Among the candidates for electrodes in future Li-based batteries, LiFePO4 (LFP) is one of the most important and most frequently studied ones. LFP belongs to the class of phase transformation cathodes (phase transformation to FePO4 (FP)). Nevertheless, there is still an extensive debate on the mechanism of the transformation. This is mainly due to the lack of in-situ observations with appreciable space and time resolution as well as due to the undefined state of the defect chemistry of most of the studied LFP materials, even though fundamentally necessary for an overall understanding of the materials behaviour. To fill this gap, we build a micrometer-sized all-solid-state thin film battery (Fig. 1) with a defect-chemically well characterized LFP single crystal as cathode material and use scanning transmission X-ray microscopy (STXM) to follow the phase evolution along the (010) orientation, which exhibits the fastest ion conductivity, during in-situ electrochemical (de)lithiation .
By this we followed the reversible LFP transformation mechanism on a micro-meter scale with a lateral resolution of 30 nm and with minutes of time resolution, disclosing the influence of the defect chemistry, in terms of ionic and electronic conductivities, as well as elastic effects on the (de)lithiation process.
Fabrication of all-solid-state micro-sized thin film battery
We use oriented LiFePO4 single crystals grown via an optical floating zone technique, which have already been carefully characterized and their defect chemistry analyzed in our group. On top we deposit a solid electrolyte layer (LiF) and a layer of aluminum, functioning as anode material. Using a focused ion beam (Ga beam) inside a scanning electron microscope we fabricate out of this layered structure an all-solid-state thin film battery with attached platinum current collectors. The oriented LFP single crystal cathode has a size of 16 x 1 x 0.2 micrometer (c x b x a direction) (Fig. 1).
STXM mapping during (de)lithiation
Upon delithiation of LFP a shift in the main absorption feature at the Fe L3 edge from ≈708 to ≈710 eV occurs, which is used to fingerprint the change in the local state-of-charge of the single crystalline sample by identifying areas containing Fe2+ and Fe3+, respectively. To visualize the delithiation kinetics, area scans at the energies of the centroids of the respective Fe2+ and Fe3+ Gaussian XANES contributions and well before the Fe L3 edge are performed and the optical densities at each point of the sample are calculated as a function of time and increasing voltage, recording the transition between the two oxidation states and therefore the (de)lithiation of LFP. Six images of compared optical densities, taken during (de)lithiation, are shown in Figs. 2(b) and 2(c). Energy spectra recorded across the phase front illustrate the transition from LiFePO4 to Li1-xFePO4/FePO4, showing a constant increase in Fe3+ concentration (Fig. 2(d)).
Two main characteristics during (de)lithiation are observed. First, upon delithiating the electrode material, the growing FP phase forms at the current collector side, while lithiation of the delithiated cathode material leads again to a growing LFP region at the electrolyte side under a retracting movement of the delithiated phase. Second, it can be seen that the delithiation of the LFP cathode material of the micro-battery does not occur homogeneously over the whole length of the sample, resulting in a flat two-phase boundary, but rather multiple filaments of delithiated material develop along c and grow along b and a. To understand these observations they have to be discussed in terms of ionic and electronic transport as well as elastic effects:
The start of delithiation and FP phase formation at the LFP-current collector contact rather than at the LFP/LiF interface indicates faster ion than electron transport σion (LFP) > σeon (LFP) along b in the studied LFP crystal (Fig. 3(a)). Owing to the defect chemistry we refer to the D-regime in Fig. 3(c). A similar rationale explains why the LFP phase starts growing in the delithiated cathode material at the electrolyte side (σeon (FP) > σion (FP), compare Fig. 3(b)). Hence, the position where the delithiation of LFP and lithiation of FP starts can well be understood in terms of defect chemistry.
Strikingly upon delithiation the FP grows filament-like in b direction. Was it the ratio of electronic to ionic conductivities in the initial phase (σeon (LFP) / σion (LFP)) that decided upon the starting position, so is it now the ratio of ionic to electronic conductivities in the two different phases LFP and FP (σion (LFP) / σeon (FP)) forming the criterion for growth morphology, if transport controlled (see figure 3a). Therefore, if transport controlled one would expect homogeneous growth of FP, as according to literature  and the above considerations the ionic conductivity along b-axis in LFP should exceed the electronic conductivity in FP. We can conclude that the growth morphology is not transport controlled. Rather, elastic effects offer a straightforward explanation: the lattice parameter change upon delithiation forces the lattice to laterally expand in c direction. Thus elastic effects favor an exclusion zone around the FP nuclei which prevents further nucleation of FP phase within the vicinity and hence causes a growth pattern which is characterized by almost regularly spaced filaments. The increased stress when the filaments grow along c is also reflected by crack formation when the sample is delithiated too far and well consistent with the finding that large single crystals of LFP disintegrate on deep delithiation . After reversal of current again the growth morphology, this time of LFP in FP, is in opposite to what is solely expected from transport criteria but in line with the above elastic considerations.
We succeeded in building a micrometer-sized all-solid-state thin film battery enabling us to follow in-situ the (de)lithiation mechanism of single crystalline LiFePO4 along the fast (010) orientation using scanning transmission X-ray microscopy. Using a defect chemical analysis it can be concluded that the growth pattern of both LFP and FP is dominated by elastic effects rather than being transport controlled. This conclusion is rather general and should not depend on the defect-chemical details. Moreover, the analysis recommends a particle size of ≈100 nm as optimum electrode grain size, as the diffusion is still rapid enough and elastic effects do not lead to mechanical failure.