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Aluminum-Assisted Densification of Cosputtered Lithium Garnet Electrolyte Films for Solid-State Batteries

  • Jordi Sastre*
    Jordi Sastre
    Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
    *E-mail: [email protected]. Phone: +41 58 765 61 10.
    More by Jordi Sastre
  • Tzu-Ying Lin
    Tzu-Ying Lin
    Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
    More by Tzu-Ying Lin
  • Alejandro N. Filippin
    Alejandro N. Filippin
    Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
  • Agnieszka Priebe
    Agnieszka Priebe
    Laboratory for Mechanics of Materials and Nanostructure, Empa—Swiss Federal Laboratories for Materials Science and Technology, Feuerwerkerstrasse 39, CH-3602 Thun, Switzerland
  • Enrico Avancini
    Enrico Avancini
    Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
  • Johann Michler
    Johann Michler
    Laboratory for Mechanics of Materials and Nanostructure, Empa—Swiss Federal Laboratories for Materials Science and Technology, Feuerwerkerstrasse 39, CH-3602 Thun, Switzerland
  • Ayodhya N. Tiwari
    Ayodhya N. Tiwari
    Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
  • Yaroslav E. Romanyuk
    Yaroslav E. Romanyuk
    Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
  • , and 
  • Stephan Buecheler
    Stephan Buecheler
    Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
Cite this: ACS Appl. Energy Mater. 2019, 2, 12, 8511–8524
Publication Date (Web):November 12, 2019
https://doi.org/10.1021/acsaem.9b01387

Copyright © 2019 American Chemical Society. This publication is licensed under CC-BY-NC-ND.

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Abstract

Garnet Li7La3Zr2O12 (LLZO) is a promising solid-state electrolyte due to its wide electrochemical stability window and high Li-ion conductivity. This electrolyte has potential to be employed in the form of thin films for solid-state batteries, a promising approach in the quest for safer batteries with higher energy densities at lower fabrication costs. In this study, we use a scalable cosputtering process to fabricate LLZO thin films with subsequent postannealing at a temperature of 700 °C, significantly below the sintering temperatures employed in ceramic pellet processing. We investigate the roles that Li excess and incorporation of Al play in the film’s crystalline phase, microstructure, phase stability, and, ultimately, ionic conductivity. Our results reveal that improving the conductivity of LLZO thin films requires not only the stabilization of the cubic phase but especially the densification of the film and the minimization of the proton exchange degradation mechanism in the presence of moisture and CO2. These issues can be mitigated by effectively controlling the amount of Li and incorporating Al as sintering agent. An ionic conductivity at room temperature of 1.9 × 10–5 S cm–1 was achieved with a 400 nm Al-substituted LLZO thin film. Finally, we prove that these LLZO thin films can be successfully deposited and crystallized on a LiCoO2 cathode.

1. Introduction

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Current state-of-the-art LIBs, based on liquid electrolytes, seem to be reaching a technological limit in terms of energy and power density. All-solid-state lithium batteries (SSLBs) have the potential to overcome this limitation and at the same time also increase the safety of the battery. (1−4) Garnet-type Li7La3Zr2O12 (LLZO) solid-state electrolyte has been drawing increasing attention since it was first synthesized by Murugan et al. in 2007, (5) due to its high ionic conductivity (around 10–3 S cm–1 at room temperature (RT) in sintered pellets), (6−8) wide electrochemical stability window (being stable against metallic lithium and up to 6 V vs Li+), (5,9−11) and robust mechanical and thermal stabilities. (12,13) Superior Li-ion conductivities are only observed in the high-temperature cubic phase of the garnet. (14−16) Nevertheless, stabilization of the high-temperature cubic phase at room temperature can be achieved by partially substituting the element sitting on one of the crystal lattice sites by an aliovalent element. The most common and successful approach consists of partially replacing the Li 24d sites by a trivalent metallic cation such as Al3+, Ga3+, or Fe3+, which increases the number of Li+ vacancies. (15,17−20)
Sub-micron-thick LLZO can be deposited using common thin film processing techniques (such as plasma sputtering, (21) pulsed laser deposition, (22) chemical vapor deposition, (23) sol–gel, (24) atomic layer deposition, (25) etc.). This fabrication approach could enable LLZO solid-state batteries with higher energy density at lower fabrication costs, in comparison to the fabrication approaches based on high-temperature sintering of ceramic pellets typically employed by the research community nowadays. Thin film electrolytes can significantly decrease the amount of nonactive material in a battery, thus reducing the stack size and increasing the energy density. Thin film processing can also enable new applications such as the fabrication of on-chip batteries or the development of flexible batteries.
However, the current best ionic conductivities measured in LLZO thin films (<1 μm), even with the addition of stabilizer elements, lag behind by about 2 orders of magnitude in comparison to that of the bulk material. Reported values for LLZO thin films range from 10–10 up to 10–5 S cm–1. (21,23,24,26−31) The work of Lobe et al. (32) shows an ionic conductivity in sputtered LLZO thin films comparable to the ones obtained with pellet samples; however, the in-plane conductivity measurements in that work were conducted using a conductive substrate, which might lead to an erroneous measurement of the ionic conductivity in the electrolyte. Very recently, Pfenninger et al. (33) reported a record value of 2.9 × 10–5 S cm–1 measured in-plane on a MgO substrate. In this work the films were fabricated using a pulsed laser deposition technique followed by a postannealing step at about 700 °C. The as-deposited films were overlithiated by introducing thin interlayers of Li3N.
The inferior ionic conductivity observed in LLZO thin films as compared to pellets and sheets may be related to the limited annealing temperature that is required to avoid a massive Li loss and the formation of nonconductive phases. Due to the high surface to volume ratio, Li evaporates more easily in a thin film. Thus, crystallization and densification are not easy to achieve with thin film processing methods, unlike in bulk pellets. Improving the ionic conductivity and achieving dense thin films under a relatively low annealing temperature are still challenges for this type of thin film electrolyte.
In this work we investigate the effects of excess lithium and incorporation of aluminum on the crystalline phase, density, and air stability of LLZO thin films, and how these two parameters can ultimately improve the ionic conductivity of LLZO thin films. For this study LLZO films were prepared using a cosputtering process followed by postannealing in oxygen atmosphere at 700 °C, about 400 °C below the typical temperatures employed in pellet processing. Our study reveals that stabilizing the cubic phase of LLZO is not the only key factor for obtaining a highly conductive LLZO thin film, but as expected, a dense and homogeneous microstructure is also very important. High porosity limits the ionic conductivity and is detrimental for the phase stability. In this sense, we studied the effect of air-induced degradation in films prepared with different amounts of Li and with Al.
The ionic conductivity of 1.9 × 10–5 S cm–1 obtained with Al-substituted LLZO thin films represents the highest value for sputtered LLZO thin films (21) and is in the same order of the value reported by Pfenninger et al. using pulsed laser deposition. (33) This value still lags about 1 order of magnitude behind the average LLZO pellet. However, taking into consideration the thickness of the film (400 nm), the cross-plane ionic resistance of this film is only 2.06 Ω cm2, about 50 times lower than the typical 1 mm thick pellet with a conductivity of 1 mS cm–1. This quantity provides evidence of the advantage of LLZO thin films in comparison to pellets: high ionic conductances can be achieved with lower processing temperatures and lower weight of the electrolyte.

2. Methods

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2.1. Thin Film Deposition

LLZO thin films of 400 nm were prepared on 10 × 10 mm2 polished MgO(100) single-crystal substrates (Jiangyin Maideli Advanced Materials Co., Ltd.). Films were prepared on MgO single-crystal substrates, which have a very low reactivity and diffusivity, to avoid possible reactions of the substrate with the film and the formation of secondary phases. This is a commonly used substrate in the investigation of LLZO thin films. (21,22,33−35) The deposition was carried out at room temperature by a radio frequency (RF) magnetron sputtering system (Orion, AJA International Inc.) with a confocal off-axis target configuration. Li7La3Zr2O12 (3N, 90% density) and Li2O (3N, 90% density) targets from Toshima Manufacturing Co. were cosputtered in order to deposit LLZO films with an excess of lithium. The sputtering process was performed at 0.3 Pa using a 50 sccm Ar + 1 sccm Ar:O2 gas flow. The deposition rate of each target was controlled independently via the sputtering power and monitored using a quartz crystal microbalance (QCM) sensor. Samples were prepared with an excess lithium ranging from 9.0(5) to 19.0(5) extra moles of Li per mole of LLZO.
Li7–3xAlxLa3Zr2O12 (Al:LLZO) thin films were obtained by following a multilayer sputtering approach, consisting of a sequence of Li-rich LLZO (cosputtered Li7La3Zr2O12 and Li2O) and aluminum metal layers. This approach was followed due to a high sputtering rate of Al, which makes it infeasible to achieve a simultaneous cosputtering together with the LLZO. The sputtering of the Al target was performed with a 25 sccm Ar flow at a pressure of 0.2 Pa.
After deposition, the amorphous LLZO films were transferred to a tube furnace (Carbolite GHA 12/300) connected to a vacuum pump on one side and a mass flow controller (MFC) on the other. The samples were annealed at 700 °C for 1 h with a heating/cooling ramp of 5 °C/min. During the annealing, an O2 flow was passed through the tube and the pressure inside the tube was monitored and regulated using the MFC and vacuum pump.
Figure 1 shows a schematic of the deposition process and the postannealing step.

Figure 1

Figure 1. Schematic of the sample preparation process. Amorphous LLZO thin films are deposited by sputtering Li7La3Zr2O12, Li2O, and Al targets. Sputtered films are postannealed in a O2 atmosphere at 700 °C for 1 h in order to crystallize the LLZO.

2.2. Characterization

The crystalline phase of the annealed LLZO thin films was assessed using in-plane grazing-incidence X-ray diffractometry (GI-XRD), impinging the sample with Cu Kα1 radiation at an incident angle (ω) of 2° and measuring the diffracted radiation (2θ) between 10° and 80°. The samples annealed in the tube furnace were measured ex situ with a Bruker D8 Discover diffractometer. For the in situ XRD measurements, a PANalytical X’Pert3 diffractometer was employed instead. In this case, XRD patterns were captured at different temperature stages during the annealing process. Phase identification and Rietveld refinements of the lattice constants were performed using the open-source software Profex.
The microstructure of the films was studied from cross-section and top-view images acquired with a scanning electron microscope (SEM) (Hitachi FEG-SEM S-4800). Cross-section samples were prepared by cleaving the MgO substrate using a diamond scriber and cleaving pliers.
The lithium-ion conductivity and its activation energy were extracted from temperature-dependent impedance spectroscopy. For this measurement, Au contacts with a thickness of ∼70 nm and a parallel spacing of 200 μm were thermally evaporated on the films after the postannealing step. Given the large ratio between the electrode separation and film thickness (t/w ∼ 500), we can neglect any fringe conductance. The impedance of the LLZO films was measured from 1 Hz to 10 MHz with an amplitude of 50 mV using a Paios all-in-one characterization system (Fluxim AG). The temperature of the sample was regulated using a temperature-controlled stage (Linkam LTSE-420-P) integrated with the measurement system. A PT100 temperature sensor was contacted on the sample’s surface to log its actual temperature. Impedance spectra were acquired for set temperatures ranging from 298 K up to 600 K. The measurement chamber was flooded with Ar to prevent degradation of the LLZO films.
The measured data was fitted using an equivalent circuit consisting of a series resistance (Rs) connected in series to a capacitance (Cgeom) in parallel with a resistance (Rion) and a constant phase element (CPEint), as proposed by Huggins et al. (36)Rs represents the series resistance introduced by the contacts, Cgeom the capacitance due to the geometry of the electrodes, Rion the ionic resistance of the electrolyte, and CPEint a constant phase element modeling the double-layer interface between the electrolyte and the electrodes. The resistance Rion is a combination of the bulk and grain boundary ionic resistances, from which the effective ionic conductivity (σion,eff) can be calculated from eq 1
(1)
where l is the electrode separation, w the electrode width, and t the film thickness. More details on this calculation are provided in the Supporting Information.
The effective ionic conductivities σion,eff at different temperatures were fitted to the Arrhenius equation:
where A is the frequency prefactor, T is the temperature measured on the film, Ea is the activation energy of the ion transport mechanism, and kB is Boltzmann’s constant.
Standard time-of-flight secondary ion mass spectrometry (ToF-SIMS) profiles were measured in a ToF.SIMS5 unit from IONTOF. Depth profiling was performed by Cs+-ion sputtering with an acceleration voltage of 2 kV on an area of 300 × 300 μm2. The primary ion source used for imaging was Bi+ with an acceleration voltage of 25 kV and a current of approximately 0.6 pA. Measurements were performed on a 100 × 100 μm2 area within the sputtering crater. A floodgun was used to avoid surface charging.
A 3D elemental distribution of the thin film was obtained using a high-resolution time-of-flight (HTOF) mass spectrometer from TOFWERK which was incorporated into a focused ion beam/scanning electron microscope (FIB/SEM) LYRA3 from Tescan. Measurements were conducted in a positive ion detection mode using a 180 pA continuous polyisotopic Ga+ beam with an energy of 20 keV (100 μm aperture, 32 μs dwell time, 512 × 512 resolution, and 2 × 2 binning). The pressure in the analytical chamber was in the range (6–9) × 10–6 mbar. The FIB scanning area was set to 10 × 10 μm2, but only a central 5 × 5 μm2 region was taken into account to prevent possible crater edge-induced artifacts. The collected 4D (x, y, z and corresponding mass spectra) data set was mass-calibrated using the TOF-SIMS signals of the main sample component, i.e., 7Li and both isotopes of the FIB beam (69Ga and 71Ga). The total signal was then decomposed in order to assess the information on single element distributions in 3D.

3. Results and Discussion

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3.1. Effects of Li Excess

3.1.1. Film Morphology

As previously shown by Rawlence et al., (21) sputtered thin films of LLZO crystallize in the nonconductive pyrochlore-type La2Zr2O7 phase when no additional Li is added to the film during preparation. This is a consequence of the loss of Li that occurs during the sputtering and especially during the postannealing process. Cosputtering of Li2O is an effective approach to overlithiate the as-deposited films and achieve a crystalline LLZO phase after postannealing. Rawlence et al. (21) reported that an excess of Li above 8 mol (4 mol of Li2O) per formula unit (p.f.u.) of LLZO was necessary to prevent the formation of nonconductive phases in LLZO thin films annealed at 700 °C for 1 h (similar conditions to the ones employed in this work). In the first part of this work we investigate the effect of further increasing the amount of Li in the as-deposited LLZO films, based on the hypothesis that an increase in lithium amount p.f.u. would have an effect on the microstructure of the films and their ionic conductivity. To decouple the effect of excess lithium from the effect of a substitutional element in the film microstructure, pristine LLZO thin films (without any substitutional element) were first investigated.
A series of pristine LLZO samples were prepared with different ratios of LLZO and Li2O deposition rates to control the excess amount of Li in the film. In addition, an interlayer of ∼15 nm Li2O was in some cases added between the MgO substrate and the LLZO film to further increase the level of lithiation. The amount of Li in the as-deposited film was calculated on the basis of the deposition rates measured with a quartz crystal microbalance (QCM) during the sputtering process. This method is not as accurate as other analysis techniques, but it offers a simple, fast, and nondestructive way to estimate the amount of Li in the as-prepared films. Given the lower accuracy of the method, wide error margins (±0.5 mol p.f.u.) are provided in the reported values. This margin of error was assessed on the basis of ICP-MS quantifications, which provide a higher level of accuracy.
Figure 2 shows cross-sectional and top-view SEM micrographs of LLZO thin films with different amounts of excess Li: 11.0(5), 13.0(5), 14.0(5), and 19.0(5) moles of excess Li p.f.u. of LLZO. The gaps between the substrate and the film observable in the cross-section SEM images appear due to the preparation of the samples. The MgO is cleaved, and therefore, the films become slightly detached due to the mechanical stress (FIB cross sections like the one in Figure S6 do not show such delamination). In all cases a layer of Li2CO3 and LiOH covers the LLZO film, which results from the exposure to air of the samples (while preparing the samples for SEM imaging) and the reaction of Li+ with moisture and CO2. In this series of samples, an increasing excess of Li has a detrimental effect on the LLZO film density. The higher the amount of excess Li is, the higher the porosity of the film. We suggest that these microstructural features appear due to the excess lithium segregating at the grain surfaces during annealing and promoting the formation of pores as a result of the vaporization of lithium oxide. In Figure S1 one can observe that the as-deposited film is dense, but it becomes porous after the annealing step. This phenomenon is in good agreement with the first-principles density functional theory study conducted by Canepa et al., (37) which predicts a spontaneous segregation of Li toward the particle boundaries during sintering. The high volatility of Li2O at the temperature range at which the samples were processed leads to the vaporization of the segregated lithium in the form of gaseous lithium (Li2O(s) → 2Li(g) + O2(g)) and gaseous lithium oxide (Li2O(s) → Li2O(g)). (38,39)

Figure 2

Figure 2. Cross-section and top-view SEM micrographs of annealed LLZO thin films on MgO prepared with different amounts of excess Li: (a) 19.0(5), (b) 14.0(5), (c) 13.0(5), and (d) 11.0(5) extra moles of Li per mole of LLZO.

3.1.2. Crystalline Phase

The X-ray diffraction (XRD) pattern of the as-deposited films shows a completely amorphous phase, as seen in Figure 3a. After the postannealing step all the prepared samples present a crystalline LLZO phase, as seen in the XRD patterns shown in Figure 3b. Rietveld refinements of the XRD measurements using as a reference the tetragonal LLZO phase (ICSD 246817) yield lattice constants of a = 13.1(1) Å and c = 12.7(1) Å for all four cases, in good agreement with previously reported values for tetragonal LLZO. The nonconductive La2Zr2O7 phase does not appear in any of the shown samples, which indicates that in all cases there is enough Li in the as-deposited film. One can observe that a peak corresponding to Li2CO3 appears around 2θ = 22°. This peak is more intense in the films with higher excess Li, but the intensity considerably drops in the films with less Li, indicating that the samples prepared with higher excess Li are more prone to form Li2CO3 when exposed to air. This difference in the amount of Li2CO3 can also be observed in the SEM images on Figure 2. At 2θ = 20° an additional peak can also be observed in the XRD patterns of the films with a higher excess of Li. This peak can be identified as Li2ZrO3, that also results from the excess Li.

Figure 3

Figure 3. Grazing-incidence X-ray diffraction patterns of (a) an as-deposited LLZO film and (b) postannealed films with different amounts of excess Li: 19.0(5), 14.0(5), 13.0(5), and 11.0(5) extra moles of Li per mole of LLZO (from top to bottom). (c) Phase ratio (tetragonal vs H+-stabilized cubic LLZO) evolution of LLZO films exposed to air at room temperature, and the top-view SEM micrographs of the film’s surfaces. Samples prepared with 9.0(5) and 11.0(5) moles of extra Li per mole of LLZO. (d) Schematic of the degradation process of porous pristine LLZO thin films when exposed to air.

In order to assess the Li loss during the processing steps, the Li amount in a postannealed thin film was quantified using ICP-MS. The amount of Li in the film, prepared with 16.0(5) moles of Li p.f.u. (9 extra moles of Li p.f.u., based on the deposition rates), was measured to be 8.92 mol p.f.u. (see Table S1), meaning that approximately 7 moles of Li p.f.u. are lost during the processing step for this given excess lithium. This can be used to draw a boundary at around 7 moles of excess lithium required to obtain a fully LLZO phase. With an initial Li amount below 14 mol p.f.u., the resulting composition after annealing would lay under the 7 moles of Li p.f.u. required for stoichiometric balance, and therefore would result in the formation of La2Zr2O7.
The amount of Li loss is strongly related to the annealing conditions (atmosphere, pressure, temperature, dwelling, etc.). For example, reducing the annealing pressure increases substantially the Li loss. In Figure S2 we show the resulting phases from two samples with a high excess of Li (19.0(5) mol p.f.u.) annealed, respectively, at a pressure of 1 mbar and 200 mbar in a pure O2 atmosphere. The LLZO film crystallized at lower pressure shows a predominant Li-deficient La2Zr2O7 phase, whereas in the sample annealed at higher partial pressure of O2, a pure LLZO phase is formed, with residual Li2CO3 and Li2ZrO3 due to an excess of Li.
These results point out the necessity of precisely controlling the amount of excess Li in the sputtered LLZO films if one seeks to fabricate compact and homogeneous films with high ionic conductivity. On one side, a deficit of Li will prevent the formation of the LLZO crystalline phase, leading to the formation of non-ionic-conductive phases, which are ineffective as an electrolyte. On the other side, excessive Li will result in less compact films, with pores and pinholes that would be very detrimental for the operation of an LLZO-based SSLB, as dendrites and shunts across electrodes could easily form.
Despite the lack of phase stabilizer, we observe that the samples prepared with different amounts of excess Li do not only show a fully tetragonal LLZO phase. As one can observe in Figure 3b, some samples show a more prominent peak splitting, characteristic of the tetragonal polymorph, whereas others show only single peaks, characteristic of the cubic polymorph. However, in all cases, phase quantification using Rietveld refinements yielded as a result a mixture of both phases, in different proportions depending on the as-deposited composition and handling history. Since the annealing process was performed in a sealed tube furnace without significant amount of air leaking in, it can be ruled out that this low-temperature cubic phase was occurring during the annealing, as is the case in the work of Xie et al. (40) and Quinzeni et al. (41) Indeed, Rawlence et al. (21) reported that nonstabilized LLZO thin films annealed in oxygen have a fully tetragonal phase prior to air exposure (as observed in the in situ XRD measurements).
We observe that samples prepared in the same batch had significantly different ratios between the cubic and tetragonal phases depending on the presence or absence of gold contacts (deposited after processing for the ionic conductivity measurements). The samples covered with gold tend to have a prominent tetragonal phase, as a result of the sealing effect provided by the gold, whereas samples without mostly show the cubic phase. Moreover, the ratio of tetragonal and cubic phases also depends on the sample history and especially on the length of time the samples have been exposed to air.
We performed systematic studies of this phase evolution using time-sequential XRD scans. Two freshly annealed samples with different amounts of excess Li in the as-deposited film (9.0(5) and 11.0(5) moles of excess Li p.f.u, respectively) were sequentially measured over the span of 10 h and after 3 days in air. Figure S3 shows the XRD patterns of the films after different air exposure periods. The phases were quantified using Rietveld refinements. Figure 3c shows the change in phase percentages obtained from the phase quantification. The cubic phase that appears in both films when exposed to air has a lattice constant above 13.0 Å, larger than the reported value of the high-temperature cubic phase. The value we obtain here agrees well with the reported lattice constants of the low-temperature cubic phase of LLZO. (42−44) With this observation we conclude that the LLZO films undergo a phase transformation (from tetragonal to cubic) at room temperature in the presence of H2O and CO2.
In both cases a clear evolution from an initially predominant tetragonal phase to a mostly cubic phase is observed after the film has been exposed to air for 3 days at RT. However, the speed of this transition is significantly different: the first sample rapidly becomes cubic whereas the second one takes several hours to become predominantly cubic. As seen in the top-view SEM images of the samples after removing residues from the surface using methanol (Figure 3c), the sample with a higher amount of excess Li presents a significantly higher number of pores as well as some pinholes, which entails a higher air-exposed surface with respect to the total volume.
As reported by Sharafi et al., (45) when LLZO pellets are exposed to air, they undergo a spontaneous reaction with moisture resulting in a Li+/H+ exchange and the formation of lithium hydroxide on the surfaces:
Subsequently, the LiOH present on the surface further reacts with CO2 to yield Li2CO3:
The cross-sectional and top-view SEM images shown in Figure 2 provide evidence of the formation of a residual layer on top of the LLZO thin films after exposure to air. The thickness of this layer is dependent on the amount of excess Li in the as-deposited film. This agrees with the proposed degradation mechanism: the higher the concentration of Li is, the more porous the film becomes during annealing and the larger the amount of Li2CO3 formed on the surface when exposed to air, due to the higher surface to volume ratio.
Figure 3d depicts the degradation mechanism of LLZO thin films in air. A higher porosity of the film implies that larger surface area will be exposed to air and therefore the faster and more extensive the degradation process will be.
ToF-SIMS was performed in two degraded samples with different amounts of deposited excess Li. Figure S4 shows the depth profiles of the Li, C, La, Zr, and MgO signals. These results confirm that a carbon-rich phase, namely, Li2CO3, is present on the upper surface of the LLZO films exposed to air. The sample with a higher amount of excess Li shows an apparently broader and less sharp interface between the Li2CO3 and LLZO layers. This can be correlated to the thicker residual layer observed in the SEM images as well as the higher porosity of the film, which allows the growth of Li2CO3 inside the film.
Energy-dispersive X-ray (EDX) spectroscopy also confirms the presence of a carbon-rich layer on the surface of the LLZO films (see Figure S5), confirming the formation of Li2CO3 on the LLZO surface after exposing the films to air.
The protonation of LLZO during this degradation process is the cause for the transition from a tetragonal phase to an expanded cubic geometry observed when the samples are exposed to air. With higher porosity, the protonation of LLZO thin films takes place faster as the surface to volume ratio is larger. The results on Figure 3c prove that a more compact LLZO film is more stable against air-induced degradation.
Similar to the substitution of Li+ sites by supervalent cations, it is still not well-understood how the presence of protons in the lattice stabilizes the cubic polymorph of LLZO. It seems though that the destabilization of the ordered Li+ sites is the cause for the rearrangement of the LLZO lattice.
The ionic conductivity of the films with a H+-stabilized cubic phase is expected to be similar or lower than the tetragonal LLZO, because the occupation of Li+ sites by H+ will reduce the amount of charge carriers and block the transport pathways. Moreover, the formation of Li2CO3 on the surfaces and particle boundaries will further increase the interfacial resistance. (46)
In the case of films, due to the large surface to volume ratio, the formation of LiOH and Li2CO3 on the surface of the thin films has a significantly higher impact than in the case of pellets, for which different approaches to remove this degraded interface exist (polishing, heat treatment, chemical etching). (47) Therefore, finding approaches to mitigate the air-induced degradation are key to improve the effective ionic conductivity of the films.

3.1.3. Li-Ion Conductivity

To study the ionic conductivity of the pristine LLZO thin films and its relation to the excess Li and film morphology, we conducted temperature-dependent impedance spectroscopy measurements on a selection of the prepared samples, as depicted in the inset in Figure 4a. The Nyquist and Bode plots of the data measured on the LLZO sample with an excess Li of 9.0(5) moles p.f.u are plotted in Figure 4a. The fitting of the data using the equivalent circuit described in the Methods section is also shown.

Figure 4

Figure 4. (a) Nyquist plot and Bode plot of the impedance spectroscopy data measured and the fitted data at temperatures ranging from 298 to 600 K on the sample with 9.0(5) excess moles of Li p.f.u. Inserts show the measurement setup. (b) Arrhenius plots of the effective ionic conductivity of four samples prepared with different amounts of excess Li.

One can observe a single semicircle due to the geometric capacitance and the ionic resistance followed by a low-frequency tail, resulting from the polarization at the ion-blocking electrodes. The presence of only one arch could misleadingly be interpreted as the lack of grain boundaries. However, the lack of a space charge between the grain interior and the grain boundary also results in a combined arch, as no capacitance appears at the boundary. The increase in the temperature results in a shrinking of the semicircles as a result of lower ionic resistances. The effect of the temperature can be better noticed in the Bode plot in Figure 4a. This behavior indicates a thermally activated ionic transport mechanism. The parameters resulting from the fitting can be found in Table S2.
The room temperature effective ionic conductivities of the pristine LLZO films with different amounts of excess Li in the as-deposited film can be found in Table 1. The error provided is based on the goodness of the fitting.
Table 1. Ionic Conductivity at Room Temperature and Activation Energy of Pristine LLZO Thin Films Prepared with Different Amounts of Excess Li p.f.u
excess Li p.f.u. in the as-deposited film [mol]ionic conductivity at RT [S cm–1]activation energy [eV]
9.0(5)2.47(10) × 10–60.494(4)
10.0(5)8.30(38) × 10–70.562(1)
11.0(5)3.32(12) × 10–70.598(6)
12.0(5)4.39(119) × 10–80.708(11)
The measured ionic conductivities show a clear correlation with the amount of excess Li added during the sputtering of the films and hence with the morphology. The higher porosity produced by the excess Li results in a poorer conductivity of the films, as the tortuosity of the Li+ migration pathways increases. The highest conductivity was measured in the samples with an excess of Li of 9.0(5) mol p.f.u., which is the densest film obtained without nonconductive phases (namely, La2Zr2O7). This value is very similar to the ionic conductivity reported for tetragonal LLZO thin films prepared with laser-assisted CVD by Loho et al. (23) and is around 1 order of magnitude lower than the highest conductivity measured for tetragonal LLZO in a sintered pellet by Wolfenstine et al. (16)
The Arrhenius plots of the LLZO films with excess Li amounts ranging from 9.0(5) to 12.0(5) mol p.f.u. are shown in Figure 4b. The activation energies range from 0.494(4) eV (in the sample with 9.0(5) extra moles of Li) to 0.708(11) eV (in the sample with 12.0(5) extra moles of Li). The best value, 0.494(4) eV, lies close to the 0.4 eV theoretically predicted by Meier et al. (48) for tetragonal LLZO.
The increase in the activation energy can be correlated with the density of the films as well. This finding is in agreement with the work by Wolfestine et al. (16) They reported an activation energy of 0.41 eV in a tetragonal LLZO pellet with a density close to the theoretical density. This value is considerably lower than in other experiments where lower density materials were studied. (14,15,49) In the case of films with lower density, the contribution of the surface diffusion to the total conductivity will be more relevant. Assuming that the transport on the surface requires a higher activation energy, the effective activation energy of the films with a larger proportion of grain boundaries therefore will be higher.

3.2. Al:LLZO Thin Films

3.2.1. Film Morphology

In the pristine LLZO, despite the improvements in the film homogeneity and compactness achieved with the regulation of the excess Li, some porosity is still visible inside the film (see details in Figure 5a). By contrast, the Al-substituted LLZO film in Figure 5b, prepared with the same amount of excess Li p.f.u. (10(5) moles of excess Li p.f.u.), presents significantly larger grains that span across the film thickness and a more compact distribution. These observations provide evidence that a higher densification of the films is achieved through the incorporation of Al.

Figure 5

Figure 5. Cross-sectional SEM micrographs of (a) a pristine LLZO film and (b) an Al:LLZO film.

Quantification of the porosity based on cross-sectional SEM images indicates a decrease in porosity of approximately 96% between the pristine LLZO film and the Al:LLZO sample annealed at higher pressure, as well as a significant reduction of the pores size (Feret diameter of porosity: 106.29 nm for the pristine LLZO and 33.52 nm for the Al:LLZO film).
We suggest that the Al incorporation plays a role as a sintering agent. The melting point of Al is around 660 °C, which is below the annealing temperature of 700 °C. The liquid phase of Al would change the edge retraction velocities from the surface diffusion; hence, it reduces the surface energy of LLZO as well as the solid–gas and solid–solid interfaces. (50−52) This role of aluminum as a sintering aid has been previously observed in the fabrication of LLZO pellets (51,53) and sheets, (54) but it had not been reported in LLZO thin films annealed at lower temperature.
In order to further investigate the distribution of Al into the LLZO thin films, FIB-TOF-SIMS measurements were carried out. This technique allows the elemental structure of a sample to be represented in a three-dimensional space with a high mass resolution while still maintaining a high spatial resolution (the best lateral resolution of this system can reach about 50 nm). See Supporting Information for more details on this technique and supporting figures (Figure S6). Figure 6a,b shows the TOF-SIMS signals of 27Al and 7Li on the analyzed area. Top views (Figure 6a) show images in the xy-plane integrated over 10 frames (i.e., FIB scans), and side views (Figure 6b) are the projections of the TOF-SIMS signals in the xz-plane integrated over 0.5 μm in the y-direction. Figure 6c shows the FIB secondary electron (SE) images collected simultaneously during the sputtering process and the acquisition of secondary ions, which allow the sample topology to be monitored.

Figure 6

Figure 6. Representation of elemental structure of the Al-incorporated LLZO film obtained using a FIB-TOF-SIMS technique. (a) Top and (b) side views of 7Li and 27Al signals. (c) Top-view secondary-electron (SE) micrographs obtained during the FIB sputtering steps. Description in the text.

On the film’s surface a stronger Li signal is observed, which spans over approximately 10 nm. This can indicate that some residual Li2O is still present on the surface and has reacted with air forming LiOH and Li2CO3, resulting in a higher positive Li-ion yield than in the LLZO film. It must be noted however that, in comparison with the thickness of the LiOH + Li2CO3 forming on top of the pristine LLZO layers, the amount appearing on the Al:LLZO films is significantly lower. This observation is consistent with the cross-sectional SEM micrographs shown in Figures 2 and 5, in which a thick residual layer (about 100 nm) is observed on top of the pristine films in comparison to the films with Al.
The top (Figure 6a) and side (Figure 6b) views of the Al and Li signals provide evidence of a correlation between the spatial distribution of both elements. The two distributions follow similar patterns but are not exactly the same, which allows topology-induced artifacts to be excluded. This similarity can indicate a chemical interaction between Li and Al, such as the formation of a Li–Al oxide at the interfaces. This phenomenon would explain the improved stability of the samples upon exposure to air as well as the higher density of the films, as it hinders the out-diffusion of Li as well as reduces the grain’s surface energy during annealing. The higher excess of Li present in the Al:LLZO film after annealing (see Table S1), in comparison to the pristine film prepared with the same excess Li, also supports this conclusion.
Spatially resolved quantification using energy-dispersive X-ray spectroscopy (EDX) of Al-substituted LLZO pellets obtained by Cheng et al. (55) also provides evidence of a tendency of Al to segregate at the grain boundaries. Pesci et al. (56) also investigated the inhomogeneous distribution of Al along grain boundaries of Al:LLZO pellets employing a similar FIB-ToF-SIMS technique. This phenomenon has also been recently reported by Wachter-Welz et al. (57) and linked to the wide distribution of conductivities measured in a set of Al:LLZO pellets prepared in the same conditions.
We can conclude that, besides the role in stabilizing the cubic phase by creating Li vacancies, Al incorporated to LLZO thin films tends to accumulate on the grain boundaries, where it acts as a sintering agent and hinders the out-diffusion of Li during annealing. Moreover, the passivation layer that forms at the interfaces limits the H+/Li+ exchange mechanism, increasing the air stability of the films.

3.2.2. Crystalline Phase

Al3+ is one of the most common supervalent ions employed in stabilizing LLZO. This element is able to partially occupy the Li sites in the LLZO lattice and generate additional Li+ vacancies. These additional vacancies are responsible for the stabilization of the high-temperature cubic phase at lower temperatures and the improved ionic conductivity. However, besides the effect on the crystalline phase, the substitutional element may also play a role in the densification of the LLZO thin films and the stability of the compound in the presence of moisture and CO2.
To investigate the overall impact of Al substitution in LLZO thin films, Al-incorporated LLZO (Al:LLZO) films were compared to the best Al-free (pristine LLZO) samples. The total amount of added Al was set to 1 mol p.f.u. of LLZO, in order to provide enough Al to stabilize the cubic phase. Studies have suggested a threshold of 0.2 moles of Al p.f.u. (58) However, in our case given the limited annealing temperature and the multilayer deposition approach, the amount of Al in the as-deposited film had to be increased substantially to guarantee the incorporation of enough Al to the LLZO grains.
The XRD patterns acquired during the in situ annealing shown in Figure 7a show the evolution of the phase from the amorphous state at room temperature until the full crystallization into cubic LLZO at 700 °C. At about 200 °C, some peaks start to appear at angles around 2θ = 30°. These peaks can be assigned to precursor phases like ZrO2, Li6Zr2O7, and La2O3. At 500 °C, tetragonal LLZO first appears, and at around 550 °C it undergoes a phase transition from tetragonal to cubic phase. Above 650 °C the LLZO peaks reach their maximum intensity, indicating a complete crystallization. The XRD measurement at 700 °C reveals a fully cubic LLZO phase. The presence of Al in the form of amorphous oxides cannot be determined from the XRD measurements, but it is likely that the excess Al accumulates at the grain interfaces in this form.

Figure 7

Figure 7. (a) Color map plot of the ambient in situ GI-XRD measurement of an Al:LLZO sample during the annealing process, starting at RT and heating up to 700 °C. (b) GI-XRD diffraction patterns of the pristine LLZO and the Al:LLZO samples after postannealing. Reference patterns of tetragonal LLZO (ICSD 246817) and cubic LLZO (ICSD 422259) are included. (c) GI-XRD patterns of the Al:LLZO sample exposed to air over 18 h, measured in 1 h steps and a final measurement of 3 h, and plot of the refined lattice constant of the cubic LLZO phase extracted from these XRD patterns.

Figure 7b shows the XRD diffraction patterns of the postannealed pristine LLZO and Al-substituted LLZO (Al:LLZO) films. XRD patterns of cubic (ICSD 422259) and tetragonal (ICSD 246817) LLZO phases were used as reference for phase quantification and lattice constants determination using Rietveld refinements. The pristine LLZO film used as a reference presents a prominent tetragonal phase (76(2)% tetragonal LLZO and 24(2)% cubic LLZO). The residual cubic phase present in the unsubstituted LLZO appears due to the Li+/H+ exchange degradation mechanism described in the previous section. By contrast, the Al-doped samples show a predominant cubic phase (87(3)% cubic LLZO and 13(3)% tetragonal LLZO). The residual tetragonal phase can be explained as a result of an incomplete diffusion of Al throughout the film.
Figure 7c shows the LLZO XRD patterns measured over 18 h (1 h per scan for the 15 first scans and 3 h for the last scan) of an Al:LLZO sample exposed to air. Rietveld refinements of the patterns do not reveal any relevant change in phases and lattice constants during the exposure to air. Figure 7c also shows the refined cubic LLZO lattice constant over time. After 18 h of exposure to air, the lattice constant remains stable at 12.98(1) Å. This value is in good agreement with previous reports for Al-substituted LLZO, (55,59) which provides evidence that the phase achieved is the desired high-temperature cubic phase, and not the H+-stabilized cubic phase previously described. Unlike the pristine LLZO thin films, Al:LLZO thin films show improved air stability, at least for the time spans investigated.

3.2.3. Li-Ion Conductivity

The results of temperature-dependent impedance measurements, performed following the same approach used with the pristine LLZO samples, were fitted using the same equivalent circuit previously described. The Nyquist and Bode plots of the measured data and the fitting for the Al:LLZO sample annealed at atmospheric pressure are shown in Figure 8a. The fitted parameters can be found in Table S3. Figure 8b shows the Arrhenius plot of the Al:LLZO sample in comparison to the best pristine LLZO sample. The effective ionic conductivity at RT of the Al:LLZO sample is 1.94(6) × 10–5 S cm–1, and the fitted activation energy is 0.481(3) eV.

Figure 8

Figure 8. (a) Nyquist plot and Bode plot of the impedance spectroscopy data measured and fitted at temperatures ranging from 298 to 600 K on the Al:LLZO sample. (b) Arrhenius plots of the effective ionic conductivity of the Al:LLZO and pristine LLZO samples.

In comparison to the pristine LLZO films, one can observe that an improvement in the effective ionic conductivity is achieved through the incorporation of Al. The value of 1.94(6) × 10–5 S cm–1 is higher than any of the previously reported conductivities in sputtered LLZO thin films (excluding the measurements carried out on conductive substrates by Lobe et al. (32)). However, this effective ionic conductivity still lags about 1 order of magnitude behind the best conductivities reported in highly dense Al:LLZO pellets processed at temperatures above 1000 °C. (58)
The activation energy improves only very slightly with respect to the pristine samples. Theoretical calculations estimate an energy barrier of ∼0.1–0.3 eV for Li+ conduction in bulk cubic LLZO. (48) Experimental measurements in Al-substituted LLZO pellets have demonstrated values as good as 0.26 eV, (58) in good agreement with the theoretical predictions.
To further investigate the ionic transport in the LLZO thin films, we also analyzed the complex electric modulus following the formalism developed by Almond and West (60) (see Supporting Information for more details on the calculations). By using this quantity, the contribution from the low-frequency tail due to the electrode polarization and other interfacial effects can be effectively minimized. (60) In this sense, it allows us to better distinguish hidden time constants in the Nyquist plot.
The imaginary parts of the electric modulus Mimag obtained from the complex impedance data of both samples are respectively plotted in Figure S7a,b. On both plots only a single peak is observable for all the temperatures, indicating that the charge carrier transport is characterized by a single Maxwell time constant (τ = ϵ0ϵr/σ, ϵ0, where ϵr is the vacuum permittivity, τ = 1/(2πfmax) the relative permittivity, and σ the conductivity). The frequency at which the imaginary part of the modulus is maximum can be related to this time constant by fmax.
Figure S7c,d shows the imaginary part of the electric modulus normalized to Mimag,max and to the maximum of the imaginary part of the modulus n = 1 – 0.7 = 0.3. One can observe that the slope of the lower frequencies in both samples (with and without Al) is close to 1, which indicates an ideal nondispersive long-range conduction as predicted by Debye theory. (60) However, the plots show some asymmetry in the high-frequency region. In this region the slope is about 0.7, which can be related to a dispersive exponent of fmax. This dispersion may be a consequence of the polycrystallinity of the material and the polarization at the electrodes.
The frequency at which the imaginary part of the electric modulus is maximized also follows a thermally activated behavior, similar to the ionic conductivity. Figure S7e shows the Arrhenius plots of Li+ for each sample. The fitted activation energies of the pristine LLZO and Al:LLZO films are 0.491(4) and 0.485(4) eV, respectively. These values are very similar to each other despite the different crystalline phases present in each sample. This reveals that in both cases the ionic transport is dominated by a similar conduction mechanism. The fact that the conductivities at room temperature present an order of magnitude difference may indicate that other factors beside the stabilization of the highly conductive phase play a role in the conductivity. We suggest these factors are the film density and air-induced protonation of the LLZO films. The incorporation of Al successfully enhances the film densification and improves the stability against air, which leads to an improvement of the conductivity by reducing the tortuosity of the conduction pathways, the increase in the charge carrier density (reduced Li-loss during annealing and Li+/H+ exchange upon exposure to air), and the reduction of the interfacial resistances due to the growth of Li2CO3.

3.3. Deposition of LLZO Films on LiCoO2

In order to demonstrate the feasibility of employing this processing approach on a real battery architecture, we deposited Al:LLZO films on top of a half-cell stack consisting of a Ni–Al–Cr current collector film and a thin film LiCoO2 cathode. Details on the current collector and the cathode films are described by Filippin et al. (61) To reduce the interdiffusivity of Co, La, and Zr between the cathode and the electrolyte during the postannealing step, a 25 nm layer of LiNbO3 was introduced as a diffusion barrier.
Figure 9a shows the SEM cross-section of the half-cell stack with the crystallized Al:LLZO film on top. The same density and homogeneity of the LLZO film previously observed on MgO substrates are also achieved when deposited on a cathode material. Also, the ToF-SIMS measurement in Figure 9b demonstrates that a sharp interface is present between the cathode and the electrolyte materials. As seen in the XRD pattern shown in Figure 9c, the high-temperature cubic LLZO phase is obtained after postannealing. The incorporation of a LiNbO3 interlayer seems to prevent the chemical reactions between the LLZO film and the cathode material that have been previously reported by Vardar et al. (62) Our future efforts will be focused on the fabrication of full battery stacks by depositing metallic lithium films as anode and investigating the electrochemical properties of such thin-film batteries.

Figure 9

Figure 9. (a) SEM cross-section image of a half-cell stack (Si/MgO/Ni–Al–Cr/LiCoO2/LiNbO3/Al:LLZO). (b) ToF-SIMS depth profile of the half-cell stack. (c) XRD pattern of the half-cell stack, including the reference patterns of cubic LLZO, LiCoO2, and the Ni–Al–Cr superalloy.

4. Conclusions

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The fabrication of dense and highly ionic-conductive lithium garnet films for all-solid-state lithium batteries is demonstrated by employing a cosputtering method followed by a postannealing step at 700 °C, which is substantially lower than the typical temperatures employed in processing powder LLZO into thick pellets. This method offers flexibility for the incorporation of excess lithium and phase stabilizers in the as-deposited LLZO films and can be transferred to the processing of other electrolyte materials.
Excess of Li in the sputtered films needs to be effectively regulated to compensate Li loss during the processing steps and achieve a crystalline and dense film. Too low an amount of Li in the as-deposited films leads to the formation of lithium-deficient nonconductive phases after postannealing, whereas a high excess of Li promotes the formation of porosity and inhomogeneities in the film as a consequence of Li evaporation. A porous morphology not only harms the ionic conductivity but also has a detrimental effect on the air stability due to the higher surface to volume ratio.
Incorporation of Al in the films not only acts as cubic phase stabilizer but also plays a role as sintering agent during the annealing of the films, which leads to a significant improvement in the density without requiring higher annealing temperatures. The superior compactness and the formation of an Al-rich passivation layer on the grain surfaces improve the stability of the films upon exposure to H2O and CO2.
We demonstrate that following this approach dense and highly conductive LLZO films can be fabricated at low processing temperatures, with an effective ionic conductivity of up to 1.94 × 10–5 S cm–1. This thin film electrolyte can be potentially integrated in an all-solid-state battery architecture, combining metallic Li as anode and a high-potential cathode material. In order to prove this, we show that the LLZO films investigated on MgO substrates can be deposited on a cathode material like LiCoO2 retaining their phase and morphological properties.

Supporting Information

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The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acsaem.9b01387.

  • Cross-sectional SEM images before and after annealing, ICP-MS measurement data, XRD measurements of films annealed with different pressures, XRD measurements of pristine films after different exposure times to air, ToF-SIMS depth profiles, EDX measurement of the LLZO surface, effective ionic conductivity calculations, fitted equivalent circuit parameters, FIB-ToF-SIMS measurement information and supplementary figures, and complex electric modulus analysis (PDF)

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Most electronic Supporting Information files are available without a subscription to ACS Web Editions. Such files may be downloaded by article for research use (if there is a public use license linked to the relevant article, that license may permit other uses). Permission may be obtained from ACS for other uses through requests via the RightsLink permission system: http://pubs.acs.org/page/copyright/permissions.html.

Author Information

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  • Corresponding Author
  • Authors
    • Tzu-Ying Lin - Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, SwitzerlandOrcidhttp://orcid.org/0000-0002-3428-9944
    • Alejandro N. Filippin - Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
    • Agnieszka Priebe - Laboratory for Mechanics of Materials and Nanostructure, Empa—Swiss Federal Laboratories for Materials Science and Technology, Feuerwerkerstrasse 39, CH-3602 Thun, SwitzerlandOrcidhttp://orcid.org/0000-0003-0425-246X
    • Enrico Avancini - Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, SwitzerlandOrcidhttp://orcid.org/0000-0003-1309-1290
    • Johann Michler - Laboratory for Mechanics of Materials and Nanostructure, Empa—Swiss Federal Laboratories for Materials Science and Technology, Feuerwerkerstrasse 39, CH-3602 Thun, Switzerland
    • Ayodhya N. Tiwari - Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
    • Yaroslav E. Romanyuk - Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland
    • Stephan Buecheler - Laboratory for Thin Films and Photovoltaics, Empa—Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, SwitzerlandOrcidhttp://orcid.org/0000-0003-0942-9965
  • Notes
    The authors declare no competing financial interest.

Acknowledgments

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This work was supported by the Swiss National Science Foundation (Grant 200021_172764) and Taiwan’s Ministry of Science and Technology (Grant 106-2917-I-564-001-A1). We acknowledge the Laboratory for Advanced Analytical Technologies at Empa for the ICP-MS measurements, the Laboratory for Energy Conversion Materials for the access to XRD equipment, and the Laboratory for Nanoscale Materials Science for the access to ToF-SIMS equipment.

References

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  • Abstract

    Figure 1

    Figure 1. Schematic of the sample preparation process. Amorphous LLZO thin films are deposited by sputtering Li7La3Zr2O12, Li2O, and Al targets. Sputtered films are postannealed in a O2 atmosphere at 700 °C for 1 h in order to crystallize the LLZO.

    Figure 2

    Figure 2. Cross-section and top-view SEM micrographs of annealed LLZO thin films on MgO prepared with different amounts of excess Li: (a) 19.0(5), (b) 14.0(5), (c) 13.0(5), and (d) 11.0(5) extra moles of Li per mole of LLZO.

    Figure 3

    Figure 3. Grazing-incidence X-ray diffraction patterns of (a) an as-deposited LLZO film and (b) postannealed films with different amounts of excess Li: 19.0(5), 14.0(5), 13.0(5), and 11.0(5) extra moles of Li per mole of LLZO (from top to bottom). (c) Phase ratio (tetragonal vs H+-stabilized cubic LLZO) evolution of LLZO films exposed to air at room temperature, and the top-view SEM micrographs of the film’s surfaces. Samples prepared with 9.0(5) and 11.0(5) moles of extra Li per mole of LLZO. (d) Schematic of the degradation process of porous pristine LLZO thin films when exposed to air.

    Figure 4

    Figure 4. (a) Nyquist plot and Bode plot of the impedance spectroscopy data measured and the fitted data at temperatures ranging from 298 to 600 K on the sample with 9.0(5) excess moles of Li p.f.u. Inserts show the measurement setup. (b) Arrhenius plots of the effective ionic conductivity of four samples prepared with different amounts of excess Li.

    Figure 5

    Figure 5. Cross-sectional SEM micrographs of (a) a pristine LLZO film and (b) an Al:LLZO film.

    Figure 6

    Figure 6. Representation of elemental structure of the Al-incorporated LLZO film obtained using a FIB-TOF-SIMS technique. (a) Top and (b) side views of 7Li and 27Al signals. (c) Top-view secondary-electron (SE) micrographs obtained during the FIB sputtering steps. Description in the text.

    Figure 7

    Figure 7. (a) Color map plot of the ambient in situ GI-XRD measurement of an Al:LLZO sample during the annealing process, starting at RT and heating up to 700 °C. (b) GI-XRD diffraction patterns of the pristine LLZO and the Al:LLZO samples after postannealing. Reference patterns of tetragonal LLZO (ICSD 246817) and cubic LLZO (ICSD 422259) are included. (c) GI-XRD patterns of the Al:LLZO sample exposed to air over 18 h, measured in 1 h steps and a final measurement of 3 h, and plot of the refined lattice constant of the cubic LLZO phase extracted from these XRD patterns.

    Figure 8

    Figure 8. (a) Nyquist plot and Bode plot of the impedance spectroscopy data measured and fitted at temperatures ranging from 298 to 600 K on the Al:LLZO sample. (b) Arrhenius plots of the effective ionic conductivity of the Al:LLZO and pristine LLZO samples.

    Figure 9

    Figure 9. (a) SEM cross-section image of a half-cell stack (Si/MgO/Ni–Al–Cr/LiCoO2/LiNbO3/Al:LLZO). (b) ToF-SIMS depth profile of the half-cell stack. (c) XRD pattern of the half-cell stack, including the reference patterns of cubic LLZO, LiCoO2, and the Ni–Al–Cr superalloy.

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    • Cross-sectional SEM images before and after annealing, ICP-MS measurement data, XRD measurements of films annealed with different pressures, XRD measurements of pristine films after different exposure times to air, ToF-SIMS depth profiles, EDX measurement of the LLZO surface, effective ionic conductivity calculations, fitted equivalent circuit parameters, FIB-ToF-SIMS measurement information and supplementary figures, and complex electric modulus analysis (PDF)


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