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Journal of Materials Science

, Volume 51, Issue 2, pp 854–860 | Cite as

Conduction below 100 °C in nominal Li6ZnNb4O14

  • Yunchao Li
  • Mariappan Parans ParanthamanEmail author
  • Lance W. Gill
  • Edward W. Hagaman
  • Yangyang Wang
  • Alexi P. Sokolov
  • Sheng Dai
  • Cheng Ma
  • Miaofang Chi
  • Gabriel M. Veith
  • Arumugam Manthiram
  • John B. Goodenough
Original Paper

Abstract

The increasing demand for a safe rechargeable battery with a high energy density per cell is driving a search for a novel solid electrolyte with a high Li+ or Na+ conductivity that is chemically stable in a working Li-ion or Na-ion battery. Li6ZnNb4O14 (LZNO) has been reported to exhibit a σ Li > 10−2 S cm−1 at 250 °C, but to disproportionate into multiple phases on cooling from 850 °C to room-temperature. An investigation of the room-temperature Li-ion conductivity in a porous pellet of a multiphase product of a nominal LZNO composition is shown to have bulk σ Li ≈ 3.3 × 10−5 S cm−1 at room-temperature that increases to 1.4 × 10−4 S cm−1 by 50 °C. 7Li MAS NMR spectra were fitted to two Lorentzian lines, one of which showed a dramatic increase with increasing temperature. A test for water stability indicates that Li+ may move to the particle and grain surfaces to react with adsorbed water as occurs in the garnet Li+ conductors.

Keywords

Solid Electrolyte LiNbO3 Scanning Transmission Electron Microscopy Solid Electrolyte Interphase Li3PO4 
These keywords were added by machine and not by the authors. This process is experimental and the keywords may be updated as the learning algorithm improves.

Introduction

Since the introduction by the SONY Corp. in 1991 of a wireless cell phone powered by a rechargeable Li-ion battery (LIB), the LIB has been used extensively to power portable computers, numerous portable electronic devices from cell phones to digital cameras, and hand-held tools. In all of these applications, the LIB does not compete with the energy stored in a fossil fuel. With an increasing awareness of the societal costs of extraction of fossil fuels and of the gaseous emissions from their combustion, the storage of electrical energy generated by alternative energy sources such as wind or solar energy has become a global priority. The development of rechargeable batteries that can power electric vehicles competitively with the internal-combustion engine and can store and supply, competitive with coal or gas, power to the electric grid in a stationary battery has become a global priority. Li-ion rechargeable batteries are receiving world-wide attention because their cell-specific energy density can, theoretically, achieve over 400 Wh kg−1 [1]. The Li–air and Li–sulfur batteries can, theoretically, reach specific energy densities per cell of 11400 [2] and 2510 Wh kg−1 [3], respectively. However, LIBs face major hurdles of safety, cost, and cycle life. Conventional LIBs contain an organic–liquid electrolyte that is flammable and is reduced by an anode that maximizes the possible voltage of a cell. The voltage can be increased where a passivating solid electrolyte interphase (SEI) layer is formed on the anode surface by the addition of a suitable chemical to the liquid electrolyte; but the SEI layer introduces two fundamental problems: (1) it must be permeable to the Li+ ion, and the Li+ of the layer comes from the cathode on the initial charge of a cell fabricated in a discharged state to introduce an irreversible capacity loss of the cathode, and (2) the SEI layer adds to the resistance to charge transfer between the anode and the electrolyte. Too fast a charge may result in a plating of metallic lithium, Li0, on the SEI layer, and a Li0 anode forms dendrites on plating that grow during charge; if the dendrites grow across a thin electrolyte to the cathode, an internal short-circuit gives rise to thermal runaway with incendiary consequences [4]. However, even where dendrites are blocked by a separator from reaching the cathode, they continuously create new anode surface area that needs to be pacified by new SEI formation, which leads to a capacity fade that limits the cycle life of the battery. Therefore, there is interest in the development of a solid Li+ electrolyte that can allow realization of a safe, low-cost LIB that approaches the theoretical limit of specific energy density [5]. However, the catalysts for the reactions at an air cathode are more stable and active in an aqueous electrolyte, which is incompatible with a Li0 or other anode giving the high voltage needed for a Li–air or Li–S cell; and the soluble intermediates of the cathode reactions in a Li–sulfur cell need to be blocked from reaching the anode. Therefore, there is a strong motivation to find a solid Li+ electrolyte that is stable in water and can be interfaced with a Li0 anode [6].

Although the advantages of a solid electrolyte are widely acknowledged, their practical application is still limited by a low ionic conductivity, too small an energy gap, and/or poor chemical/electrochemical stability in a battery. Since the advent of the Na-sulfur battery in 1967 that operates with β-alumina as the solid electrolyte and molten electrodes at over 300 °C, many material systems have been investigated for good lithium-ion conductivity. This search has included crystalline, glassy, polymer gels, and polymer–gel/oxide composites. In the 1970s, Li3N was shown to have a room-temperature conductivity as high as 6 × 10−3 S cm−1 [7]. However, this material has a small energy gap and, therefore, a poor electrochemical stability unless used only as an anode SEI layer. Many other Li+ conductors have since been discovered as is illustrated by the NASICON-structured LISICON Li1+x Al x Ti2−x (PO4)3 (LATP) [8], the perovskite-structured La3x Li(2/3)−x TiO3 (LLTO) [9], the thio-NASICON-structured LISICONS Li1.25Ge0.25P0.75S4 [10] and Li10GeP2S12 [11]. However, each of these materials has a fatal flaw that prevents their use as a solid electrolyte in a LIB. LATP and LLTO are unstable on contact with a Li0 anode and in an acidic aqueous electrolyte because the Ti4+/Ti3+ redox reaction is, respectively, at 1.8 and 2.5 V versus Li0 [12]. The NASCICON-structured LISICON and thio-LISICON compounds are chemically unstable on contact with an aqueous and an organic–liquid Li+ electrolyte and they do not provide an all-solid-state battery of adequate cathode capacity. Lithium-phosphorous-oxinitride (LIPON) films can be prepared in situ by a sputtering technique with a Li3PO4 target in a controlled N2 atmosphere [13]; it is stable on contact with Li0, but its room-temperature Li+ conductivity is only 10−6 S cm−1. The garnet-structured Li7La3Zr2−x Ta x O12 (LLZO) can have a bulk Li+ conductivity as high as 10−3 S cm−1 and it is stable against Li0 [14, 15, 16]; but on exposure to air, Li+ moves to the surface of particles and grains to react with adsorbed water and CO2. The water sensitivity limits their application.

Previously, Kanovalova et al. [17] reported a Li6ZnNb4O14 (LZNO) phase with an interesting Li+ conductivity at 250 °C. However, this temperature is too high for a room-temperature LIB with solid electrodes and the compound is metastable. In this paper, we reinvestigate the Li+ conductivity of nominal LZNO to lower temperatures; the bulk conductivity is 3.28 × 10−5 S cm−1 at room-temperature and increases to 1.43 × 10−4 S cm−1 at 50 °C, typical of the good crystalline oxide Li+ conductors.

Experimental details

The composition of LZNO was prepared by a conventional solid-state reaction. Stoichiometric amounts of dried Li2CO3 (Alfa Aesar, 99 %), Nb2O5 (Alfa Aesar, 99.9 %), and ZnO (Alfa Aesar, 99.99 %) were mixed thoroughly in an Al2O3 mortar. 10 % excess Li and Zn were used to compensate for the loss of lithium and zinc during the high-temperature firing process. The mixture of starting materials was first fired at 600 °C for 3 h, and then pelletized and fired at 1080 °C for 20 h with a ramping rate of 5 °C/min, followed by grinding with a mortar and pestle, then pelletized and fired at 1080 °C for another 20 h. Phase purity and stability were characterized with a PANalytical Empyrean diffractometer equipped with Cu Kα radiation (λ = 1.5406 Å). The scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDX) measurements of the sintered samples were performed with a Hitachi S4800 SEM operated at 20 kV. The scanning transmission electron microscopy (STEM) was performed with a JEOL 2200FS TEM/STEM operated at 200 kV to examine the particle morphology of the ground powders.

For the impedance measurement, both parallel surfaces of the pellet were sputtered with a layer of gold contacts. The experiment was performed with a Novocontrol Alpha-A impedance analyzer in the frequency range of 0.1–107 Hz. Temperature was ramped from room-temperature (21.5) to 250 °C in increments of 50 °C; about 30 min was allowed for temperature equilibration after each temperature ramp. The temperature of the sample was controlled by a Novocontrol Quatro Cryosystem with flowing nitrogen gas.

In order to estimate roughly the scale of the Li+/H+ exchange rate, 1 g of LZNO powder was dispersed in 15 mL of de-ionized water at ambient conditions. The pH values were monitored by a VWR symphony pH meter (Model SB70P). The SEM images for the pellets after 50 h in water were taken with a JEOL JSM-6060 SEM microscope operating at 10 kV.

7Li MAS NMR measurements were carried out on ground powders with a 9.4T Bruker Avance NMR spectrometer. The spectrum was examined over the temperature range of –43 to 87 °C in a 3.2 mm MAS rotor spinning at 15 kHz. 7Li spectra were acquired with an echo sequence (90°–180°) having a 90° pulse length of 2 µs. 7Li chemical shifts are given with respect to 1 M LiCl (δ = 0 ppm) with solid LiCl as the secondary standard (−1.06 ppm).

Results and discussion

The X-ray diffraction (XRD) pattern for the as-synthesized sample is presented in Fig. 1. The XRD pattern shows the presence of a major LZNO phase (PDF# 00-053-0343) and other impurity peaks of LiNbO3 and Li3NbO4. The phase of LZNO was first reported by Konovalova et al. [17]. To be metastable at room-temperature, it undergoes an irreversible phase decomposition below about 843 °C. Novel synthesis methodology is necessary to prepare LZNO single phase and needs further exploration to study its detailed crystal structure [18]. The impurities of LiNbO3 [19] and Li3NbO4 [20] formed along with the LZNO are considered to be a slow lithium ion conductor. These impurities would mostly segregates in the grain boundaries and hinder the lithium ion conduction. These impurities will also accelerate lithium/proton exchange rate when contacted with water, which will promote lithium migration to the grain boundaries and react with CO2 and form Li2CO3.
Fig. 1

Powder XRD diffraction pattern of the as-synthesized Li–Zn–Nb–O sample with a nominal composition of LZNO

The surface morphologies of the nominal LZNO pellet sintered at 1080 °C is shown in Fig. 2a. The grain size varies from 2 to 25 µm and the average is about 13.5 µm. Pores were visible on the surface of the pellet, which indicates a relatively low density of the sample. Such pores could be caused by the loss of zinc and lithium due to their high vapor pressures at high temperatures during sintering. Such a low-density pellet with pores would lead to a relatively large impedance since the active conducting area will be reduced. Further densification, e.g., by hot pressing, could be performed to improve the material performance. A bright-field STEM image of a typical as-synthesized (LZNO) grain is shown in Fig. 2b. The particle size ranges from 250 to 1000 nm. Figure 2c shows the EDX spectra taken from various positions of the sintered pellet. The pellet was broken apart to compare the elemental distribution along the original surface and the cross-section. The EDX spectra show a Nb-rich composition on the pellet surface, which indicates more Zn was lost from the surface than inside of the pellet during sintering.
Fig. 2

a A SEM image of the sintered LZNO pellet. b A bright-field STEM image of the LZNO particle. c EDX spectra of the LZNO pellet surface and interior

Figure 3 presents impedance spectroscopy characterization (in Nyquist format) of the LZNO pellets with Au-blocking electrodes at six selected temperatures: 21.5, 50, 100, 150, 200, and 250 °C. At lower temperatures (21.5, 50, 100 °C), two depressed semi-circles are clearly resolvable and the near vertical line at low frequencies associated with the double layer capacitance of the blocking electrode is present. The resistance component on the high-frequency side could be attributed to bulk conduction; the other was due to the grain-boundary contribution [14]. The experimental impedance data were initially fitted by considering only bulk and grain-boundary conduction with blocking-electrode effects. However, a large mismatch was observed between the measured and the calculated plots. Hence, the overall behavior of the impedance plots were fitted by an equivalent circuit of (R b Q b) (R gb Q gb) (R low Q low) (Q el) proposed by Tenhaeff et al. [21], where the subscripts b, gb, low, and el refer to bulk, grain-boundary, low-frequency process, and blocking electrodes, respectively. At this point, the physical phenomenon responsible for the low-frequency process is still unknown and needs further investigation. The bulk Li-ion conductivity is σ b = 3.28 × 10−5 S cm−1 and the grain-boundary Li-ion conductivity is σ gb = 1.74 × 10−5 S cm−1 at 21.5 °C. With increasing temperature, the (LZNO) pellet conductivity goes up very quickly and reaches σ b = 1.43 × 10−4 S cm−1 for bulk and σ gb = 9.30 × 10−5 S cm−1 for grain-boundary at 50 °C. Such results are close or only one order of magnitude lower than some of the intensively studied garnet and NASICON-type solid electrolytes [22, 23, 24, 25, 26, 27]. The conductivity of the pellet increased rapidly to 1.31 × 10−2 S cm−1 at 250 °C.
Fig. 3

Nyquist plots of AC impedance spectra for the as-synthesized Li–Zn–Nb–O pellet at a 21.5 °C, b 50 °C, c 100 °C, d 150 °C, e 200 °C, f 250 °C. The equivalent circuit (R b Q b) (R gb Q gb) (R low Q low) (Q el) was used to model the experimental data

Figure 4 shows the Li-ion conductivity as a function of 1000/T for a sintered sample of the nominal LZNO pellet. The bulk and grain-boundary contributions to conductivity could not be distinguished above 100 °C. Hence, the total conductivity was plotted above 100 °C and bulk conductivities are used at lower temperatures. The temperature dependence of the conductivity can be expressed by the Arrhenius equation,
Fig. 4

Temperature dependence of the Li-ion conductivity of LZNO

$$ \sigma = \frac{A}{T}\exp \left( {\frac{{ - E_{\text{a}} }}{{k_{\text{B}} T}}} \right), $$
(1)
where A is the pre-exponential factor, T the absolute temperature, E a the activation energy, and k B the Boltzmann constant. The activation energy of the (LZNO) sample is estimated to be 0.39 eV from the slope of the log (σT) versus 1/T plot in the temperature range of 294.5–523 K.

A platinum disk and lithium films of 1 µm thickness were sputtered on the as-synthesized LZNO pellet surfaces to make Li/LZNO/Pt and Li/LZNO/Li cells. The interfacial resistances were too large to allow testing of their chemical and electrochemical stability. However, excellent stability results were obtained with a study of 90 % β-Li3PS4 (LPS) and 10 % LZNO composite by Hood et al., which would suggest a possible good stability of LZNO. Cyclic voltammetry measurements show that the composite of LPS and LZNO is stable up to a potential of 5 V versus Li/Li+ and the symmetric Li/composite electrolyte/Li cell shows a long-term compatibility of the composite with metallic lithium [28].

As-synthesized LZNO powders were placed in de-ionized water and stirred at room-temperature along with a pH meter to monitor the changes in the pH values. As shown in Fig. 5, the pH value increased almost immediately upon addition of the LZNO powders and stabilized quickly. The reaction of the samples with water was not accompanied by gas evolution and a considerable amount of lithium exchange occurred rapidly after the sample was placed into the water. The pH of the solution stabilized at about 10.5, which is close to 11.2, the value of a cubic garnet sample Li7La3Zr2O12 [29] under similar test conditions (1 g sample in 15 mL de-ionized water). This comparison indicates a comparable zero-zeta-potential (ZZP) for LZNO and cubic LLZO.
Fig. 5

Time dependence of the pH value of the de-ionized water during the Li+/H+ Exchange. The time t = 0 represents the point where the powder was added

The XRD patterns of LZNO in Fig. 6 shows no obvious additional impurity peaks could be found after immersing in de-ionized water for 50 h at room-temperature. However, SEM images indicate some small particles appear on the surface after water treatment; a significant fraction of the impurity particles also appear at the grain boundaries. This observation indicates that Li+ migrates to the surface to react with adsorbed water as in the garnets [30].
Fig. 6

Comparison of the XRD patterns and SEM images of the Li–Zn–Nb–O samples before and after immersion in de-ionized water for 50 h

Figure 7a shows static 7Li NMR spectra of LZNO as a function of temperature from –43 to 87 °C. The spectrum undergoes a dramatic sharpening over this temperature range with the linewidth narrowing by more than one order of magnitude. Li mobility is greatly increased over this temperature range. The spectra display a centerband resonance that appears at 0.8 ppm in the high-temperature limit. The peak shape changes over the temperature range as the sharpest component of the line has a temperature-dependent shift (0.4–0.8 ppm). Figure 7b–e shows representative deconvolutions of the composite line shape. The experimental spectra were fitted as the sum of two Lorentzian lines. Fast Li-ion conduction is associated with the sharp component and slower Li ions with the broader components [31].
Fig. 7

a Stack plot of the static 7Li NMR spectra of LNZO as a function of temperature. b–e The line shapes at –43, 7, 27, and 87 °C are fitted as a sum of two Lorentzian components

It is notable that the chemical shift of the narrowest component in Fig. 8a increases with temperature while the broader component shows a slight diamagnetic dependence. In Fig. 8b, the linewidth of the resonance components narrow differentially with increasing temperature: the narrow component sharpens over the temperature range –30 to 50 °C while the broader components narrow over a higher temperature range: 50–90 °C. This behavior is typical for a sample that is not a pure phase in which multiple Li-ion mobility pathways occur in the material. The plot of integrated area of the components versus temperature (not shown) shows the mole fraction of the narrow and broad components are best described as independent of temperature at 0.45 and 0.55, respectively, again consistent with the multiple phase nature of the solid electrolyte.
Fig. 8

a Chemical shift dependence and b line width (FWHM) dependence with temperature in static 7Li NMR spectra of LNZO

Conclusions

A room-temperature fast Li-ion conducting solid electrolyte with a starting composition of LZNO was prepared by a solid-state reaction method. The bulk conductivity was measured to be 3.28 × 10−5 S cm−1 at room-temperature and increased to 1.43 × 10−4 S cm−1 at 50 °C. The nominal Li3ZnNb4O14 composition disproportionates by room-temperature into multiple crystalline phases on cooling from 1080 °C, but the Li+ conductivity remains comparable to that of other good crystalline oxide Li+ conductors.

Notes

Acknowledgements

The research was sponsored by the U.S. Department of Energy, Office Science, Office of Basic Energy Sciences, Materials Sciences and Engineering Division. The NMR research (L.W.G. and E.W.H.) was supported by the U.S. Department of Energy, Office of Basic Energy Sciences, Chemical Sciences, Geosciences and Biosciences Division. Scanning electron microscopy research was supported through a user project supported by ORNL’s Center for Nanophase Materials Sciences (CNMS), which is sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, U.S. Department of Energy. Dr. John B. Goodenough was supported by the Materials Sciences and Engineering Division, Office of Basic Energy Sciences, Office of Science, U.S. Department of Energy grant number (DE-SC0005397).

This manuscript has been authored by a contractor of the U.S. Government under contract DE-AC05-00OR22725. Accordingly, the U.S. Government retains a paid-up, nonexclusive, irrevocable, worldwide license to publish or reproduce the published form of this contribution, prepare derivative works, distribute copies to the public, and perform publicly and display publicly, or allow others to do so, for U.S. Government purposes.

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Copyright information

© Springer Science+Business Media New York (outside the USA) 2015

Authors and Affiliations

  • Yunchao Li
    • 1
    • 2
  • Mariappan Parans Paranthaman
    • 1
    • 2
    Email author
  • Lance W. Gill
    • 1
  • Edward W. Hagaman
    • 1
  • Yangyang Wang
    • 1
  • Alexi P. Sokolov
    • 1
  • Sheng Dai
    • 1
  • Cheng Ma
    • 3
  • Miaofang Chi
    • 3
  • Gabriel M. Veith
    • 4
  • Arumugam Manthiram
    • 5
  • John B. Goodenough
    • 5
  1. 1.Chemical Sciences DivisionOak Ridge National LaboratoryOak RidgeUSA
  2. 2.The Bredesen Center for Interdisciplinary Research and Graduate EducationThe University of TennesseeKnoxvilleUSA
  3. 3.Center for Nanophase Materials SciencesOak Ridge National LaboratoryOak RidgeUSA
  4. 4.Materials Science and Technology DivisionOak Ridge National LaboratoryOak RidgeUSA
  5. 5.Texas Materials InstituteThe University of Texas at AustinAustinUSA

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