International
Tables for
Crystallography
Volume H
Powder diffraction
Edited by C. J. Gilmore, J. A. Kaduk and H. Schenk

International Tables for Crystallography (2018). Vol. H, ch. 2.8, pp. 181-183

Section 2.8.3.2. In situ studies of electrode materials and in operando investigations of Li-ion batteries

H. Ehrenberg,a* M. Hinterstein,a A. Senyshynb and H. Fuessc

aInstitut für Angewandte Materialien (IAM-ESS), Karlsruhe Institut für Technologie (KIT), Eggenstein-Leopoldshafen, Germany,bTechnische Universität München, Garching b. München, Germany, and cTechnische Universität Darmstadt, Darmstadt, Germany
Correspondence e-mail:  helmut.ehrenberg@kit.edu

2.8.3.2. In situ studies of electrode materials and in operando investigations of Li-ion batteries

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Rechargeable energy sources in mobile electronics are mainly based on lithium-ion batteries. Their application relies on the mobility of the small Li ions, which move from the cathode through an electrolyte to the anode during charge and back during discharge. Intensive research is underway to improve the performance of such energy-storage technology. High gravimetric and volumetric energy and power densities are required. Other additional challenges are safety, lifetime, the temperature range of stable operation and production costs per unit energy at the battery level. Knowledge of the correlation between the electrochemical functionality and the structure of the electrode materials during Li exchange is essential in order to interpret the underlying mechanisms and degradation processes and to find a promising approach to better materials. The high reactivity of the cell components and the very strong interactions between materials inside an electrochemical cell require studies on complete operational devices by non-invasive in operando methods. So-called electrochemical `half cells' are often studied to follow structural changes in electrode materials. These are complete operational cells, but the electrode is cycled against an Li-metal counter electrode. Such half-cell studies are sometimes described as in situ studies. Limitations might occur with respect to fatigue studies and at very high charge and discharge rates, when the performance is determined by the Li-metal electrode. The classification of in situ and in operando methods is not unambiguous in structural studies on battery materials. Sometimes the term quasi in situ is used for studies where specific states of the materials are prepared electrochemically and handled in an Ar atmosphere with complete protection against humidity and air, but actually investigated ex situ (Oswald et al., 2009[link]).

Some early in situ setups have been described for neutron diffraction (Bergstöm, Andersson et al., 1998[link]) and transmission X-ray diffraction (Bergström, Gustafsson & Thomas, 1998[link]), and also at elevated temperatures (Eriksson et al., 2001[link]). Today, for example, good-quality full diffraction patterns can be obtained with exposure times well below 100 ms using synchrotron radiation (Herklotz et al., 2013[link]). Capillary-based micro-battery cells allow for in situ X-ray powder diffraction studies on one single electrode (Johnsen & Norby, 2013[link]). Even spatially resolved neutron diffraction studies are possible on commercial cylindrical Li-ion batteries (Senyshyn et al., 2015[link]).

The mechanism of Li extraction and insertion differs for different types of electrodes. In intercalation-type electrodes the topology of the host structure remains mainly unchanged, and suitable sites in the structure are either occupied by Li or are vacant in the delithiated state. This use of intercalation chemistry for electrochemical energy storage was established for a battery based on Li metal as the negative electrode and TiS2 as the positive electrode (Whittingham, 1976[link]). In commercial cells today the negative electrode is also based on intercalation and consists of layered graphite, which hosts Li during the charge cycle up to the composition LiC6. Another working mechanism for negative electrodes is electrochemical alloying with Li. The most promising examples involve Si, Al and Sn. These electrode concepts suffer from extreme volume changes: 100% for Al → LiAl or even 300% for Si → Li21Si5. In combination with the brittleness of these materials, the particles break down and become amorphous during successive charging and discharging, accompanied by contact losses and resulting pronounced fade in capacity. Two other mechanisms have also received considerable attention as they allow higher specific capacities. In a replacement reaction, one transition metal is replaced by Li while the topology of the structure is mainly preserved. During a conversion or displacement reaction the initial structures of transition-metal compounds, for example nanometre-sized oxides (Poizot et al., 2000[link]) or other binary compounds (Cabana et al., 2010[link]), are believed to be destroyed completely by either amorphization or phase transitions. The transition-metal ions are reduced to metallic nanoparticles, which are embedded in a complex network of Li2O and reaction products from the electrolyte. In spite of the loss of long-range order, an important short-range structure remains. This has been shown in detail for ternary Cu–Fe oxides (Adam et al., 2013[link]). The reduction of Cu2+ from CuO takes place through the formation of a Cu2O/Li2O composite, in which Cu2O reacts further to form Cu metal and Li2O. Spinel-type CuFe2O4 and CuFeO2 react to form α-LiFeO2 with the extrusion of metallic Cu and Fe nanoparticles. At even lower potentials against Li+/Li between 0.8 and 1.0 V, α-LiFeO2 is further reduced into metallic Fe nanoparticles and Li2O. While most of these displacement reactions suffer from very poor reversibility, good cycling stability was shown for Cu2.33V4O11 (Morcrette et al., 2003[link]). During cell discharge Li penetrates into the well crystallized copper vanadate, forming a solid solution up to an Li content of x = 0.6, when Bragg peaks of metallic copper were observed. The end result was a composite electrode of an amorphous Li–V–O matrix with dispersed metallic copper. The essential point is the reversibility, with the disappearance of the metallic copper and the recrystallization of the initial Cu2.33V4O11.

Two more examples belong to the class of intercalation materials: graphite as used for the negative electrode, and LiNiO2 as a candidate for the positive electrode. `Positive' and `negative' electrodes are the preferred terms for secondary batteries instead of `cathode' and `anode', because anode and cathode reactions match only for discharge (interchange between the two electrodes occurs for the charge process). A comprehensive summary of structure reports on lithiated graphite can be found in Johnsen & Norby (2013[link]). From the voltage plateau in cyclovoltammograms four distinct lithiated graphite phases have been postulated. However, only for two of them have complete structure models been reported and confirmed. The first is LiC12, P6/mmm (space group No. 191), a = 4.29, c = 7.03 Å, with C on the 12n site with x = 0.33 and z = 0.25 and Li on the 1a site; the second is LiC6, also P6/mmm, a = 4.31, c = 3.70 Å with C on the 6k site with x = 0.33 and Li again on the 1a site. According to the number of graphene layers that are needed for the smallest unit repeated by translational symmetry along the sixfold rotation axis, these structures are described as stage-II (LiC12) and stage-I (LiC6) compounds, like graphites intercalated with other alkaline elements. Note that in these phases the graphene layers are not shifted with respect to each other (AA sequence), in contrast to graphite (AB sequence). For a lower Li content, a much more complex structural behaviour was observed, including incommensurate Li distributions between the graphene layers, which were described as twisted bilayers (Senyshyn et al., 2013[link]). Higher-order reflections were observed for these phases and allowed indexing with a propagation vector (α, α, 0). Different structure models were discussed, but a complete description of the Li distribution is still lacking. Therefore, it is still an open question as to where the Li atoms in lithiated graphite are at low Li contents (below 1 Li per 12 C).

LiNiO2 is considered to be a promising positive electrode material (Ohzuku et al., 1993[link]). However, it has some limitations, which are directly linked to the underlying structure. A high degree of cation disorder, i.e. Li on the Ni site and vice versa, hinders Li transport within the layers. Furthermore, Li and Ni exchange takes place rather easily, in contrast to LiCoO2, because of a more favourable transport process through a tetrahedral interstitial site for Ni than for Co. In the cases of Li excess, Li1+δNiO2, or Ni excess, (Li1−δNiδ)NiO2, some Ni2+ ions exist, which have a very similar ionic radius to that of Li+. Therefore, it is nearly impossible to prepare stoichiometric LiNiO2 with perfect separation of Li and Ni onto distinct layers. The best samples with respect to cation order are obtained from NaNiO2 by successive Na → Li ion-exchange reactions. These drawbacks have prevented the commercial application of LiNiO2, and more complex materials like Li(Ni,Co,Al)O2 (NCA) and Li(Ni,Co,Mn)O2 (NCM) are increasingly replacing LiCoO2. Fig. 2.8.12[link] shows the structural changes in LiNiO2 during the first charge and discharge. The detailed experimental conditions are the same as those described by Nikolowski et al. (2005[link]) for Li(Ni0.8Co0.2)O2. One of the most characteristic features of the structural response to Li extraction and insertion is the pronounced change in the lattice parameters, shown by the changes in the c/a ratio for the rhombohedral structure. During Li extraction the c parameter increases, because there are fewer Li ions between (repulsive) O-atom layers. However, at lower Li contents, some of the O ions become oxidized, and the repulsion between the O-atom layers is weaker, resulting in shorter c-axis parameters. As a general rule, all layered oxides LiMO2, with M = Mn, Co and/or Ni, become intrinsically unstable in the highly de­lithiated states beyond the maximum in the c/a ratio. Note that the as-prepared material has a very high degree of cation dis­order (9.3% Ni on the Li site) and very poor capacity retention in the second cycle. The initial phase (A) gradually disappears and apparently transforms into a second phase (B) with a much lower degree of cation disorder (less than 3%). Note that this phase sequence does not necessarily reflect equilibrium conditions, but depends strongly on the chemical composition (Li or Ni excess), microstructure (size and strain) and the experimental conditions (charge rate, temperature, electrode formulation and more).

[Figure 2.8.12]

Figure 2.8.12 | top | pdf |

Li1+δNiO2 during charge and discharge. From about 210 complete diffraction patterns [a small section is shown below (λ = 0.499366 Å)], the structural response to Li extraction and insertion was monitored. In addition to changes in the unit-cell metric, the distribution of Li and Ni onto layers becomes more ordered during the first highly charged state. A pronounced capacity loss is observed during discharge after a holding time of 3 h in this overcharged state. A, B and C are three successively appearing phases. SOC = state of charge.

Such in situ studies are very important for elucidating the working mechanism and degradation processes for intercalation electrodes (Senyshyn et al., 2012[link]). Nevertheless, complementary methods are also essential for providing all the necessary information, especially about the surface near-interface region between the electrode and electrolyte, which has to be studied with surface-sensitive and local methods.

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