Patent Application
20210043967
This invention relates to electrochemical devices, such as lithium battery electrodes, lithium ion conducting solid state electrolytes, and solid-state lithium ion batteries including these electrodes and solid state electrolytes. The maximum charging rate a solid-state battery can withstand without short-circuiting can be affected by defects on the microstructure of the solid electrolyte, resulting from processing, but also the variables at which the cell is tested. In this disclosure, we describe a means to achieve relevant charging rates without short-circuiting the cell by limiting the electrode area, positioning the electrode where least defect population exist and controlling the external variables for stable lithium electrodeposition.
Current state of the art lithium ion batteries comprise two electrodes (an anode and a cathode), a separator material that keeps the electrodes from touching but allows Li+ ions through, and an electrolyte (which is an organic liquid with lithium salts). During charge and discharge, Li+ ions are exchanged between the electrodes. Currently, the liquid electrolyte used in state of the art (SOA) Li-ion batteries is not compatible with advanced battery concepts, such as the use of a lithium metal anode. All-solid-state batteries (ASSB) with lithium metal as an anode with higher energy density compared to SOA Li-ion battery technology are of interest since they have the potential to meet the electrochemical energy storage demands for applications such as electric vehicles and microelectronics. Moreover, the replacement of a flammable liquid electrolyte with a solid electrolyte (SE), can address safety concerns. Despite significant progress made in achieving ionic conductivities commensurate with liquid electrolytes and achieving low Li-SE interface resistance, the maximum charge rate does not match that of Li-ion technology (≥1C or ˜3 mA/cm2). It has been observed that at charging current densities in the 0.1 to 1 mA/cm2 range, Li metal penetrates all bulk-scale solid electrolytes (polymers, sulfide-based ceramics, oxide-based ceramics), manifested as a potential of ˜0 V, or short-circuit, across the cell under galvanostatic testing. This maximum tolerable current density at and above which Li penetration occurs is also known as critical current density (CCD).
Optimizing the CCD of a solid electrolyte is one of the last major challenges impeding commercialization of ASSB since the Li propagation mechanism has not been clearly elucidated. Generally, a fundamental understanding of the mechanism is challenging given that characterizing a Li-SE buried interface and correlating it to the electrochemical behavior observed is particularly difficult. Furthermore, replacing a liquid with a solid introduces additional challenges compared to SOA Li-ion; such as non-intimate contact between electrolyte and electrode causing an inhomogeneous current distribution across the interface. Additionally, Li with a strong reduction potential, readily forms interfacial layers between the electrode and the electrolyte that not necessarily are ionic conductors impeding charge transfer across materials.
The CCD at which lithium metal penetrates the solid electrolyte is not necessarily an intrinsic property, rather, it is likely an extensive property affected by defects/variables such as porosity, grain boundary and interface resistance, temperature and pressure. These defects/variables can play a role on the creation of ionic current focusing effects or “hot spots” resulting in the localized exceeding of the CCD. Moreover, exceeding the CCD often results in fracture of the solid electrolyte, suggesting there is an existing critical flaw size that when surpassed results in crack propagation of the solid electrolyte. Thus, control over the solid electrolyte microstructure resulting from processing and the variables at which the cell is tested can play a key role in controlling CCD.
What is needed are methods for raising the critical current density for solid-state batteries.
As described above, the critical current density can be affected by defects on the microstructure of the solid electrolyte, resulting from processing, but also the variables at which the cell is tested. In this disclosure, we describe a means to achieve relevant charging rates without short-circuiting the cell by limiting the electrode area, positioning the electrode where least defect population exist and controlling the external variables for stable Li electrodeposition.
In one aspect, the present disclosure provides an electrochemical device comprising: a cathode; a solid state electrolyte including a side having an electrolyte perimeter defining a surface area of the side of the solid state electrolyte; and an anode including a surface region having an anode perimeter defining the surface region of the anode, the surface region of the anode being in contact with the solid state electrolyte wherein at least a portion of the anode perimeter is spaced inward of the electrolyte perimeter. In one version of the electrochemical device, the entire anode perimeter can be spaced inward of the electrolyte perimeter. An area of the surface region of the anode can be a percentage or less than the surface area of the side of the solid state electrolyte, wherein the percentage is an integer between 0 and 100. The area of the surface region of the anode can be 90% or less than the surface area of the side of the solid state electrolyte. The area of the surface region of the anode can be 60% or less than the surface area of the side of the solid state electrolyte. The area of the surface region of the anode can be 30% or less than the surface area of the side of the solid state electrolyte.
In one version of the electrochemical device, the solid-state electrolyte comprises a material selected from the group consisting of lithium lanthanum zirconium oxide (LLZO), aluminum doped LLZO, tantalum doped LLZO, lithium aluminum titanium phosphate (LATP), lithium aluminum germanium phosphate (LAGP), lithium phosphorous sulfides, alkali metal cation-alumina, metal halides, and mixtures thereof.
In another version of the electrochemical device, the solid-state electrolyte comprises a material having the formula LiuRevMwAxOy, wherein
Re can be any combination of elements with a nominal valance of +3 including La, Nd, Pr, Pm, Sm, Sc, Eu, Gd, Tb, Dy, Y, Ho, Er, Tm, Yb, and Lu;
M can be any combination of metals with a nominal valance of +3, +4, +5 or +6 including Zr, Ta, Nb, Sb, W, Hf, Sn, Ti, V, Bi, Ge, and Si;
A can be any combination of dopant atoms with nominal valance of +1, +2, +3 , or +4 including H, Na, K, Rb, Cs, Ba, Sr, Ca, Mg, Fe, Co, Ni, Cu, Zn, Ga, Al, B, and Mn;
u can vary from 3-7.5;
v can vary from 0-3;
w can vary from 0-2;
x can vary from 0-2; and
y can vary from 11-12.5.
In another version of the electrochemical device, the solid-state electrolyte comprises a lithium phosphorous sulfide.
In one version of the electrochemical device, the area of the surface region of the anode is less than 10 mm2. In another version of the electrochemical device, the area of the surface region of the anode is less than 2 mm2.
In one version of the electrochemical device, a critical current density of the electrochemical device can be 2 mA/cm2 or greater. In another version of the electrochemical device, the critical current density of the electrochemical device can be 1 mA/cm2 or greater.
In one version of the electrochemical device, the anode comprises lithium, magnesium, sodium, or zinc. In another version of the electrochemical device, the anode consists essentially of lithium metal.
In another aspect, the present disclosure provides a method for forming an electrochemical device. The method comprises: (a) providing a solid state electrolyte including a side having an electrolyte perimeter defining a surface area of the side of the solid state electrolyte; and (b) placing the side of the solid state electrolyte in contact with a surface region of an electrode to form the electrochemical device, wherein the surface region of the electrode has an electrode perimeter defining the surface region of the electrode, wherein at least a portion of the electrode perimeter is spaced inward of the electrolyte perimeter. In one version of the method, the entire electrode perimeter is spaced inward of the electrolyte perimeter. In one version of the method, an area of the surface region of the electrode is a percentage or less than the surface area of the side of the solid state electrolyte, and the percentage is an integer between 0 and 100. In another version of the method, the area of the surface region of the electrode can be 90% or less than the surface area of the side of the solid state electrolyte. In another version of the method, the area of the surface region of the electrode can be 60% or less than the surface area of the side of the solid state electrolyte. In another version of the method, the area of the surface region of the electrode can be 30% or less than the surface area of the side of the solid state electrolyte.
In one version of the method, the solid-state electrolyte comprises a material selected from the group consisting of lithium lanthanum zirconium oxide (LLZO), aluminum doped LLZO, tantalum doped LLZO, lithium aluminum titanium phosphate (LATP), lithium aluminum germanium phosphate (LAGP), lithium phosphorous sulfides, alkali metal cation-alumina, metal halides, and mixtures thereof.
In another version of the method, the solid-state electrolyte comprises a material having the formula LiuRevMwAxOy, wherein
Re can be any combination of elements with a nominal valance of +3 including La, Nd, Pr, Pm, Sm, Sc, Eu, Gd, Tb, Dy, Y, Ho, Er, Tm, Yb, and Lu;
M can be any combination of metals with a nominal valance of +3, +4, +5 or +6 including Zr, Ta, Nb, Sb, W, Hf, Sn, Ti, V, Bi, Ge, and Si;
A can be any combination of dopant atoms with nominal valance of +1, +2, +3 , or +4 including H, Na, K, Rb, Cs, Ba, Sr, Ca, Mg, Fe, Co, Ni, Cu, Zn, Ga, Al, B, and Mn;
u can vary from 3-7.5;
v can vary from 0-3;
w can vary from 0-2;
x can vary from 0-2; and
y can vary from 11-12.5.
In another version of the method, the solid-state electrolyte comprises a lithium phosphorous sulfide.
In one version of the method, an area of the surface region of the electrode can be less than 10 mm2. In another version of the method, the area of the surface region of the electrode can be less than 2 mm2.
In one version of the method, step (b) comprises pressing the solid state electrolyte and the electrode together using a force in a range of 0.01 MPa to 10 MPa.
In another version of the method, step (b) comprises placing the side of the solid state electrolyte in contact with the surface region of the electrode to form the electrochemical device, wherein the surface region of the electrode is at least partially melted.
In one version of the method, the electrode consists essentially of an alkali metal, and step (b) comprises pressing the solid state electrolyte and the electrode together using a load that is higher than a yield strength of the alkali metal. In one version of the method, the electrode consists essentially of lithium metal.
In another aspect, the present disclosure provides a method for visualizing metal propagation from an anode into a solid state electrolyte during cycling of an electrochemical cell comprising the anode and the solid state electrolyte. The method comprises: (a) providing an electrochemical cell comprising a cathode, an anode, and a solid state electrolyte, wherein the anode comprises a metal; (b) repeatedly discharging and thereafter charging the cell at a current density; and (c) recording metal propagation from the anode into the solid state electrolyte using video microscopy time-synchronized to applied current density. The method can further comprise: (d) quantifying metal filament propagation as a function of the applied current density. The method can further comprise: (d) recording voltage response of the cell during galvanostatic plating of the metal. The method can further comprise: (d) imaging an interface of the anode and the solid state electrolyte after galvanostatic plating of the metal.
In one version of the method, step (a) can comprise providing an electrochemical cell comprising a stack of the cathode, the anode, and the solid state electrolyte, and step (c) can comprise recording metal propagation from the anode into the solid state electrolyte using video microscopy in a viewing direction toward a cross-section of the stack of the cathode, the anode, and the solid state electrolyte. In another version of the method, step (a) can comprise providing an electrochemical cell comprising a solid state electrolyte having a surface, a cathode deposited on the surface of the solid state electrolyte, and an anode deposited on the surface of the solid state electrolyte in spaced relationship with the cathode, and step (c) can comprise recording metal propagation from the anode into the solid state electrolyte using video microscopy in a viewing direction toward the surface of the solid state electrolyte.
In another aspect, the present disclosure provides a method for charging an electrochemical device having a cathode, an anode, and a solid state electrolyte. The method comprises: selecting a charging current density based on accumulated/irreversible damage of a cell of the electrochemical device. The method can further comprise: determining the accumulated/irreversible damage by detecting deviation from ohmic behavior when applying a current density to the cell. In one version of the method, the anode comprises lithium, magnesium, sodium, or zinc. In another version of the method, the anode consists essentially of lithium metal.
The foregoing and other aspects and advantages of the invention will appear from the following description. In the description, reference is made to the accompanying drawings which form a part hereof, and in which there is shown by way of illustration an example embodiment of the invention. Such embodiment does not necessarily represent the full scope of the invention, however, and reference is made therefore to the claims and herein for interpreting the scope of the invention.
The patent or application file contains at least one drawing executed in color. Copies of this patent or patent application publication with color drawing(s) will be provided by the Office upon request and payment of the necessary fee.
The invention will be better understood and features, aspects and advantages other than those set forth above will become apparent when consideration is given to the following detailed description thereof. Such detailed description makes reference to the drawings.
Before any embodiments of the invention are explained in detail, it is to be understood that the invention is not limited in its application to the details of construction and the arrangement of components set forth in the following description or illustrated in the following drawings. The invention is capable of other embodiments and of being practiced or of being carried out in various ways. Also, it is to be understood that the phraseology and terminology used herein is for the purpose of description and should not be regarded as limiting. The use of “including,” “comprising,” or “having” and variations thereof herein is meant to encompass the items listed thereafter and equivalents thereof as well as additional items.
The term “critical current density (CCD)” as used herein refers to the maximum tolerable current density at and above which Li penetration through a solid electrolyte occurs.
One embodiment described herein relates to a method for raising the critical current density for solid-state batteries. In one non-limiting example application, a solid state electrolyte 116 can be used in a lithium metal battery 110 as depicted in
The first current collector 112 and the second current collector 122 can comprise a conductive metal or any suitable conductive material. In some embodiments, the first current collector 112 and the second current collector 122 comprise aluminum, nickel, copper, combinations and alloys thereof. In other embodiments, the first current collector 112 and the second current collector 122 have a thickness of 0.1 microns or greater. It is to be appreciated that the thicknesses depicted in
A suitable active material for the cathode 114 of the lithium metal battery 110 is a lithium host material capable of storing and subsequently releasing lithium ions. An example cathode active material is a lithium metal oxide wherein the metal is one or more aluminum, cobalt, iron, manganese, nickel and vanadium. Non-limiting example lithium metal oxides are LiCoO2 (LCO), LiFeO2, LiMnO2 (LMO), LiMn2O4, LiNiO2 (LNO), LiNixCoyO2, LiMnxCoyO2, LiMnxNiyO2, LiMnxNiyO4, LiNixCoyAlzO2, LiNi1/3Mn1/3Co1/3O2 and others. Another example of cathode active materials is a lithium-containing phosphate having a general formula LiMPO4 wherein M is one or more of cobalt, iron, manganese, and nickel, such as lithium iron phosphate (LFP) and lithium iron fluorophosphates. Many different elements, e.g., Co, Mn, Ni, Cr, Al, or Li, may be substituted or additionally added into the structure to influence electronic conductivity, ordering of the layer, stability on delithiation and cycling performance of the cathode materials. The cathode active material can be selected from the group consisting of cathode materials having a formula LiNixMnyCozO2, wherein x+y+z=1 and x:y:z=1:1:1 (NMC 111), x:y:z=4:3:3 (NMC 433), x:y:z=5:2:2 (NMC 522), x:y:z=5:3:2 (NMC 532), x:y:z=6:2:2 (NMC 622), or x:y:z=8:1:1 (NMC 811). The cathode active material can be a mixture of any number of these cathode active materials. In other embodiments, a suitable material for the cathode 114 of the lithium battery 110 is porous carbon (for a lithium air battery), or a sulfur containing material (for a lithium sulfur battery).
In some embodiments, a suitable active material for the anode 118 of the lithium metal battery 110 consists of lithium metal. In other embodiments, an example anode 118 material consists essentially of lithium metal. Alternatively, a suitable anode 118 consists essentially of magnesium, sodium, or zinc metal.
An example solid state electrolyte 116 material for the lithium metal battery 110 can include an electrolyte material having the formula LiuRevMwAxOy, wherein Re can be any combination of elements with a nominal valance of +3 including La, Nd, Pr, Pm, Sm, Sc, Eu, Gd, Tb, Dy, Y, Ho, Er, Tm, Yb, and Lu;
M can be any combination of metals with a nominal valance of +3, +4, +5 or +6 including Zr, Ta, Nb, Sb, W, Hf, Sn, Ti, V, Bi, Ge, and Si;
A can be any combination of dopant atoms with nominal valance of +1, +2, +3 or +4 including H, Na, K, Rb, Cs, Ba, Sr, Ca, Mg, Fe, Co, Ni, Cu, Zn, Ga, Al, B, and Mn;
u can vary from 3-7.5;
v can vary from 0-3;
w can vary from 0-2;
x can vary from 0-2; and
y can vary from 11-12.5.
Li7La3Zr2O12 (LLZO) materials are beneficial for use as the solid state electrolyte 116 material for the lithium metal battery 110.
Another example solid state electrolyte 116 can include any combination oxide or phosphate materials with a garnet, perovskite, NaSICON, or LiSICON phase. Another example solid state electrolyte 116 can include lithium phosphorous sulfide (e.g., 75Li2S-25P2S5(mol %), 80Li2S-20P2S5(mol %)), lithium aluminum titanium phosphate (LATP), lithium aluminum germanium phosphate (LAGP), alkali metal cation-alumina (e.g., sodium-β-alumina and sodium-β″-alumina), metal halides, or mixtures thereof. The solid state electrolyte 116 of the lithium metal battery 110 can include any solid-like material capable of storing and transporting ions between the anode and cathode, so long as the solid-like material has negligible electronic conductivity and is electrochemically stable against high voltage cathodes and lithium metal anodes.
In the lithium metal battery 110, the solid state electrolyte 116 can include a side having an electrolyte perimeter defining a surface area of the side of the solid state electrolyte 116. The anode 118 can include a surface region having an anode perimeter defining the surface region of the anode 118, the surface region of the anode 118 being in contact with the solid state electrolyte 116. In one embodiment, at least a portion of the anode perimeter is spaced inward of the electrolyte perimeter. In another embodiment, the entire anode perimeter of the anode 118 is spaced inward of the electrolyte perimeter of the solid state electrolyte 116. In another embodiment, an area of the surface region of the anode 118 is a percentage or less than the surface area of the side of the solid state electrolyte 116, and the percentage is an integer between 0 and 100. For example, the area of the surface region of the anode 118 can be 90% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 80% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 70% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 60% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 50% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 40% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 30% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 20% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 10% or less than the surface area of the side of the solid state electrolyte 116.
In one embodiment, an area of the surface region of the anode 118 is less than 10 mm2. In another embodiment, the area of the surface region of the anode 118 is less than 8 mm2. In another embodiment, the area of the surface region of the anode 118 is less than 6 mm2. In another embodiment, the area of the surface region of the anode 118 is less than 4 mm2. In another embodiment, the area of the surface region of the anode 118 is less than 2 mm2.
In one embodiment, a critical current density of the lithium metal battery 110 is 2 mA/cm2 or greater. In another embodiment, a critical current density of the lithium metal battery 110 is 1 mA/cm2 or greater. In another embodiment, a critical current density of the lithium metal battery 110 is 0.5 mA/cm2 or greater.
One example method for forming the lithium metal battery 110 comprises: (a) providing the solid state electrolyte 116 including a side having an electrolyte perimeter defining a surface area of the side of the solid state electrolyte 116; and placing the side of the solid state electrolyte 116 in contact with a surface region of the anode 118 to form the lithium metal battery 110, wherein the surface region of the anode 118 has an anode perimeter defining the surface region of the anode 118, wherein at least a portion of the anode perimeter of the anode 118 is spaced inward of the electrolyte perimeter of the solid state electrolyte 116. Optionally, the entire anode perimeter is spaced inward of the electrolyte perimeter. In the method, the area of the surface region of the anode 118 can be 90% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 80% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 70% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 60% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 50% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 40% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 30% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 20% or less than the surface area of the side of the solid state electrolyte 116; or the area of the surface region of the anode 118 can be 10% or less than the surface area of the side of the solid state electrolyte 116.
In one embodiment of the method, an area of the surface region of the anode 118 is less than 10 mm2. In another embodiment of the method, the area of the surface region of the anode 118 is less than 8 mm2. In another embodiment of the method, the area of the surface region of the anode 118 is less than 6 mm2. In another embodiment of the method, the area of the surface region of the anode 118 is less than 4 mm2. In another embodiment of the method, the area of the surface region of the anode 118 is less than 2 mm2.
In one embodiment of the method, the solid state electrolyte 116 and the anode 118can be pressed together using a force in a range of 0.01 MPa to 10 MPa. In another embodiment of the method, the surface region of the anode 118 is at least partially melted when placing the side of the solid state electrolyte 116 in contact with the surface region of the anode 118 to form the lithium metal battery 110. In another embodiment of the method, the anode 118 consists essentially of an alkali metal (e.g., lithium), and the solid state electrolyte 116 and the anode 118 are pressed together using a load that is higher than a yield strength of the alkali metal.
The invention also provides a method for visualizing metal propagation from an anode into a solid state electrolyte during cycling of an electrochemical cell comprising the anode and the solid state electrolyte. The method includes the steps of (a) providing an electrochemical cell comprising a cathode, an anode, and a solid state electrolyte, wherein the anode comprises a metal; (b) repeatedly discharging and thereafter charging the cell at a current density; and (c) recording metal propagation from the anode into the solid state electrolyte using video microscopy time-synchronized to applied current density. In one embodiment of this method, metal filament propagation is quantified as a function of the applied current density. In another embodiment of the method, the voltage response of the cell is recorded during galvanostatic plating of the metal. The interface of the anode and the solid state electrolyte may be imaged using electron microscopy after galvanostatic plating of the metal. In one embodiment of this method, the electrochemical cell comprises a stack of the cathode, the anode, and the solid state electrolyte, and metal propagation from the anode into the solid state electrolyte is recorded using video microscopy in a viewing direction toward a cross-section of the stack of the cathode, the anode, and the solid state electrolyte. In another embodiment of the method, the electrochemical cell comprises a solid state electrolyte having a surface, a cathode deposited on the surface of the solid state electrolyte, and an anode deposited on the surface of the solid state electrolyte in spaced relationship with the cathode, and metal propagation from the anode into the solid state electrolyte is recorded using video microscopy in a viewing direction toward the surface of the solid state electrolyte.
The invention also provides a method for charging an electrochemical device having a cathode, an anode, and a solid state electrolyte. The method comprises selecting a charging current density based on accumulated/irreversible damage of a cell of the electrochemical device. In the method, the accumulated/irreversible damage can be determined by detecting deviation from ohmic behavior when applying a current density to the cell.
The following Examples are provided in order to demonstrate and further illustrate certain embodiments and aspects of the present invention and are not to be construed as limiting the scope of the invention.
It has been observed that placement of the electrode in certain regions of the solid electrolyte and the area that is being cycled have a dramatic effect on the charging rate obtained under galvanostatic testing. It is believed that defect population size (flaw or pore size) is narrower the smaller the electrode is, but also microstructural defects vary as a function of radial position. Thus, the positioning of the electrode is a variable that enables higher critical current densities compared to regions near the edges of the sample where the defect population is larger even when the area is fixed (see
The most probable critical flaw size is larger going radially from the center towards the edge of the sample, ranging from 2 μm near the center of the sample, going to ˜11 μm near the edges. In agreement, the critical current density values measured decreased as the electrode covers more of the sample area towards the edges. These defects are considered to be prone to act as ion current focusing points and stress concentrators at the interface between Li and the solid electrolyte acting as failure points. Moreover, the behavior observed suggests that this is a direct consequence of the way the ceramic has been processed (in a cylindrical die) indicating that there is an optimal effective area that results in better electrochemical performance of the solid electrolyte.
As shown in
Thus, the upper bound for the Li electrode area, assuming a square geometry, to achieve a specific charging rate with a doped LLZO solid electrolyte, should be (electrode located at the center of a 12.7 millimeter in diameter disk):
As previously mentioned, when two solid materials are put in contact in contrast to a liquid against a solid, intimate contact is not guaranteed. Furthermore, if there is a flux imbalance of Li at the interface, a depletion (or loss of contact area) can be manifested as overpolarization at the stripping interface (see left graph in
Additionally, non-linear potential response under galvanostatic testing is associated with irreversible damage onto the cell as shown in
If comparing two cells, one that has gone through current density segments of 0.1 to 2 mA cm−2 up to failure (short-circuit), 0.1 mA cm−2 increments, and one that goes through 1, 5 mA cm−2, both fixing charge to 0.25 mAh cm−2 per half cycle and measured their CCD (
Solid-state electrolytes (SSEs) have attracted substantial attention for next-generation batteries, primarily due to the promise of enabling Li metal electrodes. Despite significant progress, substantial challenges remain with interfacial stability. In many solid-state electrolyte (SSE) materials, Li metal filaments nucleate at high current densities and propagate until short-circuit. In this example, quantitative analysis of synchronized electrochemical responses and operando video microscopy paired with post-mortem electron microscopy revealed key new insights into the nature of Li filaments that propagate in SSEs. Li penetration was monitored during stepped current, galvanostatic cycling, pulsed current, and extended depth of discharge experiments to study the behavior under a wide variety of battery-relevant conditions. This example probes the coupled electrochemical-morphological-mechanical evolution of Li metal-SSE interfaces, which are critical to engineer the next-generation of solid-state batteries. Operando video microscopy was used to directly observe propagation and cycling of Li filaments inside superionic inorganic solid electrolytes at high current densities (>1 mA/cm2), providing mechanistic insights into one of the most critical challenges facing solid-state Li metal batteries.
Improvements in battery technology over the past several decades have enabled transformative changes to the way we live, travel, work, play, and communicate. From grid-scale storage to electric vehicles (EVs) to medical devices, the ability to efficiently store electrical energy and use it on-demand plays a key role in myriad applications. The demand for better batteries continues to grow, with vehicle electrification rapidly beginning to dominate total demand (Ref. 1). The goal of affordable long-range EVs with sufficiently fast charging time and improved safety has driven tremendous research effort in next-generation batteries with enhanced energy and power density (Ref. 2).
In particular, one of the most promising advances is switching from organic liquid electrolytes which are volatile and flammable to solid-state electrolytes (SSEs). In addition to the immediate improvement to the safety of these cells, this could enable the use of Li metal negative electrodes and next-generation chemistries such as Li-sulfur and Li-air, which could offer dramatically improved gravimetric/volumetric energy density (Ref. 3). To this end, numerous potential SSE materials have been developed in recent years, several of which show great promise (Ref. 4). These materials can be polymers, oxides, sulfides, halides, phosphates, and hydrides or composites thereof (Ref. 4). While each class has its own set of advantages and challenges, in many material systems, Li filaments tend to nucleate and grow towards the positive electrode at high current densities, eventually causing short-circuit (Ref. 5). The current density at which this occurs varies among different material systems. For example, in state of the art ceramic electrolytes, the critical current density (CCD) at which Li penetration occurs has been increased to above 1 mA/cm2 through careful control of material processing (Ref. 6-9). However, increasing the CCD to higher current densities has been hindered by the lack of mechanistic insight into the origins of Li penetration/propagation. Understanding and overcoming this challenge is vital to the implementation of lithium metal solid-state batteries (LMSSBs) in applications where fast charging times are required, including EVs.
Several recent studies have attempted to elucidate the underlying mechanisms for the nucleation and growth of Li filaments/dendrites within SSEs, but a range of questions remained (Ref. 5, 7, 10-14). There is significant variation in the terminology used to describe these structures in literature, but we will use the general terms “Li filament” and “Li penetration”. Among the key remaining questions are when, why, and where Li penetration occurs, and what determines the path of propagation? Early work studied Li filament propagation in polycrystalline LLZO (Ref. 10). It was found that Li propagated in grain boundaries. However, whether or not filaments entered at grain boundaries could not be determined. Regardless, it is possible that grain boundaries could serve as preferential sites for nucleation and pathways for propagation (Ref. 10), and more recent computational work suggested that the localized mechanical properties at grain boundaries could be a rationale for this (Ref. 15). Porz et al. proposed a mechanical model that links the overpotential to stress within a Li-filled crack through chemical potential, and predicted a critical flaw size that would grow at a given applied overpotential (Ref. 5). This work studied several SSE systems, including single crystal Li7La3Zr2O12 (LLZO) and glassy Li2S—P2S5 (lithium phosphorous sulfide), which demonstrates that Li filaments can propagate along paths other than grain boundaries. However, challenges with the experimental setup and geometry have made further validation of the model difficult, and the top-down or angled view of the through-plane cell limited observations of Li propagation (Ref. 16).
More recently, neutron depth profiling was applied to several SSE materials (Ref. 12). This study proposed a different underlying mechanism for nucleation of Li filaments that stems from the inherent electronic conductivity of the SSE. This mechanism was used to explain uniform, isolated accumulation of Li within the bulk of the SSE. Due to the limited spatial resolution of the technique, elevated temperatures were needed to observe changes, and the current densities applied were <200 μA/cm2 at 298° K (Ref. 12). A similar mechanism was proposed by Song et al. (Ref. 17).
In addition to initiation of Li penetration, little is known about the dynamic evolution of Li metal electrodes during cycling and deep discharge conditions. Two recent studies examined the role of diffusion within the Li metal electrode on nucleation of Li filaments and on void formation at the Li/LLZO interface (Ref. 7, 13). These studies probed the diffusion of Li away from the cathodic interface and Li vacancies away from the anodic interface. Wang et al. hypothesized that if the diffusion of Li away from the interface cannot keep up with the rate of plating, there could be a build-up of Li at “hot-spots” that then act as a nucleation site for a Li filament (Ref. 7). Krauskopf et al. studied the formation of voids in the absence of stack pressure when vacancies that are created at the interface during dissolution of Li from the electrode cannot diffuse away from the interface as fast as they are being created (Ref. 13). This leads to a decrease in interfacial contact area, which corresponds with an increase in interfacial impedance and lower CCD.
It has also shown that the interface/interphase between Li and the SSE plays an important role in determining the current density at which Li penetration is observed. In LLZO, reduction of interfacial impedance through rigorous surface cleaning and/or interfacial layers, has enabled current densities up to 1-2 mA/cm2 for planar cells at room temperature (Ref. 7, 8, 18, 19). Li7 NMR chemical shift imaging was used to show that morphology evolution at both the plating and stripping interfaces is linked with the Li penetration (Ref. 20). Interphase formation between Li metal and Li1.4Al0.4Ge1.6(PO4)3 (LAGP) has also been shown to play a role in mechanical fracture of the bulk SSE during cycling (Ref. 21).
Given the breadth of proposed mechanisms that drive Li nucleation and propagation in SSEs, there is a need for additional operando investigations to elucidate the nature and dynamic evolution of Li penetration during cycling under realistic conditions (Ref. 5, 12, 22). Filling the knowledge gaps mentioned above would provide researchers and engineers with the tools and knowledge needed to design high performance LMSSBs. Towards this goal, this example implemented an operando visualization technique that enabled high-quality optical observations within the bulk of LLZO during cycling at relevant current densities (>1 mA/cm2) with low interfacial resistances (<5 Ω-cm2). The time-synchronization of voltage response (voltage trace) with video microscopy enables significant new insights into the mechanisms that lead to the coupled electro-chemo-mechanical response of the system (Ref. 23, 24).
Al-doped LLZO was used as a model system for this example, as the inherent chemical stability against Li metal simplifies the analysis, and translucent/transparent electrolyte films can be fabricated based on recent advances in material processing (Ref. 25). In addition, the mechanical and transport properties have been well-documented (Ref. 26-29), allowing for a detailed analysis. Using both through-plane and in-plane cell geometries, Li penetration was imaged in-operando during propagation and time-synchronized with the corresponding electrochemical signatures. In addition, reversible plating and stripping of the Li within these structures and the formation of “dead Li” during cycling, were directly observed. Four distinct morphologies of Li filaments were identified and described using a combination of operando optical and post-mortem electron microscopy, illustrating that a singular mechanism/mode of failure is insufficient to capture the full complexity of Li penetration in SSEs. The velocity of Li filament propagation was quantified as a function of applied current density, providing insight into the coupling between mechanical and electrochemical behavior. The behavior during deep discharge was examined, and the dynamic evolution of the Li/SSE interfacial area was quantified and correlated with changes in interface resistance. Finally, the in-plane platform was applied to a glassy Li3PS4 SSE, demonstrating the power of the in-plane architecture and the implications of the results for SSEs more broadly.
First, the through-plane cell shown in
The gradual decrease in polarization while the filaments propagate towards the counter electrode can be explained as a result of three factors: [1] an increase in active interfacial area of the cathodic electrode as the structures grow, which decreases interfacial impedance; [2] a decrease in distance between the Li filament and the anode, decreasing the impedance associated with ionic conduction in the SSE (Ref. 20); and [3] formation of a kinetically faster interface for charge transfer, similar to dendrites in liquid electrolytes (Ref. 23). This demonstrates the power of the operando video microscopy platform to provide mechanistic insight into the origins of electrochemical signatures, which can be used in real battery applications where optical access to the cell is not available. For example, this characteristic drop in cell voltage could be used to identify the propagation of a Li filament in one cell of a battery pack with on-board diagnostics during charging to avoid a catastrophic short-circuit failure.
Although the results from this experiment provide valuable insights and clearly show the propagation of Li filaments and the synchronized voltage trace, this platform is not ideal to enable high-resolution imaging and post-mortem analysis. First, this setup requires a very transparent SSE material, and while this is possible to achieve in LLZO, it requires high sintering temperatures (>1300° C.) and long sintering times to achieve large-grained, high density pellets (Ref. 25). This makes throughput very low, and it is difficult to make comparisons to SSEs with conventional processing. In addition, the geometry of the cell means that the electrode edges are preferred nucleation sites. A majority of the filaments nucleated either at the viewing edge or at the opposite edge of the active area.
To overcome challenges with the through-plane platform, an in-plane electrode architecture was developed that enables: [1] high-quality interfaces with low interfacial impedance; [2] improved optical imaging; [3] higher throughput; [4] use of representative SSE materials; and [5] quantitative operando and post-mortem analysis without the need to remove electrodes. A schematic of this cell geometry is shown in A of
This architecture allows for many cells to be fabricated on the same SSE pellet. For example, G in
To validate the in-plane geometry experimentally, CCD tests were performed that mirror those typically used to demonstrate rate capability of SSEs (Ref. 6, 18, 19, 30). Any representative platform for operando characterization should be capable of cycling Li at sufficiently high current densities (>0.5 mA/cm2) and cell polarization (<1 V) without Li penetration. To achieve this, a clean Li/LLZO surface was formed by a combination of polishing and heat treatment, which has been previously shown to eliminate interfacial impedance and improve wettability of molten Li (Ref. 6). Li metal electrodes were deposited by thermal evaporation and heated to 250° C. without stack pressure for 5 minutes, then cooled to room-temperature. This process melted the Li and improved Li-LLZO contact. The Nyquist plot from EIS analysis after this heat-treatment in F of
To compare the behavior of the in-plane cell to typical bulk SSE cells, a constant charge of 0.25 mAh/cm2 was plated in each direction, and then current was successively increased until short-circuit was observed. Under these or similar conditions, the highest reported CCDs at room temperature for planar LLZO samples without an interlayer coating is ˜1 mA/cm2 with an applied stack pressure (Ref. 7, 19). A Li/Mg alloy electrode (Ref. 8) and an MoS2 interlayer (Ref. 18) have been demonstrated to ˜2 mA/cm2. In this in-plane cell, even without any stack pressure, flat voltage plateaus and no Li filaments were observed at 1 and 5 mA/cm2 (see B,D in
Overall, the in-plane cell closely mirrors what is seen in “typical” through-plane LLZO cells (Ref. 10, 32), but occurs at significantly higher current density (˜5-10× higher). The rate capability was consistently high in cells with annealed Li, as shown in
Several different types/morphologies of Li filaments were observed when cycling above the CCD. Here, we classify them into four groups, each of which is shown in
The straight type is shown in A-C of
The second type of structure is “branching” (D-F in
The exact cause of the branching/bifurcation is unknown, however in several cases during post-mortem FIB analysis, impurity grains were observed at the same locations where the crack bifurcated or deviated. A series of SEM images as the FIB removed thin slices of material was used to create a 3D reconstruction of one such location. This appears similar to crack deflection in ceramic matrix composites (Ref. 36-38). The ˜1% of impurities may play a significant role in determining the path of Li filament propagation.
In the 18 in-plane cells examined, the branching type always occurred at high current densities, and never occurred at low current densities. As shown in
The third type, “spalling”, is named for the similarity in appearance to a piece of glass spalling off from a larger sheet (G-I in
The final type of morphology, “diffuse”, was observed only twice during the preparation of this example, and was never observed in the optimized cells which had low interfacial impedance and high CCD. The operando video of a cell exhibited diffuse degradation. During the first three current densities (0.7, 0.8, 0.9 mA/cm2), there is no visible formation of Li filaments. During the 1 mA/cm2 cycle, a subtle diffuse darkening within the SSE is observed. The strong backlighting helps to make the features more visible. As it propagates, a dark spot forms in between the two electrodes. Upon post-mortem optical inspection at higher magnification and with lighting from the top (J in
As the diffuse darkening propagates and reaches the other side, there is almost no change in the cell polarization. A brief and very small drop is observed when the structure stops growing. This suggests that either the electronic conductivity of this feature is low, or that the “fuse effect” occurred almost instantaneously. This phenomena occurs when the electronically conductive pathway between electrodes is melted (like a fuse), and the pathway is broken (Ref. 39). When the polarity is reversed at 1.5 mA/cm2, the structure does not noticeably shrink, but several more spots of Li plating in the middle are observed. Finally, during the last current pulse, a straight type Li filament grows across the cell and causes short-circuit.
K in
Li penetration can be observed in a given cell, further emphasizing that a single mechanism or explanation is insufficient to capture the full complexity of the observed phenomena. The multiple failure types are reminiscent of the multiple failure modes observed in Na β-Al2O38 SSEs (Ref. 40). Moving forward, it is important for the research community to note that the details of the SSE, interface, and test performed may impact the observed Li penetration dramatically.
In addition to studying the conditions that lead to initial Li penetration and the morphology evolution of the interface under those conditions, the in-plane operando experimental platform is ideal for studying the dynamic evolution of filament morphology during cycling. To study this behavior, a cell with smaller electrodes ˜250 μm wide and ˜500 μm tall and a larger spacing between electrodes was used. The increased distance allows more time to observe Li filament propagation before short circuit occurs. To initially nucleate Li penetration, a two second pulse of high current at a nominal current density of 75 mA/cm2 was applied. As shown in A of
When polarity is reversed at the same current for two seconds, the branching structures recede as Li is preferentially stripped from the surface of the Li filaments that grew during the first pulse. While this is happening, similar branching structures nucleate and grow on the opposite side (C in
Examination of the voltage trace in H of
The preferential plating into and stripping from the dendritic structures suggests that there is a difference between the two interfaces. This behavior was also observed in liquid electrolytes, where it was attributed to a thinner solid electrolyte interphase (SEI) on the freshly plated Li inside the dendritic structures. A similar mechanism is likely occurring in the LLZO. Despite all efforts to create a clean and pristine LLZO interface prior to Li evaporation, it is known that some contamination remains (Ref. 42). In contrast, when a dendritic structure forms, the newly formed Li/LLZO interface is truly pristine. For this reason, Li plating into the dendritic structures is the preferred reaction pathway.
As mentioned above, the initial decrease in polarization during the first pulse (B in
As the Li inside the branching structures on the left starts to be exhausted, the polarization begins to rise. The peak occurs at the same moment when the Li is exhausted from the Li filaments on the left (E,F in
Following this initial cycle at high current density, the same amount of charge was repeatedly cycled at nominal current densities of 5 mA/cm2 (5 cycles) and then 10 mA/cm2 (5 cycles), as shown in
Although the structures appeared fully reversible during the first cycle, one of the branching structures had visible Li remaining after the second cycle (D in
H-L in
Since the structures propagated slightly further each cycle, the structures from the two sides eventually met and short-circuit occurred. In an actual battery, there would only be the potential to form Li filaments from the negative electrode, so failure in this way would not occur, but rather the filament from the Li electrode would reach all the way across to the positive electrode.
One of the primary benefits of the in-plane cell architecture is that the high-quality images can be used for quantitative analysis in ways that have not been possible previously. For example, the rate of Li filament propagation was measured as a function of applied current and a linear relationship was observed (see
For each frame in the operando video, the extension of the leading edge of the Li filament was measured by fitting the abrupt increase in brightness between the darker Li filament and the brighter SSE ahead of the Li filament (D in
The Li-filament position was measured during three sets of galvanostatic pulses. Each set consisted of pulses at nominal current densities of 0.5, 1, 2, and 5 mA/cm2. The duration of the pulses was scaled to maintain the same charge for each current in each set. The rate of change in filament length was higher for pulses with higher current, as shown by the labelled slopes in E in
During open-circuit holds after the pulses above 1 mA/cm2, the Li filament receded. For example, over a period of 6 seconds after a pulse of 5 mA/cm2, the filament length shortened by Lrelax≈3 μm (F in
On polished LLZO surfaces, grain boundaries are typically not easily visible in optical or electron microscopy. In some cases, however, the different properties of the grain boundaries make them visible. In C-D in
For Li metal batteries to outperform Li-ion, the amount of excess Li must be limited, as it is extra inactive weight and space in the cell. Thus, understanding the dynamic evolution of electrode morphology during deep discharge is of critical importance. To this end, current pulses of 4 mA/cm2 (just below the measured CCD above) were applied, with a rest period and EIS measurement between each pulse. The resulting operando images and electrochemical response are shown in
As this Li was stripped (right to left in A-C in
During the first pulse, a slight decrease in interfacial resistance is observed (F in
To gain insight into the coupled electrochemical-morphological evolution of the system, the operando optical images were analyzed. Images corresponding to the end of each current pulse were converted to the binary images shown in H in
After the initial increase during the first current pulse, the capacitance exhibits a decreasing trend, as does the apparent electrode area. A decreasing capacitance with depth of charge was also observed in a recent study of Li/LLZO interfaces (Ref. 13). The apparent electrode area decreases approximately linearly with charge, while the capacitance exhibits a highly non-linear decrease. In particular, an abrupt drop in capacitance can be observed between 0.5 and 0.8 mAh/cm2. Interestingly, this is also the depth of charge at which the interfacial resistance begins to increase dramatically (D,E in
This deviation between capacitance and apparent electrode area is unexpected, as capacitance should scale linearly with area for an SSE/electrode interface (Ref. 45). This discrepancy could be a result of multiple factors, including: [1] the apparent electrode area not accurately representing the actual area of Li contact with the LLZO due to interfacial void formation; [2] chemical and/or structural heterogeneities along the Li/LLZO interface, which would affect the local capacitance along the interface. Both of these effects could also contribute to the increase in interfacial resistance with time, and will be discussed in further detail below.
As depicted in G in
Representative SEM analysis on FIB cut cross-sections of the Li/LLZO interface confirms the presence of voids at the anodic interface (
Chemical and/or structural heterogeneities along the Li/LLZO interface could also contribute to the evolution of both interfacial resistance and capacitance over time. As Li will be preferentially removed from areas of the interface with faster kinetics (lower activation barrier), these areas will be exhausted of Li more quickly than areas which have slower kinetics. For example, we have previously shown that even after the LLZO surface preparation techniques used in the present work, a non-zero fraction of the surface has Li2CO3 present (Ref. 42). Furthermore, the presence of the “peaking” voltage behavior described in the reversibility and cycling section above indicates that the kinetics of the bulk Li-SSE interface are slower than at the new interface formed as Li propagation occurs.
As stripping proceeds at the anode, kinetically faster regions along the Li/SSE interface will experience higher local current densities. This will lead to faster vacancy accumulation in the Li in these regions, increasing the likelihood of void formation (G in
This demonstrates how interfacial heterogeneities may be ultimately responsible for the non-linear increase in interface resistance and decrease in interface capacitance observed at 0.5-0.8 mAh/cm2 through the formation of voids.
The fact that stripping of Li at high rates leads to electrode morphology evolution in the form of depleted regions and voids indicates that the design of experimental conditions for evaluating rate capability should take into consideration the intended application. In any LMSSB, stripping of Li from the Li/SSE interface would occur during discharge of the battery, and plating would occur during charging. Thus, the formation of voids would occur during a fast discharge. For example, in EV or personal electronics applications, fast charge is more important than fast discharge, as the goal is to make the battery last at least several hours of use time. This suggests that an asymmetric cycling protocol may be a better indicator of performance than the symmetric tests above.
All of the experiments above were conducted in the model system of LLZO, but the in-plane architecture can be used in a range of SSE materials. For example, similar behavior to nearly all of the results in LLZO above was also observed in a glassy LPS material, where multiple morphologies (straight and spalling) nucleated in the cycle immediately following the first rapid increase in cell polarization during a current pulse. As shown in
This result shows the parallels between these two substantially different electrolyte systems. They have different anions (S vs. O), structure (glassy vs. garnet), electrochemical stability (narrow vs. wide), interphase formation (significant vs. little/none) (Ref. 3), but exhibit very similar Li propagation and electrode evolution in this case. This demonstrates the wider implications of this study for Li metal LMSSBs in general. At any Li/SSE interface, the morphology evolution of the Li metal electrode plays a key role in determining the rate capability. Whether that is Li penetration (of various types), void formation, or wetting/dewetting, it is critical to understand the coupled chemo-electrochemical-mechanical phenomena that drive interfacial performance and stability in LMSSBs.
In this example, an in-plane cell geometry was demonstrated that enables high quality optical imaging of SSEs in-operando at battery-relevant current densities. With this platform, several key observations were made:
These results provide significant new insight into the underlying mechanisms that drive failure in solid state electrolyte (SSE) materials, including how/why Li penetration and void formation occurs, what the nature of the structures is, and how they evolve during cycling. This understanding informs future work to suppress the formation of these structures and enable fast-charging of LMSSBs.
Cubic Al-doped LLZO (Al-LLZO) with a nominal composition of Li6.25Al0.25La3Zr2O12was prepared for the through-plane cell using the solid-state synthetic technique. The methodology has been explained elsewhere in more detail (Ref. 6, 26). A combination of hot-pressing and annealing was used to grow grains and prepare transparent LLZO specimen in this example. First, the calcined powder was densified using a rapid induction hot-press (RIHP) at 1100° C. under a uniaxial 62 MPa pressure for 1 hour in a graphite die under flowing argon atmosphere to achieve >97% relative density. Then, the hot-pressed pellets were embedded in LLZO mother powder with the same nominal composition and annealed in a tube furnace under flowing argon for 50 hours at 1300° C. Details about the method to prepare transparent LLZO can be found elsewhere (Ref. 26).
Al-LLZO samples for in-plane cells were synthesized from starting powders of Li2CO3, Al2O3, La2O3, and ZrO2 from a solid-state reaction method, calcined at 1000° C. for 4 hours in dry air. Densification of Al-LLZO was achieved by Rapid Induction Hot Pressing (RIHP) green bodies of calcined Al-LLZO at 1225° C., 48 MPa for 40 minutes in an argon atmosphere.
For all pellets, polishing and surface preparation to attain low interfacial resistance was done following a previously reported procedure (Ref. 26), adding a final polishing step of 0.1 μm of diamond polishing abrasive for the in-plane cells before heat treatment at 400° C. for 3 hours in an Ar atmosphere.
The amorphous 25Li2S-75P2S5 (mol %) solid electrolyte was synthesized from crystalline Li2S (99.98%, Aldrich) and P2S5 (99%, Sigma Aldrich) by mechanochemical synthesis after being mixed in an agate mortar and pestle. The mixed precursors were placed in a 45 cc zirconia pot with 10 zirconia balls of 10 mm in diameter and 10 zirconia balls of 5 mm in diameter, sealed in a dry Ar-filled glovebox (water concentration below 0.5 ppm), placed inside stainless steel vessels and transported in an inert atmosphere. The pots were spun at 510 RPM for 10 hours in a Planetary Micro Mill (Pulverisette 7, Fritsch GmbH), with 2 hour intervals of milling followed by 10 minute rest intervals at room temperature. Milled LPS powder was hot-pressed at 200° C. for 4 hours at 270 MPa with a heating rate of 0.7° C. min−1.
Immediately following the heat treatment procedure, the LLZO pellets for the through-plane cells were pressed against Li foil (99.9%, Alfa Aesar) that had been scraped with a stainless steel spatula, with areas apart from the active surfaces masked with Kapton tape. In-plane cells were loaded into a glovebox-integrated thermal evaporator (Ångstrom Engineering Inc.). Laser-cut Ni foil was gently clamped on the pellet surface to define the geometry of the electrodes. Molten Li (99.9%, Alfa Aesar) in a Mo crucible was heated to 550° C. while the sample was spinning and thickness was monitored by quartz crystal microbalance until the desired thickness was reached.
Optical imaging of the through-plane cell was performed with a 12 megapixel color camera (Amscope) and a 75 mm c-mount lens (Fujinon) in an Ar glovebox. A 3-watt white LED (Cree) with focusing optics was used for back-lighting while an LED ring light (Amscope) was used for top lighting.
In-plane cells were imaged using a Nikon LV-150 microscope in an Ar glovebox. A 5 megapixel color camera (IDS) was used to capture still images and video. 5×, 10×, 20×, and 50× objectives (Nikon) were used for various experiments. A halogen lamp was used for axial illumination with circular polarizers on the incoming light and in the imaging column for cross-polarization. A 3-watt white LED (Cree) with focusing optics was used for back-lighting while an LED ring light (Amscope) was used for top lighting. Varying the relative intensity of each light source allowed tunability to maximize image quality and contrast.
Extended depth of focus images shown in
Electrochemical measurements were conducted with an SP-200 potentiostat (Bio-Logic) connected to the sample through BNC feedthroughs into the glovebox. Through-plane cells were contacted through the Ni pins that applied stack pressure. In-plane cells were contacted with tungsten needles and xyz manipulators (Signatone).
For the crack growth measurements, optical images were synchronized with corresponding electrochemical measurements using Matlab. The Li filament was initiated by first imparting a small scratch on the LLZO surface at the edge of the cathodic electrode with a diamond scribe and then plating Li at a nominal current density of 1 mA/cm2 until the filament propagated beyond the end of the scratch by a distance (L0) of ˜60 μm (
Y
fit
=A tan h[B(x−C)]+Dx+E.
The fitted function Yfit utilized a hyperbolic tangent function with five fitting terms: brightness magnitude (A) of the hyperbolic tangent, crack position scale (B) and shift (C), linear brightness scaling (D), and brightness shift (E). The key fitting term of interest is the crack position shift (C), with the growth of the Li filament between frames given by the change in C.
For the analysis of electrode area during deep discharge, a color threshold on the R channel was applied in ImageJ. Then the image was converted to greyscale and another threshold was applied to convert the image to binary white and black. Outlier removal of both light and dark regions smaller than 20 pixels was used to reduce noise. Electrode area was then calculated by counting white pixels in each frame with Matlab.
Post-mortem electron microscopy and focused-ion-beam (FIB) cutting was performed on an FEI Helios 650 Nanolab dual beam SEM/FIB. Air exposure during transfer into the instrument was minimized, usually less than one minute between first air exposure and reaching 10−3 Torr vacuum levels, with UHV levels reached within 2 minutes after that. SEM imaging was performed at 1-2 kV accelerating voltage and 100 pA to minimize charging and other artifacts. FIB cross-sectioning and tomography was performed using Ga+ ions with an accelerating voltage of 30 kV and a beam current of 21 nA.
Reconstruction of a series of SEM images during successive FIB polishing cuts was performed using the Avizo software package (Thermo Fisher Scientific Inc.). Images were aligned and shear-corrected, and a non-local means filter was applied for noise reduction. Subsequently, segmentation based on brightness was performed to assign regions of interest.
Electrodes were defined by masking the surface of the sample everywhere but a 1.5 mm wide strip closest to the fractured surface. By doing so, any Li filaments that nucleated were close to the viewing surface and therefore provided good contrast (minimized scattering between the features and the objective), and were close to the focal plane of the lens. The cell was then placed between two nickel pins and a stack pressure of approximately 350 kPa was applied.
Backlighting was used to silhouette the Li features and make them more visible. This also resulted in contrast from flaws that resulted from hot-pressing and polishing (the dark features in D in
The post-mortem image shown in C in
1. G. Zubi, R. Dufo-López, M. Carvalho, G. Pasaoglu, The lithium-ion battery: State of the art and future perspectives. Renew. Sustain. Energy Rev. 89, 292-308 (2018).
2. Z. Gao et al., Promises, Challenges, and Recent Progress of Inorganic Solid-State Electrolytes for All-Solid-State Lithium Batteries. Adv. Mater. 30, 1-27 (2018).
3. Z. Zhang et al., New horizons for inorganic solid state ion conductors. Energy Environ. Sci. 11, 1945-1976 (2018).
4. A. Manthiram, X. Yu, S. Wang, Lithium battery chemistries enabled by solid-state electrolytes. Nat. Rev. Mater. 2, 16103 (2017).
5. L. Porz et al., Mechanism of Lithium Metal Penetration through Inorganic Solid Electrolytes. Adv. Energy Mater. 7, 1-12 (2017).
6. A. Sharafi et al., Surface Chemistry Mechanism of Ultra-Low Interfacial
Resistance in the Solid-State Electrolyte Li7La3Zr2O12. Chem. Mater. 29, 7961-7968 (2017).
7. M. Wang, J. B. Wolfenstine, J. Sakamoto, Temperature dependent flux balance of the Li/Li7La3Zr2O12 interface. Electrochim. Acta. 296, 842-847 (2019).
8. C. Yang et al., An Electron/Ion Dual-Conductive Alloy Framework for High-Rate and High-Capacity Solid-State Lithium-Metal Batteries. Adv. Mater. 31, 1-7 (2019).
9. R. Garcia-Mendez, F. Mizuno, R. Zhang, T. S. Arthur, J. Sakamoto, Effect of Processing Conditions of 75Li2S-25P2S5 Solid Electrolyte on its DC Electrochemical Behavior. Electrochim. Acta. 237, 144-151 (2017).
10. E. J. Cheng, A. Sharafi, J. Sakamoto, Intergranular Li metal propagation through polycrystalline Li6.25Al0.25La3Zr2O12 ceramic electrolyte. Electrochim. Acta. 223, 85-91 (2017).
11. T. Swamy et al., Lithium Metal Penetration Induced by Electrodeposition through Solid Electrolytes: Example in Single-Crystal Li6La3ZrTaO12 Garnet. J. Electrochem. Soc. 165, A3648-A3655 (2018).
12. F. Han et al., High electronic conductivity as the origin of lithium dendrite formation within solid electrolytes. Nat. Energy. 4, 187-196 (2019).
13. T. Krauskopf, H. Hartmann, W. G. Zeier, J. Janek, Toward a Fundamental Understanding of the Lithium Metal Anode in Solid-State Batteries—An Electrochemo-Mechanical Study on the Garnet-Type Solid Electrolyte Li6.25Al0.25La3Zr2O12. ACS Appl. Mater. Interfaces. 11, 14463-14477 (2019).
14. F. Shen, M. B. Dixit, X. Xiao, K. B. Hatzell, Effect of Pore Connectivity on
Li Dendrite Propagation within LLZO Electrolytes Observed with Synchrotron X-ray Tomography. ACS Energy Lett. 3, 1056-1061 (2018).
15. J. Wolfenstine, J. L. Allen, J. Sakamoto, D. J. Siegel, H. Choe, Mechanical behavior of Li− ion-conducting crystalline oxide-based solid electrolytes: a brief review. Ionics (Kiel)., 1-6 (2017).
16. T. Swamy et al., Lithium Metal Penetration Induced by Electrodeposition through Solid Electrolytes: Example in Single-Crystal Li6La3ZrTaO12 Garnet, J. Electrochem. Soc. 165, A3648-A3655 (2018).
17. Y. Song et al., Revealing the Short-Circuiting Mechanism of Garnet-Based Solid-State Electrolyte. Adv. Energy Mater. 1900671, 1900671 (2019).
18. J. Fu et al., In situ formation of a bifunctional interlayer enabled by a conversion reaction to initiatively prevent lithium dendrites in a garnet solid electrolyte. Energy Environ. Sci. 12, 1404-1412 (2019).
19. X. Huang et al., None-Mother-Powder Method to Prepare Dense Li-Garnet Solid Electrolytes with High Critical Current Density. ACS Appl. Energy Mater. 1, 5355-5365 (2018).
20. L. E. Marbella et al., 7 Li NMR Chemical Shift Imaging To Detect Microstructural Growth of Lithium in All-Solid-State Batteries. Chem. Mater. 31, 2762-2769 (2019).
21. J. A. Lewis et al., Interphase Morphology between a Solid-State Electrolyte and Lithium Controls Cell Failure. ACS Energy Lett. 4,591-599 (2019).
22. W. Manalastas et al., Mechanical failure of garnet electrolytes during Li electrodeposition observed by in-operando microscopy. J. Power Sources. 412, 287-293 (2019).
23. K. N. Wood et al., Dendrites and Pits: Untangling the Complex Behavior of Lithium Metal Anodes through Operando Video Microscopy. ACS Cent. Sci. 2,790-801 (2016).
24. K. N. Wood, M. Noked, N. P. Dasgupta, Lithium Metal Anodes: Toward an Improved Understanding of Coupled Morphological, Electrochemical, and Mechanical Behavior. ACS Energy Lett. 2,664-672 (2017).
25. A. Sharafi, C. G. Haslam, R. D. Kerns, J. Wolfenstine, J. Sakamoto, Controlling and correlating the effect of grain size with the mechanical and electrochemical properties of Li7La3Zr2O12 solid-state electrolyte. J. Mater. Chem. A. 5, 21491-21504 (2017).
26. W. E. Tenhaeff et al., Resolving the Grain Boundary and Lattice Impedance of Hot-Pressed Li7La3Zr2O12 Garnet Electrolytes. ChemElectroChem. 1, 375-378 (2014).
27. S. Yu et al., Elastic Properties of the Solid Electrolyte Li7La3Zr2O12 (LLZO). Chem. Mater. 28,197-206 (2016).
28. R. Murugan, V. Thangadurai, W. Weppner, Fast lithium ion conduction in garnet-type Li7La3Zr2O12. Angew. Chemie-Int. Ed. 46,7778-7781 (2007).
29. V. Thangadurai, S. Narayanan, D. Pinzaru, Garnet-type solid-state fast Li ion conductors for Li batteries: Critical review. Chem. Soc. Rev. 43,4714-4727 (2014).
30. Y. Lu et al., An: In situ element permeation constructed high endurance Li-LLZO interface at high current densities. J. Mater. Chem. A. 6,18853-18858 (2018).
31. A. Sharafi, H. M. Meyer, J. Nanda, J. Wolfenstine, J. Sakamoto, Characterizing the Li-Li7La3Zr2O12 interface stability and kinetics as a function of temperature and current density. J. Power Sources. 302, 135-139 (2016).
32. N. J. Taylor et al., Demonstration of high current densities and extended cycling in the garnet Li7La3Zr2O12 solid electrolyte, J. Power Sources. 396, 314-318 (2018).
33. K. Trustrum, Jayatilaka, Statistical approach to brittle fracture. J. Mater. Sci. 12, 1426-1430 (1977).
34. W. S. LePage et al., Lithium Mechanics: Roles of Strain Rate and Temperature and Implications for Lithium Metal Batteries. J. Electrochem. Soc. 166, A89-A97 (2019).
35. A. Masias, N. Felten, R. Garcia-Mendez, J. Wolfenstine, J. Sakamoto, Elastic, plastic, and creep mechanical properties of lithium metal, J. Mater. Sci. 54, 2585-2600 (2019).
36. Z. Xia et al., Direct observation of toughening mechanisms in carbon nanotube ceramic matrix composites. Acta Mater. 52, 931-944 (2004).
37. K. K. Chawla, in Composite Materials: Science and Engineering (Springer, New York, ed. 3rd, 2012; https://link.springer.com/book/10.1007%2F978-0-387-74365-3#toc), pp. 249-290.
38. K. M. Sevener, J. M. Tracy, Z. Chen, J. D. Kiser, S. Daly, Crack opening behavior in ceramic matrix composites. J. Am. Ceram. Soc. 100, 4734-4747 (2017).
39. M. Rosso et al., Dendrite short-circuit and fuse effect on Li/polymer/Li cells. Electrochim. Acta. 51, 5334-5340 (2006).
40. L. C. De Jonghe, L. Feldman, A. Beuchele, Slow degradation and electron conduction in sodium/beta-aluminas. J. Mater. Sci. 16, 780-786 (1981).
41. A. Gupta et al., Evaluating the Effects of Temperature and Pressure on Li/PEO-LiTFSI Interfacial Stability and Kinetics. J. Electrochem. Soc. 165, A2801-A2806 (2018).
42. A. Sharafi et al., Surface Chemistry Mechanism of Ultra-Low Interfacial Resistance in the Solid-State Electrolyte Li7La3Zr2O12. Chem. Mater. 29, 7961-7968 (2017).
43. S. Wenzel et al., Direct Observation of the Interfacial Instability of the Fast Ionic Conductor Li10GeP2S12 at the Lithium Metal Anode. Chem. Mater. 28, 2400-2407 (2016).
44. S. Yu, D. J. Siegel, Grain Boundary Softening: A Potential Mechanism for Lithium Metal Penetration through Stiff Solid Electrolytes. ACS Appl. Mater. Interfaces.
10, 38151-38158 (2018).
45. X. Ge, C. Fu, S. H. Chan, Double layer capacitance of anode/solid-electrolyte interfaces. Phys. Chem. Chem. Phys. 13, 15134-15142 (2011).
The citation of any document is not to be construed as an admission that it is prior art with respect to the present invention.
Thus, the invention provides a means to achieve relevant charging rates without short-circuiting a lithium battery cell by limiting the electrode area, positioning the electrode where least defect population exist and controlling the external variables for stable lithium electrodeposition.
Although the invention has been described in considerable detail with reference to certain embodiments, one skilled in the art will appreciate that the present invention can be practiced by other than the described embodiments, which have been presented for purposes of illustration and not of limitation. Therefore, the scope of the appended claims should not be limited to the description of the embodiments contained herein. Various features and advantages of the invention are set forth in the following claims.
This application claims priority to U.S. Patent Application No. 62/885,107 filed Aug. 9, 2019.
This invention was made with government support under grant number DE-AR-0000653 awarded by Advanced Research Projects Agency. The government has certain rights in this invention.
| Number | Date | Country | |
|---|---|---|---|
| 62885107 | Aug 2019 | US |
