Toward Understanding the Lithium Transport Mechanism in Garnet-type Solid Electrolytes: Li+ Ion Exchanges and Their Mobility atOctahedral/Tetrahedral SitesDawei Wang,† Guiming Zhong,† Wei Kong Pang,‡,§ Zaiping Guo,§ Yixiao Li,† Matthew J. McDonald,†

Riqiang Fu,*,∥ Jin-Xiao Mi,⊥ and Yong Yang*,†

†Collaborative Innovation Center of Chemistry for Energy Materials, State Key Lab of Physical Chemistry of Solid Surface andDepartment of Chemistry, College of Chemistry and Chemical Engineering and ⊥Department of Material Science and Engineering,Xiamen University, Xiamen 361005, China‡The Bragg Institute, Australian Nuclear Science and Technology Organisation, Locked Bag 2001, Kirrawee DC, New South Wales2232, Australia§Institute for Superconducting and Electronic Materials, University of Wollongong, Wollongong, New South Wales 2522, Australia∥National High Magnetic Field Laboratory, 1800 E. Paul Dirac Drive, Tallahassee, Florida 32310, United States

*S Supporting Information

ABSTRACT: The cubic garnet-type solid electrolyte Li7La3Zr2O12 with aliovalentdoping exhibits a high ionic conductivity, reaching up to ∼10−3 S/cm at roomtemperature. Fully understanding the Li+ transport mechanism including Li+ mobilityat different sites is a key topic in this field, and Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1)are selected as target electrolytes. X-ray and neutron diffraction as well as ac impedanceresults show that a low amount of aliovalent substitution of Zr with W does notobviously affect the crystal structure and the activation energy of Li+ ion jumping, but itdoes noticeably vary the distribution of Li+ ions, electrostatic attraction/repulsion, andcrystal defects, which increase the lithium jump rate and the creation energy of mobileLi+ ions. For the first time, high-resolution NMR results show evidence that the 24d,96h, and 48g sites can be well-resolved. In addition, ionic exchange between the 24dand 96h sites is clearly observed, demonstrating a lithium transport route of 24d−96h−48g−96h−24d. The lithium mobility at the 24d sites is found to dominate the totalionic conductivity of the samples, with diffusion coefficients of 10−9 m2 s−1 and 10−12

m2 s−1 at the octahedral and tetrahedral sites, respectively.

1. INTRODUCTION

Solid-state batteries promise to lead the next generation ofrechargeable, high energy density lithium batteries, because oftheir great advantages in safety and working voltage rangecompared to conventional liquid electrolyte lithium ionbatteries, in which the intrinsic properties of liquid electrolytesplague their performance.1,2 Solid electrolyte is an importantconstituent of all solid-state batteries and determines most oftheir performance.3 As a result, it has attracted great interest inrecent years.4,5 Of various solid electrolytes, the garnet-typesolid electrolyte Li7La3Zr2O12 shows great prospects withregard to the development of highly safe lithium ion batteries,because of its high ionic conductivity at room temperature, inparallel with its good chemical and electrical stability.6

Li7La3Zr2O12 with a cubic structure was first reported byMurugan et al. as a solid electrolyte with an ionic conductivityas high as 2.4 × 10−4 S cm−1 at 22 °C, achieved only in aluminacrucibles.7 It was found that aliovalent doping at 24d (Li1) and16a (Zr) sites could dramatically enhance the stability and ionic

conductivity of the cubic phase, using elements such as Al3+,Ga3+, Nb5+, Ta5+, Te6+, W6+, etc.8−11

Therefore, an important question regarding this material ishow such doping affects the ionic conductivity. To address thisissue, we need to understand the states of the lithium ions atrespective sites and how lithium migrates through the crystalstructure. Using density functional theory, Xu et al. concludedthat the Li+ ions jumped from one tetrahedral site to aneighboring octahedral site and then leaped forward to anothertetrahedral site in the Li7La3Zr2O12 structure.12 Somepreliminary studies have also been carried out to investigateLi migration and jumping via neutron powder diffraction(NPD) in Li7La3Zr2O12, yielding a similar conclusion.13,14

However, Thompson et al. drew an opposite conclusion withNPD results and claimed that the ionic conductivity ofLi7La3Zr2O12 trended with the occupancy of octahedral sites;

Received: June 24, 2015Revised: September 18, 2015Published: September 18, 2015

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they deduced that the lithium jumps from one octahedral siteto a neighboring octahedral site, bypassing tetrahedral sitesentirely.15

Explanations of the enhanced ionic conductivity of thegarnet-type electrolytes are usually correlated with the lithiumstoichiometry in the compounds, although detailed transportmechanisms are seldom investigated, in particular the sittingsites of Li+ ions as well as their mobility and rate-determiningsteps.6,12,15 Previously, we reported on the synergistic effects ofAl and Te using Al2O3 and ZrO2 crucibles, demonstrating thatthe concentration of Li sites and ion dynamics can bemodulated by aliovalent ions and accurately monitored byusing NPD and solid-state NMR techniques.16 In this work, weelaborate further the effects of binary aliovalent substitution onenhancing the ionic conductivity and bolstering the lithiumtransport mechanism in Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1)with a cubic garnet structure. Furthermore, we illustrate therelationship between composition, crystal structure, lithiumdynamics, and ionic conductivity in the garnet-type electrolytes.

2. EXPERIMENTAL SECTIONA conventional solid-state reaction was used to synthesize the garnet-type electrolyte. The starting materials, LiOH·H2O, La2O3, ZrO2, andWO3 (purities of 95.0%, 99.99%, 99.0%, and 99.99%, respectively, allfrom the Sinopharm Chemical Reagent Co., Ltd.), were mixed andground by a planetary grinder, with isopropanol used as the grindingreagent. The precursors were ground for 12 h and dried in a 120 °Cbox, then calcined at 800 °C for 12 h. After the second grinding anddrying, the material was pressed into pellets at 2 MPa by uniaxial coldpressing. Finally, the pellets with W, covered with green constituentpowder, were sintered at 1150, 1175, and 1200 °C for 12 h in Al2O3crucibles. To get a pure phase with cubic symmetry, the pellet with thecomposition of Li7La3Zr2O12 (i.e., without W doping) was sintered atdifferent temperatures for 36 h in an alumina crucible.The crystal structure of the as-prepared garnet electrolyte powder

was characterized using X-ray powder diffraction. A Panalytical X’Pert(Philip, Netherlands) instrument with Cu Kα radiation was used forthis purpose.High-resolution NPD was also used to study the material, because

of its high sensitivity for “light” atoms. The powders were stored andloaded into 6 mm vanadium cans in an Ar-filled glovebox, which weresealed properly in an airtight condition. The NPD data was collectedusing ECHIDNA, the high-resolution neutron powder diffractometerat the Australian Nuclear Science and Technology Organisation(ANSTO).17 The neutron beam wavelength was 1.62158(8) Å,determined using the La11B6 NIST standard reference material 660b.The NPD data was obtained in the 2θ angular range of 4 to 164° witha step size of 0.125°. GSAS-II was employed to perform Rietveldanalysis of the obtained NPD data.18 The refining parameters thatwere optimized included background coefficients; zero-shift; peakshape parameters; lattice parameters; the positional parameters of Li(48g), Li (96h), and O (96h); the fractional factors of all Li, Zr, W,and O; and isotropic atomic displacement parameters (Uiso). Duringthe refinement, the total occupancy of Zr and W, sharing 16a sites, wasconstrained to be unity, and the Uiso of all elements was set to beequivalent.The density of sintered pellets was measured using the Archimedes

method. Theoretical density was calculated from the lattice parametersdetermined by XRD. A Hitachi S-4800 scanning electron microscopewas used to check the cross-sectional morphology of samples.Measurements of ac impedance were undertaken to measure

conductivity, using a Solatron 1260 impedance analyzer with thefrequency range of 1 to 106 Hz and an amplitude of 100 mV. Themeasurement cell was constructed as follows: first, silver slurry wasspread onto the surfaces of pellets and cured at 600 °C for 30 min inorder to vaporize any organic solvent. Next, the silver-coated pelletswere placed between two pieces of steel for testing. Test temperatures

were varied from −50 to 150 °C at 20 °C intervals to study thetemperature-dependent relationship. The electronic conductivity wasmeasured by the dc polarization method, using a Solatron 1287electrochemical analyzer with a polarization potential of 0.1 V. Theholding time was 10 000 s.

The magic-angle spinning nuclear magnetic resonance (MASNMR) technique was used to probe the local chemical environmentof the atoms and to analyze the lithium dynamics of the materials. The6Li and 27Al MAS NMR spectra were gathered at Larmor frequenciesof 58.9 and 104.3 MHz, respectively, on a Bruker Avance 400 NMRspectrometer with a sample spinning speed of 12 kHz. The chemicalshifts of the 6Li and 27Al were calibrated by using LiCl powder (0ppm) and Al(OH)3 (0 ppm), respectively. 27Al MAS NMR spectrawere collected using a pulse length of 1 μs and a recycle delay of 0.1 s.6Li MAS NMR spectra were collected using a 1/4 π pulse and a recycledelay of 40 s. The saturation recovery method was used here tocharacterize the spin−lattice relaxation times for different lithium sites.To reduce the Li2 signal in 6Li MAS NMR spectra, a π pulse was firstapplied to invert the magnetization from the +z to −z axis; then after arecovery time of 2 s at which the strong peak (Li2) was zero-crossingwhile the other signals still remained in the −z axis, due to theirdifference in T1, a 1/2 π pulse was applied to flip those signals alongthe −z axis into the xy plane for observation, with a recycle delay of110 s. The 6Li−6Li exchange spectra were recorded after suppressionof the strong peak (Li2), with a mixing time of 10 s. The diffusioncoefficient was detected by a static 8 mm 7Li probe. 7Li PFG NMRspectra were collected at the Larmor frequency of 158 MHz using a 1/2 π pulse and a recycle delay of 2 s. A series of pulse field gradientsbetween 50−400 G/cm was applied, with a duration time of 1.42 s anda diffusion time of 10 μs.

3. RESULTS AND DISCUSSION

3.1. Crystal Structure and Morphologies. Figure 1shows the XRD patterns of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x≤ 1) sintered at different temperatures. The sample with thecomposition of Li7−3yAlyLa3Zr2O12 (y = 0.26), sintered at 1150°C for 36 h, shows a pure cubic structure.16 The samples withthe compositions of x = 0.25 and 0.5 also show pure cubicstructures. It was very difficult to synthesize a pure cubic phasewhen the W doping reached x = 0.75. In most circumstances, asmall amount of W salts can be detected in the samples (withpeaks visible at 18°). Higher sintering temperatures reduced theamount of the impurity phase, suggesting that high sinteringtemperatures favored the formation of solid solutions. Finally,at the composition of x = 1, more impurities were detected inthe samples. As the substitution content increases, the latticeparameters decrease monotonically, as seen in Figure S1. Inoctahedral coordination, the ionic radius of Zr4+ is 0.72 Å andthe ionic radius of W6+ is 0.60 Å.19 Assuming no phasetransition, the monotonic decrease in lattice parameters withincreasing W content indicates the successful substitutiono f Z r 4 + b y W 6 + a n d t h e f o rm a t i o n o f t h eLi7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) solid solutions. Thesamples with the same composition showed a decrease in latticeparameters when the sintering temperature increases, indicatingthat more Li+ vacancies and O defects were formed in thegarnet structure at higher sintering temperatures.As alumina crucibles were used during the sintering process,

it was necessary to know the content of Al and its localenvironment in the samples. Our NMR and EDS results showthat a fraction of Al occupies the 24d sites, and the otherrandomly exists along grain boundaries in the form of LaAlO3in some batches, with an Al content of 0.2 mol for both x =0.25 and x = 0.5. When the W content reaches x = 0.75, there isno Al detectable in the samples. A similar “synergistic effect” of

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Al and Te could be found in our previous study.16 Detailedinformation is listed in Supporting Information, Figure S2 andTable S1.The model proposed by Goodenough and co-workers was

used for the Rietveld refinement of the NPD data for theLi7−2x−3yAlyLa3Zr2−xWxO12 (0.25 ≤ x ≤ 0.75) samples.20 The W

atoms were considered to occupy the Zr (16a) sites, due to theclose values of their ionic radii. However, the presence ofrelatively heavy La and Zr atoms in the structure complicatedthe refinement of the “light” aluminum and lithium atoms.Furthermore, the neutron scattering length for Al is rather low,and thus it was difficult to quantify a small concentration in thecompound. With the NMR and EDS results, the Al occupancyis 0.0667 and located at the 24d sites for both x = 0.25 and 0.5,while the Al occupancy is 0 for x = 0.75. The best fits of theNPD data were obtained with the parameters listed in Table 1,and the results of Rietveld refinement are shown in Figure 2and Figure S3.

The lithium number is slightly higher than expected from thestoichiometry of the formula unit but still smaller than thetolerance of 7.5.21 The occupancy of lithium in the 24d sitesranges from 32 to 41%, and the smallest value, with x = 0.25,might have been due to the occupancy of Al at these sites. TheLi occupancy in the 48g sites ranges from 48% to 35%, whichdecreases along with the total lithium content. Finally, the Lioccupancy in the 96h sites remains constant. The partialoccupancy of lithium at different sites might have been thereason for the observed high ionic conductivity of theLi7−2x−3yAlyLa3Zr2−xWxO12 (0.25 ≤ x ≤ 0.75) system. Inaddition, oxygen defects were also detected from the Rietveldanalysis.F i g u r e 3 s h ow s t h e mo r p h o l o g i e s o f t h e

Li7−2x−3yAlyLa3Zr2−xWxO12 samples sintered at different tem-

Figure 1. XRD patterns of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1)sintered at different temperatures: (a) x = 0, (b) x = 0.25, (c) x = 0.5,(d) x = 0.75, and (e) x = 1.

Table 1. Refined Structural Parameters of Li7−2x−3yAlyLa3Zr2−xWxO12 (Space Group Ia-3d, No. 230)

stoichiometry formula Li5.9Al0.2La3Zr1.75W0.25O12 Li5.4Al0.2La3Zr1.5W0.5O12 Li5.5La3Zr1.25W0.75O12

refinement results Li6.36(12)Al0.2La3Zr1.62(4)W0.38(4)O11.30(8) Li6.15(12)Al0.2La3Zr1.46(6)W0.54(6)O11.31(8) Li5.76(16)La3Zr1.12(6)W0.88(6)O11.20(9)

Rwp 6.59% 6.70% 6.60%lattice parameter a 12.954(1) 12.924(1) 12.900(1)Li1 (24d) 0.32(3) 0.41(3) 0.39(3)Li2 (48g) 0.48(2) 0.4(2) 0.35(2)Li3 (96h) 0.21(1) 0.21(1) 0.21(1)Al 0.067a 0.067a 0La 1 1 1Zr (R/S) 0.81(2)/0.875 0.73(3)/0.75 0.56(3)/0.625W(R/S) 0.19(2)/0.125 0.27(3)/0.25 0.44(3)/0.375O 0.94(7) 0.942(7) 0.933(8)

aFixed to the same value as from the EDS results. R = refinement results; S = stoichiometry formula.

Figure 2. Rietveld refinement of the structural models, based on NPDdata collected at room temperature from Li7−2x−3yAlyLa3Zr2−xWxO12samples with x = 0.25 and sintered at 1150 °C.

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peratures. For x = 0, the grains can be clearly seen in the crosssections and become larger as the sintering temperatureincreases. For x = 0.25, the grains contact well with eachother, such that it is difficult to distinguish the grain boundariesin the cross sections. For x = 0.5, the grain boundaries can beclearly distinguished. With an increase in sintering temperature,the grains become larger, but the size of the grain is distributedinhomogeneously. For x = 0.75, at the sintering temperature of1150 °C, the large grains are surrounded by many smallparticles. An increase in sintering temperature causes theselarge grains to increase in size and the small particles to vanish.By combining the XRD patterns with the SEM results, a likelyexplanation for this behavior is that the small particles wereregions of the W-salt impurities. As the sintering temperatureincreased, the W-salts reacted with the samples such that theimpurity content started to decrease. However, for x = 1, theimpurities became overwhelming in quantity and thus severelyinterfered with the crystallization of the garnet phase. At thiscomposition, in Figure 3, no obvious grains can be seen in thecross sections. Overall, from the morphologies of the samples, asmall amount of W substitution significantly increased thesintering properties of the material. Excessive W substitution,on the other hand, gave rise to poor contact between grains. Onthe basis of the Li content in the samples and theirmorphologies, the composition of x = 0.25 appears to haveexhibited the best ionic conductivity.21

Figure 3 also describes the elemental distribution at x = 0.25.The elements O, Al, La, Zr, and W are distributedhomogeneously in the cross section, indicating the formationof solid solutions. The relative density is a key factor whenevaluating the sintering properties of ceramics. Here, therelative densities of the Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1)

samples are between 93% and 95%, which are slightly higherthan that of the sample without W substitution (Table S3). Thesubstitution of W is an effective way to decrease the sinteringrequirements and increase the sintering properties.

3.2. Electrical Properties of Li7−2x−3yAlyLa3Zr2−xWxO12(0 ≤ x ≤ 1) Solid Solutions. Measurements of ac impedancewere used to calculate the impedance and conductivity of thelithium ion conductors. Figure 4 shows the Nyquist plots ofLi7−2x−3yAlyLa3Zr2−xWxO12 (x = 0.25). In the case of testingtemperature below 25 °C, there exists a semicircle and a curveat the high- and low-frequency sides. With an increase of testingtemperature, the semicircle at the high-frequency side graduallydecreases and eventually vanishes after 50 °C. Thisphenomenon was previously explained in the Te system.16

Two long conductive lines were used during the impedancetesting, and the inductance was considered in the fittingprocess. When the testing temperature was below 25 °C,L(R1CPE1)(R2CPE2) as an equivalent circuit was used to fitthe Nyquist plots, whereas for a testing temperature higher than50 °C, LR1(R2CPE2) was applied, because of the vanishing ofthe semicircle at the high-frequency side. In these equivalentcircuits, L, R1, and CPE1 represent the conductance, the totalresistance of the sample, and the constant phase element of thesample, respectively, while R2 and CPE2 refer to the electrodeprocesses. The Nyquist plots were well-fitted, and the resultsare listed in Table S2.Other compositions of Li7−2x−3yAlyLa3Zr2−xWxO12 (x = 0.5,

0.75, and 1) showed similar behavior (Figure S4) and wereanalyzed in the same way. Of particular relevance is the fact thatin the Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) system, thecapacitance value of CPE1 was on the order of 10−11 to 10−12 F,which is the characteristic value of the bulk capacitance. This

Figure 3. Morphologies of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1): (a) x = 0, with sintering temperatures of 1000, 1100, and 1230 °C from top tobottom; (b) x = 0.25; (c) x = 0.5; (d) x = 0.75; and (e) x = 1. Sintering temperatures were 1150, 1175, and 1200 °C, from top to bottom. The EDSmappings show the elemental distribution in the sample with composition of x = 0.25 sintered at 1150 °C for 12 h: red, O; dark green, Al; blue, La;yellow−green, Zr; purple, W.

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indicates that the total resistance R1 was dominated by the bulkresistance and the grain boundaries played a minor or negligiblerole in determining the total resistance. Thus, the bulkconductivity is believed to exhibit behavior similar to thetotal conductivity. No information regarding the grainboundary resistance could be gathered from the Nyquistplots, although the morphologies of different compositionsvaried. This suggests that the grain boundary properties couldnot be judged solely from the apparent morphologies.Figure 5 shows the Arrhenius plots of the samples. The

e l e c t r i c a l c o n d u c t i v i t i e s o f t h e W - d o p e dLi7−2x−3yAlyLa3Zr2−xWxO12 (0.25 ≤ x ≤ 1) are much higherthan that of the pristine Li7−3yAlyLa3Zr2O12 (y = 0.26),

indicating that the substitution of W greatly improves theconductivity of the garnet phase. When x = 0.25, the sampleshows the highest ionic conductivity, 4.9 × 10−4 S cm−1 at 25°C, and lowest activation energy, 34.1 kJ/mol, although theionic conductivity is slightly lower than that reported in theliterature.11 With an increase in substitution content, theconductivity decreases while the activation energy increases, atrend that also occurred previously in the Te-doped system.16 Itwas found that the Te played a more prominent role than Aldid in determining the ionic conductivity of Al and Te codopedsystem in our previous study; here, the change in ionicconductivity and activation was believed to be determined bythe doping of W rather than Al.16 Meanwhile, at 25 °C, theArrhenius plots show an inflection point, where the activationenergy at the low-temperature side is higher than that at thehigh-temperature side, which might be due to the existence ofan “ion-trapping effect” at low temperatures. Table S3 detailsthe conductivities and activation energies of theLi7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) solid solutions.Activation energies listed in Table S3 are from the high-temperature side, with the electrical properties at lowtemperatures being discussed in the next section. In addition,the electrical conductivities and activation energies of thegarnet phases with the same composition appear to not beaffected by the sintering temperature.As a solid electrolyte used in lithium ion batteries,

Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) is required to be apurely lithium ionic conductor with a negligible electronicconductivity. The dc polarization method was applied tomeasure the electronic conductivities of the samples, asdocumented in Supporting Information, Figure S5. Theelectronic conductivities of samples are on the order of 10−8

to 10−9 S cm−1, and the transference number of lithium isnearly unity. This demonstrates that Li7−2x−3yAlyLa3Zr2−xWxO12(0 ≤ x ≤ 1) were pure lithium ionic conductors and thereforesuitable for lithium ion batteries.

3.3. Local Environment and Diffusion of Li inLi7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) Solid Solutions. Inthe cubic garnet structure, there are 3 lithium sites, namely, thetetrahedral 24d sites and the octahedral 48g and 96h sites. The96h sites are evolved from the distortion of 48g sites. Becausethe NMR technique is sensitive to the local environment, thecoordination and local environment of lithium can be deducedby detecting the chemical shift of lithium. NMR measurementscan also provide information on diffusion through the influenceof atomic movement on the width of nuclear resonance linesand on relaxation times. In addition, an external magnetic fieldgradient can be used for a direct determination of diffusioncoefficients.22

Here, the 6Li MAS NMR technique was used to check theLi+ environment and dynamics in Li7−2x−3yAlyLa3Zr2−xWxO12 (0≤ x ≤ 1). Figure 6 shows the 6Li MAS NMR spectra ofLi7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1). Two partiallyoverlapped peaks are observed in the spectra, which areattributable to the lithium ions occupying the tetrahedral (0.8ppm) and octahedral (1.9 ppm) sites.23

The saturation recovery method was used to characterize thespin−lattice relaxation time (T1). Figure S6 shows thesaturation recovery of the 6Li signals. A biexponential functionfitted the data very well, as shown in the red curves, yieldingtwo different T1 values, as listed in Table 2. Apparently, thelarger component with the shorter T1 value (between 1.0−5.0s) belongs to the lithium in the octahedral sites, while the

Figure 4. Nyquist plots of Li7−2x−3yAlyLa3Zr2−xWxO12 (x = 0.25)sintered at 1150 °C for 12 h: (a) collected at low temperature and (b)collected at high temperature.

Figure 5. Arrhenius plots of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1)sintered at 1150 °C for 12 h.

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smaller component with the longer T1 value (in the range of14−28 s) represents the lithium in the tetrahedral sites. It wasalso indicated that the Li at the octahedral (48g) sites with theshorter T1 value should exhibit a higher mobility than the Li atthe tetrahedral (24d) sites with the longer T1 in our case.22

Here we utilized the difference in T1 to suppress the strongsignals from the octahedral sites. This was done byincorporating the inversion recovery with a given recoverytime, in such a way that the shorter T1 component (i.e., the 6Lisignals from the octahedral sites) was relaxed to the null pointwhile the longer T1 component (i.e., the 6Li signals from thetetrahedral sites) still remained in the −z axis.24 Figure 7 showsthe 6Li spectrum after suppressing the strong signals at 1.9ppm. Interestingly, there exists an additional peak at 1.6 ppm,which is completely buried in Figure 6. This peak should have asimilar T1 value to 0.8 ppm. Because of its very weak signalintensity, this newly discovered peak at 1.6 ppm most likelybelongs to those lithium ions occupying the distorted

octahedral sites (96h), and the peak at 1.9 ppm should belongto the lithium at the 48g sites.With the appearance of the new peak at 1.6 ppm (96h), we

deconvoluted the experimentally observed peaks with threemixed Lorentz−Gaussian peaks (Figure 6). The deconvolutionsfitted well with the experimental data, yielding relative Lipopulations in the 24d, 48g, and 96h sites, as listed in Table 2.The integrated area corresponds to the number of lithiumoccupying corresponding sites; thus, the occupancies atrespective sites can be obtained. It is worth noting that thepeak at 1.6 ppm corresponds only to a portion of the distortedoctahedral 96h sites that have a longer T1 value. There mayexist another fraction that is close to the center of theoctahedral sites and has a T1 value similar to that for theoctahedral 48g sites. Those fractions would be lost during theinversion recovery time, which would have affected the fittingresults, especially the occupancy of the 48g/96h sites.The occupancy of Al at the 24d sites reduced the occupancy

of Li (24d). Because the Al occupancies at the 24d sites variedwith the composition x, it is not logical to discuss theoccupancy of Li+ ions at the 24d sites with x. Instead, theoccupancies at the octahedral sites were directly associated withx, as Al does not occupy these sites. From the data, it is clearthat, as W content increases, the Li occupancy at the octahedral48g sites decreases, i.e., a decrease in the total Li content of thegarnet structure predictably leads to a decrease in the lithiumcontent at the octahedral sites. Both the NPD and NMR resultsindicated that the Li occupancy at the tetrahedral 24d sites isbetween 30−40% and the total amounts of Li at the octahedralsites (48g and 96h) are similar. However, the difference in 48g/

Figure 6. 6Li MAS NMR spectra of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x≤ 1) sintered at 1150 °C for 12 h.

Table 2. Fitting Results of 6Li MAS NMR Spectra of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) Solid Solutions

fit peak 1 (Li1, 24d sites) fit peak 2 (Li2, 48g sites) fit peak 3 (Li3, 96h sites)

x temp.c.s.ppm

L.W.(Hz)

area(%) occu T1

c.s.ppm

L.W.(Hz)

area(%) occu T1

c.s.ppm

L.W.(Hz)

area(%) occu T1

0 0.88 46.85 22.3 0.52 1.90 51.74 77.7 0.910.25 1150 0.82 35.0 15.8 0.34 23.8 1.97 35.03 79.5 0.86 4.2 1.41 36.3 0.31 0.025

1175 0.81 35.0 15 0.33 24.8 1.94 36.8 80.5 0.87 4.9 1.47 36.3 0.29 0.0241200 0.82 36.7 16.6 0.36 22.5 1.94 40.6 81.6 0.88 4.9 1.53 36.3 0.12 0.010

0.5 1150 0.84 31.7 17.8 0.36 28.0 1.92 33.4 73.6 0.74 2.5 1.54 34.5 0.52 0.0431175 0.82 34.5 20.8 0.42 21.4 1.91 35 72.3 0.72 2.0 1.54 34.5 0.41 0.0351200 0.79 31.7 15.3 0.31 17.7 1.92 33.4 75.3 0.75 2.4 1.53 34.5 0.56 0.047

0.75 1150 0.8 32.8 17.8 0.33 24.8 1.89 33.5 76.7 0.70 1.0 1.48 38.07 0.31 0.0261175 0.77 31.2 14.2 0.26 18.8 1.87 33.5 80.9 0.74 1.1 1.40 38.1 0.27 0.0221200 0.79 32.8 17 0.31 14.6 1.87 31.9 76.9 0.70 0.97 1.41 32.9 0.34 0.028

Figure 7. 6Li MAS NMR spectra of Li7‑2x‑3yAlyLa3Zr2−xWxO12 (0 ≤ x≤ 1) gathered by incorporating the inversion recovery with a givenrecovery time.

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96h occupancy determined from Rietveld refinement and NMRfitting results might be due to the error associated with eachmethod as well as the loss of a portion of the 96h sites signalfrom NMR fitting.An important question regarding the lithium migration

pathway is whether the Li+ ions jump from the octahedral siteto its neighboring octahedral sites, bypassing the tetrahedralsites? Or are the 24d sites in fact involved in the lithiummigration pathway? In NMR, two-dimensional chemicalexchange experiments can be used to measure the slowexchange between the lithium ions at the tetrahedral andoctahedral sites. Unfortunately, it is extremely difficult toobserve any exchange signals between the less populatedlithium at the tetrahedral sites and the much more populatedlithium at the octahedral sites, especially when the latter has amuch shorter T1 than the former. Again, we incorporated theinversion recovery with a given recovery time to suppress thestrong signals from the octahedral sites at 1.9 ppm before doingthe chemical exchange experiment. Figure 8 shows the two-

dimensional 6Li−6Li exchange spectrum after suppressing thestrong signals at 1.9 ppm. Clearly, there exists a cross-peak at(ω1 = 0.8, ω2 = 1.6) ppm, indicating the correlation betweenthe octahedral (96h) and tetrahedral (24d) sites: that is, the Li+

ions jumped from one tetrahedral site to neighboringoctahedral (96h) sites and vice versa. In the crystal structureof the garnet phase, the 96h sites are between 48g and 24dsites; this is in agreement with our deduction of lithiumattribution. These results confirm that the Li+ ions in thetetrahedral 24d sites are involved in the lithium migrationpathway and support the Li+ ions jump along 24d−96h−48g−96h−24d. Furthermore, the 24d sites exhibiting a lowermobility were expected to determine the overall conductivity.The diffusion coefficients of Li7−2x−3yAlyLa3Zr2−xWxO12 (x =

0.5) were obtained at different temperatures using 7Li diffusionmeasurements. Figure 9 and Figure S7 show the decays of 7Liintensities as a function of the applied pulse field gradient.Figure S7 clearly shows a sharp peak and a broad peak in thespectra, where the sharp peak is attributed to the Li+ ions

occupying the octahedral sites, and the broad peak is assignedto the Li+ ions occupying the tetrahedral sites based on theirintegrated areas. Clearly, the broad peak shows almost no decaywithin the range of the applied pulse field gradient, indicatingthat the Li+ ions occupying the tetrahedral sites diffuse ratherslowly. The sharp peak decays, indicating that the Li+ ionsoccupying the octahedral sites diffuse fast. Fitting the data usinga Stejskal−Tanner function yielded the diffusion coefficients of1.30 × 10−9 m2 s−1, 1.47 × 10−9 m2 s−1, and 1.89 × 10−9 m2 s−1

at temperatures of 9.4, 25.6, and 50.6 °C, respectively.25 Thediffusion coefficients obtained here characterize the diffusionrate for those Li at the octahedral sites, more precisely the 48gsites. The inset in Figure 9 shows the Arrhenius plot of thediffusion coefficient at the 48g sites, giving rise to the activationenergy of 6.12 kJ/mol, which is much smaller than that fromthe ionic conductivity and is also much less than the activationenergy of 31 kJ/mol from the average lithium jump rate thatwill be calculated in the next section. The average diffusioncoefficient from the Nernst−Einstein relationship is on theorder of 10−12 m2 s−1, much lower than that obtained throughthe diffusion measurements, 10−9 m2 s−1. Thus, we believe thatthe Li+ ions should have jumped only along the 24d−96h−48g−96h−24d pathway, because any Li+ ions jumping fromone octahedral site to a neighboring octahedral site would haveyielded a much higher ionic conductivity. Taking into accountthe lithium transport route of 24d−96h−48g−96h−24d andthe fact that the mobility at the 24d sites determined the wholeionic conductivity of the sample, it is reasonable to concludethat the diffusion coefficient of Li at the 24d sites was on theorder of 10−12 m2 s−1.

3 . 4 . L o w - T em p e r a t u r e P r o p e r t i e s o fLi7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) Solid Solutions.Figure 5 shows an inflection point of the Arrhenius plot at 25°C, where the activation energy of the ionic conductivity at thelow-temperature side is much higher than that at the high-temperature side, a phenomena stemming from the ion-trapping effect. For fast ion conductors, Nyquist plotting at lowtemperatures can show a more complete semicircle, which isnecessary for the collection of microinformation.Usually, the ac conductivity and the frequency follow

Jonscher’s universal power law, eq 1,26

σ ω σ ω ω ω ω= + = +A K( ) [1 ( / ) ]n ndc p p (1)

where σdc is the dc ionic conductivity, ω is the angularfrequency, ωp represents the jump rate of the charge carriers, A

Figure 8. 6Li−6Li exchange spectra of Li7−2x−3yAlyLa3Zr2−xWxO12 (x =0.5) sintered at 1150 °C for 12 h.

Figure 9. 7Li PFG NMR spectra of Li7−2x−3yAlyLa3Zr2−xWxO12 (x =0.5) sintered at 1150 °C for 12 h. The inset shows the Arrhenius plotof the diffusion coefficient.

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and K are the prefactors, and n is a dimensionless frequencyexponent that lies in the range 0 < n < 1. Figure 10 shows the

frequency-dependent σ(ω). There are two platforms in thedata, with the platform at the high-frequency side representingthe dc conductivity σdc and the decrease in ionic conductivity atthe low-frequency side being attributed to the electrodeprocess. Here, the bulk properties are the only focus of thefitting, and so the Jonscher relationship was used to fit the dataat the high-frequency side. The intercept of the fitting curve onthe y axis represents the dc conductivity, while the jump rate ωpcould be directly taken from the fitting results. The mobile ion

concentration was calculated from the Nernst−Einstein−Smoluchowski relationship, eq 2,

σω α

π=

c q

kT2dcp

2 2

(2)

where c is the mobile ion concentration, q is the charge, α is thejump distance and set to be 2 Å, k is the Boltzmann constant,and T is the absolute temperature in degrees Kelvin. The resultsare listed in Table S4.σdc, ωp, and c are important factors when studying solid

electrolytes, and their relationships with test temperatures aregiven in Figure 11 and Figures S8 and S9. The data fit theArrhenius plots well, indicating a thermal activation process.The activation energies were calculated and are listed in Table3. The sample with the composition of x = 1, however, shows alarge amount of impurities, which greatly affects the shape ofthe Nyquist plot and the fitting results. For that reason, onlythe data gathered from compositions with x = 0.25, 0.5, and0.75 are discussed.The activation energy of the ionic conductivity at the high-

temperature side in this data shows a clear pattern, but becauseof the ion-trapping effect, the difference in activation energy ofvarious compositions is not as obvious at the low-temperatureside. This ion-trapping effect is treated as the creation energy ofmobile Li+ ions, that is, it increases the activation energy thatactivates Li+ from fixed sites to a mobile status.With an increase in x from 0.25 to 0.75, σdc decreases from

1.7 × 10−4 to 2 × 10−5 S cm−1 (10 °C), the jump rate ωpdecreases from 77 256 768 to 13 434 467 rad/s (10 °C), andthe mobile Li+ concentration c decreases from 5.25 × 1021 to3.01 × 1021 cm−3 (10 °C). ωp shows a more obvious decrease

Figure 10. Frequency-dependent σ(ω) of Li7−2x−3yAlyLa3Zr2−xWxO12(x = 0.25) sintered at 1150 °C and associated fitting results.

Figure 11. Low-temperature properties of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) sintered at 1150 °C. σtotal values are the fitting results of Nyquistplots while σdc values are the fitting results of Figure 10.

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compared to that of c, indicating that the decrease in ionicconductivity is mainly affected by the decrease in ωp. However,by examining different compositions, the activation energy ofthe jump rate at between 30 and 32 kJ/mol shows negligibledifferences, indicating that the barrier for lithium jumping is notobviously changed. The decrease in the crystal lattice size isthus apparently too small to affect the activation energy of thelithium jump rate. The decrease in lithium jump rate with x ismainly due to the decrease in attempt frequency, although thestructural reasons for this phenomenon are not yet very clear.At 10 °C, only about 1−10% of the total lithium is mobile inthe garnet structure: most of the Li+ ions are fixed in the latticeand do not participate in lithium transport.27 With an increasein x, the creation energy of mobile Li+ ions increases from 2.9to 11.6 kJ/mol at the high-temperature side. The dissolutionenergy of W−O is 653.2 kJ/mol, which is smaller than the 760kJ/mol of the Zr−O band.19 With an increase in W content, theaverage dissolution energies of the M−O band decrease, the Oshows a stronger attraction to Li+, and the dissolution energy ofthe Li−O band increases. This causes the creation energy ofmobile Li+ ions to increase. The activation energy of the jumprate ωp is nearly constant. Therefore, the difference inactivation energies of ionic conductivity is mainly affected bythe creation energy of mobile Li+ ions. However, at lowtemperatures, the ion-trapping effect decreases from 8.1 to 1.1kJ/mol, behavior that is mainly due to a decrease in Li+ ioncontent and crystal defects in the garnet structure. These twoopposite behaviors cause the creation energy of mobile Li+ ionsto exhibit a minor increase with x at the low-temperature side,from 11 to 13 kJ/mol. Because the lithium mobility at 24d sitesdetermines the total conductivity, the obtained parameters aremainly attributed to the Li+ ions at 24d sites.

4. CONCLUSIONA series of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) solidsolutions were synthesized, and their properties were studied.The sample with the composition of x = 0.25 showed thehighest ionic conductivity, with σrt = 4.9 × 10 −4 S cm−1, andthe lowest activation energy, 34.1 kJ/mol. The ionicconductivities of the solid solutions decreased with the furthersubstitution of W, while the activation energies increased. Thedecrease in ionic conductivity was due to a decrease in jumprate at the 24d sites, which determined the total ionicconductivity of the samples. The increase of activation energy

was due to the heightened creation energy of mobile Li+ ions,and these mobile ions made up only a small amount of the totallithium. In addition, the ion-trapping effect decreased with anincrease in substitution content. The aliovalent substitution ofW did not seriously affect the crystal structure or the activationenergy of Li+ ion jumping, but it noticeably varied thedistribution of Li+ ions, electrostatic attraction/repulsion, andcrystal defects, which increased the lithium jump rate and thecreation energy of mobile Li+ ions. This is the mechanism ofthe previously observed effect that aliovalent substitution hason ionic conductivity.Three lithium sites, the tetrahedral 24d as well as the

octahedral 48g and 96h, were successfully resolved by high-resolution solid-state NMR, with the diffusion coefficients oflithium found to be 10−9 m2 s−1 at the octahedral sites and10−12 m2 s−1 at the tetrahedral sites. The lithium ions exchangebetween the 24d sites and 96h sites was observed, whichdirectly demonstrated a lithium transport route of 24d−96h−48g−96h−24d. The tetrahedral sites were the key pointsand de termined the ion i c conduct i v i ty o f theLi7−2x−3yAlyLa3Zr2−xWxO12 structure.The work carried out here sets up a direct relationship

between the macroscale ionic conductivity and microscalelithium dynamics, i.e., lithium jump rates and/or lithiumdiffusion coefficients at different sites, and provides a deeperunderstanding of lithium transport in the lithium-stuffedLi7−2x−3yAlyLa3Zr2−xWxO12 structure. It is anticipated that acombination of ac impedance testing (at high- and low-temperature ranges) and the use of the solid-state NMRtechnique would be a powerful way to study the lithiumtransfer performance and could be helpful in understanding thelithium dynamics in all solid electrolytes.

■ ASSOCIATED CONTENT

*S Supporting InformationThe Supporting Information is available free of charge on theACS Publications website at DOI: 10.1021/acs.chemma-ter.5b02429.

Lattice parameters of Li7−2xLa3Zr2−xWxO12 (0 ≤ x ≤ 1);27Al MAS NMR spectra of Li7−2x−3yAlyLa3Zr2−xWxO12 (0≤ x ≤ 0.75); EDS results of Li7−2x−3yAlyLa3Zr2−xWxO12

(0 ≤ x ≤ 0.75); Rietveld refinement of the structuralm o d e l s , b a s e d o n N P D d a t a f r o m

Table 3. Low-Temperature Properties of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1)

Ea/HTa(kJ/mol) Ea/LT

a (kJ/mol)

x S.T. σdc/10 °C (S cm−1) ωp/10 °C (rad s−1) c/10 °C (cm−3) ω0 (rad s−1) Ea,σ Ea,c Ea,σ Ea,σ/dc Ea,ω Ea,c Etp

0.25 1150 1.7 × 10−4 7.73E7 5.25E21 4.47E13 34.1 2.9 42.4 42.2 31.1 11 8.11175 1.37 × 10−4 6.72E7 4.89E21 3.89E13 34.4 3.4 42.5 42.2 31 11.2 7.81200 1.31 × 10−4 1.38E8 2.27E21 2.69E13 34.8 4.4 42.7 41.9 30.2 11.5 7.1

0.5 1150 4.72 × 10−5 2.97E7 3.81E21 1.44E13 38.9 7.9 44.5 43.5 31 12.5 4.61175 4.47 × 10−5 2.22E7 4.84E21 8.9E12 38.4 8 44.4 43.4 30.4 13 51200 3.79 × 10−5 2.29E7 3.96E21 1.2E13 39.4 8.3 44.5 43.6 31.1 12.5 4.2

0.75 1150 2.52 × 10−5 1.75E7 3.45E21 7.2E12 41.93 11.3 45.2 43.1 30.5 12.5 1.21175 2.02 × 10−5 1.60E7 3.01E21 7.2E12 42 11.6 45.6 43.1 30.7 12.7 1.11200 2.75 × 10−5 1.34E7 4.52E21 3.8E12 41.7 10 46.5 44.9 31.6 13.2 3.2

1 1150 7.91 × 10−6 6.35E6 2.89E21 1.6E11 43.7 46.5 40.6 24 16.51175 5.27 × 10−6 3.22E6 3.96E21 1.3E10 44.8 47.4 40.7 19.4 21.31200

aEa,σ is the activation energy of ionic conductivity; Ea,c is the creation energy of mobile Li+ ions; Ea, ω is the activation energy of jump rate; and Etp isthe ion-trapping effect.

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Li7−2x−3yAlyLa3Zr2−xWxO12 samples; Nyquist plots ofLi7−2x−3yAlyLa3Zr2−xWxO12 (x = 0.5 and 0.75); fittingresults of Nyquist plots of Li7−2x−3yAlyLa3Zr2−xWxO12 (x= 0.25); fitting results of Arrhenius plots ofLi7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1); dc polarizationcurves of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) solidsolutions; spin−lattice relaxation curves and fitting curvesof Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1); 7Li PFGNMR spectra of Li7−2x−3yAlyLa3Zr2−xWxO12 (x = 0.5);fitting results of σ(ω) of Li7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤x ≤ 1); and low-temperature properties ofLi7−2x−3yAlyLa3Zr2−xWxO12 (0 ≤ x ≤ 1) (PDF)

■ AUTHOR INFORMATIONCorresponding Authors*E-mail: rfu@magnet.fsu.edu.*Phone: 86 0592 2185 753. Fax: 86 0592 2185 753. E-mail:yyang@xmu.edu.cn.Author ContributionsThe manuscript was written through contributions of allauthors. All authors have given approval to the final version ofthe manuscript.

FundingThe National Basic Research Program of China (973 program,Grant no. 2011CB935903) and the National Natural ScienceFoundation of China (Grant nos. 21233004 and 21021002, andin part 21428303).

NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTSThe authors are grateful to Dr. Vanessa K. Peterson of BraggInstitute, ANSTO, Australia, for her technical support on high-resolution NPD data collection. R.F. is also indebted to being aPCOSS fellow supported by the State Key Lab of PhysicalChemistry of Solid Surfaces, Xiamen University, China.

■ ABBREVIATIONSXRD, X-ray diffraction; EDS, energy-dispersive spectrometry;NPD, neutron powder diffraction; SEM, scanning electronmicroscopy; PFG NMR, pulse field gradient nuclear magneticresonance; c.s., chemical shift; L.W., line width; occu,occupancy

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