The comparative influences of structural ordering, grain size, Li-content, and bulk density on the Li+-conductivity of Li0.29La0.57TiO3
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- Sutorik, A.C., Green, M.D., Cooper, C. et al. J Mater Sci (2012) 47: 6992. doi:10.1007/s10853-012-6650-5
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The lattice and total Li+-ionic conductivity of Li0.29La0.57TiO3 ceramic (LLTO) sintered at 1200 °C were determined as functions of powder calcination temperature and sintering duration, and these results were correlated with the relative degrees of Li+-ordering, Li-content, grain size, and bulk density to assess the relative impact of these parameters on material performance. Under all conditions, LLTO formed with a high degree of tetragonal superstructure to its perovskite related framework, and the lattice conductivity closely followed the relative amounts of the superstructure, as evaluated via determination of the sample ordering parameter from X-ray diffraction data. LLTO powders that were calcined at 900 °C for 1 h and sintered at 1200 °C for 6 h gave lattice conductivity values (~1.14 × 10−3 S cm−1) comparable within the highest ranges reported in the literature. This coincided with the lowest degree of tetragonal superstructure formation, and it was also found to be largely independent of the values of Li-content measured on sintered ceramic despite significant Li2O volatilization at longer sintering times (up to 23 % after 12 h at 1200 °C). Samples of LLTO powder that were calcined at 1100 °C and sintered at 1200 °C for 12 h resulted in the highest total Li-ion conductivity value ~6.30 × 10−5 S cm−1. The total conductivity of LLTO varied inversely with grain size when the grains were <20 μm but was insensitive to that parameter above that size threshold. The strongest influence on total conductivity was primarily the bulk ceramic density. It was estimated from measured values that as the bulk ceramic density approached the full theoretical value for LLTO the total conductivity could near the lattice conductivity of ~1.2 × 10−3 S cm−1.
Aqueous lithium-air batteries offer the potential for higher energy density compared to the current battery technologies [1, 2]. However, many technical challenges remain until their full potential can be realized. One of the most significant issues is that aqueous Li-air battery technology requires a solid electrolytic membrane that possesses fast lithium ion transport, low electronic conductivity, high mechanical strength, and excellent chemical stability . Ceramic electrolytes offer several of the needed characteristics, and lithium lanthanum titanate (LLTO) (stoichiometry Li3xLa(2/3)−xTiO3) is currently one of the materials under consideration for this application.
One reason for the interest in LLTO is that among Li+ conducting ceramics it has some of the highest reported values for room temperature lattice conductivities at ~1 × 10−3 S cm−1 (R.T.) [3, 4]. However, the total conductivity of polycrystalline LLTO is a function of both the lattice and grain boundary contributions. Li-ion conduction through the grain boundaries still limits the total conductivity of LLTO to the typical range of 1 × 10−5–1 × 10−4 S cm−1 [4, 5]. A primary requirement for optimized conductivity of the phase pure ceramic is that the final product needs to be near full density to maximize grain-to-grain contact and, ideally, remove porosity [4, 5].
Previous studies have examined the correlation between LLTO synthesis, sintering, and ionic conductivity. Ceramics of LLTO are most commonly sintered in the temperature ranges of 1100–1350 °C from powder which has been previously prepared from direct combination, solid-state reaction between Li2CO3, La2O3, and TiO3 [5, 6]. There is general agreement in the open literature that the maximum lattice conductivity (σL) of ~1 × 10−3 S cm−1 occurs for Li3xLa(2/3)−xTiO3 with x close to 0.1 . Some groups have reported a range of high conductivity for x = 0.11–0.13 [5, 6, 8] but a maximum value at compositions as high as x = 0.21 has also been reported . In most of the literature, it is also believed that decreased Li-content in LLTO ceramics, resulting from Li2O volatilization on prolonged thermal exposure, leads to a decrease in lattice conductivity [2, 5–8]. Due to possible Li2O volatility at sintering temperatures, precautions are typically taken to retard Li-loss during powder synthesis and sintering. These include over batching of the Li-content in the starting materials, minimizing the temperature of powder reaction (i.e., the calcination temperature), packing of the green ceramic parts in loose LLTO powder to provide a Li2O-vapor-rich local atmosphere around the parts, and minimizing the time and temperature of final sintering [2, 6–9]. Once LLTO is fully sintered, the Li-content appears stable as the material can undergo repeated annealing cycles in the temperature range 600–1350 °C and yet still recover its original conductivity [8, 9].
In addition to the Li-content, another key parameter impacting the lattice conductivity of LLTO is the crystal structure polytype which is stabilized on sintering. These are all variations on a perovskite (ABO3) framework. At its most idealized is the cubic structure (PDF# 46–465; Space Group: Pm-3m) in which Li+- and La3+-cations are randomly distributed among the A sites. This structure can be promoted by values of x close to 0.11, but regardless of its composition, it is a high temperature phase and must be stabilized by quenching from ≥1150 °C [8, 9]. With slow cooling or low temperature annealing, a tetragonal superstructure forms as the Li+- and La3+-cations partially order into alternate A-site layers in the lattice (PDF# 53–109; Space Group: P4/mmm). This is accompanied by a slight shrinking of the ao lattice parameter, which distorts the [TiO6] octahedron positions and forms “bottlenecks” around the A-site cations . Li+-mobility is consequently lowered, as evidenced by both a reduction in ionic conductivity and an increase in its activation energy [8, 10]. The extent of tetragonal superstructure ordering can vary with processing conditions, and so to maximize Li+-conductivity procedures which retard its formation are favored. Attempts to increase the stability of the cubic phase and, concomitantly, the Li+-conductivity, through doping of different sized cations have not proven successful . The kinetic equilibrium between the cubic and tetragonal superstructures is actually one of the simplest models for LLTO structure. Not only do other polymorphs form for other values of x , but also the existence of two different orthorhombic superstructures has been reported using either synchrotron X-ray diffraction  or neutron diffraction . Because of the rapid diffusion of cations through a network of A-site vacancies, a precise structural model of LLTO is difficult to attain, and it is clearly high-influenced by synthetic conditions. Nevertheless, a general distinction between an A-site cation-disordered structure (i.e., more cubic character) and an A-site cation ordered structure (i.e., more tetragonal) can readily be distinguished by X-ray diffraction and so can serve as a valuable guiding principle toward the synthesis of LLTO with optimized Li+-lattice conductivity.
While lattice conductivity is tied to the atomic scale properties of composition and crystal structure, representing, in effect, a sample’s maximum Li+-conductivity, the total conductivity is how the sample actually performs macroscopically as this value includes resistance from the grain boundaries. This is tied to characteristic of the ceramic part such as residual porosity and grain size . To maximize total conductivity, near full density should be attained for the ceramic but care should be made that the determination of sample density with the Archimedes method include evaluation of the sample open porosity, without which a ceramic’s apparent density (i.e., the density of the contiguously connected ceramic) may be confused with the bulk density (the density of the ceramic piece which includes both open and closed porosity).
The synthesis and sintering of LLTO ceramics thus present an interesting challenge because processing conditions which promote one aspect of the materials conductivity can simultaneously degrade conductivity by a different mechanism. For example, sintering temperatures >1150 °C would work to achieve high density and the more conductive cation disordered, cubic superstructure, but such temperatures also could cause Li+-loss thru Li2O volatilization, which can lower ionic conductivity. The interpretation of experimental results in the open literature is further complicated when incomplete characterization is presented, thereby introducing doubt as to the actual cause of variation in both lattice and total conductivity. To elucidate some of these issues, experiments have been performed in which ceramics of LLTO have been prepared and extensively characterized for ionic conductivity (lattice and total), electronic conductivity, phase content, bulk and apparent density, Li-content remaining after sintering, and grain size. By this approach the relative impacts of the competing factor which could influence Li+-conductivity can be compared and prioritized so as to guide the consistent optimization of the material’s performance. Since the experimental range of such an undertaking would be considerable, the current report is limited in scope to LLTO with x = 0.097 (that is, Li0.29La0.57TiO3) and to sintering at 1200 °C for 1, 6, and 12 h. Future reports will detail studies on other compositions and processing conditions with the eventual goal of preparing LLTO with optimized and consistent properties suitable for evaluation in Li-air battery test cells.
The following powders were used in the synthesis of LLTO: Li2CO3 (Alfa Aesar, 99.0 %), La2O3 (99.99 %; Alpha Aesar), and TiO2 (99.8 % Alfa Aesar). Materials were used as received. The amount of moisture and other volatile impurities for each powder was determined by a loss on ignition (LOI) procedure in which 1–2 g of sample were loaded into alumina crucibles, and sample masses were determined before and after firing under a ramp rate of 5 °C min−1–900 °C with a 1 h hold. These correction factors were then used in the stoichiometric calculations.
Synthesis and sintering
Starting powders of Li2CO3, La2O3, and TiO2, were mixed via magnetic stirring in ethanol in amounts necessary to form Li0.29La0.57TiO3. In a typical procedure for a target yield of 100 g of LLTO, 6.051 g Li2CO3, 52.439 g La2O3, and 45.112 g TiO2 were added to stirring ethanol. No other dispersants or additives were used. The slurry was stirred for approximately 1 h. After mixing, the slurry was poured into a Labrota 4000 rotary evaporator where the bulk of the ethanol was volatilized under 23 kPa of pressure and 60 rpm rotation in a 60 °C oil bath. Once the ethanol was visibly removed, the remaining powder was coarsely broken and transferred to a clean, loosely covered, watch-glass. The powder was left to stand, loosely covered for a period of 1–2 days to allow for complete solvent volatilization. After drying, the powder was calcined in a high-alumina crucible under ambient atmosphere in a closed box furnace. Samples were heated with a ramp rate of 5 °C min−1 to either 900 or 1100 °C (hereafter denoted C900 or C1100), held isothermally for 1 h, and cooled back to room temperature at 5 °C min−1.
Powders were pressed into pellets in a 13 mm diameter stainless steel die with a uni-axial hydraulic press. Each sample was pressed under 68 MPa for approximately 10 s. Samples were also subjected to further consolidation in the green state using a cold isostatic press under 207 MPa pressure. Duplicate samples for each calcination condition were sintered in ambient atmosphere in a closed box furnace to 1200 °C at 5 °C min−1. Isothermal hold durations were varied from 1, 6, and 12 h, and all samples were cooled at a rate of 5 °C min−1.
Dynamic light scattering
Particle size distributions for synthesized powders were measured using a Horiba LA-910 Light Scattering Particle Size Distribution Analyzer. Samples were prepared by pre-dispersion of 0.1 g powder into 10 g of deionized water with the container immersed in an ultrasonic bath (Li2CO3 was dispersed in ethanol). Samples were added drop wise to a blanked sample holder containing the necessary medium under magnetic stirring until instrument response fell into the optimum measurement range. Values of D50 and D90 are reported on a volume distribution basis.
Surface area analysis
The surface area of starting and synthesized powders was measured using a computer controlled Micromeritics ASAP 2010 using an 11-point BET analysis protocol. One gram of sample was loaded into a dry sample tube and stored open in a dynamic vacuum furnace at 75 °C overnight to insure physisorbed species were removed from the samples surface. Nitrogen gas was used for the adsorbing agent.
Powder X-ray diffraction (XRD)
Powder X-ray diffraction analysis was employed to determine the crystalline phases in the starting powders, calcined LLTO, and sintered LLTO ceramics. Powders of sintered pellets were prepared by grinding fragments in a high-alumina mortar and pestle. Measurements were taken using a computer controlled Rigaku minflex operating at 30 kV and 15 mA with a step size of 0.03° 2θ and scan rate of 2° 2θ/min, scanning over 5–60°. Powder samples were loaded directly into low-background Si sample holders.
Field emission scanning electron microscopy (FESEM)
A Hitachi 4700 Scanning Electron Microscope was used to observe powder morphology of starting powders, synthesized LLTO powders, and sintered ceramic grain structure. An accelerating voltage of 2 kV and beam current of ~10 mA were used throughout this procedure to minimize charging. Samples were applied to double sided carbon tape and attached to an aluminum sample holder.
Grain size measurements
Inductively coupled plasma mass spectroscopy measurements (ICP-MS)
ICP-MS was conducted to determine the Li-, Ti-, and La-content of the uncalcined, calcined, and sintered LLTO. Analysis was performed by NSL Analytical (Cleveland, OH).
Samples were prepared for electrochemical characterization by applying a gold coating of ~50 nm to the electrode contact surfaces with a Denton Desk IV vacuum cold-sputtering/etch unit. Dimensions of the samples were measured using a Vernier caliper. Averages of five diameter and five width measurements from opposing points around the pellets were used in subsequent calculations.
A Solatron SI 1287 Electrochemical Interface was used to measure the AC impedance of the LLTO samples. All samples were measured under a 100 mV AC potential, from an initial frequency of 1.0 × 106 Hz to a final frequency of 1.0 Hz using a logarithmic scale in a Teflon sample holder. All measurements were conducted in a dry-room with negligible humidity at room temperature. Electronic resistance of each of the samples was measured using a Keithly 6517A High Resistance Electrometer with 8009 Resistivity Test Fixture. All samples were left overnight in the test fixture and final electronic resistivity was reported from the equilibrated value.
Results and discussion
Starting powder surface areas and particle diameters
Surface area (m2 g−1)
28.66 ± 0.02
50.47 ± 0.02
1.03 ± 0.04
14.07 ± 1.27
37.88 ± 10.18
3.95 ± 1.50
1.13 ± 1.59
2.37 ± 2.61
4.63 ± 0.80
The starting powders were also investigated with powder X-ray diffraction. No crystalline impurities were found in Li2CO3. In addition to being free from impurities, the TiO2 proved to be 100 % in the rutile structure type; the relative amounts of anatase and rutile structure types in TiO2 starting powders is not typically reported in papers on LLTO synthesis, so the potential impact on reactivity and ceramic properties is not known. The La2O3 starting material proved to be a mixture of the oxide and La(OH)3; LOI analysis confirmed the presence of 0.7 wt % volatile species, on average. Note that if the hydroxide content is not taken into account, the intended composition of LLTO would be in error.
Formulated powder surface areas and particle diameters for uncalcined (UC), C900, and C1100 powders
Surface area (m2 g−1)
4.35 ± 0.20
9.33 ± 0.32
5.31 ± 0.52
10.52 ± 0.25
26.62 ± 8.05
1.10 ± 0.87
10.25 ± 0.86
22.53 ± 4.64
1.47 ± 0.25
Sintered ceramic characterization
The density of LLTO ceramics relied more on the temperature of powder calcination than the duration of sintering at 1200 °C. The C900 powder sintered to a high apparent density of 4.856 ± 0.0168 g cm−3 after 1 h at 1200 °C. This sample, however, has a much lower bulk density of only 4.369 ± 0.571 g cm−3, indicating that significant open porosity remains in the sample. Note that if the bulk density had not been determined, then the Archimedes density measurement, using only the ceramic dry mass, would have erroneously indicated a much higher level of ceramic density. At a 6 h isotherm, the apparent and bulk density of C900 converge to similar values, which indicates that the ceramic is approaching a closed porosity stage. Sintering for 12 h brings the two density values even closer. The situation is very different for the C1100 samples which exhibit near identical values for the bulk and apparent densities for all sintering times. Closed porosity is achieved very readily with this formulation, and the remaining residual porosity is thus trapped within the ceramic. As noted in Table 2, the size distribution characteristics of C900 and C1100 are very similar, whereas the XRD analysis indicates that C1100 has a greater degree of conversion to LLTO. This modest difference may be enough to allow for a more active and uniform densification during sintering.
Average ordering parameters S, for sintered ceramics with varied calcination temperature and sintering times
0.49 ± 0.04
0.43 ± 0.06
0.57 ± 0.04
0.44 ± 0.04
0.53 ± 0.03
0.53 ± 0.03
AC impedance measurements
Contributions of grain boundary and lattice resistance to total resistance as a function of calcining temperature and sintering isothermal hold time
Lattice ionic conductivity
The σL data in Fig. 5 does not appear to trend with the decreasing Li-content as opposed to the apparent inverse relationship between σL and S noted earlier. This implies that once the structure reaches a thermodynamically stable endpoint, as in the case of C1100, changes to the Li+-content have reduced impact on σL. In other words, crystal structure appears more important than Li+-content to σL. Because of its calcination at a lower temperature, C900 has not advanced as far to the thermodynamic endpoint, and so its sintering at 1200 °C for 6 h represents a kinetically stabilized outcome in which the greater structural disorder (i.e., lower degree of tetragonal superstructure ordering) combines with the Li+-content (which is the highest of the sintered ceramics) to result in the highest σL among the present samples. The outcome, however, is metastable in the case of C900, and longer sintering times drive the ceramic to a more thermodynamically stable outcome and closer to the properties of C1100 for the same sintering conditions. This is a logical scenario in that the earlier stages of a solid-state reaction are more likely to exhibit residual levels of structural disorder before coming to full thermodynamic equilibrium. In the case of LLTO, disorder represented by randomization of Li+ and La3+ in the A-sites promotes higher Li+-conductivity. Hence, what is reported as a high temperature structural model can exist at low temperatures under kinetic conditions.
Figure 8 highlights that there is only a small change in σT as a result of increased isothermal hold time. For the C1100 samples, no changes in σT were observed by increasing the sintering duration from 6 to 12 h (a plateau in conductivity is seen within measurement error). Reasons for the differences between the C900 and C1100 samples’ σT may be a result of grain size and/or density [4, 13], and so both have been examined.
Sintered LLTO average grain size (μm) as a function of the calcination temperature of the starting powder and the sintering hold time
6.87 ± 0.22
6.75 ± 1.26
45.28 ± 15.69
12.23 ± 1.21
19.53 ± 0.75
25.73 ± 2.25
Ceramic powders of LLTO with compositions of Li0.29La0.57TiO3 were prepared under varied calcination temperatures and sintering durations to assess the relative impacts of phase content, Li+-content, bulk density, and grain size on lattice and total conductivity. Samples of LLTO powder that were calcined at 900 °C for 1 h and sintered at 1200 °C for 6 h exhibited σL of 1.14 × 10−3 ± 0.11 × 10−3 S cm−1, a value comparable to the highest reported in the literature. This also coincided with the lowest amount of tetragonal superstructure ordering as represented by the sample ordering parameter, S, and the inverse relationship between S and σL is observed, consistent with the work of Harada et al. Structural disorder appears to exert a greater influence on σL than Li+-loss (which was significant after 12 h of sintering at 1200 °C), and the more ordered tetragonal superstructure appears to have preferred thermodynamic stability.
It was also observed that polycrystalline LLTO σT is more dependent on bulk sample density than on average grain sizes. LLTO samples with grain sizes >20 μm were found to influence σT to a significantly lesser degree than samples with grains smaller than 20 μm. Moreover, it was determined that σT is primarily a function of the samples’ bulk densities. At densities approaching the theoretical value for LLTO ceramics, σT’s dependence on density rapidly increased; it was estimated that a value of 1.2 × 10−3 S cm−1 for σT can be possible if density is optimized. If both A-site cation ordering (or the lack thereof) and density can be optimized for the LLTO material, rapid improvements can be made in its overall σT.
The authors of this paper would like to acknowledge and thank the Oak Ridge Institute for Science and Education (ORISE) and the United States Army Research Laboratory (ARL), for their unwavering support in this research effort.