electrolyte was realized, due to high crystallization temperature required for the garnet phase leading to detrimental reactions with electrode material. Garnet thin films were deposited from the gas phase by pulsed laser deposition (PLD) (Kim et al., 2013d; Reinacher et al., 2014), sputtering (Lobe et al., 2016, 2018), CVD (Loho et al., 2017a,b), and by wet-chemical techniques like sol-gel processes (Bitzer et al., 2016). In general, all processes suffer from the high surface to volume ratio, so that lithium is lost during sintering, which can lead to decomposition of the material or insufficient conductivity values. Nevertheless, garnet thin films were produced with sufficient ionic conductivities. The highest ionic conductivities at room temperature of garnet thin films are 4.2  10
6 S cm
1 (Ea ¼ 0.5 eV) for tetragonal Li7La3Zr2O12 (Loho et al., 2017a) and 3.8  10
5 for cubic Li5La3Ta2O12 (Loho et al., 2017b) by CO2-laser-assisted CVD, 1.0  10
5 S cm
1 for epitaxially grown Li7La3Zr2O12 on Gd3Ga5O12 by PLD (Kim et al., 2013d), 2  10
6 S cm
1 for Li6BaLa2Ta2O12 deposited by PLD (Reinacher et al., 2014), and 1.2  10
4 S cm
1 (Ea ¼ 0.47 eV) (Lobe et al., 2016) by sputter deposition. Films deposited by aerosol deposition show only low conductivities after deposition (Ahn et al., 2015; Hanft et al., 2017), a subsequent annealing step leads to an increase in conductivity up to 7  10
5 S cm
1 (Hanft et al., 2017).

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The thickness-dependent phase formation in tape-casting processes of Al-doped Li7La3Zr2O12 was investigated by Yi et al. Using calcined powder, they showed that a green film thickness of around 50 μm gives an ideal sintered tape with a thickness of <30 μm and conductivities between 8  10
5 (at
35°C) and 1.5  10
3 S cm
1 (at 85°C) (Yi et al., 2016). Ga-LLZ tapes produced by a similar process showed ionic conductivities of 1.3  10
3 S cm
1 at room temperature (Yi et al., 2017). Xu et al. (2018a) described tape-casted layered structures consisting of a dense garnet layer between two porous layers. On the cathode side, carbon nanotubes were first introduced into the porous structure followed by melting of sulfur, which is acting as cathode. On the anode side, a thin ZnO film was deposited on the surface to reduce the surface tension of liquid lithium, which was introduced in a second step. These structures were used in lithium-sulfur battery prototypes. For full functionality, liquid electrolyte was added to enhance the contact of the cathode side. The discharge capacity was close to 1200 mAh g
1, which shows the suitability to use garnet electrolytes as separators in lithium-sulfur batteries (Xu et al., 2018a).

Perovskites Perovskites with the general formula Li3xLa2/3
x□1/3
2xTiO3 got a lot of attention from the battery community when Inaguma et al. discovered high bulk ionic conductivity values for x ¼ 0.11 at room temperature in 1993. The bulk conductivity was given to be 1  10
3 S cm
1 and the grain boundary conductivity 7.5  10
5 S cm
1, leading to a total conductivity of 7  10
5 S cm
1. The activation energy of the bulk conductivity was 0.40 eV at room temperature showing a change at around 123°C– 0.15 eV, whereas the activation energy for conduction across grain boundaries was constant over the whole temperature range (from
50°C to 400°C) (Inaguma et al., 1993). The structure is described as a typical ABO3 perovskite structure with lithium, lanthanum, and vacancies (indicated by □) sitting on the A-site and Ti sitting on the B-site. The bottleneck of the ionic conductivity is formed by four adjacent BO6-octahedra. In addition to the cubic perovskite structure, where lithium, lanthanum, and vacancies are randomly distributed on the A-site, other structure types can occur due to ordering of the A-site atoms and vacancies as well as distortion of TiO6-octahedra. Tetragonal, orthorhombic, and hexagonal structure types were described for Li3xLa2/3
x□1/3
2xTiO3. Fig. 46 shows four unit cells of the cubic structure. The metal-oxygen polyhedra are visualized, as well as the bottleneck for Li-ion transport, which is formed by the rectangle of oxygen atoms between two adjacent A-sites (Stramare et al., 2003). Even though the bulk conductivity shows a remarkable high value, the total conductivity of Li3xLa2/3
x□1/3
2xTiO3 is low compared to other ceramic ion conductors, because it suffers from a low conductivity across the grain boundaries. Ma et al. (2014) highlighted two types of grain boundaries by high-resolution transmission electron microscopy (HR-TEM). Type I shows a large lattice mismatch and a depletion of lanthanum as well as a change of the atomic structure at the grain boundary. Furthermore, an additional depletion of lithium was detected. Type II shows a smaller

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649

Fig. 46 Structure and bottleneck of perovskite Li3xLa2/3
x 1/3
2xTiO3, shown at a cubic modification, generated by ICSD254044 (Stramare et al., 2003).

lattice mismatch so that the perovskite structure is preserved at the grain boundary. In addition, electron energy loss spectroscopy (EELS) also revealed La-loss. The phase changes and the depletion of lanthanum and lithium are suggested to be the reasons for low grain boundary conductivity. In order to increase the total conductivity of the material, several substitutions at the A-site as well at the B-site of the perovskite structure were investigated. Exchanging the lanthanum atom on the A-site by heavier rare earth metals, like Sm, Pr, or Nd leads to lower ionic conductivity. The reason is the decreased size of the bottle neck due to lower ionic radii of Sm3+, Pr3+, and Nd3+ (Itoh et al., 1994). Substitution by a larger element, like Sr2+, leads to a slightly increased bulk conductivity up to 2.54  10
3 S cm
1 for Li0.36La0.53Sr0.03□0.08TiO3 at 22°C due to a slight expansion of the bottleneck (Morata-Orrantia et al., 2003). As possible B-site substitution, Nb instead of Ti was investigated. In the solid-solution LixLa(1
x)/3NbO3 for x ¼ 0.1 the highest bulk conductivity of 4.7  10
5 S cm
1 is observed at 25°C (Kawakami et al., 1998). The Nb-based solid-solution (Li0.25La0.25)1
xSr0.5xNbO3 has a conductivity maximum of 7.3  10
5 S cm
1 for x ¼ 0.125 at room temperature (Kawakami et al., 1998). Furthermore, small amounts of Al can be substituted to the

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Advanced Ceramics for Energy Conversion and Storage

B-site leading to a bulk conductivity of 2.95  10
3 S cm
1 for La0.56Li0.36Ti0.97Al0.03O3 at 22°C. However, the grain boundary conductivity of 4.94  10
5 S cm
1 is still low (Morata-Orrantia et al., 2003). In contact with Li-metal, the Ti4+ in Li3xLa2/3
x□1/3
2xTiO3 is reduced leading to electronic conductivity, which makes it unsuitable as electrolyte in all-solid-state lithium batteries with metallic Li anodes. This process was investigated by sputtering lithium on top of a perovskite sample in a special in situ XPS setup. Ti3+, Ti2+, and Ti0 were observed as reduced species, while La3+ was not affected. At the same time, no lithium thin film was observed at the surface, so it can be expected that lithium is directly incorporated to the structure (Wenzel et al., 2015). The electrochemical window of Li3xLa2/3
x□1/3
2xTiO3 was calculated to be between 1.75 and 3.71 V vs lithium, limited in the low-voltage range by the reduction of Ti4+ and in the high-voltage range by decomposition to TiO2 and La2Ti2O7 as well as oxygen release (Zhu et al., 2015). Therefore, several elements were tested to replace the reducible titanium to increase the stability in contact with Li-metal. First investigations in the LiSr1.65□0.35Zr1.3Ta1.7O0.9 showed a total conductivity of 1.3  10
5 S cm
1 at 30°C (Thangadurai et al., 1999). Based on this result, stoichiometry and processing was optimized leading to Li3/8Sr7/16Ta3/4Zr1/4O3 with a total conductivity of 2.7  10
4 S cm
1 when sintered in a powder bed (Inada et al., 2014). Cyclic voltammetry of this material with liquid electrolyte and a lithium metal anode shows an onset voltage of 1.0 V which is attributed to the reduction of Ta5+ to Ta4+ (Chen et al., 2004). Several all-solid-state batteries with Li3xLa2/3
x□1/3
2xTiO3 electrolyte were described in the literature. The concept proposed by Kotobuki et al. uses a 3D-structured electrolyte in order to increase the area for the active material. LiMn2O4 and Li4Mn5O12 were used as cathode and anode material, respectively. The electrodes were applied by using a sol-based impregnation process, including a calcination step at 600°C. In one charge-discharge cycle, only 2% of theoretical capacity was utilized. Most likely, this low performance is related to processing issues (Kotobuki et al., 2011c). Another battery having a SnO2 anode and a LiMn2O4 cathode was demonstrated by Trong et al. The cathodes were applied by a thin film process in order to facilitate a good interface with the ceramic electrolyte. The discharge capacity of the cell reached 3 μAh cm
2 μm
1 at a current density of 2 μA cm
2, which is <7% of theoretical capacity (Trong et al., 2015). Li3xLa2/3
x□1/3
2xTiO3 (LLTO) shows good stability in contact with water and LiOH solutions, so it could be used as separator in Li-air batteries between liquid catholyte (LiOH) and anolyte (organic electrolyte). A proof of concept without optimized air electrode was published by Inaguma and Nakashima (2013). Their battery used a process-optimized LLTO separator with large grains, so that a total conductivity of 5.7  10
4 S cm
1 was reached. The LLTO tube separated the aqueous LiOH solution (catholyte) and the organic anolyte (1 M LiClO4 in EC/DMC). Carbon and lithium were used as cathode and anode, respectively. The cell showed a stable charge-discharge behavior at discharge currents of 60 and 120 μA cm
2 (Inaguma and Nakashima, 2013).

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651

Li-NASICON Another material subclass that includes potentially high Li- as well as Na-ion conductors consists of the so-called NASICON-type materials. The naming NASICON is short for NA Super Ion CONductor and it goes back to reports by Goodenough et al. (1976) and Hong (1976) from 1976, where Na-ion conductors derived from the basic formula Na1+ xZr2SixP3
xO12 (0  x  3) were presented. In case of the Li-ion-conducting NASICON-type materials, the general composition can be written as Li1+xM2
xMx0 (PO4)3, where the M-site is usually occupied by a tetravalent cation like, for example, Ge4+, Ti4+, Hf4+, or Zr4+ and the M0 -site can be partially occupied by an aliovalent substituent, for example, Al3+ and Ga3+ (Thangadurai and Weppner, 2006; Anantharamulu et al., 2011; Aono et al., 1990b). While in the unsubstituted case of, for example, LiTi2(PO4)3 a total conductivity of only around 2  10
6 S cm
1 (Aono et al., 1990b) at room temperature can be achieved, the substitution of the material can greatly enhance the total ionic conductivity. Among the various reported possible compositions, some of the most promising Li-ion conductors of the NASICON family originate from the material systems of Li1+ xAlxTi2
x(PO4)3 (LATP) and Li1+ xAlxGe2
x(PO4)3 (LAGP). The substitution of the tetravalent cation (Ti4+, Ge4+) by trivalent Al3+ ions, which comes along with an increase in Li concentration, improves the densification of the ceramic and the ionic conductivity of the material (Aono et al., 1990b). In this manner, a roomtemperature ionic conductivity of 7  10
4 S cm
1 (Aono et al., 1990b) and 2.4  10
4 S cm
1 (Aono et al., 1992) is achieved for Li1.3Al0.3Ti1.7(PO4)3 and Li1.5Al0.5Ge1.5(PO4)3, respectively. Despite their high ionic conductivity, one considerable disadvantage of NASICON Li-ion conductors with Ti4+ cations is the incompatibility with lithium metal. The Ti4+ is reduced in contact with Li metal and a mixed ion/electron-conducting interphase is formed. Unfortunately, due to its ionic and electronic conductivity, the growth of this interphase is not self-limited (Hartmann et al., 2013; Wenzel et al., 2015). On the other hand, the presence of redox-active cations in the Li-NASICON system does also imply that this material class can as well be used as active electrode material in solid-state batteries ( Jian et al., 2017). Exemplary for Li-NASICON-type electrolytes, Fig. 47 displays the rhombohedral crystal structure (R-3c) for Li1.3Ti1.7 Al0.3(PO4)3. The rigid framework of the crystal structure is formed by the MO6 octahedra, which are occupied by either Ti4+ or Al3+ cations and connected to the PO4 tetrahedra at the corners. Within this framework there are two Li sites, which form a 3D conduction path through the structure. While in case of NASICON materials with low Li concentration [e.g., LiTi2(PO4)3] only the Li(1) site is fully occupied, increasing the Li concentration in the structure leads to partial occupation of the Li(2) site, for example, in Li1+ xTi2
xAlx(PO4)3 (Adachi et al., 1996). As already mentioned for oxygen-based ion conductors in general, the total ionic conductivity of NASICON-type electrolytes is mainly controlled by its grain boundary contribution (Adachi et al., 1996). One approach to reduce the impact of the grain boundaries on the total ionic conductivity is the processing of the LATP and LAGP material systems as a glass ceramic. In this way, increased room-temperature Li-ion

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Advanced Ceramics for Energy Conversion and Storage

Fig. 47 R-3c NASICON structure for Li1.3Ti1.7Al0.3(PO4)3 (ICSD 257190). The three-dimensional conduction path through the structure is build up by the connection of fully occupied Li(1) and partially occupied Li(2) sites.

conductivities of 1.3  10
3 S cm
1 (Fu, 1997b) and 5.08  10
3 S cm
1 (Thokchom and Kumar, 2008) could be achieved for glass-ceramic LATP and LAGP, respectively. Another inherent processing challenge for NASICON-type solid electrolytes arises from the highly anisotropic coefficient of thermal expansion (CTE) (Oota and Yamai, 1986). While for the densification of the material sintering temperatures of around 1000°C are needed, during the subsequent cooldown of the material the anisotropy of the CTE causes the formation of microcracks if the grain size in the material surpasses a critical value ( Jackman and Cutler, 2012; Ma, 2016b). Therefore, much effort is put into the development of feasible synthesis routes that provide a small grain size, a proper densification, and a good conductivity. A critical factor for the processing of complete working solid-state batteries or mixed composite electrodes, as already discussed for the garnet Li-ion conductors, is the thermal stability of the Li-NASICON electrolytes with active electrode materials. As an example, Miara et al. (2016) published a study where they combined DFT calculations and experimental data from differential thermoanalysis and thermogravimetry in order to investigate the thermal stability of Li1.5Al0.5Ti1.5(PO4)3 with different high-voltage spinel cathode materials. While the LATP on its own shows a high thermal stability up to a temperature of around 1200°C, in combination with the spinel cathode materials the decomposition temperature is reduced to temperatures of around 600°C (Miara et al., 2016). In addition, Gellert et al. (2018) investigated the thermal stability of Li1.5Al0.5Ti1.5(PO4)3 in combination with different phosphate-based cathode materials like LiCoPO4, LiFePO4, and LiMn0.5Fe0.5PO4. LiFePO4 exhibits the highest stability among the investigated phosphates up to 700°C, whereas the LiCoPO4 did already show signs of decomposition at around 500°C (Gellert et al., 2018). Overall, these experimental results highlight that the combination of Li-NASICON-type electrolytes (requiring sintering temperatures of 1000°

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653

C or even higher for an acceptable densification) with state-of-the-art electrode materials is challenging. In literature, glass-ceramic sheets of LATP and LAGP were already used as substrate material for setting up solid-state batteries in combination with thin film electrodes. For example, in case of the LATP electrolyte (Dokko et al., 2007), glass-ceramic sheets prepared by the company Ohara Inc. were coated with a LiMn2O4 cathode via sol-gel method and Li-metal was used anode. The Li-metal was separated from the LATP by a polymethyl methacrylate gel-polymer electrolyte. In a similar approach, LAGP glass-ceramic sheets were combined with a TiO2 working electrode and Li metal (Kotobuki et al., 2011a). Furthermore, there are also reports on bulk-type solid-state batteries that make use of the glass ceramic as well as of the crystalline LATP or LAGP as solid electrolyte in combination with different electrode materials (Yu et al., 2018a; Robinson et al., 2017; Aboulaich et al. 2011). Recently, Yu et al. reported, for example, a monolithic all-phosphate solidstate Li battery that uses a Li1.3Al0.3Ti1.7(PO4)3 sintered pellet as electrolyte in combination with a thick screen-printed composite LiTi2(PO4)3 anode and Li3V2(PO4)3 cathode. They were able to cycle this cell setup at 30°C in the potential range between 0.5 and 2.2 V vs Li/Li+ maintaining 84% of the initial discharge capacity after 500 cycles (Yu et al., 2018a). LISICON Inspired by the high conductivities obtained in the Na1+ xZr2SixP3
xO12 (0  x  3) NASICON system, Hong also reported in 1978 on a material class named as LISICON that exhibit reasonable Li-ion conductivities (Hong, 1978). Despite the similar naming the LISICON material class does considerably differ from the previously discussed Li-NASICON. In general, the LISICON electrolytes are based on the general formula Li16
2xDx(TO4)4 with divalent (D) and tetravalent (T) cations like, for example, Mg2+ or Zn2+ and Si4+ or Ge4+. Li14Zn(GeO4)4 shows an ionic conductivity of 0.125 S cm
1 at 300°C, decreasing down to 2  10
6 S cm
1 at 50°C (Alpen et al., 1978). The structure of LISICON, shown in Fig. 48, is similar to γ-Li3PO4, so that all structures based on Li3PO4 are considered as LISICON-like materials (Abrahams et al., 1988). Substitutions inside the polyanion group in Li3PO4 leads to a maximum conductivity of 1  10
3 S cm
1 at 200°C and an activation energy of 0.53 eV for the composition Li4Al0.33Si0.33P0.33O4 (Deng et al., 2017). For the composition Li4Al1/3Si1/6Ge1/6P1/3O4 a room-temperature conductivity of 0.9 mS cm
1 was predicted by molecular dynamics simulation but not verified by experimental work, yet (Deng et al., 2017). Based on thermodynamic calculations, the stability window for Li14Zn(GeO4)4 was determined to be between 1.4 and 3.71 V (Zhu et al., 2015), which confirms the instability of this composition in contact with Li-metal (Alpen et al., 1978). Furthermore, materials with LISICON structure show a significant aging during time. Bruce and West (1982) described a decrease of factor 100 of the ionic conductivity of a LISICON pellet after 28 days storage in vacuum. The reduced conductivity of the aged sample was assumed to be due to an intragranular effect. These both detrimental properties make LISICON materials up to now unpracticable in all-solid-state batteries.

654

Advanced Ceramics for Energy Conversion and Storage

Fig. 48 Crystal structure of the LISICON material Li14Zn(GeO4)4 with space group Pnma (ICSD 100169). Only the Li(1) and Li(2) contribute to the conduction while the Li(3) and Li(4) site, which are also partially occupied by Zn, are part of the rigid crystal framework.

LISICON materials got famous when the sulfide analogs, called thio-Lisicon, were first described by Kanno and Murayama (2001). This discovery paved the way for the development of the large class of sulfide containing Li-ion conductors which are further described in the following section.

Glasses While in the previous subsections, the materials were classified according to their crystal structure, this subsection focuses on glassy ion conductors that lack of a long-range crystalline order. One of the most prominent glassy lithium-ion conductor is LiPON, the so-called LiPON, which was first reported by Bates et al. (1992). Whereas the previously discussed crystalline ion conductors are still under development, the amorphous LiPON electrolyte is already used in commercially available thin film solid-state batteries. Therefore, thin film processing of LIPON electrolyte layers with thicknesses of a few micrometers in combination with, for example, LiCoO2 cathode thin films and Li metal anodes is already quite elaborated. Thin film cells can be operated at room temperature and withstand thousands of charge-discharge cycles with reasonable C-rates and good capacity retention (Wang et al., 1996; Dudney, 2005; Lacivita et al., 2018). On a laboratory scale, electrolyte layers of LiPON and Li3PO4 were also already used for thin film batteries operated close to or even above 5 V vs Li+/Li. Yada et al., for example, combined a LiPON electrolyte with a LiCr0.05Ni0.45Mn1.5O4
δ high-voltage spinel cathode material and a Li metal anode to setup a thin film battery operating at around 4.7 V vs Li+/Li (Yada et al., 2014). In case of the pure Li3PO4

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655

electrolyte layers without nitrogen doping, Kuwata et al. already demonstrated a thin film battery with the high-voltage spinel cathode material LiCoMnO4 and Li metal that was cycled up to 5.5 V vs Li+/Li with a capacity retention of 99.4% after 20 cycles (Kuwata et al., 2014). This demonstrates quite nicely the wide electrochemical window of LiPON well above 5 V vs Li+/Li (Yu et al., 1997). Unfortunately, due to its relatively low conductivity of around 2–3  10
6 S cm
1 at room temperature (Yu et al., 1997; Bates et al., 1993a), the utilization of LiPON is limited to thin film applications. However, due to its stability in contact with Li metal, the material is nowadays not only considered as electrolyte in thin film batteries, but also as possible thin film functional layer in bulk-scale batteries, for example, for dendrite inhibition or for stabilizing the Li metal/SEI (Chung et al., 2004; Dudney, 2000; Kozen et al., 2015a,b). Whereas it was so far assumed that there is no reaction at all occurring at the LiPON/Li metal interface (Yu et al., 1997; Bates et al., 1992), Schw€ obel et al. recently demonstrated by an in situ X-ray photoemission study that LiPON indeed decomposes into products like Li3PO4, Li3P, Li3N, and Li2O in contact with Li metal. These interface reactions are limited by the formation of a passivation layer which maintains the good cyclability of LiPON thin film batteries (Schw€ obel et al., 2015). In addition to its reasonable ionic conductivity, LiPON also exhibits a very high electronic resistivity of around 1013 Ω cm (Bates et al., 1993a), which leads to negligible self-discharge in the thin film batteries. A considerable advantage of the LiPON in comparison to oxygenbased crystalline ion conductors is the low-temperature processing (e.g., via sputter deposition without active heating), which diminishes the thermal compatibility issues occurring when the material is implemented into a solid-state battery. Furthermore, the glassy electrolytes provide an isotropic Li-ion conduction network, which is not dictated by a rigid crystal framework, nor limited by grain boundary conductivities like in their polycrystalline analogs (Zhang et al., 2018b). Most commonly, LiPON thin films are prepared by a reactive radio frequency magnetron sputtering process using a Li3PO4 sputter target in combination with a pure or partial nitrogen sputter plasma (Bates et al., 1992, 1993a; Yu et al., 1997). Besides, there are also publications on the thin film preparation of LiPON by ALD (Kozen et al., 2015a,b), metal-organic CVD (Kim et al., 2013a), or PLD (Zhao et al., 2002). Independently of the processing technique, a drastic increase in ionic conductivity upon nitrogen incorporation in the Li3PO4 thin films is reported in literature. Bates et al. (1993a), for example, observed an almost 50-fold increase in ionic conductivity of sputter-deposited thin films starting from around 7  10
8 S cm
1 for pure Li3PO4 thin films up to 3.3  10
6 S cm
1 for a thin film composition of Li2.9PO3.3N0.46. Up to now the underlying mechanism behind the conductivity improvement by nitrogen incorporation is not fully understood and different theories are discussed in literature (Mun˜oz, 2012). Alongside LiPON, there is also some research and development on other oxygenbased glassy Li-ion conductors. Some of these electrolytes originate, for example, from the Li2O-SiO2-P2O5 (Bates et al., 1993b) or the Li2O-V2O5-SiO2 (Kuwata et al., 2004) system. But so far none of this other glassy material can surpass the overall performance of the LiPON solid electrolyte.

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Advanced Ceramics for Energy Conversion and Storage

4.2.2 Na-ion conductors Fig. 49 shows an overview of conductivities of some oxygen-based Na-ion conductors that are considered as possible candidates for Na solid-state batteries. In the following, these different material systems are discussed in more detail.

β-Alumina

β-Alumina is an ionic conductor which is obtained after high-temperature treatment of γ-Al2O3 with Na2O (provided by, e.g., Na2CO3) at 1600°C, but also other techniques, like wet-chemical processes (Lu et al., 2010). Two types of β-alumina are known, the hexagonal β-phase which has the chemical formula (Na2O)(Al2O3)11 and the rhombohedral β00 -phase with the chemical formula (Na2O)(Al2O3)6. During the synthesis process, Na2O slabs are introduced in the Al2O3 structure, like it can be seen in Fig. 50, leading to a high unidirectional sodium-ion conductivity perpendicular to the c-axis. The low activation energy for the Na-ion transport process is remarkable. For single-crystalline Na-β-alumina, conductivities up to 3.6  10
2 S cm
1 (Armstrong et al., 1976) were measured. The β00 -phase exhibits a higher conductivity due to the increased amount of Na-conduction paths. Single crystals of Na-β00 -alumina

Fig. 49 Conductivity overview for some solid-state Na-ion conductors of the material classes β-alumina (Hooper et al., 1976) and Na-NASICON (Ma et al., 2016a; Vogel et al., 1984; Naqash et al., 2017) plotted against inverse temperature. For comparison, the conductivity of a Na-ion liquid electrolyte (Ponrouch et al., 2012) at room temperature is also given.

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657

Fig. 50 Crystal structure of hexagonal Na-β-alumina. Generated from ICSD66561 (Edstr€om et al., 1991).

show conductivities up to 0.1 S cm
1 (Ea ¼ 0.2 eV) (Briant and Farrington, 1980), whereas polycrystalline mixtures of β-Na-alumina and β00 -Na-alumina with conductivities up to 1.2  10
3 S cm
1 (Ea ¼ 0.27 eV) at 25°C were demonstrated (Hooper, 1977). Na-β00 -Alumina is used as electrolyte in Na-sulfur-batteries, due to its high stability toward liquid Na-metal as well as to liquid sulfur. In the so-called ZEBRA battery, Na-β00 alumina is mainly used as separator, a liquid electrolyte is added at the cathode side to improve the contact between the metal chloride cathode and the Na-ion conducting separator (Sudworth, 2001). The Li-analog of Na-β-alumina can be prepared by ion exchange in melted salts. Due to kinetic issues, a two-step process is necessary. In a first step, Na-β-alumina is transformed to Ag-β-alumina by using a silver nitrate bed at 350°C. The second step is the ion exchange between Ag+ and Li+ in a mixture of LiCl and LiNO3 (Yung-Fang Yu and Kummer, 1967). Further processes, like direct and indirect incorporation of Li from the gas phase were also described (Koh et al., 2012). Measurements on single-crystal material showed a conductivity of 2.7  10
3 S cm
1 and activation energy of 0.24 eV. Even these values are very high for solid Li-ion

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conductors, the strong hygroscopic behavior prevent a practical use of the material up to now (Farrington et al., 1981).

Na-NASICON As already mentioned in the discussion of Li-ion electrolytes, the material class of Na super ion conductors (NASICON) is going back to two publications of Goodenough et al. (1976) and Hong (1976). In their original publications, they investigated the crystal structure and the Na-ion transport in Si substituted Na1+ xZr2SixP3
xO12 (NZSP) in the range 0  x  3. In a more general way, the formula of the Na-NASICON electro3+ 5+ 4+ lytes can be written as Na1+2w+ x
y+ zM2+ w Mx My M2
w
x
y(SiO4)z(PO4)3
z, which clearly highlights the huge variety of possible substituents in this material system (Guin and Tietz, 2015). The various substituents have a great influence on the ionic conductivity and the crystal structure. With respect to the crystal structure, Goodenough and Hong already showed for the NSZP system that the material can adopt a rhombohedral (R-3c) or monoclinic (C2/c) symmetry depending on the Si concentration. In the concentration range of 1.8  x  2.2 of the general formula Na1+ xZr2SixP3
xO12, the NSZP exhibits a monoclinic distortion, while in the residual concentration regime a rhombohedral crystal structure is obtained (Goodenough et al., 1976; Hong, 1976). In both cases, the rigid framework is built by corner-sharing metal octahedra (MO6) and by the SiO4/PO4 tetrahedra (Fig. 51). In the rhombohedral symmetry, two different lattice sites are occupied by Na ions, while in the monoclinic symmetry an additional third Na site

Fig. 51 Na-NASICON electrolytes (A) Na1.1Zr2P2.89Si0.11O12 (ICSD67014) and (B) Na3Zr2Si2PO12 (ICSD473) with (A) rhombohedral R-3c and (B) C2/c monoclinic crystal structure.

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contributes to the conduction pathway. Within the conduction pathway, the Na ions have to pass through two triangles (T1, T2), which are made up by oxygen atoms of two edge-sharing MO6 octahedra, and considered as bottleneck for the transport process. An effective mean ionic radius (reff) can be defined by reff ¼ ðrM2 + w + rM3 + x + rM5 + y + rM4 + ð2
w
x
yÞÞ=2 that combines the ionic radii of the different cations (rM) in one geometrical parameter. By comparing literature data of >100 different compositions of Na-NASICON electrolytes, Guin and Tietz (2015) pointed out that an optimal conductivity is ˚ . Furthermore, the optimal obtained with a mean effective radius of 0.72 0.1 A amount of charge carriers seems to be given for a concentration of around 3.3 mol of Na per f.u. (Guin and Tietz, 2015). Overall, the ionic conductivities of the different Na-NASICON ion conductors vary over several orders of magnitude in the range from 6  10
13 S cm
1 for NaGe0.5Ti1.5(PO4)3 (Carrasco et al., 1993; Guin and Tietz, 2015) up to 4  10
3 S cm
1 for Na3.4Sc0.4Zr1.6(SiO4)2(PO4) (Ma et al., 2016a). While Goodenough already reported a conductivity of 6.7  10
4 S cm
1 for x ¼ 2 in Na1+ xZr2SixP3
xO12 at room temperature (Goodenough et al., 1976; Hong, 1976), recently the so far highest room-temperature ionic conductivity within the Na-NASICON system of 4.0  10
3 S cm
1 was reported by Ma et al. for Na3.4Sc0.4Zr1.6(SiO4)2(PO4) (Ma et al., 2016a). The substitution of Zr by Sc given by the formula Na3+ xScxZr2
x(SiO4)2(PO4) was investigated in the concentration range 0  x  0.6. The trivalent Sc3+ was chosen as substituent for the tetravalent Zr4+, because the aliovalent substitution increases the Na content per f.u. while the difference in the ionic radii is with only 2.5 pm the smallest among all possible trivalent substituents (Ma et al., 2016a). In addition, they elaborated a new processing route for the material, the so-called solution-assisted solid-state reaction (SASSR). With this SASSR, they could achieve small (<50 nm) and nearly spherically shaped particles. This fine powder morphology promotes the sintering behavior of the material and the small grain size (2 μm) after sintering impedes the formation of microcracks, which usually occur in Na-NASICON-type materials during cooling after sintering due to the anisotropic CTE (Ma et al., 2016a). Ma et al. (2016a) also demonstrated the good electrochemical stability of Na3.4Sc0.4Zr1.6(SiO4)2(PO4) up to 6 V vs Na+/Na by testing the electrochemical window of the material in combination with Na metal and a gold thin film electrode. Another group of Na-NASICON electrolytes with quite high room-temperature ion conductivities originate from the material system Na1+ xHf2SixP3
xO12 where Zr4+ is fully substituted by Hf4+. By this isovalent substitution, a room-temperature conductivity of 1.0  10
3 S cm
1 was achieved for Na2Hf2Si2PO12, which is already around 50% higher than the one of the Zr analogous composition (Cava et al., 1982). In addition, Vogel et al. reported in 1984 that the conductivity can be further improved up to 2.3  10
3 S cm
1 at room temperature by increasing the mobile ion concentration in the structure to Na3.2Hf2Si2.2P0.8O12 (Vogel et al., 1984).

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Advanced Ceramics for Energy Conversion and Storage

Despite the exceptional high Na-ion conductivities of Na-NASICON materials so far there are only a few reports about solid-state sodium batteries based on this electrolyte material. In 2014, Lale`re et al. reported a sodium solid-state battery composed of a Na3Zr2SiPO12 electrolyte and Na3V2(PO4)3 as positive and negative electrode material. In order to provide a good electronic as well as ionic conduction composite electrodes with 25% of Na3V2(PO4)3, 60% of Na3Zr2SiPO12, and 15% of carbon additive were used for the cell testing. The complete battery cell stack was 560 μm in thickness and was prepared in a single step by spark plasma sintering. Due to the poor densification of the electrolyte during this co-sintering process, the battery cell showed a quite high internal resistance and needed to be operated at 200°C. At this operating temperature, the cell achieved 85% of its theoretical capacity at a C-rate of C/10 and an operating potential of 1.8 V (Lale`re et al., 2014). Another sodium battery that uses a La-substituted NASICON of the formula Na3.3Zr1.7La0.3Si2PO12 as electrolyte, Na3V2(PO4)3 as active material in a composite cathode and Na metal as anode, was published by Zhang et al. (2016d). Although they cycled their cell at an elevated temperature of 80°C with a slow C-rate of C/10, they could only achieve inferior electrochemical performance which they attributed to the poor contact between electrolyte and electrodes. In order to improve this poor contact and to be able to cycle the cell at room temperature a small amount of an ionic liquid was added. With this modification, it was possible to cycle the cell at a C-rate of 10C for 10,000 cycles with a nearly constant capacity of around 90 mAh g
1.

4.3 Sulfur-based solid electrolytes Sulfur-based solid electrolytes gained a lot of attention in the last years due to their outstanding Li-ion conductivity. Room-temperature conductivity values of 1 mS cm
1 and higher are usual for the different kinds of sulfur-based electrolytes (see Fig. 52). A further advantage, which makes the materials interesting for industrial use, is that the material can be synthesized by mechanical milling methods and densified by pressure under moderate heating (Sakuda et al., 2013), so that battery assembly is possible without high-temperature sintering steps. Nevertheless, a very precise control of structure is required, as it directly impacts the properties—that’s why up scaling of synthesis in a reproducible way is still a challenging issue. In addition, the sulfur-based electrolyte materials are highly reactive with oxygen and moisture, which leads to the formation of toxic H2S (Muramatsu et al., 2011). An inert atmosphere and dry environment are thus required for their synthesis, processing, and integration in cells, leading to higher infrastructure costs. First sulfide-based solid electrolytes were reported for glass systems. The glasses were synthesized by melting the binary sulfides in and subsequent cooling or quenching. Some glass systems and compositions are summarized in Table 2. The thio-Lisicon family is a family of crystalline Li-conductors with the formula LixM1
yM0 yS4 (M ¼ Si, Ge and M0 ¼ P, Al, Zn, Ga). The structure is similar to the (oxide)-Lisicons, which are based on the γ-Li3PO4 structure. In the pseudobinary Li4GeSe4-Li3PS4-system, three regions were identified with different monoclinic superstructures, depending on the cation ordering. The best conductivity can be found

Ceramics for electrochemical storage

661

Temperature (°C)

Ionic conductivity (S cm–1)

150 125

100

75

50

25

0

10 –1

10 –2

10 –3

10 –4 2.25

Li9.54Si1.74P1.44S11.7Cl0.3 Li7P3S11 Li6PS5Cl Li10GeP2S11 Li3.25Ge0.25P0.75S4 Na11Sn2PS12

2.50

2.75

3.00

3.25

3.50

3.75

4.00

1000/T (K–1) Fig. 52 Comparison of ionic conductivities of different sulfur-based electrolytes (Kato et al., 2016; Seino et al., 2014; Yu et al., 2018; Kamaya et al., 2011; Kanno and Murayama, 2001; Zhang et al., 2018a).

Table 2 Overview of some early investigated sulfur-based glassy Li-ion conductors. System

Author

Composition

Li2S-SiS2-LiI

Kennedy and Zhang (1988) Sakamoto et al. (1999) Mercier et al. (1981) Zhang and Kennedy (1990)

0.42 Li2S 0.25 SiS2 0.3 LiI 55.5Li2S 40SiS2 3Li3N (Li2S P2S5)0.55 (LiI)0.45 0.67Li2S 0.10 P2S5 0.23 B2S3

Li2S-SiS2-Li3N Li2S-P2S5-LiI Li2S-P2S5-B2S3

l

Conductivity at 25°C (S cm21)

Activation energy (eV)

1.3  10
3

0.31

1.5  10
3

0.28

l

l

l

l

l

10
3 1.41  10
4

0.4

l

for 0.6 < x < 0.8 in Li4
xGe1
xPxS4. Li3.25Ge0.25P0.75S4 exhibits the highest conductivity of 2.17  10
3 S cm
1 at 25°C and an activation energy of 20 kJ mol
1. This value is remarkably higher than for the oxidic analogs. The thio-Lisicon materials were described to be stable in contact with molten lithium and having a high electrochemical stability up to 5 V vs Li/Li+ (Kanno and Murayama, 2001). At the interface

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Advanced Ceramics for Energy Conversion and Storage

of Li3.25Ge0.25P0.75S4 with LiCoO2 high interfacial resistances were observed. Li4Ti5O12 and LiNbO3 coatings applied on LiCoO2 before cell processing reduced the contact resistance. Especially the LiNbO3-coating improved the cycling properties of all-solid-state cells (Ohta et al., 2006, 2007). Highly conducting compositions without additional metal ions like Ge were investigated because a higher electrochemical stability was expected. Glass-ceramic materials with the compositions 75Li2S 25P2S5 and 80Li2S 20P2S5 were obtained after high-energy ball milling at room temperature. The as-prepared room-temperature conductivities are 1.8  10
4 S cm
1 and 1.7  10
4 S cm
1. After a heating cycle to 500°C, the conductivities increase to 2.8  10
4 and 7.2  10
4 S cm
1 respec, tively. As a reason, the intermediate formation of a phase similar to the thio-Lisicon II phase in 80Li2S 20P2S5 during heating was proposed (Hayashi et al., 2003). These results also show that the conductivity of glass-ceramic materials is higher than for the pure crystalline phases obtained by solid-state reaction. Crystalline γ-Li3PS4 shows a Li-ion conductivity of 3  10
7 S cm
1 at room temperature, and the extrapolation of β-Li3PS4 gives only around 9  10
7 S cm
1 (Homma et al., 2011; Liu et al., 2013b). The importance of processing and its impact on properties can also be shown for 70Li2S 30P2S5. Using a solid-state reaction process, a conductivity of only 10
8 S cm
1 is observed. In contrast, mechanical milling and heating to 240°C leads to a room-temperature conductivity of 2.2  10
3. Heating to 360°C increases the room-temperature conductivity to 3.2  10
3 S cm
1 coupled with a low activation energy of 12 kJ mol
1 (Tatsumisago et al., 2006). By improving the processing of 70Li2S 30P2S5, Seino et al. obtained Li7P3S11 glass-ceramic material with a conductivity of 1.7  10
2 S cm
1 and a low activation energy of 17 kJ mol
1, which was the highest conductivity known for a solid Li-ion conductor in 2014. An electrochemical stability between 0.1 and 5 V vs Li/Li+ was also reported. A disadvantage is the reaction with oxide-based cathode materials during the heat treatment which is necessary for the consolidation process (Seino et al., 2014). Halide-containing, argyrodite-structured Li-ion conductors with the general formula Li6PS5X (X ¼ Cl, Br, I) were first reported in 2008. Solid-state NMR measurements and conductivity measurements suggested very high Li-ion conductivities in the range of 10
2–10
3 S cm
1. The cubic structure is consisting of sulfur and halide atoms, which are forming 136 tetrahedral sites in the unit-cell. PS4 tetrahedra are formed by occupation of P in four tetrahedral sites which are only formed by the S1 position. The remaining tetrahedral sites are partially filled with Li+ ions. This structure is similar to the structure of Ag9AlSe6, which is a well-known Ag+-conductor (Deiseroth et al., 2008). High values of 3  10
3 and 7  10
3 S cm
1 for Li6PS5Cl and Li6PS5Br were indeed confirmed by impedance spectroscopy (Rao and Adams, 2011). Synthesis by solid-state reaction at 550°C also led to a high conductivity of 4.96 mS cm
1 at 26.2°C and an activation energy of 0.33 eV for Li6PS5Cl (Yu et al., 2018). Argyrodite Li-ion conductors are decomposed in contact with oxide cathode materials, like LiCoO2, LiNi1/3Co1/3Mn1/3O2, and LiMn2O4, forming compounds like sulfur, lithium polysulfides, P2Sx, phosphates, and LiCl (Auvergniot et al., 2017). In 2011, Kamaya et al. reported the first Li-superionic conductor showing conductivity similar to liquid analogs. Li10GeP2S12 (LGPS), synthesized from the metal and phosphor sulfides by reaction at 550°C, showed an ionic conductivity of 12 mS cm
1 l

l

l

l

l

Ceramics for electrochemical storage

663

at 27°C and an activation energy of 24 kJ mol
1 in the range from
110°C to 110°C. The tetragonal structure consists of (Ge0.5P0.5)S4 tetrahedra, PS4 tetrahedra, LiS4 tetrahedra, and LiS6 octahedra and shows a preferred Li-ion pathway along the c-axis (Kamaya et al., 2011). Initially, it was reported to be stable with Li-metal, but later the formation of an interface layer was verified by in situ XPS studies (Wenzel et al., 2016). The stability of the LGPS/LiCoO2 interface during battery cycling was investigated in all-solid-state cells. To avoid a chemical reaction between LGPS and the positive active material, a LiNbxTa1
xO3-coated LiCoO2 was used. After 100 cycles, a mechanical failure of LiCoO2 as well as reactions between the LGPS and the cathode material were observed (Zhang et al., 2018b). In 2016, Li9.54Si1.74P1.44S11.7Cl0.3 (LiSiPSCl) with a conductivity of 25 mS cm
1 and an activation energy of 23 kJ mol
1 was reported by the Kanno group. The structure determined by Rietveld-refinement shows a 3D Li-ion pathway, explaining the increased conductivity in comparison to the isostructural LGPS. LiSiPSCl was demonstrated in batteries delivering either high current or high voltage. In the latter case, Li9.6P3S12 was used for the anode because of its higher stability with Li-metal compared to LGPS and LiSiPSCl. Due to the reactivity of electrolyte materials with LiCoO2, LiNbO3-coated LiCoO2 was used as positive active material in the cells. The batteries showed an outstanding performance, for example, discharge rates up to 150C at 25°C. The batteries were also cycled with high-capacity retention at 100°C, state-of-the-art Li-ion batteries cannot be cycled under these conditions due to the thermal instability of the liquid electrolyte (Kato et al., 2016). Similar to the oxides, sodium-ion conductors based on sulfur are also known. One example is the Na3SP4 glass ceramic which can reach, in dependence of the processing, conductivities up to 2  10
4 S cm
1 at room temperature with an activation energy of 27 kJ mol
1 (Hayashi et al., 2012). Another example is Na11Sn2PS12, which synthesis was inspired by the LGPS material. It has a high ionic conductivity of 1.4 mS cm
1 and an activation energy of 0.25 eV (Zhang et al., 2018a). However, so far sulfide-based sodium-ion conductors were not able to outperform the values of NASICON materials, which are still higher (see Section 4.2.2).

5

Outlook

Battery technology advancements, discussed in detail in this chapter, have enabled energy storage to meet the performance requirements for a range of applications. The recent development of battery technologies is mainly driven by the transition to sustainable intermittent renewable energy sources and emission-free electric mobility. The requirements for battery systems for grid-scale energy storage include low lifetime cost, long cycle life, high specific power, and capacity. For mobility applications, such as BEVs and airplanes, high-energy density is the primary requirement in order to increase the range while cycle life must allow battery use throughout the projected lifetime (e.g., ca. 10 years for BEV). Specific power and energy density of the modern battery technologies, as well as of various capacitors, are compared in the so-called Ragone plot (Fig. 53) (Kato et al., 2016). Due to the superior combination of the specific power and the specific energy,

664

103

Supercapacitors

All-solid-state batteries

at 100 °C at 100 °C at 100 °C

102 SC2

SC1

Li-S batteries LiS1

101

LiB2

Mg battery

1

2

2 SiB

B3

LiS2

–1

10

100

101

102

103

LiO2

Na-ion batteries

Li–O2 batteries

104

Specific energy, E (Wh kg ) –1

Li4Ti5O12/LiNbO3-coated LiCoO2 Li4Ti5O12 + LGPS/LGPS/LiNbO3-coated LiCoO2 + LGPS Li4Ti5O12 + LSiPSCI/LSiPSCI/LiNbO3-coated LiCoO2 + LSiPSCI Graphite + LPS/LPS | LGPS/LiNbO3-coated LiCoO2 + LGPS

SiB1: Na3V2(PO4)3 (NVP)+graphene/NVP+graphene SiB2: NVP+CNT/NVP+CNT SiB3: NVP+activated carbon (AC)/NVP + AC Al-ion battery: AI/graphite Mg battery: Mg/V2O5

SC1: activated carbon/activated carbon SC2: reduced graphene oxide RuO2/RuO2-polyaniline SC3: activated carbon/activated carbon

LiS1: Li/S (graphene+single-walled CNT) LiS2: Li/S

Fig. 53 The Ragone plots of different battery and capacitor technologies (Kato et al., 2016).

LiB1: graphite/LiCoO2 LiB2: Li/LiFePO4 LiB3: Li4Ti5O12/LiNi0.5Mn1.5O4

LiO1: Li/O2 (graphene) LiO2: Li/O2 (carbon nanofibres) LiO3: Li/O2 (carbon nanotubes(CNT))

Advanced Ceramics for Energy Conversion and Storage

LiB1

Al-ion battery

Si

100

LiB3

Si B

Li-ion batteries

LiO1 LiO

Specific power, P (kW kg–1)

SC3

Ceramics for electrochemical storage

665

Li(-ion) ASSBs still represent the most promising technology for electric mobility applications (Kato et al., 2016). For the stationary grid-scale storage of renewable energy, various competitive or complementary technologies have been considered depending on particular application characteristics (Malhotra et al., 2016). While the well-established LIB technology will hold its position on the electric mobility market for the next years, new beyond LIBs technologies, such as all-solid-state and alternative metal ion batteries, like SIBs will eventually be optimized for specialized applications, especially where higher-energy and power density, faster charging, and lower cost are required (Wernecke and Morgenroth, 2018). For all these applications, specific inorganic materials for cathodes, anodes, and electrolytes functional material are required and highly reproducible, cheap processing into reliable cells is needed.

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