In Situ Neutron Diffraction Monitoring of Li7La3Zr2O12 Formation

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In situ Neutron Diffraction Monitoring of Li7La3Zr2O12 formation: Towards a Rational Synthesis of Garnet Solid Electrolytes R. Prasada Rao, Wenyi Gu, Neeraj Sharma, Vanessa K. Peterson, Maxim Avdeev, and Stefan Adams Chem. Mater., Just Accepted Manuscript • Publication Date (Web): 01 Apr 2015 Downloaded from http://pubs.acs.org on April 2, 2015

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Chemistry of Materials

In situ Neutron Diffraction Monitoring of Li7La3Zr2O12 formation: Towards a Rational Synthesis of Garnet Solid Electrolytes. R. Prasada Rao†, Wenyi Gu†, Neeraj Sharma‡, Vanessa K. Peterson‡, Maxim Avdeev‡, and Stefan Adams†, §,* †

Department of Materials Science and Engineering, National University of Singapore, Singapore 117575 (Singapore)

‡

Bragg Institute, Australian Nuclear Science and Technology Organisation, Locked Bag 2001, Kirrawee DC, NSW 2232, Australia. §

Institute of Materials Research and Engineering (IMRE), Agency for Science, Technology and Research (ASTAR), Singapore 117602 ABSTRACT: The favourable combination of fast-ionic conductivity and high electrochemical stability of Li-stuffed garnet type Li7La3Zr2O12 (LLZ) makes this material a promising candidate for applications as a solid-state electrolyte in high energy density batteries. However, a wide-spread technical use of LLZ is impeded by difficulty in reliable formation and densification of the pure fast-ion conducting phase. The present study of the phase-formation process enables rational fabrication-procedures to be devised based on a thorough understanding of the complex phase-formation of LLZ. In situ neutron powder diffraction monitoring of the phase formation revealed an influence of the partial melting of precursors on the formation of the fast-ion conducting phase, indicating that in the typical synthesis route LLZ is not formed in a solid-state reaction but from a partial carbonate melt that decomposes on further heating. The cooling rate critically influences lithium ordering and ionic conductivity.

INTRODUCTION High conductivity solid electrolytes (SEs) that are stable in contact with lithium are required for safe high-energy batteries, e.g. Li-air and Li-sulphur batteries in which the SE protects the lithium anode from chemical attack by the liquid electrolyte. Additionally, SEs open the way for the complete replacement of flammable organic liquid electrolytes in all-solid-state lithium batteries, where the fast ion-conducting solid acts as both the electrolyte and as a mechanical barrier preventing short-circuiting by Li dendrite growth. To improve the power performance of all-solid-state lithium secondary-batteries, SEs need to exhibit a combination of fast lithium ionic conductivity of ~ 10-3 Scm-1 at room temperature with negligible electronic conductivity, high electrochemical stability against reactions with lithium anodes and the cathode, as well as be suitable for industrial fabrication of dense solid electrolyte membranes. Despite intense research on solid lithium ionconductors over the past decades, a material that fulfils all the above-mentioned requirements remains to be realised. As an example, NASICON-related glass-ceramic SEs such as lithium aluminium germanium phosphate, Li1.5Al0.5Ge1.5(PO4)3, exhibit lithium-ion conductivity as high as 5×10-3 Scm-1 at room temperature1,2,3 but are unstable against lithium metal, which rapidly leads to a

drastic increase of the interfacial resistance when in contact with lithium.4 LixPOyN(5+x-2y)/3, is stabile in contact with lithium metal and cathode materials (with a stability window of 0 - 5.5 V vs. Li/Li+) and can be fabricated into dense thinfilms by sputtering in a straightforward yet costly process. However, its very low ionic conductivity on the order of only 10-6 Scm-1 at 298 K limits application of LixPOyN(5+xto low-power thin-film batteries.5 Kanno and co2y)/3 workers reported high room-temperature Li-ion conduction of up to 2.2 ×10-3 Scm-1 in Li4-xGe1-xPxS4 with the thioLISICON structure, and Li10GeP2S12 with an unprecedented room-temperature conductivity of 1.2×10-2 Scm-1;6,7 and more recently, a room temperature ionic conductivity of 1.7×10-2 Scm-1 was reported by Seino et al. for Li2S-P2S5 glass-ceramic containing fast-ion conducting Li7P3S11.8 Deiseroth et al.9 and our group10,11 reported that lithium argyrodytes (Li6PS5X, X = Cl, Br) exhibit ionic conductivity on the order of 7×10-4 Scm-1. These sulphides are stabile in contact with Li metal, but are extremely hygroscopic, producing toxic H2S on contact with water, and their long-term stability in contact with high-voltage cathode materials is not proven. The Weppner group reported the highly stabile garnettype solid Li-ion conductor Li7La3Zr2O12 (LLZ)12. X-ray diffraction data of LLZ and its chemical properties, such as stability against molten lithium, matched well with

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earlier findings concerning the related garnet-type cubic material Li5La3M2O12 (M=Nb, Ta)13,14,15 although for the Zr-based garnet the structural stability of the fast ion conducting cubic phase at room temperature requires partial-replacement of Zr by dopants (such as Al3+ impurities from crucibles,16 intentional doping by tri- or pentavalent cations such as Nb5+, Ta5+ 17,18). Pure Li7La3Zr2O12 crystallizes at room temperature in the tetragonal space group I41/acd. The bulk ionic-conductivity of a sintered pellet (92% theoretical density) of the cubic phase was 5×10-4 Scm-1 at 25 °C, which is attractive for application as a lithium-battery electrolyte, while the room-temperature conductivity of the Li-ordered tetragonal phase is about three orders of magnitude lower.19 Shimonishi et al.20,21 reported LLZ to be stable in aqueous electrolyte, and a water-stable lithium-metal containing electrode can therefore be produced using LLZ without a polymerelectrolyte interface. While a minor reversible H+ for Li+ exchange is to be expected in aqueous solutions,22,23,24,25,26 it does not significantly affect the dominance of Li-ionic conductivity. Such a water-stable lithium-metal electrode may become a key component of a ‘‘protected’’ anode system in high energy-density lithium-air secondary batteries. 27,28 In contrast to atomistic18 and ab initio29 simulations clarifying that the stable low-temperature phase of pure stoichiometric LLZ is tetragonal, Xie et al. report that sintering at 1023 K yields the cubic high-temperature phase of LLZ in the absence of known dopants30 and Awaka et al.31 prepared a nominally-pure single crystal of cubic LLZ by sintering at 1530 K. Recent in situ hightemperature X-ray diffraction data of LLZ suggest that the formation of cubic or tetragonal LLZ depends on the evaporation of CO2 from the precursors that may act as a heterogeneous dopant.32 Another study of undoped LLZ using neutron diffraction at 4K and 300K suggested that at 0.9 wt% Al is necessary for stabilisation of LLZ in cubic structure.33 In another study by Thompson et al., using a combination of Raman and neutron diffraction studied indicated that 0.4–0.5 Li vacancies per formula unit are required to stabilize the cubic LLZ.34 Yet another study using TGA indicated the 0.6 Ta is the minimum substitution level to stabilise the cubic structure35. These studies are contradicting each other. No temperature dependent in situ neutron studies were reported to understand the conflicting results in our knowledge. Another MD simulation study indicated that a few types of lithium clusters dominate in both phases of Li7La3Zr2O12, which leads to highly correlated motion of lithium atoms36. The local symmetry of these clusters dictates a ‘‘center-pass’’ mechanism as lithium goes through the bottleneck.

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synthesis of Al-doped LLZ followed by annealing at 900 °C.38 Sol-gel synthesis of Nb-doped LLZ had to be complemented by annealing at 1423 K to achieve acceptable conductivity.39 We previously reported the preparation of LLZ as well as Ta (x = 0.25) and Nb (x = 0.25) doped LLZ at 1223 K by an 18 h sintering of ball-milled precursors (Fritsch P6, milling for 5 h with 15 balls of 10 mm diameter in an 80 ml ZrO2 bowl).18 Our molecular-dynamics (MD) simulations revealed a temperature-dependent Li redistribution amongst the tetrahedral and octahedral coordination sites and partial ordering on cooling below the transition temperature. This implies that the temperature and rate of cooling to room temperature will affect the distribution of Li on the tetrahedrally and octahedrally coordinated sites as well as the degree of intermediaterange Li ordering and, consequently, the transport behaviour. Here we use in situ neutron powder diffraction to gain a deeper understanding of the formation and phasetransformation characteristics of LLZ. EXPERIMENTAL SECTION Li7-xLa3Zr2-xMxO12 (M = Nb, Ta, x = 0, 0.25) was prepared from stoichiometric Li2CO3, La2O3 (pre-dried at 1173 K for 12 h), ZrO2 and for the doped samples Nb2O5 or Ta2O5. Stoichiometric ratios of the starting materials were mixed by mortar grinder (Retsch RM200) for 15 min and then transferred into 9 mm V cans (undoped LLZ) or Pt cans (Nb and Ta doped LLZ). The sample container cans were equipped with a small hole preventing a build-up of pressure in the event of CO2 gas evolution. In situ neutron powder diffraction (NPD) data were collected for the Li7-xLa3Zr2-xMxO12 precursor mixture for 16° < 2ϴ < 136° using the high-intensity diffractometer, WOMBAT, at the Australian Nuclear Science and Technology Organisation (ANSTO) with a wavelength of 2.5191(1) Å (determined using the NBS SRM 676 Al2O3). WOMBAT features a 120 ° area detector and data were acquired every min during heating of the sample in a high-vacuum furnace from 384 to 1050 °C with a ramp rate of 2.15 °C/min where the sample was held (±15 °C) for 6 h before the furnace being switched off. Data continued to be acquired during cooling to room temperature. Samples for ex situ neutron powder diffraction analysis were annealed at 850 °C, 950 °C, or 1050 °C for 18 h in a Pt crucible. This temperature is significantly-lower than the commonly-applied annealing temperature of 1200 °C that leads to a partial decomposition of LLZ due to the evaporation of Li. The annealed samples were characterized using the high-resolution neutron powder diffractometer, ECHIDNA,40 at the ANSTO (λ = 1.6220 Å) in the 2ϴ range 10 – 165° and by X-ray powder diffraction (Cu Kα, 10° < 2ϴ < 100°, nominal scan rate of 120 s-1, step size 0.016°).

The above LLZ synthesis and densification methods typically require high temperature (~ 1530 K) and/or longtime sintering (~48 h), so that preventing losses of lithium and oxygen from the sample becomes a technical issue. Attempts to produce cubic Al-free LLZ by a sol-gel method and annealing at only 700 °C lead to products with low total conductivity.37 The same applies to sol-gel

Rietveld refinements structural models using the X-ray and neutron powder diffraction data were performed with the General Structure Analysis System (GSAS),41,42 along with the graphical user interface EXPGUI. In the refine-

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ments, the distribution of the n = 56 Li per unit cell in LLZ or n = 54 Li per unit cell in LLZT on the two crystallographic site-types Li(1) (24d) and Li(2) (96h) the siteoccupancy factors (SOFs) are constrained to be correlated such that SOF(Li(2)) = n/96 – SOF(Li(1)/4). Thus, only the variation in the SOF of Li(1) is discussed.

amounts of the intermediate products La2O2CO3 and La2Zr2O7, in addition to the reactants Li2CO3, La2O3, and ZrO2. Therefrom it may also be concluded that the degree of amorphisation should not be a major influencing factor, as the sample anyway completely crystallizes again before the LLZ formation sets in. Further heating to 720 °C causes a decrease in the phase fractions of the La2O2CO3 and La2Zr2O7 intermediates. Key LLZ formation processes occur on reaching the Li2CO3 melting point (720 °C). The presence of a liquid phase is evidenced by the sudden appearance of broad background features in the diffraction data at ~ 170 min. The mass fractions of crystalline La2O3 and ZrO2 rapidly decrease so that La2Zr2O7 becomes the dominant crystalline phase, with excess La in the carbonate melt leading to a resurgence of the La2O2CO3 phase fraction. Above 800 °C LLZ is observed, with its phase fraction rapidly increasing above ~ 850 °C. At 950 °C all other crystalline phases disappeared and at 980 °C the evaporation of the carbonate melt is completed (as seen from both the constant sample mass and constant absolute scaling factors for the diffraction intensities), leaving pure LLZ.

It should be noted that during the in situ NPD measurement the amount of crystalline sample in the beam varies with temperature due to the presence of amorphous phases, including the carbonate melt, as well as the evaporation of CO2. Therefore the mass fractions are based on a model assuming a slight continuous decrease of the sample mass above 560 °C and a more pronounced drop around 840 °C before the sample mass stabilizes and remains constant above 900 °C and throughout the subsequent cooling. Computational geometry optimisations and molecular dynamics simulations were performed for 2×2×2 superstructures of Li6.75La3Zr1.75Ta0.25O12 using our dedicated forcefield that is based on Morse-type interactions derived from our softBV bond valence parameters,43,44 and the software package GULP as implemented in Materials Studio 6.0. Details of the forcefield are published in our earlier work.18

It should be noted that the indirect approach that is required here for the refinement of phase fractions due to the variable sample mass and the presence of amorphous or molten states will limit the achievable precision of phase fractions. Deviations of the observed elemental distribution from the expected one for stoichiometric phases were found to be within 10%, of which a part may be due to the mutual doping of the co-existing phases e.g. by La3+-doping of ZrO2 etc.

The samples are pressed into pellets of about 13 mm diameter and 2-3 mm in thickness. The density of the pellets are measured using gas pycnometry. The density of the samples are about 89-90%. Ionic conductivities of samples were determined by impedance spectroscopy (Schlumberger Solartron SI1260, frequency range 1 Hz - 1 MHz) using sputtered gold (≈ 100 nm thick) electrodes in the temperature range 300 K – 673 K. The samples were kept at each temperature for 20 min for thermal equilibration before recording impedance spectra.

In situ neutron powder diffraction monitoring of the heating of undoped LLZ from ball-milled precursors (see Fig. 1) in combination with a thermogravimetric study of the loss of CO2 from the carbonate precursors permits the determination of the complex sequence of reactions that occurs during the formation of LLZ. To this end the thermogravimetrically determined CO2 loss is translated into an expected variation of the unit-cell-mass-weighted sum of the scaling factors if all phases were crystalline (filled symbols in Figure 1(b)) with the help of the assumption that the sample is completely crystalline at the end of the heating cycle. The deviation of the results from the experimentally observed weighted scaling factor sum (open symbols in Figure 1b) yielded the amount of amorphous phase near room temperature (from the ballmilling of precursors) or of the carbonate melt around 800 °C.

On passive cooling the tetragonal room-temperature phase transition occurs at ~ 650 °C (Fig. 1); higher than expected. This is consistent with the 650 °C transition temperature recently reported by Matsui et al.32 after complete evaporation of CO2 (as expected in our open system), in comparison to the transition at ~ 150 °C for the CO2-containing system. Despite this agreement, we cannot completely rule out the possibility that the observed shift in transition temperature for undoped LLZ in our study was affected by unintentional V3+ doping as the vanadium sample container used in the experiment was attacked by the sample. Even usually-inert materials such as α-Al2O3 are attacked by this melt, with tri- and pentavalent dopants being incorporated into the LLZ. Due to the low scattering cross section of vanadium (which makes it a suitable container material for neutron diffraction experiments) a possible vanadium doping is difficult to probe directly. Trials to refine Zr and La site occupancies did not show any significant deviations from full occupancy. To safely prevent potential influences from vanadium reactivity, Pt cans were used in subsequent experiments, though this leads to Pt peaks in the diffraction patterns.

In the temperature range up to ca. 540 °C the initial glass-ceramic ball-milled precursor mixture fully crystallizes. At this stage the sample contains besides substantial

A mortar-ground precursor mixture for Nb-doped Li6.75La3Zr1.75Nb0.25O12 was studied during an analogous heating process. Although in this measurement the phase

RESULTS AND DISCUSSION

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formation was monitored over the entire temperature range from room temperature to 1000 K, the small weight fraction of dopants and its amorphisation and/or reaction with the ZrO2 during mechanical treatment did not permit us to identify the dopant as a separate phase. Hence, in our analysis of this sample the ball-milled precursor mixture for the Nb-doped LLZ was considered to consist of the same phases as for the undoped LLZ (crystalline ZrO2, La2Zr2O7, La2O3, La2O2CO3, La2Zr2O7, Li2CO3, and an amorphous sample fraction) and the nominal compositions ZrO2 and La2Zr2O7 refer to Nb-doped phases with unknown Nb content.

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with that stipulated by mutual Li-Li repulsion, as indicated by the dashed line in Figure 3. A higher occupancy of the tetrahedral sites would block so many of the adjacent Li(2) sites that there are no longer enough accessible sites for all the Li in the structure. Hence it is to be expected that the increase of SOF(Li(1)) has to stop when the blocking limit marked by the broken horizontal line is reached. For a more detailed discussion on the influence of mutual site blocking on the Li distribution can be found in our earlier work.18 At lower temperature data accuracy is reduced (as the cooling process is no longer controlled) and it cannot be conclusively determined from these measurements whether the high Li(1)occupancy approaching room temperature was due to lattice strain caused by fast cooling below 500 °C, or if this would also have been observed on continued slow cooling. In contrast to these results, separate furnace-cooled Ta- and Nb-doped, Ta/Ga and Nb/Ga co- doped, as well as undoped LLZ exhibited the low Li(1) site occupancy found also in the MD simulations. Hence, we conclude that the rate of cooling (following final annealing) strongly affects the distribution and ordering of Li+ ions in LLZ and is an essential parameter for controlling ionic conductivity at room temperature.

It should also be noted that the mortar-grinding of this sample mixture expectedly reduced the extent of amorphisation again. The more prominent formation of the intermediate product La2Zr2O7 in the neutron diffraction experiment at low temperatures might be traced back to the influence of the Nb-doping. The lower fraction of La2O2CO3 found in this sample indicates that the oxycarbonate formation requires the (mechanochemical) decomposition of Li2CO3. Again only minor changes of the phase composition occurred in the solid state, most notably the reduction of the phase fractions of La2Zr2O7. In this experiment the partial melting occurs at a significantly lower temperature T ~ 750°C, i.e. at a ~ 50 °C lower temperature than for the undoped LLZ and in two stages. Up to 40% of the sample is molten by ~ 750 °C. Again it is only in the presence of the melt that the phase composition of the when compared to the ball-milled undoped sample, and heating to 335°C was sufficient to completely crystallise the sample precursor mixture changes drastically: La2Zr2O7 becomes the dominant crystalline phase and the phase fraction of La2O2CO3 increases slightly before the formation of Nb-doped LLZ occurs temperature than for the undoped LLZ. The carbonate melt vanishes at ~ 950 °C where the pure garnet phase is obtained. In this experiment the rest time at 1000 °C was limited to 1 h and on cooling the Nb-doped sample remained cubic down to room temperature.

To shed more light onto the cooling rate dependence of the LLZ properties, we systematically studied ionic conductivity and lattice parameters of samples that differ only in the cooling rate after the final heat-treatment step. The impedance plot was fitted as bulk (Rb), grainboundary resistance (Rgb) and constant phase elements (CPE). As an example, Figure 4 presents Nyquist plots of impedance data for Li6.75La3Ta0.25Zr1.75O12. The solid line in Figure 4 represents the fit to the experimental data based on the equivalent circuit (RbQb)(RgbQgb)(Qel) using the ZVIEW program. For measurements above 150°C it appears difficult to segregate bulk and grain boundary contributions, so we have to consider the total (bulk + grain-boundary) contribution for determining the conductivity over the temperature range investigated. Figure 5 shows the temperature-dependent ionic conductivities of samples of Li6.75La3Nb0.25Zr1.75O12 after having been kept at 950 °C for 18 h to equilibrate the lithium distribution and subsequently cooled at different rates. The results reveal that samples cooled at a slower rate exhibit higher Li+-ionic conductivity, and lead to the conclusion that the cooling rate following the final heat-treatment should be sufficiently slow to allow for the Li+ distribution amongst the two site-types to remain in equilibrium during cooling. It may be noted in passing that for samples containing Ga3+ as a sintering agent the cooling rate dependence becomes more pronounced. Details on the influence of such sintering agents will be discussed separately.24 Local cation-ordering processes also appear to be the most straightforward explanation for the pronounced nonlinear variation of the cubic lattice parameter a for different pentavalently-doped LLZ samples that are consistent-

Rietveld refinements of the in situ data for undoped LLZ in the high-temperature section of the heating cycle (where the temperature varied linearly with time) yielded Li distributions that are in agreement with those obtained from our previous MD simulations (see Figure 3). Due to the controlled cooling in the second experiment and the non-reactive Pt vessel we could, for the Nb-doped sample, analyse the thermal expansion and the temperaturedependence of the lithium distribution of the garnet-type structure over a wider temperature range. We find that the lithium distribution follows the previously-assumed temperature-dependence only in the high-temperature regime. At ~ 750 °C a lithium-ion ordering process sets in as evidenced by the continuous increase of the tetrahedral Li(1) site occupancy up to the limit imposed by the mutual-occupancy exclusion of adjacent octahedral and tetrahedral sites. At 500 °C the site occupancy factor of Li(1) reaches this maximum value (x ≈ 0.583), consistent

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ly observed in our X-ray and neutron in situ studies as well as our MD simulations (see Figure 6). Analogous lithium clustering has been observed in simulations of the garnet phase Li5La3Ta2O12.45 Analogous non-linearities also occur on cooling cubic undoped-LLZ as a precursor to the cubic-to-tetrahedral phase transition that is triggered by complete lithium site-ordering.

magnitudes below those for the equilibrated structure. Though the pressure dependence of the ionic conductivity can only be estimated roughly from the available simulation data, at ambient pressure the conductivity of such cubic phases with high concentration of Li on the tetrahedral sites (and hence inevitably high degree of Li ordering) will hardly exceed the conductivity of the tetragonal undoped phase.

To gain a deeper understanding into the correlations between the lithium distribution, lattice constants and Li+ ion mobility we have conducted a series of atomistic simulations using the same bond-valence based atomistic forcefield that has been developed and used by our group previously to reproduce structural phase transitions and ionic conductivity.18 To this end we constructed and geometry-optimised a series of local structure models in which the lithium occupancies of tetrahedrally and octahedrally coordinated Li sites was systematically varied within the limitations set by the mutual Li+- Li+ repulsion. Examples of such configurations are shown in Figure 7. A visual inspection of the migration pathway networks already shows that increasing the occupancy of the tetrahedral sites (each blocking four neighbouring octahedral sites) requires an ordering of the octahedral Lithium on the distant octahedral sites and is thus detrimental to the local mobility. The formation energies as determined from the geometry-optimised simulated local structure models (see Figure 8) suggest that the structure with a site occupancy factor SOF(Li(1))= 1/3 is the most stable structure at ambient pressure. Both the geometry optimisations and subsequent room temperature molecular dynamics simulations of 2×2×2 supercells (cf. Figure 8) demonstrate that the variation of the room temperature lattice parameters is closely linked to the formation energy reaching again a minimum for SOF(Li(1))=1/3 and significantly higher values for configurations that maximise the occupancy of the tetrahedral site. Hence, it may be expected that tensile stresses on fast cooling the sample trigger the lithium ordering process.

It may be noted that a redistribution of lithium by hops between octahedrally and tetrahedrally coordinated sites will be possible with moderate activation energy and hence reasonably fast as long as the lack of long-range ordering renders all tetrahedral sites energetically (nearly) equivalent. This changes when intermediate-range order affect the energy levels of the tetragonal sites as well as of the octahedral sites, which can also be observed as a change in chemical shift in NMR studies.46 Within the ordered domain the energy landscape for ion transport looks more like in the tetragonal phase with significantly enhanced activation barriers that may be understood as defect association energies. Thereby the Li+ ions trapped in these ordered domains not only show a drastically reduced mobility (so that the concentration of effectively mobile Li+ drops) but also cannot redistribute sufficiently fast as they should to stay in thermal equilibrium when the temperature drops, which promotes more pronounced deviations from the equilibrium Li+ distribution and further Li ordering. Thus it is to be expected that the practically observed phase transition temperature for the cubic to tetragonal cubic to tetragonal phase transition temperature does not only depend on the dopant concentration but also on the cooling rate. This of course also translates into a cooling-rate dependence for the dopant concentration required to achieve a pure cubic phase at room temperature. CONCLUSIONS In summarising recommendations for the more-reliable synthesis of phase-pure fast ion-conducting (doped) LLZ that arise from this detailed in situ study, we conclude that the common addition of excess Li2CO3 will not only mitigate lithium loss, but may provide a flux to effectively accelerate reactions that would otherwise be slower in the solid state. Heating beyond 1000 °C appears to be detrimental to fast ion-conduction in LLZ as it leads to a partial decomposition of the sample. Higher temperatures are also not necessary to achieve a complete conversion to LLZ as long as the particle size of the reactants is sufficiently small and sufficient time for the completion of the garnet phase formation and the equilibration of the Li site distribution is spent at ~ 900 °C.

We also attempted to analyse the effect of the initial local lithium ionic conductivity on the ionic conductivity of Li6.75La3Zr1.75Ta0.25O12. Results of simulations at temperatures of 300 – 400 K for initial values of SOF(Li(1)) around 1/3 are consistent with our earlier reported MD simulations and experimental data of the fast-ion conducting phase, but a quantitative comparison with conductivities in systems with higher values of SOF(Li(1)) ≈ 0.5 is prevented by the limited stability. If the simulated period is sufficiently long to allow for an extraction of the Li+ diffusion coefficient, then this diffusion will substantially alter the Li distribution towards the equilibrium configuration. Higher Li concentrations on tetrahedral sites can be stabilized by fixing the cubic lattice parameters at sufficiently large values (e.g. 13.00 Å at 350 K corresponding to a pressure of -1.4 GPa for the case of a structure with SOF(Li(1))≈ 0.58), but despite the extra free volume for Lithium migration the mean square displacements obtained in such simulations remained one to two orders of

The study also showed that even for pure cubic garnet samples the observed room temperature conductivity will vary substantially depending on the cooling rate after the last high temperature treatment step. The synopsis of neutron diffraction data and atomistic simulations then

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clarifies the role of cooling-rate dependent lithium ordering for the ionic conductivity. In order to maximize conductivity the cooling rate needs to be sufficiently slow to allow for the distribution of Li+ among the two site-types to remain in equilibrium.

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D. Safanama, D. Damiano., R. Prasada Rao, S. Adams, Solid State Ionics, 2014,262,211-215. 4 P. Hartmann, T. Leichtweiss, M. R. Busche, M. Schneider, M. Reich, J. Sann, P. Adelhelm, J. Janek, J. Phys. Chem. C, 2013, 117 (41), 21064–21074. 5 J. B. Bates, N. J. Dudney, G. R. Gruzalski, R. A. Zuhr, A. Choudhury and C. F. Luck, Solid State Ionics, 1992, 53–56, 647–654. 6 R. Kanno and M. Murayama, J. Electrochem. Soc., 2001, 148, A742–A746. 7 N. Kamaya, K. Homma, Y. Yamakawa, M. Hirayama, R.Kanno, M. Yonemura, T. Kamiyama, Y. Kato, S. Hama, K. Kawamoto and A. Mitsui, Nat. Mater., 2011, 10, 682– 686. 8 Y. Seino, T. Ota, K. Takada, A. Hayashi, M. Tatsumisago, Energy Environ. Sci., 2014,7, 627-631. 9 H.-J. Deiseroth, S.-T. Kong, H. Eckert, J. Vannahme, C. Reiner, T. Zaiß and M. Schlosser, Angew. Chem., Int. Ed., 2008, 47, 755–758. 10 R. Prasada Rao and S. Adams, Phys. Status Solidi A, 2011, 208, 1804–1807. 11 R Prasada Rao, N. Sharma, V.K. Peterson, S. Adams, Solid State Ionics, 2013,189, 385-390.

AUTHOR INFORMATION Corresponding Author * S. Adams, Department of Materials Science and Engineering, National University of Singapore, Singapore 117575 (Singapore); E-mail: [email protected].

Present Addresses ‡ Neeraj Sharma’s current address: School of Chemistry, UNSW Australia, Sydney, NSW 2052, Australia.

Author Contributions R.P.R. synthesized and characterised the materials and analyses the neutron diffraction data. N. S., V.K.P. contributed to the neutron diffraction data collection and interpretation at WOMBAT. M.A. contributed to data collection and interpretation at ECHIDNA. W.G. studied the cooling rate dependence of the conductivity. S.A. was the project coordinator and conducted the Molecular Dynamics simulations. R.P.R. and S.A. prepared the manuscript.

NOTES

12 R. Murugan, V. Thangadurai and W. Weppner, Angew. Chem., Int. Ed., 2007, 46, 7778–7781.

The authors declare no competing financial interest.

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V. Thangadurai, S. Adams and W. Weppner, Chem. Mater., 2004,16, 2998–3006. 14 V. Thangadurai, H. Kaack and W. Weppner, J. Am. Ceram. Soc., 2003, 86, 437–440. 15 V. Thangadurai, S. Narayanan, D. Pinzaru, Chem. Soc. Rev., 2014, 43, 4714-4727. 16 C. A. Geiger, E. Alekseev, B. Lazic, M. Fisch, T. Armbruster, R. Langner, M. Fechtelkord, N. Kim, T. Pettke and W. Weppner, Inorg. Chem., 2011, 50, 1089–1097. 17 J. Awaka, A. Takashima, K. Kataoka, N. Kijima, Y. Idemoto and J. Akimoto, Chem. Lett., 2011, 40, 60–62. 18 S. Adams, R. Prasada Rao, J. Mater. Chem., 2012, 22, 1426-1434. 19 J. Wolfenstine, E. Rangasamy, J. L. Allen, J. Sakamoto, J. Power Sources, 2012, 208, 193–196. 20 Y. Shimonishi, A. Toda, T. Zhang, A. Hirano, N. Imanishi, O. Yamamoto and Y. Takeda, Solid State Ionics, 2011, 183, 48–53. 21 N. Imanishi, M. Matsui, Y. Takeda, O. Yamamoto, Electrochemistry, 2014, 82, 938-945. 22 L. Truong, M. Howard, O. Clemens, K. S. Knight, P. R. Slater, V. Thangadurai, J. Mater. Chem. A, 2013, 1, 1346913475. 23 C. Galven, J. Dittmer, E. Suard, F. Le Berre, M.-P. Crosnier-Lopez, Chem. Mater. 2012, 24, 3335−3345. 24 W. Gu, R. Prasada Rao, S. Adams, Solid State Ionics, 2015, 274, 100–105. 25 C. Ma, E. Rangasamy, C. Liang, J. Sakamoto, K.L. More, M. Chi, Angew. Chem. Int. Ed., 2015, 54, 129 – 133.

ASSOCIATED CONTENT Supporting Information. TGA-DSC data of Li6.75La3Zr1.75Nb0.25O12 and Li6.75La3Zr1.75Ta0.25O12; Details of Rietveld refinement results in Li6.75La3Zr1.75Nb0.25O12 and Li7La3Zr2O12 determined by in situ neutron powder diffraction. This material is available free of charge via the Internet at http://pubs.acs.org.

ACKNOWLEDGMENT This research was supported by the National Research Foundation, Prime Minister’s Office, Singapore under its Competitive Research Programme (CRP awards No. NRF-CRP 102012-6 and NRF-CRP 8-2011-4).

REFERENCES

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J. S. Thokchom, N. Gupta and K. Kumar, J. Electrochem. Soc., 2008, 155, A915–A920. 2 J. Fu, Solid State Ionics, 1997, 104, 191-194.

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N. Janani, S. Ramakumar, L. Dhivya, C. Deviannapoorani, K. Saranya, R. Murugan, Ionics 2011, 17, 575580. 38 K. Tadanaga, H. Egawa, A. Hayashi, M. Tatsumisago, J. Mosa, M. Aparicio, A. Duran, J. Power Sources, 2015, 273, 844-847. 39 K. Ishiguro, Y. Nakata, M. Matsui, I. Uechi, Y. Takeda, O. Yamamoto, N. Imanishi, J.Electrochem. Soc. 2013, 160, A1690-A1693. 40 K.-D. Liss, B. A. Hunter, M. E. Hagen, T. J. Noakes, S. J. Kennedy, Physica B, 2006, 385-386, 1010-1012. 41 A.C. Larson, R.B. von Dreele, General Structure Analysis System (GSAS); Report LAUR 86-748; Los Alamos National Laboratory: Los Alamos, NM, 2000. 42 B. Toby, J. Appl. Crystallogr. 2001, 34, 210–213. 43 S. Adams, R. Prasada Rao, Phys. Status Solidi A, 2011, 208, 1746-1753. 44 S. Adams, “Practical Considerations in Determining Bond Valence Parameters”, in “Bond Valences” (I. D. Brown, K.R. Poeppelmeier eds.), Springer Berlin Heidelberg, Structure and Bonding 2014, 158, 91-128.(DOI 10.1007/430_2013_96) 45 Y. Wang, M. Klenk, K. Page, W. Lai, Chem. Mater. 2014, 26, 5613−5624. 46 G. Larraz, A. Orera, J. Sanz, I. Sobrados, V. DiezGómez, M. L. Sanjuán, J. Mater. Chem. A, 2015, 3, 56835691.

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Temperature / °C 0

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(c) Figure 1. (a) A two-dimensional intensity contour-map of diffraction patterns during the formation of Li7La3Zr2O12 from a ball-milled partially-amorphous phases during heating at a constant rate of 2.15 °C/min, followed by isothermal annealing at 1050 °C and passive cooling to room temperature. The solid magenta line indicates the temperature during the measurement. (b) Open squares: variation of the unitcell-mass-weighted sum of scale factors for all crystalline phases as derived from the Rietveld refinements. Filled squares: variation of the expected sum of unit-cell-mass-weighted scale factors based on the thermogravimetric analysis of mass loss. The insert shows the thermogravimetry analysis of a parallel sample of the same precursor mixture. (c) Weight fractions of the phases present in the LLZ sample during the heating determined using Rietveld analysis of the data shown in (a) averaged over 20 min taking into consideration the change in scattering mass according to the thermogravimetric results shown in (b).


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(a)

(b) Figure 2. (a) A two-dimensional intensity contour-map of diffraction patterns for the formation of Li6.75La3Zr1,75Nb0.25O12 from a mortar-ground partially-amorphous phase during heating at a constant rate of 2 °C/min followed by isothermal annealing at 1000 °C for 1 h and controlled cooling to 500°C followed by natural cooling till room temperature. The solid magenta line indicates the temperature during the measurement. (b) Weight fractions of the phases in the sample during heating run as determined by Rietveld analysis of the data shown in (a) averaged over 20 min taking into consideration the change in scattering mass.


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Figure 3. Temperature dependence of the Li(1) site occupancy factor in undoped and pentavalentlydoped LLZ as predicted by our previous MD simulations18 (open symbols) and determined experimentally (filled symbols) using room-temperature ex situ neutron powder diffraction and at high temperature using in situ neutron powder diffraction. The site occupancy of the Li(2) site is constrained to be SOF(Li(2)) = (7 - 3*SOF(Li(1)))/12 in undoped LLZ and SOF(Li(2)) = (6.75 – 3*SOF(Li(1))/12 for the pentavalently doped phases. Data for undoped LLZ are diamonds, for Nb-doped LLZ are squares, and for Tadoped LLZ are triangles. Horizontal dashed lines indicate the predicted limits of Li(1) occupancy based on the mutual exclusion of Li at neighbouring tetrahedral and octahedral sites at the chosen Nb (or Ta) dopant concentration, while horizontal dashed-dotted lines show the corresponding limits for the Li(1) occupancy in undoped LLZ. Solid lines represent polynomial fits to the MD simulation results.


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8.0E+02 7.0E+02 6.0E+02

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-Z"

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Figure 4. The Nyquist plots of Li6.75La3Ta0.25Zr1.75O12 for selected temperatures. The solid dots are observed data. The solid line in Figure represents fitted data with an equivalent circuit consisting of (RbQb)(RgbQgb)(Qel) for 50 °C data as an example.

Figure 5. Arrhenius plot of Li6.75La3Nb0.25Zr1.75O12 cooled at 2, 5, 20°C/min or quenched to room temperature after keeping the respective pelletized sample at 950 °C for 18 hours.


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Figure 6. Temperature dependence of the lattice parameter a for cubic Nb and Ta-doped samples, determined from MD simulations and Rietveld refinement using X-ray and neutron diffraction data. Solid lines are polynomial fits. Broken lines indicate linear extrapolations of the high-temperature behaviour (above 500 °C) to lower temperature during the cooling of the Nb-doped LLZ obtained from neutron diffraction and Ta-doped LLZ obtained from MD simulations. All samples and methods yield nearly identical thermal-expansion in the high-temperature range and an according non-linear thermal expansion at lower temperatures that implies local ordering processes.


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Figure 7. Ball-and-stick representation of lithium distribution in local structure models of pentavalently doped LLZ. Sticks mark the possible migration pathways alternating between octahedral and tetrahedral Li sites. Large spheres refer to the tetrahedrally coordinated Li(1) sites and are connected to 4 small spheres representing the octahedrally coordinated Li(2) sites. Light grey colour is used to mark unoccupied sites, while coloured spheres mark occupied sites. Li ions that can perform local hops without the + need to push neighbouring Li ions from their site are marked in green, while tetrahedral (octahedral) + sites that are immobilised by the neighbouring Li are shown in red (blue). In the l.h.s. model, where 8 of the 54 Lithium occupy tetrahedral sites, in total 16 of the lithium are locally mobile, while in the r.h.s. model with 14 Li are on tetrahedral sites the number of locally mobile Li is reduced to 2.

Figure 8. Average pseudocubic lattice parameter a (black squares, lower graph) and potential energy per atom (grey triangles, upper graph) for local structure models with different Li site distributions. For each data point the axis below the graphs indicates the occupancy of the tetrahedral 24d site Li(1), while the axis above indicates the corresponding occupancy of the octahedrally coordinated 96h site. The lattice parameters are derived from the cube root of the time-averaged unit cell volume in short room temperature molecular dynamics simulations of 2×2×2 supercells. Formation energies are derived from geometry optimisations fixing the pressure to 1 atm.


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241x137mm (96 x 96 DPI)


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In Situ Neutron Diffraction Monitoring of Li7La3Zr2O12 Formation

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