Manufacturing of metal-ionic conductor layers via solution combustion synthesis

Abstract

Solution combustion synthesis of amorphous metal-ionic conductor layers is provided using an ambient processing environment having 5% or less relative humidity. The resulting layers are ionic conductors having an ionic conductivity at least 10x their electronic conductivity, and preferably higher in some applications such as solid-state battery electrolytes. The amorphous nature of these layers is a significant advantage for battery applications, since the amorphous morphology inhibits dendrite formation.

Description

[0001] Manufacturing of metal-ionic conductor layers via solution combustion synthesis

[0002] FIELD OF THE INVENTION

[0003] This invention relates to the fabrication of metal-ion conductor layers .

[0004] BACKGROUND

[0005] Metal oxide layers that contain alkaline metals for batteries are very di f ficult to deposit via conventional methods . Sensitivity to carbon dioxide and water in the atmosphere , high required sintering temperatures ( often over 400 ° C over long duration) , and expensive precursors have made scalable manufacturing of solid oxide electrolytes di f ficult and posed a critical hurdle for the field .

[0006] The standard approaches using vacuum-based deposition ( such as sputtering or CVD) struggle with low throughputs , high costs , and expensive equipment . Solution-based methods are limited generally to the sol-gel approach which is very sensitive to atmospheric conditions , relative humidity, and solution pH to deliver consistent films . Other methods like flame or spray pyrolysis can demand high temperatures to sinter particles of oxide species or use halogenated precursors which can contaminate resulting films .

[0007] Particular focus is directed toward a need for amorphous metal oxide electrolyte layers due to their ability to block dendrites in batteries . In order to achieve amorphous character, high temperature syntheses with long annealing times have to be avoided excluding many processing techniques available for traditional metal oxide fabrication . Accordingly, it would be an advance in the art to provide improved fabrication of amorphous metal-ion conductor layers .

[0008] SUMMARY

[0009] In one example , we present a method to rapidly generate amorphous lithiated metal oxide layers with controllable lithium content for use as solid state ionic electrolytes . The method was able to generate an amorphous layer of lithium lanthanum zirconium oxide ( LLZO) with a Li carbonate- free bulk in less than 5 minutes at 350 ° C . Furthermore , blade casting was used to generate the layers on a variety of substrates at up to 10 mm / sec .

[0010] The method leverages solution combustion synthesis ( SCS ) which is a next-generation metal oxide fabrication technique that can generate metal oxides at substantially lower temperatures ( often hundreds of degrees Celsius lower and / or at much higher speeds ) than conventional oxide forming counterparts owing to the presence of both oxidative and reductive moieties complexed to the metal cation . Our precursor solution includes a blend of three or more metals with one metal being a variable lithium content of up to 60% lithium oxide content . In a one step process , the precursor is deposited onto a substrate , heated to T > Tignition and the oxide layer is generated .

[0011] We have started by demonstrating success to form LLZO, one of the most used and popular Li battery solid state electrolytes and have shown variable composition compatibility between the lanthanum oxide , zirconium oxide , and lithium components . Additionally, we have demonstrated that this synthesis is compatible with extremely dry air as a processing environment , thereby eliminating the need for inert atmospheres and vacuum-based processing often used in traditional battery manufacturing .

[0012] The SCS-based method for metal oxide formation opens the door to the commerciali zation of various next generation solid state battery technologies currently gated by the lack of scalable processing techniques . The roughly doubled energy density of solid-state batteries would allow for the production of electric vehicles with higher ranges than currently available , as well as open the path to electri fy commercial airplanes , drones , and other areas such as safer consumer electronics .

[0013] Current competing technologies for depositing battery layers tend to be poorly scalable or expensive processes including spin coating and vacuum sputtering methods . Furthermore , the deposited chemistry of such technologies is typically formed through inherently unstable processes such as sol-gel methods . These processes are high temperature and relatively slow, often leading to lithium loss during sintering and requiring expensive excesses of lithium- containing reagents . This loss in material is also associated with the creation of pinholes in solid state layers , which leads to di f ficulty in creating large-scale layers without signi ficant shorting .

[0014] Due to the low temperature required and fast speed of the synthesis , we can form lithium-containing layers cheaply and without substantial loss in materials . We are able to form amorphous oxide layers which have been shown to have improved dendrite blocking ability . Additionally, dopant additives can be introduced simply by co-dissolving the low- cost metal nitrate or organo-metallic salt corresponding to the desired oxide product . BRIEF DESCRIPTION OF THE DRAWINGS

[0015] FIG . 1 schematically shows solution combustion synthesis ( SCS ) of metal-ion conductor layers .

[0016] FIGs . 2A-C show characteri zation results for a SCS- fabricated LLZO layer .

[0017] FIGs . 3A-B shows scanning electron microscopy and atomic force microscopy results of a SCS- f abricated LLZO layer .

[0018] FIG . 4 shows electrochemical impedance spectroscopy results for SCS- f abricated LLZO layers .

[0019] FIG . 5 is an Arrhenius plot for ion conductivity activation energy in SCS- f abricated LLZO layers .

[0020] DETAILED DESCRIPTION

[0021] Section A describes general principles relating to embodiments of the invention . Section B describes detailed experimental examples .

[0022] A. General principles

[0023] An exemplary embodiment of the invention is a method of making a metal-ionic conductor having the following steps :

[0024] 1 ) Depositing a precursor on a substrate , where the precursor includes two or more metal species , at least one fuel species , and at least one oxidant species . In the example of FIG . 1 , 102 schematically shows the precursor chemistry, and 104 is the precursor layer deposited on substrate 106 . The detailed example below relates to LLZO, a three-metal compound . However, amorphous layers of two- metal ion-conductor compounds can also be made with the present approach, with exemplary materials including, but not limited to: lithium zirconate, sodium zirconate, lithium aluminate, and sodium aluminate.

[0025] 2) Combusting the precursor in an ambient processing environment having relative humidity 5% or less to provide an amorphous layer of metal-ionic conductor that is a compound of the two or more metal species and at least one bridging heteroatom species. In the example of FIG. 1, 108 is this amorphous layer of metal-ionic conductor, and the lower left of FIG. 1 schematically shows the larger free energy of combustion synthesis than that of sol-gel synthesis. At least one of the two or more metal species provides mobile cations in the layer of metal-ionic conductor, and the ionic conductivity of the layer of metalionic conductor is at least lOx the electronic conductivity of the layer of metal-ionic conductor. Such materials can be suitable for use in batteries 110.

[0026] Preferably, the ionic conductivity of the layer of metal-ionic conductor is at least lOOx the electronic conductivity of the layer of metal-ionic conductor. The ionic conductivity of the layer of metal-ionic conductor is preferably at least 10“6S / cm. The electronic conductivity of the layer of metal-ionic conductor is preferably no more than 10“9S / cm.

[0027] The ambient processing environment preferably has a relative humidity of 2% or less, and more preferably this relative humidity is 1% or less.

[0028] Suitable metal species include, but are not limited to: Li, Na, K, Al, La, Zr, and Ti. Suitable mobile cations include, but are not limited to ions of: Li, Na, K, and Al. Suitable bridging heteroatom species include, but are not limited to: 0, S, P, N, and halogens. B . Detailed examples

[0029] Bl . Introduction

[0030] Lithium lanthanum zirconium oxides ( LLZO) are critical materials of interest for solid-state battery electrolyte . LLZO can be used for battery applications thanks to its high ionic conductivity and high stability against lithium metal , enabling high charge rates and energy densities . LLZO' s amorphous phase ( a-LLZO) has a high conductance in the submicron thin film case while still maintaining a low electronic conductivity in the 10-11range .

[0031] Thin film solid-state batteries have the potential to enable substantial increases in both energy density and safety for energy storage . Compared to traditional device architectures with polymeric membranes , liquid electrolytes , or thick ( >20 pm) ceramic pelleti zed separators , solid-state thin film batteries of fer potential for both gravimetric and volumetric savings by relying on anode and cathode separation by a film thinner than 1 pm . However, manufacturing approaches for critical thin film energy storage materials can be laborious and limit end applications due to high costs , requirements for vacuums , limitations on throughput to batch-to-batch processes , and lack of stable precursors .

[0032] We demonstrate a rapid, facile method for producing low-def ectivity a-LLZO metal oxide thin films using solution combustion synthesis . Solution combustion synthesis has its origins in solid-state combustion reactions to form nanomaterials , largely for the formation of refractory ceramics . While the original principles were purely solid- state processes ( e . g . , the reaction of Zr particles with elemental boron) , the principles developed are widely applicable to solution-based approaches . Solution combustion synthesis ( SCS ) has held a growingly influential method of generating nanostructures at substantially reduced process temperatures ( often below 300 ° C ) . Solution-phase combustion synthesis opens the door to the liquid precursor' s solution equilibria with the chelating fuel and solvent properties being an additional source of both complexity and tunability for the synthetic pathway .

[0033] Using SCS , we success fully fabricated LLZO with excellent material properties , including <2 nm roughness with low electronic conductivities . The films are combusted and subsequently annealed for three minutes on a hotplate in an ambient pressure dry-air environment , negating the need for expensive inert gases or slow vacuum steps used extensively in vacuum sputtering and sol-gel processes . The rapid curing step is among the fastest recorded in the literature for any phase of LLZO .

[0034] B2 . Results and Discussion

[0035] B2 . 1 Optimi zation of Combustion Reaction for Thin Films

[0036] SCS has increased in influence for thin film synthesis over the last 15 years . The thin film geometry poses a number of unique opportunities : the substrate acts as a heat sink largely limiting the thermal run-away observed in bulk combustions , the metal-organic complex must be easily deposited as a uni form layer on a substrate , and the film has to have low defectivity for device applications . Taking these thin film coating requirements in mind, the SCS precursor has three important components : the metal cation, a non-volatile oxidant , and a complexing fuel which influence the film combustion temperature , film morphology, and final film chemistry . The chemical identity of the complexing fuel is flexible , though some constraints apply that will influence the combustion reaction . Generally for thin film synthesis , the metal organic complex formed must remain soluble in the polar organic solvent of choice , the fuel should not be heavily oxidi zed such that the combustion is suf ficiently energetic, and the complex should inhibit premature metal hydrolysis . In previous work, we have applied these principles to success fully generate thin film oxides of indium, tin, nickel and aluminum .

[0037] Zirconium presents an additional challenge with a 4+ cationic charge and the lack of a stable metal nitrate . Despite the existence of ZrO (NOa ) 2 , the limited solubility of the partially oxidi zed zirconium cation requires alternative approaches . Rather than start with a metal nitrate and form a metal complex in this work, we begin with the fully chelated metal complexes of acetylacetone (AcAc ) . AcAc is ubiquitously used as a complexing fuel for combustion owing to its ability to form stable complexes with a large number of cations . Given the presence of stable AcAc complexes of La, Zr, and Li , the fully complexed cations were dissolved with nitric acid as the oxidant in a polar solvent rather than the standard approach of forming the complex in solution . The chemical principles of SCS for forming LLZO are outlined in FIG . 1 .

[0038] Upon heating the precursor to above the ignition temperature , SCS occurs with a rapid exothermic event fully forming oxide in <1 minute . Subsequent annealing for 3 minutes results in removal of carbon species . As a result , SCS has substantial potential in scalable manufacturing applications for thin film devices . Moreover, the ability to arrest crystalli zation due to the substrate acting as a heat sink opens the door to forming amorphous oxides at industrially relevant scales .

[0039] Combustion reactions were first optimi zed for ideal combustion temperature and ef ficiency by modulating the fuel / oxidant ratio . Because of starting with fully AcAc- complexed cations and the use of nitric acid as an oxidant to form soluble , combustible metal nitrate / AcAc complexes , solution equilibrium was a critical factor in driving the fuel / oxidant ratio around the cation . For both La and Zr, the amount of nitric acid was modulated to yield fuel / oxidant ratios between 1 and 10 . The resulting solution solubility and stability were assessed along with the combustion in TGA. When optimi zed, the resulting solutions combust at <200 ° C in the bulk ( FIG . 2A) resulting in white or light-yellow powders .

[0040] B2 . 2 Chemical Properties of a-LLZO films

[0041] A time-temperature sweep was then implemented using XPS measurements to optimi ze the curing protocol while removing residual combustion contaminants for LLZO . While the bulk combustion reaction for lanthanum and zirconium occurs below 200 ° C, excess un-reacted nitrogen and carbon remain in the film post-combustion in the LLZO film due to the di f ficulty in combusting lithium . We found a temperature of 425 ° C was capable of fully removing carbon and nitrogen contaminants from the bulk of the lithiated LLZO film in three minutes of total curing time . A residual layer of carbonate on the surface is present in the lithiated LLZO likely due to exposure to trace moisture post-combustion .

[0042] XPS depth profiling in FIG . 2B reveals consistent stoichiometry to the precursor solutions with equivalent concentrations of Zr and La in the LLZO . The film-interface layer with air shows adventitious carbon for LLZO which decreases in the film bulk . The LLZO precursor yields nearly carbon- free films .

[0043] GIXRD ( grazing incidence X-ray di f fraction) was used to determine the phase of the material . The X-ray di f fraction results of FIG . 2C show the fabricated LLZO layer . This film was formed at 425 C and is largely amorphous with small peaks indicating the presence of nano-crystallites that could not be directly assigned to a known phase . Lower temperatures for film curing are expected to give fully amorphous layers . The LLZO film coated in polystyrene has three small peaks which proved di f ficult to index due to their si ze and being within the noise of the measurement . These peaks were compared with references in the ICDD database to no avail , and we postulate they may belong to a very small quantity of a crystalline lithium zirconate phase . High resolution GIWAXS ( grazing incidence wide-angle X-ray scattering) with synchrotron radiation may be required to fully characteri ze the small observed peaks .

[0044] B2 . 3 Surface and Morphological Characteri zation a-LLZO films

[0045] A thorough characteri zation sweep was carried out on combustion synthesi zed a-LLZO using top-down SEM, cross- sectional SEM, and AFM . From top-down SEM we determine our films are highly smooth and conformal , with no phase inhomogeneity detectable from back scattered electrons ( FIG . 3A) . The films were inspected at various magni fications , but no pinholes or negative defects were visible . An analysis of the surface roughness was carried out using tapping AFM and determined the presence of a low detectivity film with root mean square (RMS ) roughness below 2 nm for LLZO ( FIG . 3B ) . Based on these measurements , we confirm the surface roughness created by the process is not enough to cause inherent shorting in the LLZO films , which have a thickness of approximately 50 nm.

[0046] B2 . 4 Electrochemical Properties of LLZO films

[0047] The electrochemical properties of a-LLZO films were also investigated for use as solid state electrolytes . We measured the ionic conductivity, conductance and activation energy of the a-LLZO across a range of temperatures using EIS ( electrochemical impedance spectroscopy) measurements ( FIG . 4 ) . FIG . 5 is an exemplary activation energy measurement .

[0048] Although cubic LLZO is typically the phase of interest for its high ionic conductivity, in the thin film case the conductivity of a-LLZO is able to approach the conductivity of cubic LLZO, while maintaining a lower electronic conductivity than the cubic form . The ionic conductivity can be measured by fitting an equivalent circuit model to the PEIS (potentiostatic electrochemical impedance spectroscopy) curve . We measure an ionic conductivity of 1 • 10“7S / cm at 30 ° C, which improves to 4 • 10“4S / cm at 95 ° C . This matches or surpasses the ionic conductivity of other solid- state materials such as LiPON, ( 10“6- 10“5S / cm) , or cubic LLZO ( 10“4- 10“3S / cm) , especially at higher temperatures . Here LiPON is lithium phosphorous oxynitride .

[0049] This becomes even more signi ficant when considering conductance , which combines film thickness and ionic conductivity into a device-relevant property . Conductance can be calculated using the equation (Equation 1 ) where A is the electrode area, cr is the ionic conductivity, and L is the electrode spacing ( electrolyte thickness ) . Based on the 5-pm spacing of our test structure and assuming a normali zed electrode area, the conductance of our thin a- LLZO films can be calculated to be 2 • 10“4S at 30 ° C, which improves to 1 S at 95 ° C . This high conductance is one of the highest in the literature . Due to the sub- 100 nm thickness of our films as compared to the test structure , we expect through-plane measurements to lead to signi ficant further improvements .

[0050] The signi ficant improvement in both ionic conductivity and conductance with temperature scaling is attributable to the high activation energy, which can be calculated using the Arrhenius equation (Equation 2 ) where cr is ionic conductivity, Eais the activation energy, kb is the Boltzmann constant , cr0is the ionic conductivity pre-exponential factor, and T is temperature . We calculate the activation energy of our films as 1 . 2 eV ( FIG . 5 ) . This activation energy is in line with previously reported tetragonal and amorphous LLZO . While crystalline cubic LLZO is typically generated to minimi ze Li motion activation energy, ionic conductivity is governed by a log cr(T) oc 1 / T relationship, where the slope is directly proportional to the activation energy . As such, our high activation energy leads to excellent ionic conductivity at high temperatures .

[0051] B3 . Conclusions

[0052] Thin film lithium lanthanum zirconate was demonstrated using high-throughput solution combustion synthesis . As a result of the exothermic reaction to form the oxide , in-line and roll-to-roll strategies can be employed to fabricate thin film LLZO battery electrolytes requiring as little as three total minutes of processing time , thereby greatly mitigating energy consumption for manufacturing . The precursors were shown to combust at <200 ° C . Of particular importance is our demonstrated ability to generate a ternary ( Li , La, and Zr ) mixed oxide via combustion synthesis with largely amorphous character . The LLZO film shows a carbon- free interior and no nitrogen contamination indicating the formation of the desired oxide phase . The LLZO films have RMS roughness <2 nm opening the door to generating multilayer devices necessary for thin film dielectrics , membranes , and solid-state separators . We believe these strategies open the door to imminently manufacturable , ultra-compact , and high energy density solid-state device applications in energy storage .

[0053] B4 . Methods

[0054] B4 . 1 Combustion LLZO synthesis and substrates

[0055] Amorphous LLZO thin films were synthesi zed on current collectors from metal-ion precursors at atmospheric conditions by a solution combustion synthesis process in a controlled humidity dry box . The a-LLZO was synthesi zed with the target stoichiometry of Li7LaaZr20i2 . Lithium acetylacetonate ( Sigma-Aldrich) , lanthanum acetylacetonate ( Sigma-Aldrich) , and zirconium acetylacetonate ( Sigma- Aldrich) were generated in three separate precursors , mixed, and filtered using a 0 . 22 pm PVDF syringe filter . Due to the precursor salt hydrophilicity, lithium acetylacetonate was stored under nitrogen in a glovebox . 2-methoxyethanol ( 2ME , Sigma-Aldrich) ( dried over prepared 3A molecular sieves ) and 70% nitric acid ( 99 . 999+% , redistilled) were added . The concentrations of nitric acid were experimentally determined to promote full dissolution and combustion of the precursors. Upon 2ME and nitric acid addition, the solutions were stirred for 1 hour at room temperature and combined in the desired end film stoichiometric ratio before use. In a preferred embodiment, the final solution concentrations for the lithium solution were: 261 mg LiAcAc (2.5 mmol) , 267 pL 70% HNOa (4.2 mmol) and 1.733 mL 2ME . The lanthanum solution included 487.3 mg La(AcAc)s (1 mmol) , 495 pL 70% HNOa (7.8 mmol) , and 1.834 mL 2ME . The zirconium solution included 488 mg Zr(AcAc)4 (1 mmol) , 267 pL 70% HNO3 (4.2 mmol) , and 1.773 mL 2ME . The resulting fuel oxidant ratios were 2.8, 2, and 4.6 for Li, La, and Zr, respectively.

[0056] The solution was deposited via spin-coating onto polished steel current collectors or silicon wafers coated with AI2O3. The AI2O3 was used for XPS measurements and was deposited on (100) single side polished silicon wafers via e-beam evaporation (AJA International, Inc.) using a target thickness of 2000 A and with a rate of 2 angstroms per second (A / s) . Electronic conductivity measurements were taken with a through-plane shadow-masked aluminum symmetric design, whereas ionic conductivity measurements were made using a commercial platinum-coated interdigitated electrode manufactured by Micrux Technologies.

[0057] B4.2 Combustion Synthesis Spin Coating and Blade Deposition

[0058] Lithiated LLZO combustion solution was cured in low humidity (<2% RH) air to prevent the formation of unwanted phases such as lithium hydroxide and lithium carbonate. Spin coated films (24 second, 2000 max rpm, 500 rpm / second acceleration) were layered with no intermediate curing steps and a number of layers between one and six. The stack was combusted and annealed in a single curing step on a hotplate at 450 °C for three minutes. Blade coated films were deposited using a MXI Fixed-gap Blade Coater (MTI Corporation) with a 10 pm gap at room temperature and were subsequently cured using the same curing procedure as the spin-coated films . Lithiated LLZO thin films were subsequently transported using an air- free vessel to an inert nitrogen glovebox where they were stored to minimi ze carbonate formation .

[0059] B4 . 3 X-Ray Photoemission Spectroscopy (XPS ) a-LLZO thin films were deposited as described above on AI2O3 coated silicon wafers . The chemical composition was analyzed using XPS ( PHI XPS Versaprobe 4 ) with a monochromati zed Al source at a vacuum of <1 • 10“6Pa . The samples were transported in a sealed vacuum trans fer system (ULVAC-PHI ) from the dry air box to the XPS without humid air exposure . The samples were depth profiled via Ar+sputtering in a 2 mm x 2 mm spot si ze by a 5 kV, 3 pA in 1- minute intervals until the full film thickness was sputtered . Data was analyzed in MultiPak software .

[0060] B4 . 4 X-Ray Di f fraction (XRD) a-LLZO films were deposited as described above on alumina coated fused silica ( LLZO) . Thin films were measured in grazing incidence using a Cu K-alpha source XRD ( PANalytical Empyrean 2 ) with a mirror / PPC geometry using a PIXcel detector with a 1-degree incidence angle and a 20 range of 20 to 80 deg . The LLZO films were capped in inert polymer using a 2-3 mM polystyrene solution ( 35 kDa ) in anhydrous toluene to prevent air exposure during the measurement . B4 . 5 Scanning Electron Microscopy ( SEM) a-LLZO thin films were deposited as described above on silicon wafers . The surface and cross-sectional morphology were analyzed using scanning electron microscopy ( FEI Magellan 400 XHR) . Samples were cleaved and transported in a vacuum trans fer system (MAST ) from the dry air box to the SEM and were never exposed to humid air to prevent exposure to moisture and subsequent carbonate contamination . Samples were imaged in immersion mode using a low acceleration voltage of 500 V was used with a current of 25 pA and a voltage bias of 1500 V . Parameters were chosen to optimi ze thin film viewing by minimi zing the interaction volume and minimi zing film reactivity with the electron beam reported by other groups in the literature .

[0061] B4 . 6 Atomic Force Microscopy (AFM) a-LLZO thin films were deposited as described above on alumina coated silicon wafers and surface roughness was measured using AFM (Asylum AFM MFP-3D) . A conductive tapping probe (NANOANDMORE ) was used in tapping mode with a rated frequency of 325 Hz . The AFM measurement was carried out within an argon glovebox enclosure at <1 PPM oxygen and humidity levels .

[0062] B4 . 7 Thermogravimetric Analysis ( TGA)

[0063] TGA measurements were performed in a Waters TGA-5500 at a constant heating rate of 10 ° C / min to 400 - 450 ° C with platinum pans under air as the purge gas . The precursors were initially prepared using a vacuum oven under 50- 60 ° C overnight and then packed into a small 5-25 mg pellet in a consumable aluminum DSC pan before being loaded on the Pt

[0064] TGA analysis pan .

[0065] B4 . 8 Impedance spectroscopy and voltage transient measurements

[0066] EIS and steady-state voltage measurements were performed with in-plane LLZO measurements , using a commercial interdigitated Pt back electrode design (Micrux Technologies ) . EIS was recorded between 7 MHz and 100 MHz , with 6 points per decade and a sinusoidal amplitude of 40 mV (BioLogic VSP-300 ) . Measurements were taken immediately after coating - samples were trans ferred air- free to a pure nitrogen glovebox and PEIS was measured on a temperature calibrated hot plate for temperatures from 30 ° C to 100 ° C . Temperature calibration was validated with an external 4- channel k-type data logger thermometer ( Channel Thermometer ) with attached K-type thermocouple .

[0067] Ionic conductivity, activation energy, and electronic conductivity were calculated from EIS measurements . The ionic conductivity was calculated by fitting the semicircular portion of the PEIS measurements to an R-R / Q circuit . The resistance in ohms was calculated from the di f ference between the bulk and charge trans fer resistance which was then converted to ionic conductivities at each temperature . To calculate the activation energies , conductivities were plotted on a natural log plot vs 1 / T , and the slope of the line was used to calculate the activation energy by way of the Arrhenius equation (Equation 2 ) . Lastly, the electronic conductivity was calculated by applying a steady state voltage current for 1 hour . The steady-state resistance was calculated from Ohm' s law, then converted into an electronic conductivity value averaged across all measured voltages .

Claims

CLAIMS1 . A method of making a metal-ionic conductor, the method comprising : depositing a precursor on a substrate , wherein the precursor includes two or more metal species , at least one fuel species , and at least one oxidant species ; combusting the precursor in an ambient processing environment having relative humidity 5% or less to provide an amorphous layer of metal-ionic conductor that is a compound of the two or more metal species and at least one bridging heteroatom species ; wherein at least one of the two or more metal species provides mobile cations in the layer of metal-ionic conductor, and wherein an ionic conductivity of the layer of metalionic conductor is at least l Ox an electronic conductivity of the layer of metal-ionic conductor .2 . The method of claim 1 , wherein the ionic conductivity of the layer of metal-ionic conductor is at least l O Ox the electronic conductivity of the layer of metal-ionic conductor .3 . The method of claim 1 , wherein the ionic conductivity of the layer of metal-ionic conductor is at least 10“6S / cm .4 . The method of claim 1 , wherein the electronic conductivity of the layer of metal-ionic conductor no more than 10“9S / cm .

5. The method of claim 1, wherein the ambient processing environment has a relative humidity of 2% or less.

6. The method of claim 1, wherein the two or more metal species are selected from the group consisting of: Li, Na, K, Al, La, Zr, and Ti.

7. The method of claim 1, wherein the mobile cations includes ions from one or more species selected from the group consisting of: Li, Na, K, and Al.

8. The method of claim 1, wherein the at least one bridging heteroatom species includes one or more species selected from the group consisting of: 0, S, P, N, and halogens.

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